EDITOR IN CHIEF Rudy J. M. Konings European Commission, Joint Research Centre, Institute for Transuranium Elements, Karlsruhe, Germany SECTION EDITORS Todd R. Allen Department of Engineering Physics, University of Wisconsin, Madison, WI, USA Roger E. Stoller Materials Science and Technology Division, Oak Ridge National Laboratory, Oak Ridge, TN, USA Shinsuke Yamanaka Division of Sustainable Energy and Environmental Engineering, Graduate School of Engineering, Osaka University, Osaka, Japan Elsevier Radarweg 29, PO Box 211, 1000 AE Amsterdam, The Netherlands The Boulevard, Langford Lane, Kidlington, Oxford OX5 1GB, UK 225 Wyman Street, Waltham, MA 02451, USA Copyright © 2012 Elsevier Ltd. All rights reserved The following articles are US Government works in the public domain and not subject to copyright: Radiation Effects in UO2 TRISO-Coated Particle Fuel Performance Composite Fuel (cermet, cercer) Metal Fuel-Cladding Interaction No part of this publication may be reproduced, stored in a retrieval system or transmitted in any form or by any means electronic, mechanical, photocopying, recording or otherwise without the prior written permission of the publisher Permissions may be sought directly from Elsevier’s Science & Technology Rights Department in Oxford, UK: phone (þ44) (0) 1865 843830; fax (þ44) (0) 1865 853333; email:
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Alternatively you can submit your request online by visiting the Elsevier web site at http://elsevier.com/locate/permissions, and selecting Obtaining permission to use Elsevier material Notice No responsibility is assumed by the publisher for any injury and/or damage to persons or property as a matter of products liability, negligence or otherwise, or from any use or operation of any methods, products, instructions or ideas contained in the material herein, Because of rapid advances in the medical sciences, in particular, independent verification of diagnoses and drug dosages should be made British Library Cataloguing in Publication Data A catalogue record for this book is available from the British Library Library of Congress Catalog Number: 2011929343 ISBN (print): 978-0-08-056027-4 For information on all Elsevier publications visit our website at books.elsevier.com Cover image courtesy of Professor David Sedmidubský, The Institute of Chemical Technology, Prague Printed and bound in Spain 12 13 14 15 16 10 9 8 7 6 5 4 3 2 1 Editorial : Gemma Mattingley Production : Nicky Carter EDITORS BIOGRAPHIES Rudy Konings is currently head of the Materials Research Unit in the Institute for Transuranium Elements (ITU) of the Joint Research Centre of the European Commission. His research interests are nuclear reactor fuels and actinide materials, with particular emphasis on high temperature chemistry and thermodynamics. Before joining ITU, he worked on nuclear fuel-related issues at ECN (the Energy Research Centre of the Netherlands) and NRG (Nuclear Research and Consultancy Group) in the Netherlands. Rudy is editor of Journal of Nuclear Materials and is professor at the Delft University of Technology (Netherlands), where he holds the chair of ‘Chemistry of the nuclear fuel cycle.’ Roger Stoller is currently a Distinguished Research Staff Member in the Materials Science and Technology Division of the Oak Ridge National Laboratory and serves as the ORNL Program Manager for Fusion Reactor Materials for ORNL. He joined ORNL in 1984 and is actively involved in research on the effects of radiation on structural materials and fuels for nuclear energy systems. His primary expertise is in the area of computa- tional modeling and simulation. He has authored or coauthored more than 100 publications and reports on the effects of radiation on materials, as well as edited the proceedings of several international conferences. Todd Allen is an Associate Professor in the Department of Engineering Physics at the University of Wisconsin – Madison since 2003. Todd’s research expertise is in the area of materials-related issues in nuclear reactors, specifi- cally radiation damage and corrosion. He is also the Scientific Director for the Advanced Test Reactor National Scientific User Facility as well as the Director for the Center for Material Science of Nuclear Fuel at the Idaho National Laboratory, positions he holds in conjunction with his faculty position at the University of Wisconsin. v vi Editors Biographies Shinsuke Yamanaka is a professor in Division of Sustainable Energy and Environmental Engineering, Graduate School of Engineering, Osaka University since 1998. He has studied the thermophysics and thermochem- istry of nuclear fuel and materials. His research for the hydrogen behavior in LWR fuel cladding is notable among his achievements and he received the Young Scientist Awards (1980) and the Best Paper Awards (2004) from Japan Atomic Energy Society. Shinsuke is the program officer of Japan Science and Technology Agency since 2005 and the visiting professor of Fukui University since 2009, and he is also the associate dean of Graduate School of Engineering, Osaka University since 2011. PREFACE There are essentially three primary energy sources for the billions of people living on the earth’s surface: the sun, radioactivity, and gravitation. The sun, an enormous nuclear fusion reactor, has transmitted energy to the earth for billions of years, sustaining photosynthesis, which in turn produces wood and other combustible resources (biomass), and the fossil fuels like coal, oil, and natural gas. The sun also provides the energy that steers the climate, the atmospheric circulations, and thus ‘fuelling’ wind mills, and it is at the origin of photovoltaic processes used to produce electricity. Radioactive decay of primarily uranium and thorium heats the earth underneath us and is the origin of geothermal energy. Hot springs have been used as a source of energy from the early days of humanity, although it took until the twentieth century for the potential of radioactivity by fission to be discovered. Gravitation, a non-nuclear source, has been long used to generate energy, primarily in hydropower and tidal power applications. Although nuclear processes are thus omnipresent, nuclear technology is relatively young. But from the moment scientists unraveled the secrets of the atom and its nucleus during the twentieth century, aided by developments in quantum mechanics, and obtained a fundamental understanding of nuclear fission and fusion, humanity has considered these nuclear processes as sources of almost unlimited (peaceful) energy. The first fission reactor was designed and constructed by Enrico Fermi in 1942 in Chicago, the CP1, based on the fission of uraniumby neutron capture. AfterWorldWar II, a rapid exploration of fission technology took place in theUnited States and the Union of Soviet Socialist Republics, and after the Atoms for Peace speech by Eisenhower at the UnitedNations Congress in 1954, also in Europe and Japan. Avariety of nuclear fission reactors were explored for electricity generation and with them the fuel cycle. Moreover, the possibility of controlled fusion reactions has gained interest as a technology for producing energy from one of themost abundant elements on earth, hydrogen. The environment to which materials in nuclear reactors are exposed is one of extremes with respect to temperature and radiation. Fuel pins for nuclear reactors operate at temperatures above 1000 �C in the center of the pellets, in fast reactor oxide fuels even above 2000 �C, whereas the effects of the radiation (neutrons, alpha particles, recoil atoms, fission fragments) continuously damage the material. The cladding of the fuel and the structural and functional materials in the fission reactor core also operate in a strong radiation field, often in a dynamic corrosive environment of the coolant at elevated temperatures. Materials in fusion reactors are exposed to the fusion plasma and the highly energetic particles escaping from it. Furthermore, in this technology, the reactor core structures operate at high temperatures. Materials science for nuclear systems has, therefore, been strongly focussed on the development of radiation tolerant materials that can operate in a wide range of temperatures and in different chemical environments such as aqueous solutions, liquid metals, molten salts, or gases. The lifetime of the plant components is critical in many respects and thus strongly affects the safety as well as the economics of the technologies. With the need for efficiency and competitiveness in modern society, there is a strong incentive to improve reactor components or to deploy advanced materials that are continuously developed for improved performance. There are many examples of excellent achievements in this respect. For example, with the increase of the burnup of the fuel for fission reactors, motivated by improved economics and a more efficient use of resources, the Zircaloy cladding (a Zr–Sn alloy) of the fuel pins showed increased susceptibility to coolant corrosion, but within a relatively short period, a different zirconium-based alloy was developed, tested, qualified, and employed, which allowed reliable operation in the high burnup range. xxi xxii Preface Nuclear technologies also produce waste. It is the moral obligation of the generations consuming the energy to implement an acceptable waste treatment and disposal strategy. The inherent complication of radioactivity, the decay that can span hundreds of thousands of years, amplifies the importance of extreme time periods in the issue of corrosion and radiation stability. The search for storage concepts that can guarantee the safe storage and isolation of radioactive waste is, therefore, another challenging task for materials science, requiring a close examination of natural (geological) materials and processes. The more than 50 years of research and development of fission and fusion reactors have undoubtedly demonstrated that the statement ‘technologies are enabled by materials’ is particularly true for nuclear technology. Although the nuclear field is typically known for its incremental progress, the challenges posed by the next generation of fission reactors (Generation IV) as well as the demonstration of fusion reactors will need breakthroughs to achieve their ambitious goals. This is being accompanied by an important change in materials science, with a shift of discovery through experiments to discovery through simulation. The progress in numerical simulation of the material evolution on a scientific and engineering scale is growing rapidly. Simulation techniques at the atomistic or meso scale (e.g., electronic structure calculations, molecular dynam- ics, kinetic Monte Carlo) are increasingly helping to unravel the complex processes occurring in materials under extreme conditions and to provide an insight into the causes and thus helping to design remedies. In this context, Comprehensive Nuclear Materials aims to provide fundamental information on the vast variety of materials employed in the broad field of nuclear technology. But to do justice to the comprehensiveness of the work, fundamental issues are also addressed in detail, as well as the basics of the emerging numerical simulation techniques. R.J.M. Konings European Commission, Joint Research Centre, Institute for Transuranium Elements, Karlsruhe, Germany T.R. Allen Department of Engineering Physics, Wisconsin University, Madison, WI, USA R. Stoller Materials Science and Technology Division, Oak Ridge National Laboratory, Oak Ridge, TN, USA S. Yamanaka Division of Sustainable Energy and Environmental Engineering, Graduate School of Engineering, Osaka University, Osaka, Japan FOREWORD ‘Nuclear materials’ denotes a field of great breadth and depth, whose topics address applications and facilities that depend upon nuclear reactions. The major topics within the field are devoted to the materials science and engineering surrounding fission and fusion reactions in energy conversion reactors. Most of the rest of the field is formed of the closely related materials science needed for the effects of energetic particles on the targets and other radiation areas of charged particle accelerators and plasma devices. A more complete but also more cumbersome descriptor thus would be ‘the science and engineering of materials for fission reactors, fusion reactors, and closely related topics.’ In these areas, the very existence of such technologies turns upon our capabilities to understand the physical behavior of materials. Performance of facilities and components to the demanding limits required is dictated by the capabilities of materials to withstand unique and aggressive environments. The unifying concept that runs through all aspects is the effect of radiation on materials. In this way, the main feature is somewhat analogous to the unifying concept of elevated temperature in that part of materials science and engineering termed ‘high-temperature materials.’ Nuclear materials came into existence in the 1950s and began to grow as an internationally recognized field of endeavor late in that decade. The beginning in this field has been attributed to presentations and discussions that occurred at the First and Second International Conferences on the Peaceful Uses of Atomic Energy, held in Geneva in 1955 and 1958. Journal of Nuclear Materials, which is the home journal for this area of materials science, was founded in 1959. The development of nuclear materials science and engineering took place in the same rapid growth time period as the parent field of materials science and engineering. And similarly to the parent field, nuclear materials draws together the formerly separate disciplines of metallurgy, solid-state physics, ceramics, and materials chemistry that were early devoted to nuclear applications. The small priest- hood of first researchers in half a dozen countries has now grown to a cohort of thousands, whose home institutions are anchored in more than 40 nations. The prodigious work, ‘Comprehensive Nuclear Materials,’ captures the essence and the extensive scope of the field. It provides authoritative chapters that review the full range of endeavor. In the present day of glance and click ‘reading’ of short snippets from the internet, this is an old-fashioned book in the best sense of the word, which will be available in both electronic and printed form. All of the main segments of the field are covered, as well as most of the specialized areas and subtopics. With well over 100 chapters, the reader finds thorough coverage on topics ranging from fundamentals of atom movements after displacement by energetic particles to testing and engineering analysis methods of large components. All the materials classes that have main application in nuclear technologies are visited, and the most important of them are covered in exhaustive fashion. Authors of the chapters are practitioners who are at the highest level of achievement and knowledge in their respective areas. Many of these authors not only have lived through a substantial part of the history sketched above, but they themselves are the architects. Without those represented here in the author list, the field would certainly be a weaker reflection of itself. It is no small feat that so many of my distinguished colleagues could have been persuaded to join this collective endeavor and to make the real sacrifices entailed in such time-consuming work. I congratulate the Editor, Rudy Konings, and xxiii xxiv Foreword the Associate Editors, Roger Stoller, Todd Allen, and Shinsuke Yamanaka. This book will be an important asset to young researchers entering the field as well as a valuable resource to workers engaged in the enterprise at present. Dr. Louis K. Mansur Oak Ridge, Tennessee, USA Permission Acknowledgments The following material is reproduced with kind permission of Cambridge University Press Figure 15 of Oxide Dispersion Strengthened Steels Figure 15 of Minerals and Natural Analogues Table 10 of Spent Fuel as Waste Material Figure 21b of Radiation-Induced Effects on Microstructure www.cambridge.org The following material is reproduced with kind permission of American Chemical Society Figure 2 of Molten Salt Reactor Fuel and Coolant Figure 22 of Molten Salt Reactor Fuel and Coolant Table 9 of Molten Salt Reactor Fuel and Coolant Figure 6 of Thermodynamic and Thermophysical Properties of the Actinide Nitrides www.acs.org The following material is reproduced with kind permission of Wiley Table 3 of Properties and Characteristics of SiC and SiC/SiC Composites Table 4 of Properties and Characteristics of SiC and SiC/SiC Composites Table 5 of Properties and Characteristics of SiC and SiC/SiC Composites Figure 5 of Advanced Concepts in TRISO Fuel Figure 6 of Advanced Concepts in TRISO Fuel Figure 30 of Material Performance in Supercritical Water Figure 32 of Material Performance in Supercritical Water Figure 19 of Tritium Barriers and Tritium Diffusion in Fusion Reactors Figure 9 of Waste Containers Figure 13 of Waste Containers Figure 21 of Waste Containers Figure 11 of Carbide Fuel Figure 12 of Carbide Fuel Figure 13 of Carbide Fuel Figure 4 of Thermodynamic and Thermophysical Properties of the Actinide Nitrides Figure 2 of The U–F system Figure 18 of Fundamental Point Defect Properties in Ceramics Table 1 of Fundamental Point Defect Properties in Ceramics Figure 17 of Radiation Effects in SiC and SiC-SiC Figure 21 of Radiation Effects in SiC and SiC-SiC Figure 6 of Radiation Damage in Austenitic Steels Figure 7 of Radiation Damage in Austenitic Steels Figure 17 of Ceramic Breeder Materials Figure 33a of Carbon as a Fusion Plasma-Facing Material Figure 34 of Carbon as a Fusion Plasma-Facing Material i 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SiC/SiC Composites Figure 10 of Properties and Characteristics of SiC and SiC/SiC Composites Figure 11 of Properties and Characteristics of SiC and SiC/SiC Composites Figure 12 of Properties and Characteristics of SiC and SiC/SiC Composites Figure 22d of Fission Product Chemistry in Oxide Fuels Figure 3 of Behavior of LWR Fuel During Loss-of-Coolant Accidents Figure 14a of Irradiation Assisted Stress Corrosion Cracking Figure 14b of Irradiation Assisted Stress Corrosion Cracking Figure 14c of Irradiation Assisted Stress Corrosion Cracking Figure 25a of Irradiation Assisted Stress Corrosion Cracking Figure 25b of Irradiation Assisted Stress Corrosion Cracking Figure 1 of Properties of Liquid Metal Coolants Figure 5b of Fast Spectrum Control Rod Materials Figure 3 of Oxide Fuel Performance Modeling and Simulations Figure 8 of Oxide Fuel Performance Modeling and Simulations Figure 10 of Oxide Fuel Performance Modeling and Simulations Figure 11 of Oxide Fuel Performance Modeling and 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Carbides Figure 28b of Thermodynamic and Thermophysical Properties of the Actinide Carbides Figure 2 of Physical and Mechanical Properties of Copper and Copper Alloys Figure 5 of Physical and Mechanical Properties of Copper and Copper Alloys Figure 6 of The Actinides Elements: Properties and Characteristics Figure 10 of The Actinides Elements: Properties and Characteristics Figure 11 of The Actinides Elements: Properties and Characteristics Figure 12 of The Actinides Elements: Properties and Characteristics Figure 15 of The Actinides Elements: Properties and Characteristics Table 1 of The Actinides Elements: Properties and Characteristics Table 6 of The Actinides Elements: Properties and Characteristics Figure 25 of Fundamental Properties of Defects in Metals Table 1 of Fundamental Properties of Defects in Metals Table 7 of Fundamental Properties of Defects in Metals Table 8 of Fundamental Properties of Defects in Metals www.springer.com The following material is reproduced with kind permission of Taylor & Francis Figure 9 of Radiation-Induced Segregation Figure 6 of Radiation Effects in Zirconium Alloys Figure 1 of Dislocation Dynamics Figure 25 of Radiation Damage Using Ion Beams Figure 26 of Radiation Damage Using Ion Beams Figure 27 of Radiation Damage Using Ion Beams Figure 4 of Radiation-Induced Effects on Material Properties of Ceramics (Mechanical and Dimensional) Figure 7 of The Actinides Elements: Properties and Characteristics Figure 20 of The Actinides Elements: Properties and Characteristics Figure 18a of Primary Radiation Damage Formation Figure 18b of Primary Radiation Damage Formation Figure 18c of Primary Radiation Damage Formation Figure 18d of Primary Radiation Damage Formation Figure 18e of Primary Radiation Damage Formation Figure 18f of Primary Radiation Damage Formation Figure 1 of Radiation-Induced Effects on Microstructure Figure 27 of Radiation-Induced Effects on Microstructure Figure 5 of Performance of Aluminum in Research Reactors 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Irradiated Metals Figure 16d of Atomic-Level Dislocation Dynamics in Irradiated Metals Figure 16e of Atomic-Level Dislocation Dynamics in Irradiated Metals Figure 17a of Atomic-Level Dislocation Dynamics in Irradiated Metals Figure 17b of Atomic-Level Dislocation Dynamics in Irradiated Metals Figure 17c of Atomic-Level Dislocation Dynamics in Irradiated Metals Figure 17d of Atomic-Level Dislocation Dynamics in Irradiated Metals www.taylorandfrancisgroup.com http://www.taylorandfrancisgroup.com 4.01 Radiation Effects in Zirconium Alloys F. Onimus and J. L. Béchade Commissariat à l’Energie Atomique, Gif-sur-Yvette, France � 2012 Elsevier Ltd. All rights reserved. 4.01.1 Irradiation Damage in Zirconium Alloys 2 4.01.1.1 Damage Creation: Short-Term Evolution 2 4.01.1.1.1 Neutron–zirconium interaction 2 4.01.1.1.2 Displacement energy in zirconium 2 4.01.1.1.3 Displacement cascade in zirconium 2 4.01.1.2 Evolution of Point Defects in Zirconium: Long-Term Evolution 4 4.01.1.2.1 Vacancy formation and migration energies 4 4.01.1.2.2 SIA formation and migration energies 4 4.01.1.2.3 Evolution of point defects: Impact of the anisotropic diffusion of SIAs 6 4.01.1.3 Point-Defect Clusters in Zirconium Alloys 7 4.01.1.3.1 hai Dislocation loops 7 4.01.1.3.2 hai Loop formation: Mechanisms 8 4.01.1.3.3 hci Component dislocation loops 9 4.01.1.3.4 hci Loop formation: Mechanisms 9 4.01.1.3.5 Void formation 10 4.01.1.4 Secondary-Phase Evolution Under Irradiation 10 4.01.1.4.1 Crystalline to amorphous transformation of Zr-(Fe,Cr,Ni) intermetallic precipitates 10 4.01.1.4.2 Irradiation effects in Zr–Nb alloys: Enhanced precipitation 13 4.01.2 Postirradiation Mechanical Behavior 14 4.01.2.1 Mechanical Behavior During Tensile Testing 14 4.01.2.1.1 Irradiation hardening: Macroscopic behavior 14 4.01.2.1.2 Irradiation hardening: Mechanisms 14 4.01.2.1.3 Post-yield deformation: Macroscopic behavior 16 4.01.2.1.4 Post-yield deformation: Mechanisms 16 4.01.2.2 Effect of Postirradiation Heat Treatment 17 4.01.2.3 Postirradiation Creep 18 4.01.3 Deformation Under Irradiation 19 4.01.3.1 Irradiation Growth 19 4.01.3.1.1 Irradiation growth: Macroscopic behavior 19 4.01.3.1.2 Irradiation growth: Mechanisms 21 4.01.3.2 Irradiation Creep 24 4.01.3.2.1 Irradiation creep: Macroscopic behavior 24 4.01.3.2.2 Irradiation creep: Mechanisms 25 4.01.3.3 Outlook 26 References 27 Abbreviations BWR Boiling-water reactor CANDU Canadian deuterium uranium DAD Diffusion anisotropy difference EAM Embedded atom method EID Elastic interaction difference FP-LMTO Full-potential linear muffin-tin orbital GGA Generalized gradient approximation hcp Hexagonal close-packed HVEM High-voltage electron microscope LDA Local density approximation MB Many body MD Molecular dynamics NRT Norgett–Robinson–Torrens 1 2 Radiation Effects in Zirconium Alloys PKA Primary knocked-on atom PWR Pressurized water reactor RXA Recrystallization annealed SANS Small-angle neutron scattering SIA Self interstitial atom SIPA Stress-induced preferential absorption SIPA-AD Stress preferential induced nucleation- anisotropic diffusion SIPN Stress preferential induced nucleation SRA Stress-relieved annealed TEM Transmission electron microscopy Tm Melting temperature UTS Ultimate tensile strength YS Yield stress 4.01.1 Irradiation Damage in Zirconium Alloys 4.01.1.1 Damage Creation: Short-Term Evolution 4.01.1.1.1 Neutron–zirconium interaction Zirconium alloys are used as structural components for light and heavy water nuclear reactor cores because of their low capture cross section to thermal neutrons and their good corrosion resistance. In a nuclear reactor core, zirconium alloys are subjected to a fast neutron flux (E> 1 MeV), which leads to irradiation damage of the material. In the case of metallic alloys, the irradiation damage is mainly due to elastic interaction between fast neutrons and atoms of the alloy that displace atoms from their crystallo- graphic sites (depending on the energy of the incom- ing neutron) and can create point defects without modifications of the target atom, as opposed to inelastic interactions leading to transmutation, for instance. During the collision between the neutron and the atom, part of the kinetic energy can be trans- ferred to the target atom. The interaction probability is given by the elastic collision differential cross sec- tion1,2 which depends on both the neutron kinetic energy and the transferred energy.3 For a typical fast neutron of 1 MeV, the mean transferred energy ð �TÞ of the Zr atom is �T � 22keV. For low value of the transferred energy, the target atom cannot leave its position in the crystal, leading only to an increase of the atomic vibrational amplitude resulting in simple heating of the crystal. If the transferred energy is higher than a threshold value, the displacement energy (Ed), the knocked-on atom can escape from its lattice site and is called the primary knocked-on atom (PKA). For high transferred energy, as is the case for fast neutron irradiation, the PKA interacts with the other atoms of the alloy along its track. On average, at each atomic collision, half of its current kinetic energy is transferred to the collided atom, since they have equal masses. The collided atoms can then interact with other atoms, thus creating a displacement cascade within the crystal. 4.01.1.1.2 Displacement energy in zirconium In the case of zirconium, the displacement energy has been measured experimentally using electron irra- diations performed at low temperatures ( Radiation Effects in Zirconium Alloys 3 centimeters), it can be considered that only one PKA is created by the incoming neutron going through the Zr cladding used in pressurized water reactors (PWRs) (with a thickness of 0.6 mm). There- fore, if the PKA creation rate per unit volume within the cladding is known for a typical fuel assem- bly in a PWR (with typical fast neutron flux is 5� 1017 nm�2 s�1 (E> 1 MeV)), the number of dis- placed atoms per unit volume and per second can be computed. From this value, the overall number of displacements per atom (dpa) and per second can be simply computed. This calculation can be achieved, as described by Lunéville et al.,3 by taking into account the PWR neutron spectrum as well as the neutron–atom differential cross section. It can be shown that a typical damage rate for a cladding in a PWR core is between 2 and 5 dpa year�1, depending on the neutron flux history. This means that each atom of the cladding has been displaced 2–5 times per year! A more accurate correspondence between the fast fluence and the damage for a cladding in a PWR is provided by Shishov et al.12 These authors evaluate that a fluence of 6� 1024 nm�2 (E> 1 MeV) corresponds to a damage of 1 dpa. This simple approach gives a good description of the number of displaced atoms during the creation of the cascade, but does not consider intracascade elas- tic recombinations that occur during the cascade relaxation or cooling-down phase.11,13,14 In addition, this approach does not give any information on the form of the remaining damage at the end of the cascade, such as the point-defect clusters that can be created in the cascade. In order to have a better understanding of the created damage in a-zirconium, several authors have 1 0 2 4 6 8 10 12 14 2 3 4 5 6 7 8Cluster size 9 1011 12131415 25 1 1 5 5 1 0 10 20 20 0.5 0.5 (a) (b N um b er o f i nt er st iti al s p er c as ca d e PKA ene rgy (ke V) 100 K Figure 1 Number of single and clustered (a) interstitials and (b Adapted from Gao, F.; Bacon, D. J.; Howe, L. M.; So, C. B. J. N performed MD computations also using different types of interatomic potentials. It is shown that, at the end of the cascade creation ( 4 Radiation Effects in Zirconium Alloys The form of these small clusters is also of major importance since it plays a role on the nucleation of dislocation loops. Wooding et al.16 and Gao et al.8 have shown that the small SIA clusters are in the form of dislocation loops with the Burgers vector 1=3 11�20h i. The collapse of the 24-vacancy cluster to a dislocation loop on the prism plane was also found to occur. 4.01.1.2 Evolution of Point Defects in Zirconium: Long-Term Evolution After the cascade formation and relaxation, which last for a few picoseconds, the microstructure evolves over a longer time. The evolution of the microstruc- ture is driven by the bulk diffusion of point defects. For a better understanding of the microstructure evolution under irradiation, the elementary proper- ties of point defects, such as formation energy and migration energy, have first to be examined. 4.01.1.2.1 Vacancy formation and migration energies Concerning the vacancy, all the atomic positions are identical in the lattice and so there is only one vacancy description leading to a unique value for the vacancy formation energy. Due to the rather low a–b phase transformation temperature, the measurement of vacancy formation and migration energy in the Zr hexagonal close-packed (hcp) phase is difficult. The temperature that can be reached is not high enough to obtain an accurately measurable concentration and mobility of vacancies.18 Nevertheless, various experi- mental techniques (Table 1), such as positron annihila- tion spectroscopy or diffusion of radioactive isotopes, have been used in order to measure the vacancy formation and migration energies or the self-diffusion Table 1 Experimental determination formation (Ef), migration (Em) and self diffusion activation (Ea) energies for vacancy (in eV) Experimental methods Ef Em Ea Reference Semiempirical 1.8–1.9 1.3–1.6 3.3 [18] Self-diffusion – – 1.2–3.5 [18] Diffusion behavior of various solutes in Zr 1.4–2.1 1.1–1.5 3.2–3.5 [19] Self-diffusion – – 2.85 [20] coefficient.18–26 The values obtained by the various authors are given in Table 1. It is pointed out by Hood18 that there is great discrepancy among the vari- ous results. It is particularly shown that at high tem- perature, the self-diffusion activation energy is rather low compared to the usual self-diffusion activation energy in other metals.18 However, as the temperature decreases, the self-diffusion activation energy increases strongly. According to Hood,18 this phenomenon can be explained assuming that at high temperature the vacancy mobility is enhanced by some impurity such as an ultrafast species like iron. At lower tempera- ture, the iron atoms are believed to form small pre- cipitates, explaining that at low temperatures the measured self-diffusion energy is coherent with usual intrinsic self-diffusion of hcp crystals. It is also shown that the self-diffusion anisotropy remains low for normal-purity zirconium, with a slightly higher mobil- ity in the basal plane than along the hci axis.22,26,27 For high-purity zirconium, with a very low iron con- tent, the anisotropy is reversed, with a higher mobility along the hci axis than in the basal plane.27 The vacancy formation and migration energies have also been computed either by MD methods, where the mean displacement distance versus time allows obtaining the diffusion coefficient, or by static computation of the energy barrier corresponding to the transition between two positions of the vacancy using either empirical interatomic potential7,28–34 or the most recent ab initio tools.35–38 Since the different sites surrounding the vacancy are not similar, due to the non-ideal c/a ratio, the migration energies are expected to depend on the crystallographic direction, that is, the migration energies in the basal plane E == m and along the hci direction E?m are different. The results are given in Table 2. The atomistic calculations are in agreement with the positron annihilation spectroscopy measurement but are in disagreement with the direct measure- ments of self-diffusion in hcp zirconium.20 As dis- cussed by Hood,18 and recently modeled by several authors,39,40 this phenomenon is attributed to the enhanced diffusion due to coupling with the ultrafast diffusion of iron. 4.01.1.2.2 SIA formation and migration energies In the case of SIAs, the insertion of an additional atom in the crystal lattice leads to a great distortion of the lattice. Therefore, only a limited number of configurations are possible. The geometrical descrip- tion of all the interstitial configuration sites has been Radiation Effects in Zirconium Alloys 5 proposed for titanium by Johnson and Beeler41 and is generally adopted by the scientific community for other hcp structures (Figure 2). � T is the simplest tetrahedral site, and O is the octahedral one, with, respectively, 4 and 6 coordi- nation numbers. � BT and BO are similar sites projected to the basal plane with three nearest neighbors, but with dif- ferent numbers of second neighbors. � BC is the crowdion extended defect located in the middle of a segment linking two basal atoms. T BT BC O BS BO C S Figure 2 Interstitial sites configuration: (a) static localizations 249–265) and (b) relaxed configurations (adapted from Willaime Table 2 Computation determination formation (Ef), migration (Em), and self-diffusion activation (Ea) energies for vacancy (in eV) Computation methods Ef E == m E ? m Ea Reference Pair potential 1.59 1.21 1.10 [28] Finnis–Sinclair MB potential 1.79 0.93 0.93 [33] Finnis–Sinclair MB potential 1.79 – – – [7] Finnis–Sinclair MB potential 1.79 0.84 0.88 2.64 [34] EAM potential 1.74 0.57 0.59 2.32 [31] Ab initio FP-LMTO 2.07 – – – [30] Ab initio GGA 1.86 – – – [36,37] Ab initio GGA 2.17 0.51 0.67 2.76 [38] Ab initio LDA 2.29 0.23 0.43 2.78 [38] MB: many body; EAM: embedded atom method; FP-LMTO: full-potential linear Muffin-Tin orbital; GGA: generalized gradient approximation; LDA: local density approximation. � C is the interstitial atom located between two adjacent atoms of two adjacent basal planes in the 20�23h i direction. This direction is not a close- packed direction, and allows easier insertion of the SIA. � S is the split dumbbell position in the hci direction. The only way to have access to the SIA formation energy is from atomistic computations taking into account the different configurations of the SIA given previously. In their early work on titanium, Johnson and Beeler41 found that the most stable SIA configuration was the basal-octahedral site (BO). Several other sites were also found to be metastable, like asymmetric variants of the T and C sites. As reviewed by Willaime,35 the relative stabilities of the various SIA configurations were observed to depend strongly on the interatomic potential used (Table 3). The mobility of SIAs can be estimated experimen- tally using electron irradiation at very low tempera- tures (4.2 K), followed by a heat treatment. During the recovery, the electrical resistivity is measured. The main recovery process was found around 100–120K and analysis of the kinetics gives the SIA migration energy of Em� 0.26 eV.4 Atomistic computations have also brought results (Table 3) concerning the SIA migration energy. Sev- eral authors7,28–31,33–37 have found that the mobility of SIAs is anisotropic, with low migration activation energy for the basal plane mobility (E==m � 0.06 eV) and a higher migration activation energy in the hci direction (E?m � 0.15 eV). In the temperature range of interest for the power reactors (T� 600K), the diffusion coefficients obtained are the following: D == i ¼ 8� 10�9m2 s�1 (in the basal plane) and O S C BCBSBO (adapted from Bacon, D. J. J. Nucl. Mater. 1993, 206, , F. J. Nucl. Mater. 2003, 323, 205–212). Table 3 Computation of SIAs formation (Ef) andmigration (Em) energies in Zr by ab initio, MD, or MS (molecular statics) (in eV) Method Ef Em Reference O BO BS/BC C S T E == m E ? m Pair potential – 3.83 – 4.01 – – BO: 0.8 BO: 0.49 [28] – C: 0.49 C: 0.29 EAM potential 2.8 2.63 2.5 2.78 3.04 _ 0.05 0.14 [31] Finnis-Sinclair MB potential – 3.97 3.76 3.97 4.32 _ – – [7] Finnis–Sinclair MB potential – – – – – – 0.06 0.15 [33] Ab initio GGA 2.84 2.88 2.95 3.08 3.01 4.03 – – [36,37] Ab initio LDA 2.79 2.78 2.90 3.07 2.80 – – – [35] Ab initio GGA 3.04 3.14 3.39 3.52 3.28 – – – [35] Finnis–Sinclair 4.13 3.97 3.75 3.96 3.77 3.98 – – [34] MB potential MB: many body; EAM: embedded atom method; FP-LMTO: full-potential linear Muffin-Tin orbital; GGA: generalized gradient approximation; LDA: local density approximation. 6 Radiation Effects in Zirconium Alloys D?i ¼ 10�9 mm2 s�1 (along the hci direction). These authors have also shown that the anisotropy depends on the temperature. Computing the effective diffu- sion rate of SIAs in all directions, taking into account the multiplicity of the jump configurations for each type of migration, Woo and co-workers34,42 have obtained the anisotropy for self-interstitial diffusion as a function of temperature. It is shown that the SIA mobility is higher in the basal plane than along the hci axis and that the anisotropy decreases when the temperature increases. 4.01.1.2.3 Evolution of point defects: Impact of the anisotropic diffusion of SIAs In zirconium alloys, as in other metals, under irradia- tion both vacancies and SIAs (Frenkel pairs) are created within the cascade leading to an increase of the point-defect concentration with the irradiation dose. However, even at very low temperature, the Frenkel pair concentration saturates at values about 1% due to the mutual recombination of vacancies and SIAs.43 At higher temperatures, point defects migrate and can therefore disappear because of a large variety of defects/defects reactions. Three major mechanisms contribute to defect elimination: vacancy–SIA recombination, point-defect elimina- tion on defect sinks (dislocation, grain boundaries, free surface, etc.), and agglomeration in the form of vacancy dislocation loops and interstitial dislocation loops. It has to be noted that, because of the rapid migration of SIAs compared to the slow migration of vacancies, at steady state the vacancy concentration is several orders of magnitude higher than the SIA concentration. Because of the elimination of point defects on point-defect clusters, the clusters can grow under irradiation depending on their relative capture effi- ciency. In the case of cubic metals, since the relaxa- tion volume of SIAs is usually much larger than that of vacancies, edge dislocations eliminate SIAs with a higher efficiency than vacancies (positive bias toward SIAs). Assuming an isotropic diffusion of point defects, this phenomenon leads to a preferred absorp- tion of SIAs by dislocations, provided that there is another type of sink within the material. Because of this preferential absorption of SIAs, the intersti- tial loops tend to grow under irradiation and the vacancy loops tend to shrink. However, in hcp zirconium, the point-defect diffusion is usually considered to be anisotropic although there is little experimental evidence of this phenomenon. From the experimental results, it is believed that vacancy migration is only slightly anisotropic but the SIA migration is believed to be significantly anisotropic, as shown by atomistic com- putations. This diffusional anisotropy difference (DAD) has a strong impact on capture efficiency of point defects by sinks.44 Indeed, assuming SIAs to have a higher mobility in the basal plane than along the hci axis and that the vacancies have an isotropic diffusional behavior, it can be seen that grain bound- aries perpendicular to the basal plane absorb more SIAs than vacancies. On the other hand, grain bound- aries parallel to the basal plane absorb more vacancies (a) (b) 0.5mm Figure 3 hai dislocation loops obtained in EBR-II at 700 K: (a) 1.1�1025 n m�2 and (b) 1.5�1026 n m�2. Diffracting vector g¼10�11 and beam direction B¼ 0�111� � Griffiths, M. J. Nucl. Mater. 1988, 159, 190–218. Radiation Effects in Zirconium Alloys 7 than SIAs. Similarly, a line dislocation parallel to the hci axis absorbs more SIAs than vacancies and a line dislocation in the basal plane absorbs more vacancies than SIAs. As discussed by Woo,44 this geometrical effect due to the DAD can overwhelm the conven- tional bias caused by the point-defect/sink elastic interaction difference (EID). Thus, contrary to the implications of the conventional rate theory, edge dislocations in a-zirconium are not necessarily biased toward SIAs, and grain boundaries are no longer neutral sinks. As will be described in the following, this phenomenon can explain some anomalous irra- diation-induced microstructural features as well as the growth phenomenon of zirconium alloys. 50 nm Figure 4 Typical hai loop microstructure observed on recrystallized Zy-4 irradiated at 280 �C in Siloé up to a fluence of 6�1024 n m�2. + + + + + + + + 1019 1020 Neutron fluence (cm–2) Not visible (b) (a) 4 5 6 7 0 2 N (1 02 2 m –3 ) 4 1021 1022 d (n m ) Figure 5 Evolution with dose of the dislocation loops characteristics: (a) density and (b) mean size of defects for Zy-2 irradiated at 300 �C. Adapted from Northwood, D. O. Atomic Energy Rev. 1977, 15, 547–610. 4.01.1.3 Point-Defect Clusters in Zirconium Alloys In the case of zirconium alloys, many authors have studied the postirradiation microstructure by using transmission electron microscopy (TEM). In 1979, an international ‘round robin’ was undertaken consisting of TEM observations of neutron-irradiated recrys- tallized zirconium alloys45 in order to determine the nature of the point-defect clusters. A more recent compilation of observations is given by Griffiths.46 It has been now proved by numerous authors that in zirconium alloys mainly dislocation loops with hai Burgers vector can be found. Only for high fluence, the hci component dislocation loops appear. Cavities are observed only in very specific cases. 4.01.1.3.1 hai Dislocation loops It is now clearly established by numerous authors45–57 that for commercial neutron-irradiated zirconium alloys (e.g., annealed Zircaloy-2 described in Northwood et al.45) at temperatures between 250 and 400 �C and for irradiation dose lower than 5� 1025 nm�2, the point-defect clusters that can be observed by TEM (>2 nm) consist of perfect dislocation loops, either of vacancy or interstitial nature, with Burgers vector ah i ¼ 1=3 11�20h i, situated in the prismatic planes with typical diameter from 5 to 20 nm, depending on the irradiation temperature (Figures 3 and 4). These loops are found in very high density, typically between 5� 1021 and 5� 1022 m�3 depending on the irradiation temperature (Figure 5).45,51 The three hai Burgers vectors are equally represented. Thorough studies of neutron damage in zirconium using the high-voltage electron microscope (HVEM) have also been given.53,58,59 8 Radiation Effects in Zirconium Alloys The proportion of vacancy loops to interstitial loops depends on the irradiation temperature. Indeed, it is observed that for an irradiation temper- ature of 350 �C approximately 50% of observed loops are vacancy loops, whereas for an irradiation temper- ature of 400 �C, 70% of loops are vacancy loops.45,46 For a low irradiation temperature (below 300 �C), the majority of loops present in the material are of the interstitial type. The loop habit plane is close to the prismatic plane, but accurate determination proves that the loops are not pure edge but their habit plane is usually closer to the first-order prismatic plane 10�10f g. The authors have also observed that for loop diameters lower than 40 nm the loops are circu- lar but for diameters larger than 40 nm the vacancy loops become elliptical with the great axis along the hci axis, the interstitial loops remaining circular. The hai loops also appear to be aligned in rows parallel to the trace of the basal plane.46,50 For an irradiation temperature of 300 �C, no dislocation loop can be observed below a neutron fluence of 3� 1023 nm�2 in the case of annealed Zy-2 (Zircaloy-2) irradiated at 300 �C.51 However, from this fluence, the loop density increases rapidly with increasing fluence but saturates at a density of 3� 1022 m�3, from a relatively low fluence of approx- imately 1� 1024 nm�2 (Figure 5). The loop density saturation has been confirmed by X-ray analysis.60 The loop size exhibits a parabolic increase with fluence but no clear saturation in the evolution of the loop size is seen even after a fluence of 1� 1026 nm�2.51,67 Increasing the irradiation temperature leads to a decrease in the loop density and to an increase of the loop size.45,55,61 Indeed, it was shown by Northwood et al.45 that neutron irradiation performed at 350 �C of annealed Zy-2 up to a fluence of 1� 1025 nm�2 leads to a mean loop diameter between 8 and 10 nm and a loop density between 8� 1021 and 5� 1022 m�3; whereas a neutron irradiation of the same alloy per- formedat400 �Cuptoa fluenceof1� 1025nm�2 leads toamean loopdiameterbetween16 to23nmanda loop density between 4� 1021 and 2� 1022 m�3.45 Above 500 �C, no irradiation damage is formed.52 The hai loop microstructure is found to be very sensitive to alloying elements such as oxygen. Indeed, for high- purity zirconium with very low oxygen content, the hai loops are large and in low density, whereas for commercial zirconium alloys (with oxygen content between 1000 and 1500 ppm) the growth speed of loops is considerably reduced yielding smaller loops in much higher density.45,55 It was also reported from TEM observations that a particular band contrast of alternative black and white was superimposed on the usual radiation damage nor- mally visible on thin foils of irradiated materials. This phenomenon has been connected to the alignment of the loops in the same direction and is believed to be a thin-foil artifact. It has been named ‘corduroy’ contrast by Bell.62 The commonly accepted explanation of this artefact is based on the local elastic relaxation of the internal stresses in TEM thin foils, in areas where pronounced alignment of hai loops is present.63 4.01.1.3.2 hai Loop formation: Mechanisms The origin for the stability of the hai loops in zirco- nium is attributed to the relative packing density of the prismatic plane compared to the basal plane, which depends on the c/a ratio of the hcp lattice. Foll and Wilkens64 have proposed that when the c/a ratio is higher than ffiffiffi 3 p , loops are formed in the basal plane with Burgers vector 1=6 20�23h i, whereas if c/a is lower than ffiffiffi 3 p , then loops are formed in the prismatic plane with Burgers vector ah i ¼ 1=3 11�20h i. For all hcp metals, this means that loops are formed in the prismatic plane except for Zn and Cd. This is not the case for Zr, Ti, and Mg where loops are also formed in the basal planes, depending on the irradiation dose, irradiation temperature, and purity of the metal.56,57 MD computations for a-zirconium have also shown that most of the small interstitial clusters produced in the cascade have the form of a dislocation loop with Burgers vector ah i ¼ 1=3 11�20h i. The small vacancy clusters are also found in the prismatic plane.8,28,65 For larger point-defect clusters,66 it is shown that the point-defect clusters in the prismatic plane always relax to perfect dislocation loops with Burgers vector ah i ¼ 1=3 11�20h i. On the other hand, vacancy clusters in the basal plane form a hexagonal loop enclosing a stacking fault with 1=2 0001h i Burgers vector. The simultaneous observation of vacancy and interstitial hai loops in zirconium alloys45,48,50,54,61 is a rather surprising feature.53,57 Indeed, as discussed for usual cubic metals, interstitial loops tend to grow under irradiation and the vacancy loops tend to shrink since the edge dislocations are biased toward SIAs due to the EID. According to Griffiths,57 the coexistence of these two types of loops in zirconium can be explained by a modified SIA bias in zirconium due to (i) a rela- tively small relaxation volume of SIA relative to vacancy (low bias), (ii) interaction with impurities, and (iii) spatial partitioning of vacancy loops and interstitial loops as a result of elastic interactions or Radiation Effects in Zirconium Alloys 9 anisotropic diffusion. Other authors53,68 think that this phenomenon is due to a subtle balance of the bias factors of the neighboring point-defect sinks that lead to an increasing bias as the loop size increases if the loop density is high. Woo44 considers that the coexistence of both types of hai loops can be explained in the frame of the DAD model, which induces a strong DAD-induced bias. Indeed, in this model, the hai type loops are shown to be relatively neutral and may therefore receive a net flow of either interstitials or vacancies, depending on the sink situ- ation in their neighborhood. Finally, recent computations,69 using the Monte Carlo method, that take into account the large vacancy and interstitial point-defect clusters created inside the cascade as an input microstructure show that both vacancy and interstitial loops are able to grow simultaneously, the proportion of vacancy loops increasing with increasing irradiation temperature. This last phenomenon can be related to the so-called production bias discussed previously.14 4.01.1.3.3 hci Component dislocation loops At the time of the thorough review by Northwood,51 no hci component loops had been observed yet. The ‘round robin’ work45 also established that up to an irradiation fluence of 1� 1025 nm�2 no hci component dislocation loop is observed. As highly irradiated Zircaloy samples became available, for fluence higher than 5� 1025 nm�2, evidence of hci component loops arose.46,54,70–73,189 The hci component loops have been analyzed as being faulted and of the vacancy type.They are located in the basal plane with a Burgers vector (a) Figure 6 Comparison of neutron damage in Zr at 700 K followi purity (500 wt ppm) with no c-component loops. (b) Sponge pur edge-on orientation (arrowed). Only hci component defects are is [10�10] for each micrograph. Adapted from Griffiths, M. Philos 1=6 20�23h i having a component parallel to the hci axis (Figure 6). The hci component loops are much larger than the hai loops but their density is much lower. For instance, for recrystallized Zy-2 and Zy-4 irradiated at 300 �C, after 5.4� 1025 nm�2, hci component loops are found with a diameter of 120 nm and with a density between 3 and 6� 1020 m�3. Whatever the irradiation conditions, these hci component loops are always present in conjunction with more numerous and finer hai loops. The hci component loops can therefore only be observed edge-on by TEM by using the g¼ 0002 diffraction vector, which leads to invisible hai type defects. The hci loops thus appear as straight-line segments. There is considerable evidence to show that their formation is dependent on the purity of the zirconium used (Figure 6).46,74–76,190 It is also observed that at the beginning of their formation, these dislocation loops appear to be located close to the intermetallic precipi- tates present in the Zircaloy samples46,76 (Figure 7). By using an HVEM on iron-doped samples, it has been possible to prove that iron enhances the nucle- ation of the hci loops, the loop density increasing as a function of the iron content. Moreover, ironwas found to have segregated in the plane of the loops.76 4.01.1.3.4 hci Loop formation: Mechanisms It is rather surprising that although the most stable loops are the prismatic loops, basal loops are also observed in zirconium alloys. Moreover, these loops are of the vacancy character. According to the usual rate theory, vacancy loops should not grow as a result of the bias of edge dislocation toward SIAs. 0.5mm(b) ng irradiation to a fluence of 1.5� 1026 n m�2. (a) Crystal bar ity (2000 wt ppm) containing basal hci component in an visible with diffracting vector of [0002]. The beam direction . Mag. B. 1991, 63(5), 835–847. 500 nm Figure 7 High density of c-component loops in the vicinity of the precipitates in a Zy-4 sample irradiated to 6�1025 nm�2; at 585 K. The arrow shows the diffracting vector [0002]. Adapted from De Carlan, Y.; Régnard, C.; Griffiths, M.; Gilbon, D. Influence of iron in the nucleation of hci component dislocation loops in irradiated zircaloy-4. In Eleventh International Symposium on Zirconium in the Nuclear Industry, 1996; Bradley, E. R., Sabol, G. P., Eds.; pp 638–653, ASTM STP 1295. 10 Radiation Effects in Zirconium Alloys The reason for the nucleation and growth of the hci component loops in zirconium alloys has been ana- lyzed and discussed in great detail by Griffiths and co-workers.46,56,57,74 The most likely explanation for their appearance46 is that they nucleate in collision cascades, as shown recently by De Diego.66 Their stability is dependent to a large extent on the pres- ence of solute elements, which probably lower the stacking-fault energy of the Zr lattice, making the basal hci component loops more energetically stable. It is also possible that small impurity clusters, espe- cially iron in the form of small basal platelets, could act as nucleation sites for these loops.74,76 However, according to Griffiths,46 this cannot account for the very large vacancy hci component loops observed, since the growth of vacancy loops is not favorable considering the EID discussed previously. In order to understand the reason for the important growth of the hci component loops, another mechanism must occur. As discussed by Woo,44 the growth of hci component loops is well understood in the frame of the DADmodel. Indeed, because of the higher mobil- ity of SIAs in the basal plane rather than along the hci axis (and the isotropic diffusion of vacancies), dislo- cations parallel to the hci axis will absorb a net flux of SIAs whereas dislocations in the basal plane will absorb a net flux of vacancies. This can therefore explain why the basal vacancy loops can grow. The incubation period before the appearance of hci component loops can be explained, according to Griffiths et al.,73 by the fact that the hci loop formation is dependent on the volume of the matrix containing a critical interstitial solute concentration. This volume increases as the interstitial impurity concentration is gradually supplemented by the radiation-induced dis- solution of elements such as iron from intermetallic precipitates (or b-phase in the case of Zr–Nb alloys). 4.01.1.3.5 Void formation Early studies failed to showany cavity in Zr alloys after irradiation.77 From all the obtained data, it is seen that zirconium is extremely resistant to void formation during neutron irradiation (Figure 8).46,52 The effect of very low production of helium by (n, a) reactions during irradiation was mentioned as a possible reason for this absence of voids. But most probably, the fact that in zirconium alloys vacancy type loops are easily formed can be the reason for the absence of void.52 To favor the formation of voids, various studies per- formed, especially on model alloys, have shown that stabilization of voids can occur when impurities are present in the metal. Helium coming from transmu- tation of boron on Zr sponge67 as well as impurities located near Fe-enriched intermetallics are found to favor the stability of voids.54 Irradiations with elec- trons give better conditions to stabilize voids: the main reason is that irradiation doses can be very high – hundreds of displacements per atom can be reached after few hours.190 Moreover, electron irradi- ation on Zr samples preimplanted with He at various concentrations showed the nucleation and growth of voids only for the samples doped with at least 100 ppm of He.78 4.01.1.4 Secondary-Phase Evolution Under Irradiation 4.01.1.4.1 Crystalline to amorphous transformation of Zr-(Fe,Cr,Ni) intermetallic precipitates In addition to point-defect cluster formation, irra- diation of metals can affect the precipitation state as well as the solid solution. In the case of zirco- nium alloys, while investigating the effect of irradi- ation on corrosion, TEM observations revealed that for Zircaloy, irradiated at temperatures typical for commercial light water reactors (lower than 600 K), Zr(Fe,Cr)2 precipitates began to become amorphous after a fluence of about 3� 1025 nm�2. Interestingly, the other common precipitate in Zy-2, Zr2(Fe, Ni), remained crystalline up to higher irra- diation doses.77 The instability of these precipitates under irradiation is of great importance since the secondary-phase precipitate plays a major role on 0.1mm (a) (b) (c) (d) Figure 8 Examples of radiation-induced cavities in zirconium alloys. (a) Annealed crystal-bar zirconium, prism foil, 673K, 1.2�1025n/m2; (b) annealed zircaloy-2, prism foil, 673K, 1.2�1025n/m2; (c) annealed Zr-2.5 wt% Nb, basal foil, 923K, 0.7�1025n/m2; (d) typical cavity attached to inclusion on a grain boundary, material (c). Adapted from Gilbert, R. W.; Farrell, K.; Coleman, C. E. J. Nucl. Mater. 1979, 84(1–2), 137–148. (b) (a) 0.1mm Figure 9 Crystalline to amorphous transformations of Zr (Cr, Fe)2 particle in Zy-4 irradiated in a BWR at 560 K: (a) 3.5�1025 n m�2 and (b) 8.5�1025 n m�2. Adapted from Griffiths, M.; Gilbert, R. W.; Carpenter, G. J. C. J. Nucl. Mater. 1987, 150(1), 53–66. Radiation Effects in Zirconium Alloys 11 the corrosion resistance of Zircaloy (see Chapter 5.03, Corrosion of Zirconium Alloys). The effect of temperature on the crystalline to amorphous transformation has been studied by vari- ous authors.75,79–83 It is shown that at low tempera- tures (353 K), under neutron irradiation, both Zr(Fe, Cr)2 and Zr2(Fe, Ni) undergo a rapid and complete crystalline to amorphous transformation. As the irra- diation temperature increases, a higher dose is required for amorphization. It is indeed seen that, at 570 K, Zr(Fe,Cr)2 precipitates undergo only a partial amorphous transformation and Zr2(Fe,Ni) particles remain crystalline (Figure 9). It is also observed that the crystalline to amor- phous transformation starts at the periphery of par- ticles, and then the amorphous rim moves inward until the whole precipitate becomes fully amor- phous. The chemical concentration profile within the precipitates also exhibits two distinct zones corresponding to the two different states: the crystal- line core and the amorphous periphery. It is observed that the amorphous layer exhibits a much lower iron 500 400 300 200 100 0 500 400 300 200 100 0 350 450 550 Energy (eV)�10–1 Energy (eV)�10–1 650 750 850 350 450 550 650 750 850 CrKa CrKa FeKa FeKa C ou nt s � 10 1 C ou nt s � 10 1 0.1mm Figure 10 Crystalline to amorphous transformations of Zr(Cr, Fe)2 particle in Zy-4 irradiated at 560 K at 3.5�1025 n m�2. EDX spectrum shows that the amorphous volume is coincident with a depletion of Fe. Adapted from Griffiths, M.; Gilbert, R. W.; Carpenter, G. J. C. J. Nucl. Mater. 1987, 150(1), 53–66. 12 Radiation Effects in Zirconium Alloys content than the precipitate, the iron profile showing a local drop from the standard value of 45 at.% to below 10 at.% (Figure 10). At higher temperatures (T> 640 K), amorphiza- tion was not detected and the precipitates remain crystalline, but some authors79 have nevertheless observed loss of iron and even total dissolution of Zr2(Fe, Ni) and Zr(Fe, Cr)2 precipitates and redistri- bution of alloying elements. The crystalline to amorphous transformation is eas- ily understood in terms of ballistic radiation-induced disordering at a temperature where recombination of point defects or recrystallization within the interme- tallic precipitate is too slow to compensate for the rate of atomic displacement (at 350K).79 The dissolution of alloying elements remains limited at this low tempera- ture and the amorphization is mainly due to sputtering, that is, transfer of material from the particle because of atomic displacements by neutrons. When the point- defect concentration becomes too high and/or when the chemical disordering is too high, the crystalline structure is destabilized and undergoes a transforma- tion to an amorphous phase.75,79 The fact that the Zr2(Fe, Ni) phase remains crys- talline at intermediate temperatures (520–600 K) is presumably due to a more rapid reordering than the disordering in this structure (Zintl phase structure). Concerning the Zr(Fe, Cr)2 (Laves phase structure), it is seen that the amorphization starts at the precipi- tate–matrix interface forming a front that gradually moves into the precipitate. The amorphization is believed to happen by a deviation from stoichiometry due to a ballistic interchange of iron and zirconium atoms across the precipitate–matrix interface. It also agrees with the observed kinetics of amorphization, predicting an amorphous thickness proportional to fluence and the absence of an incubation period for the transformation to start.84 The reason for the depletion of iron from the precipitates is not clearly understood yet, according to Griffiths et al.79 It is suggested that iron may be in some form of irradiation-induced interstitial state in irradiated Zr-alloys and may then diffuse intersti- tially out of the intermetallic particles. At high temperatures (640–710 K), corresponding to 0.3Tm, the thermal activation is sufficient to induce dynamic recrystallization impeding the amorphiza- tion of the precipitates. However, depletion and some precipitate dissolution would still occur, but the level of damage necessary for amorphization would not be reached due to the absence of cascade damage.84 Because of the high mobility of Fe and Cr, redistribution of solute can occur, leading to secondary-precipitate formation. Radiation Effects in Zirconium Alloys 13 4.01.1.4.2 Irradiation effects in Zr–Nb alloys: Enhanced precipitation In binary Zr–Nb alloys (Zr–1% Nb and Zr–2.5% Nb), the microstructure is usually in a metastable state due to the thermomechanical processing in the upper a range or in the aþ b domain. Indeed, at this relatively low temperature (around 580 �C), the atomic mobility is low and the equilibrium state cannot be reached in reasonable time. After cooling, the matrix is therefore supersaturated in Nb and the composition of secondary phases (Nb rich) still corresponds to the high-temperature chemical com- position. It is indeed shown by Toffolon-Masclet et al.85 that a Zr–1% Nb–O alloy that has undergone a final heat treatment at 580 �C for a few hours can still evolve toward its thermodynamic equilibrium after 10000 h of heat treatment at 400 �C. Under irradiation, it is observed that the micro- structure of Zr–Nb alloys is not stable and very fine Nb-rich precipitates, with diameter of a few nanometers, are observed in very high density (Figure 11). This precipitation of Nb from the super- saturated matrix is observed in any type of binary alloys: in Zr–1% Nb such as M5™(86) and E110(12,87) as well as Zr–2.5% Nb.88 This needle-like precipita- tion has been studied mainly by TEM, and also by small angle neutron scattering (SANS) analyses.86 100 n (a) (b) (e)(d) Figure 11 Micrographs of needle-like radiation-enhanced pre 2.8�1025 n m�2, (c) M5™ 3.6�1025 n m�2, (d) Zr–1% NbO 5.7 (f) M5™ 13.1� 1025 n m�2. Reprinted, with permission, from J. Drive, West Conshohocken, PA 19428. Simultaneously, a noticeable decrease of Nb content in the matrix occurs.89 This precipitation is due to an enhanced mobility of Nb atoms under irradiation due to the very high vacancy concentration created by irradiation. This enhances the Nb mobility and allows the rapid evolu- tion of the microstructure toward its thermodynamic equilibrium, leading to precipitation of very fine Nb- rich precipitates in Zr–Nb binary alloys. In Zr–Nb alloys, the Nb-rich phases also undergo chemical changes under irradiation. Indeed, it is shown that the o phase, obtained in Zr–2.5Nb by transformation of the b-Nb after extrusion, disap- pears and transforms into b-Nb.60 For the b-Nb phase and in the case of M5™ alloys, an evolution of the chemical composition under irradiation has also been observed, but the b-Nb precipitates still remain fully crystalline even after six PWR cycles of irradiation (70 GWd t�1). Only a decrease in Nb content with a small increase in the size of the precipitates has been noticed86 (Figure 11). The same has been obtained for E110 and E635 Russians alloys, where b-Nb precipitates are altered in com- position to reduce the Nb content from 85–90% to 50%.12 Moreover, for the Zr(Nb, Fe)2 Laves phases with hcp structure found in E635 and E110 alloys, it seems that a release of iron atoms into the matrix from the m (c) (f) cipitation: (a) M5™ 2.1� 1025 n m�2, (b) Zr–1% NbO � 1025 n m�2, (e) Zr–1% NbO 8.2� 1025 n m�2, and ASTM Int., copyright ASTM International, 100 Barr Harbor 14 Radiation Effects in Zirconium Alloys precipitates has occurred after irradiation, leading to the transformation into b-Nb particles with bcc structure.12,89 4.01.2 Postirradiation Mechanical Behavior 4.01.2.1 Mechanical Behavior During Tensile Testing 4.01.2.1.1 Irradiation hardening: Macroscopic behavior As for many other metals, zirconium alloys exhibit strong hardening after neutron irradiation. It is indeed observed by numerous authors90–99 and reviewed21,77,100 that the yield stress (YS), as well as the ultimate tensile strength (UTS), of both recrystallization-annealed (RXA) and stress-relieved annealed (SRA) zirconium alloys is strongly increased by neutron irradiation (Figures 12 and 13). Micro- hardness tests also prove this phenomenon.101–105 The irradiation-induced hardening increases rapidly for fluences below 1� 1024nm�2 (E> 1MeV), at irra- diation temperatures between 320 and 360 �C, but saturates above 1� 1024 nm�2 (E> 1 MeV) and little change occurs from 1� 1024 up to 1.5� 1025 nm�2 (E> 1 MeV).92 It is however to be noticed that some authors do not find a clear saturation of the irradiation-induced hardening for fluences up to 1.5� 1025 nm�2 and irradiation temperatures between 320 and 360 �C.92,97 Although the YS (and UTS) of SRA Zr alloys is significantly higher than the YS of Strain rate 2.5% min–1 0.25 0.5 0.025 0.025 Irradiated (~3 � 1024n m–2) Unirradiated 0 0 200 400 600 800 1000 0.02 0.04 True strain (mm mm–1) 0.06 0.08 0.1 Tr ue s tr es s (M P a) Figure 12 Stress–strain curves indicating the effect of irradiation and strain rate of RXA Zy-2 measured during uniaxial tensile test at 616 K. Reprinted, with permission, from Seventh International Symposium on Zirconium in the Nuclear Industry, Strasbourg, France, June 24–27, 1985, copyright ASTM International, 100 Barr Harbor Drive, West Conshohocken, PA 19428. RXA Zr alloys before irradiation, the YS of both alloys, measured after high irradiation doses, at saturation, become close.21,90,100 According to Higgy and Hammad,92 and reviewed by Douglass,21 as the irradiation temperature increases from temperatures below 100 �C up to temperatures between 320 and 360 �C, the irradiation-induced hard- ening decreases. According to these authors, this shows that the accumulation of damage decreases as the irra- diation temperature increases, presumably due to recovery during irradiation. The chemical composition seems to play a second- ary role in the irradiation-induced hardening com- pared to the effect of the metallurgical state (SRA vs. RXA). The oxygen content is nevertheless shown to have a slight effect on the irradiation-induced harden- ing. Indeed, Adamson and Bell101 have shown using microhardness tests that the irradiation-induced hard- ening is higher for RXA Zy-2 alloy with high oxygen content (1800 ppm) than in the case of an RXA Zy-2 alloy with low oxygen content (180 ppm). It can also be noticed that the test temperature seems to have only a small influence on the irradiation- induced hardening, for a given irradiation temperature, up to a test temperature of 400 �C. Indeed, as reported by Onchi et al.96 (Figure 14), the YS of both irradiated and unirradiated RXA Zy-2 decreases with the test temperature, the decrease being only slightly lower for the irradiated specimens between 20 and 300 �C. However, beyond a test temperature of 400 �C, a strong decrease of the irradiation hardening occurs due to the recovery of the irradiation damage. 4.01.2.1.2 Irradiation hardening: Mechanisms It is widely agreed77,100 that the irradiation-induced hardening in zirconium alloys results mainly, as for many other metals, in the creation of a high density of small point-defect clusters that act as obstacles for dislocation glide. As described earlier, the point- defect clusters in zirconium alloys consist mainly of small prismatic loops, with Burgers vector lying in the hai direction and the habit plane close to the prismatic plane of the hcp crystal lattice. Several authors have discussed that dislocations interact with irradiation-induced dislocation loops through their long-range stress field106,107 and also through contact interactions, which can lead to junction creation that are strong obstacles to dislocation motion.108–110 Several authors have investigated in more detail the junction formation between dis- locations and loops in zirconium alloys. Particularly, Carpenter111 has considered the mechanism YS Uniform elongation Fast fluence (E > 1 MeV) UTS 80 60 40 S tr en gt h (k g m m –2 ) E lo n g at io n 20 0.2 � 1021 0.4 � 1021 0.6 � 1021 0.8 � 1021 1.0 � 1021 1.2 � 1021 1.4 � 1021 4 8 12 16 20 24 1.6 � 1021 Figure 13 The effect of fast fluence (given in n cm�2, E>1 MeV) on the room temperature tensile properties of RXA Zircaloy-4 for irradiation temperature between 320 and 360 �C. Adapted from Higgy, H. R.; Hammad, F. H. J. Nucl. Mater. 1972, 44, 215–227. 800 700 600 500 400 300 200 100 300 400 500 600 Test temperature (K) 700 Proportional limit, yield and ultimate tensile stress 3.2 � 1019nvt (0.1% offset) (lower yield point)Unirrad spl s p l, s y a nd s U TS (M P a) sy sUTS Figure 14 Proportional limit, yield, and ultimate tensile stress as a function of temperature for unirradiated and irradiated annealed (RXA) Zircaloy-2, tested at a strain rate of 1.1�10�4 s�1. Adapted from Onchi, T.; Kayano, H.; Higashiguchi, Y. J. Nucl. Mater. 1980, 88(2–3), 226–235. Radiation Effects in Zirconium Alloys 15 proposed by Foreman and Sharp109 and he applied it to the prismatic glide in zirconium alloys. He has shown that an edge dislocation gliding in the pris- matic plane that is pinned by a loop can annihilate the loop. More recently, it has been discussed that the junctions between the loops and the dislocations gliding in the basal plane are always glissile, whereas they are sessile when the dislocations glide in the prismatic plane.112,113 This phenomenon could then lead to a lower hardening of the basal slip system compared to the other slip systems. Lately, MD com- putations114 have been undertaken in order to gain a better understanding of the interaction mechanisms between dislocation and loops in zirconium alloys. It is shown that all the slip systems are not affected in the same way by the presence of the hai type loops, the basal slip system being less hardened than the prismatic slip system, for instance. 700 nm Figure 15 Propagating basal channels observed after tensile testing at 350 �C. Adapted from Onimus, F.; Monnet, I.; Béchade, J. L.; Prioul, C.; Pilvin, P. J. Nucl. Mater. 2004, 328, 165–179. 16 Radiation Effects in Zirconium Alloys 4.01.2.1.3 Post-yield deformation: Macroscopic behavior Concerning the mechanical behavior beyond the YS, it is pointed out by several authors97,115,116 that for RXA zirconium alloys, the strain hardening rate is higher after irradiation at the onset of plastic flow but decreases rapidly with the plastic strain, more rapidly than before irradiation, resulting in a low strain hard- ening capability, and therefore in little difference between YS and UTS.21 This strong decrease of the strain hardening rate is believed to be the cause of the early localization of the plastic strain at the specimen scale, observed particularly in RXA zirconium alloys, which leads to a strong decrease of the uniform elon- gation, as reported by numerous authors.92–94,96–98,117 Several authors112,118–120 have shown that, for RXA zirconium alloys, this apparent or macroscopic loss of ductility is related to the early localization of the plastic strain inside shear bands, the failure mode remaining ductile with dimples.97,112,117,121,122 The material does not become brittle considering the frac- ture mode but localizes all the plastic strain in a limited part of the specimen, which leads, at the specimen scale, to a very low, uniform elongation (Figures 12 and 13). As the irradiation-induced hardening increases with the fluence, the uniform elongation decreases rap- idlywith the fluence from 10% tovalues lower than 1% for RXA alloys at 350 �C, and saturates from a fluence of 5� 1024 nm�2.92 As for the irradiation-induced hardening, the SRA and RXA zirconium alloys exhibit similar uniform elongation at saturation.100 Some authors96,117 suggest that there is a minimum of uniform elongation for RXA zirconium alloys for testing tem- peratures between 300 and 400 �C. This loss of ductility could be due to an additional hardening that can occur in this temperature range because of the trapping of oxygen atoms by the loops,117 as already observed using microhardness tests.101 For testing temperatures above 400 �C, the ductility is progressively recovered as shown by Garde.117 4.01.2.1.4 Post-yield deformation: Mechanisms Several authors96,112,113,119–121,123–125 have studied the deformation mechanisms using TEM by taking thin foils out of the specimens after testing. They have observed that, as for many other irradiated metals, after testing, numerous cleared bands free of irradiation defects are present in the material (Figure 15). These cleared bands are the consequence of the dislocation channeling mechanism reviewed in detail byHirsch,110Wechsler,126 and Luft.127 According to several authors,128–130 the irradiation-induced loops, which are obstacles to dislocation glide, can be overcome by dislocations when a sufficient stress is applied, the loops being subsequently annihilated or dragged by dislocations following different pos- sible mechanisms.108–110,131,132 This process of removal of irradiation loops by moving dislocations produces a cleared zone free of defects inside the grain. These obstacle-free channels or swaths will therefore constitute preferred areas for further dis- location gliding, leading to plastic strain localiza- tion at the grain scale with regions of very high local plastic strain surrounded by regions of almost zero plastic strain. According to Williams et al.118 and Adamson et al.,119 the local plastic strain could reach up to 100% inside these bands. Some dis- agreement on the activated slip systems seems to remain in the case of zirconium alloys. Indeed, some authors have observed channels along the prismatic planes101,119 for tests performed at 250 and 327 �C on a Zircaloy-2 containing 1500 ppm oxygen, whereas more recently other authors113,124,125 have observed channels along the basal plane as well as along the prismatic plane depending on the loading conditions. This discrepancy could probably be explained by the differences in the texture or test temperature used by the different authors. Nevertheless, it is now clearly proved113,124 that for materials with texture characteris- tic of RXA tubing or rolled sheets, with hci axes ori- ented in the (r, y) plane with an angle between 20� to 45� to the radial (r) direction, and for internal pressure tests or transverse tensile test performed at 350 �C, only Radiation Effects in Zirconium Alloys 17 basal channels are observed for low plastic strain level. Therefore, most of the plastic strain is believed to occur by basal slip inside the channels. However, it is shown that, for an axial tensile test, basal slip is not active because of its very poor orientation and only prismatic and maybe pyramidal channels can be observed. The fact that the basal slip becomes the easy glide slip system at 350 �C after irradiation constitutes a major change in the deformation mechanisms since, before irradiation, for the same test temperature it is the prismatic slip system that is the easy glide slip system. This change in the deformation mechanisms can be explained by the difference in the interaction between the irradiation-induced loops and the dis- locations gliding either in the basal plane or in the prismatic plane, as pointed out previously. Indeed, the junction created between a dislocation gliding in the basal plane and a loop is always glissile, whereas it is sessile when the dislocation is gliding in the pris- matic plane. Therefore, when the dislocation glides in the basal plane and encounters a loop, the loop can be dragged along the slip plane, leading to a progres- sive clearing of the basal channel. Since the loops are cleared by gliding dislocations inside the channels, it is usually assumed133 that within the channels a strain softening occurs. This phenomenon is believed to be the cause of the decrease of the strain-hardening rate with irradiation and thus to the early localization of the deformation at the specimen scale, explaining the dramatic decrease of the uniform elongation after irradia- tion.96,133 According to several authors,119,127 the strong texture of the rolled sheets or tubing leads to an even stronger localization of the plastic strain. Indeed, due to the texture, the hci axis of the hcp grains is along the (r, y) plane in the case of a tube. Since for internal pressure test or transverse tensile tests the channels are along the basal plane, the basal channels can easily propagate from grain to grain, as has been shown by Onimus et al.113,124 When the entire section of the specimen is crossed by disloca- tion channels, a strong necking is observed on the specimen. As was pointed out by Franklin et al.,134 the RXA alloys are more susceptible to the plastic insta- bility since the dislocation tangles that remain in SRA alloys are believed to inhibit the easy glide and the plastic flow localization. As discussed by Onimus and Béchade,135 the polycrystalline nature of the material is also believed to play an important role in the overall macroscopic response of irradiated zirconium alloys after irradiation. Indeed, the intergranular stresses that develop because of strain incompatibilities between grains can balance the local microscopic softening occurring in the dislocation channels up to the UTS. Based on various mechanical data such as Knoop hardness test136 or plane strain and plane stress tensile tests, several authors93,122 have shown that the irradia- tion decreases the plastic anisotropy of the RXA zirconium alloys. Concerning the SRA zirconium alloys, the mechanical behavior is already more iso- tropic before irradiation than RXA zirconium alloys137 and the relative decrease of the anisotropy is therefore lower.122 According to these authors,122,136 this decrease of the anisotropy of RXA zirconium alloys is due to the fact that the basal slip is more activated after irradiation than before irradiation. 4.01.2.2 Effect of Postirradiation Heat Treatment A heat treatment performed at a temperature higher than the irradiation temperature on vari- ous zirconium alloys results in a recovery of the radiation-induced hardening90,138 (Figure 16). This recovery can also be measured using microhardness tests.101,102,105,139–142 The recovery of the hardening is always associated with the recovery of the ductility and the fracture properties.138 Howe and Thomas90 have shown that in a cold- worked zirconium alloy most of the recovery occur- ring between 280 and 450 �C appears to be the annealing out of radiation damage rather than cold work. In the case of strongly cold-worked zirconium alloys such as SRA Zy-4, radiation hardening recov- ery is also observed. The hardness of the material can even become lower than the initial hardness of the SRA Zy-4(105) owing to the recovery of the disloca- tions, in addition to the recovery of the loops. Some authors,101,140,143 on the basis of various experimental results, have suggested that there is an interaction between oxygen and irradiation- produced dislocation loops, which increases the dislocation–defect barrier interaction. During the recovery, this phenomenon can lead to an additional hardening, as shown by Snowden and Veevers.140 Several authors48,101,105,141,144,145 have shown that during a heat treatment performed on a RXA zirco- nium alloy, the hai loop density strongly decreases and the loop size increases. This decrease of the obstacle density to dislocation motion has been clearly correlated to the decrease of the radiation- induced hardening.101,105 40 100 200 300 400 500 Postirradiation annealing temperature (�C) 600 700 50 60 70 80 90 100 P ile t em p er at ur e (2 80 �C ) P L YS UTS UTS ANN YS ANN PL ANN N or m al s tr es s (p si � 10 –3 ) Figure 16 Recovery curves for irradiated annealed Zy-2. PL: Proportional limit, YS: 0.2% offset yield stress, UTS: ultimate tensile strength. Adapted from Howe, L.; Thomas, W. R. J. Nucl. Mater. 1960, 2(3), 248–260. 18 Radiation Effects in Zirconium Alloys Concerning the nature of the loops, Kelly and Blake48 have studied 240 loops in a zirconium alloy sample heat-treated at 490 �C during 1 h after irradi- ation up to a fluence of 1.4� 1024 nm�2. These authors show that, although the initial microstructure is composed of both interstitial and vacancy loops in equal amount, after the heat treatment, two-thirds of the analyzed loops are vacancy loops and only one- third are interstitial loops. This implies that the interstitial loops undergo a more rapid recovery than the vacancy loops. These observations have been recently confirmed by Ribis et al.,105 who stud- ied the evolution of the proportion of the vacancy loops and interstitial loops with heat treatment for various temperatures. These authors have shown that after 960 h at 450 �C, only large vacancy loops in low density are observed. In the literature, several mechanisms are proposed in order to explain the irradiation damage recovery. The most commonly agreed mechanism is based on bulk diffusion of vacancies during the recovery and their exchange between loops of various size.105,146–148 Indeed, the smaller vacancy loops emit vacancies that diffuse toward larger vacancy loops, which absorb more vacancies than they emit, leading to a growth of the larger loops at the expense of the smaller loops. On the other hand, interstitial loops always absorb vacancies whatever their size, since the vacancies are in supersaturation during the heat treatment, exp- laining the rapid disappearance of the interstitial loops.105,146 4.01.2.3 Postirradiation Creep There are relatively few data in the literature concerning the postirradiation creep behavior of zir- conium alloys as pointed out by Peehs and Fleisch.149 Even in the thorough review by Franklin et al.,134 very few results concerning the postirradiation creep are given. In the case of the SRA zirconium alloys142,150–155 or RXA Zy-2,142,156 several authors have shown that irradiation leads to a strong decrease of the creep rate (Figure 17). This phenomenon is attributed to the presence of a high density of irradi- ation defects that harden the material. However, according to Ito et al.142 and Schäffler et al.,152 irradi- ation does not seem to affect strongly the stress sensitivity coefficient of SRA Zy-4 (Zircaloy-4), at least for the high stress range. However, for low applied stress, Ito et al.142 have shown that the stress sensitivity coefficient is lower after irradiation than before irradiation. They have also shown that irradi- ation has a weak effect on the creep activation energy of SRA Zy-4 for temperatures from 330 to 600 �C and for stresses from 77 to 384 MPa. Murty and Mahmood157 have suggested that the creep anisot- ropy of RXA Zy-2 is decreased by irradiation. According to these authors, this phenomenon is due to the activation of other slip systems than the pris- matic slip system after irradiation, such as the basal and the pyramidal slip systems. Cappelaere et al.154 and Limon and Lehmann155 have shown that for low applied stress, a ‘tertiary 0.05 0.04 0.03 0.02 0.01 0 0 50 000 100 000 t (s) D ia m et ra l c re ep e q (–) 92.41 � 1024 n m–2 20.80 � 1024 n m–2 8.25 � 1024n m–2 350 �C 445 MPa unirr. 4.39 � 1024n m–2 45.03 � 1024 n m–2 150 000 200 000 0.06 0.07 0.08 0.09 Figure 17 Effect of fluence on thermal creep behavior at 350 �C of irradiated SRA Zy-4 cladding tubes. Reprinted, with permission, from Thirteenth International Symposium on Zirconium in the Nuclear Industry, 2002, copyright ASTM International, 100 Barr Harbor Drive, West Conshohocken, PA 19428. Radiation Effects in Zirconium Alloys 19 creep’ occurs for SRA Zy-4, even though the creep strain level remains low. This phenomenon cannot be explained by the increase of the stress due to the thinning of the wall of the tube. This phenomenon is therefore interpreted as a result of the recovery of the irradiation defects during the creep test and also due to the beginning of the recrystallization that can occur for high-temperature creep tests. Tsai and Billone158 have come to the same conclusions by analyzing their own long-term creep tests. The recovery of irradiation loops during creep tests has been observed, using TEM, by several authors on SRA Zy-4(154) or RXA Zr–1% Nb–O alloy,124 but it is the recent work by Ribis et al.105 that gives the most detailed study of the microstructure evolution during creep tests of the above alloy. The microstructure is compared to that observed after postirradiation heat treatment or after creep of the nonirradiated mate- rial. In this study, it is clearly shown that in RXA zirconium alloys, the irradiation loops are progres- sively annealed during the creep test, as for a heat treatment without an applied stress, the magnitude of the recovery being similar in both cases. Moreover, these authors show that other mechanisms associated with the deformation occur. Indeed, it is noticed that for tests performed at 400 �C and for low applied stress (130 MPa), in addition to the recovery of loops, the microstructure observed after creep tests exhibits a high dislocation density, much higher than the dislocation density observed in the nonirradiated material deformed up to the same plastic strain. According to these authors, this phenomenon results from the irradiation loops that act as obstacle to dislocation motion, especially in the prismatic planes, and limit their mean free path. This leads to an important multiplication of dislocations in order to accommodate the plastic strain. This high dislocation density can then lead to a significant hardening in addition to the hardening due to loops. This could explain that for long-term creep test performed at 400 �C under an applied stress of 130 MPa, although a significant recovery of the irradiation damage occurs, the creep strain remains limited. Additional hardening due to the high density of small b-Nb needles can also occur in the case of Zr–Nb alloys. For higher applied stress, higher than 200 MPa, these authors suggest that a sweeping of loops probably occurs. This mechanism is believed to be similar to the dislocation channeling mechanism that is observed for burst tests and tensile tests.113,124 This mechanism therefore allows the deformation of the material for high applied stress, despite the high loop density. 4.01.3 Deformation Under Irradiation 4.01.3.1 Irradiation Growth 4.01.3.1.1 Irradiation growth: Macroscopic behavior One of the most specific macroscopic effects of irra- diation on materials is the dimensional change with- out applied stress. In the case of zirconium alloys, it is known that under neutron irradiation, a zirconium single crystal undergoes an elongation along the hai C-axis crystals A-axis crystals 553 K dpa Carpenter et al.164 Iodide zirconium (6,7) Zone-refined zirconium (2,5) 0 8 6 4 2 0G ro w th s tr ai n (1 0– 4 ) –2 –4 0 1 2 3 4 Neutron fluence (1025n m–2) 5 6 7 8 1 2 3 4 5 6 7 8 9 Figure 18 High-fluence growth strain as a function of fluence for annealed zirconium single crystals at 553 K. Adapted from Carpenter, G. J. C.; Zee, R. H.; Rogerson, A. J. Nucl. Mater. 1988, 159, 86–100. 20 Radiation Effects in Zirconium Alloys axis and a shortening along the hci axis without significant volume evolution. Thorough reviews of this phenomenon have been given.72,150,159–163 It is observed that the elongation along the hai axis is rapid at the beginning of the irradiation and slows down until reaching a low stationary growth rate (Figure 18). The growth strain remains small ( 0.5), exhibit a negative growth in this direction and a positive growth in the direction with low Kearns factor (fd< 0.2). In the case of highly textured products such as cold-worked tubing, in SRA or RXA metallurgical state, a large majority of the grains exhibit their hci axis close to the radial direction (hci axes oriented in the (r, y) plane with an angle between 20� and 45� to the radial direction, the Kearns factor along the radial direction being fr� 0.6). The directions h112�0i or h10�10i are along the rolling direction (low Kearns factor along the rolling, or axial direction fa� 0.1–0.16.167,168) Due to this strong texture, an elongation of the tube along the rolling direction is observed159,169,168 as well as a decrease of the thickness as shown on rolled sheet,159 the strain along the diameter of the tube remaining low.153 In the case of pressure tube for Canadian deuterium uranium (CANDU) reactors, made of cold-worked Zr–2.5Nb, since the hci axes are mainly along the transverse direction (fr� 0.3, fa� 0.05, ft� 0.6, respectively for radial, axial, and transverse Kearns factors), the irradiation growth leads to an increase of the length in the axial direction and a decrease of the diameter.163 As for the zirconium single crystal, textured RXA Zy-4 or Zy-2 products, for instance, in the form of tubing, exhibit first a rapid elongation along the roll- ing direction, and then a decrease in the growth rate, reaching a low stationary growth rate.159 It can be noticed that the stationary growth strain of the polycrystal is higher than that for the Zr single crys- tal.161 This demonstrates the role of the grain bound- aries on the growth mechanisms. For higher fluence, higher than 3–5� 1025 nm�2, a growth breakaway is observed, yielding a high growth rate. It is reported150,160,166 that for polycrystalline zir- conium alloys, the grain size affects the growth rate 0 0 50 10 G ro w th s tr ai n � 10 4 15 20 25 30 35 40 45 353 K 553 K Annealed zircaloy–2.20mm 25% C.W. zircaloy 2.5–8 mm 50 55 20 40 60 80 100 120 140 160 180 Fast neutron dose (�1024n m–2, E > 1 MeV) FL= 0.1 Figure 19 Irradiation growth in annealed and 25% cold-worked Zircaloy-2 at 353 and 553 K. Rogerson, A. J. Nucl. Mater. 1988, 159, 43–61. Radiation Effects in Zirconium Alloys 21 of RXA zirconium alloys during the initial growth transient at 553 K, the growth rate increasing when the grain size decreases. On the other hand, the stationary growth is not affected by the grain size. This phenomenon is also observed for Zircaloy-2.159 Ibrahim and Holt170 and Holt171 have also suggested that the grain shape, especially in the case of Zr–2.5% Nb material, can play a role on the growth behavior. It is shown that for cold-worked materials (e> 10%) the growth rate increases as the cold work- ing increases150,159,160 (Figure 19). For the extreme case of SRA zirconium alloys, which could undergo up to 80% cold working followed by a SRA treat- ment, the growth rate is so high that the stationary growth rate is not observed, and from the beginning of the irradiation, the growth rate is comparable to the growth rate measured for RXA zirconium alloys after the breakaway growth. Several authors, as reviewed by Fidleris et al.159 and Holt,72 have clearly correlated the increase of the growth rate with the increase of the dislocation density due to the cold working. This also proves the importance of the ini- tial dislocations network in the growth mechanisms. Several authors have studied the effect of the impu- rity and alloying elements on the growth rate and especially on the growth acceleration. At 280 �C, for a high-purity zirconium single crystal obtained by the melting zone method, no growth breakaway is observed. On the other hand, for a lower purity zirco- nium single crystal obtained by using the iodine puri- fication method161 the breakaway growth is observed. Similarly, for polycrystalline RXA zirconium alloys, irradiated at elevated temperature (390–430 �C), the growth rate is higher than that of pure zirconium.73,160 It is particularly noticed by Griffiths et al.73 that RXA zirconium alloys exhibit accelerated growth contrary to pure zirconium. It is believed that minor elements (Fe, Cr), and especially iron, play a major role on the breakaway.54,160 On the other hand, it appears that the tin content, in solid solution, has no effect on the stationary growth rate at high temperatures (280 �C)150,160 but that the niobium leads to a reduced growth rate compared to RXA Zy-4.168 The irradiation temperature has a complex influ- ence on the growth behavior72,150 (Figure 20). For SRA zirconium alloys, it is shown that the growth rate increases as the temperature increases. On the other hand, for RXA zirconium alloys the prebreakaway growth rate has a very low temperature sensitivity, the growth rate increasing very slowly with increasing temperature. A growth peak is even observed around 570 K, the growth rate decreasing rapidly above 620 K. However, for postbreakaway growth, the tempera- ture sensitivity is high, as high as for SRA zirconium alloys.150 It is also shown that the breakaway fluence decreases with increase in the temperature.72 4.01.3.1.2 Irradiation growth: Mechanisms The mechanisms proposed in the literature in order to explain the growth under irradiation of zirconium and its alloys have progressively evolved as the obser- vations of the microstructure have progressed. 700 10–27 10–28 10–29 10–30 1.4 � 10–3 1.8 � 10–3 2.2 � 10–3 1/T(K) Temperature (K) 2.6 � 10–3 3.0 � 10–3 600 500 400 f = 0.10 p = 5 � 1014m–2 p = 1 � 1014m–2 350 Recrystallized (prebreakaway) Recrystallized (postbreakaway) G ro w th r at e (m 2 n– 1 ) Q»150 kJ mol–1 Q»3 kJ mol–1 Cold worked Figure 20 Generalized representation of the temperature dependence of irradiation growth of Zircaloy. Adapted from Fidleris, V. J. Nucl. Mater. 1988, 159, 22–42. (a) (b) (d) (c) Figure 21 (a–c) The three phases of growth of recrystallized zirconium alloys. (d) Growth mechanisms of stress relieved zirconium alloys. 22 Radiation Effects in Zirconium Alloys Several comprehensive reviews of these mechanisms have been given,44,46,72,163 and a nice history of the various mechanisms for irradiation growth of zirco- nium alloys is provided by Holt.162 Some of these mechanisms are not compatible with all the observa- tions. For instance, the fact that both vacancy and interstitial hai loops are present in the polycrystalline material, as described in the first part, shows that the model proposed by Buckley172 described in Northwood173 and Holt162 for the growth of zirconium alloys is not correct. The most promising model that gives the best agreement with the observations is the model based on the DAD, first proposed by Woo and Gösele174 and described in detail by Woo.44 This last model is based on the assumption that the diffusion of SIAs is anisotropic, the vacancy diffusion anisotropy being low. Indeed, as reported in the first part of this chap- ter, several authors28,33,34,175 have shown, using atom- istic simulations, that the mobility of the SIAs is higher in the basal plane than along the hci axis and that the vacancy diffusion is only slightly anisotropic. The growth mechanism proposed by Woo44 is the most convincing model, since every feature of the growth phenomenon is understood in its frame unlike in the previous models. According to this mechanism, during the first stage of the irradiation of RXA zirconium alloys, with low initial dislocation density, the grain boundaries are the dominant sinks. Due to the rapid mobility of SIAs in the basal plane, the grain boundaries perpendicular to the basal plane are preferential sinks for SIAs. In contrast, grain boundaries parallel to the basal plane constitute preferential sinks for vacancies. This leads to a fast initial growth of polycrystalline zirconium alloys, in agreement with the model first proposed by Ball176 (Figure 21(a)). This mechanism explainswhy the initial growth transient is sensitive to the grain size. As the irradiation dose increases, the hai loop density increases and the hai loops become the dom- inant sink for point defects. In the absence of hci component dislocation (as is the case in RXA zirco- nium alloys), calculations of DAD-induced bias Radiation Effects in Zirconium Alloys 23 show that linear hai type dislocations parallel to the hci axis are preferential SIA sinks while hai type loops are relatively neutral and may receive a net flow of either interstitials or vacancies, depending on the sink situation in their neighborhood. This explains why both interstitial and vacancy hai type loops can be observed. This also explains why in the neigh- borhood of prismatic grain boundaries, or surfaces, which experience a net influx of SIAs, there will be a higher vacancy supersaturation leading to a predom- inance of vacancy loops towards interstitial loops as shown by Griffiths.46 It has to be pointed out that the simultaneous growth of interstitial and vacancy hai type loops in the prismatic plane does not induce strain of the crystal although they are the dominant sinks (Figure 21 (b)). This explains the low station- ary growth rate observed. For irradiation doses higher than 5� 1025 nm�2, vacancy hci component dislocation loops in the basal plane are observed in RXA zirconium alloys (Figure 21 (c)). The origin of the nucleation of hci component loops remains unclear. Nevertheless, it has been shown, as described previously, that it is favored by the iron dissolution in the matrix coming from the precipitates.57,73,75,76 The appearance of hci component defects has been clearly correlated to the breakaway growth71 (Figure 22). The fact that these vacancy hci component basal loops are able to grow in zirconium alloys, whereas it is the hai prismatic loops that are the most stable, is easily explained in the frame of the DAD model. Indeed, it can be shown that it is due to the DAD that vacancies are elimi- nated preferentially on the hci component loops and on the grain boundaries parallel to the basal plane. The SIAs are eliminated on hai type dislocations A2 G2D2LD3T 0.15 0.10 0.05 Ir ra d ia tio n gr ow th s tr ai n (% ) 0 1 2 3 4 Fluence (n m–2) (E No component dislocations Som dislo Growth specimens in DIDO (553 K) Figure 22 Irradiation growth in annealed (RXA) Zircaloy at 550 (E>1 MeV). Adapted from Fidleris, V. J. Nucl. Mater. 1988, 159 and grain boundaries parallel to the prismatic plane. This partitioning of the point defects on these differ- ent sinks leads to the growth of the vacancy hci component loops and therefore to the accelerated growth of RXA zirconium alloys. However, as pointed out by Griffiths et al.,73 although there is a clear correlation between the occurrence of the breakaway and the appearance of hci loops, the strain induced by the loops observed is much lower than the growth strain measured. The fast and continuous growth of cold-worked or SRA zirconium alloys can also be easily explained by this model. Indeed, since in these materials the hcþ ai line dislocations are already present before irradiation, under irradiation, the vacancies are prefer- entially eliminated on the dislocations with hcþ ai Burgers vector in the basal plane,72,162,163 leading to the climb of these dislocations. On the other hand, the SIAs are eliminated on hai type dislocations, leading to the climb of these dislocations. This parti- tioning of point defects therefore leads to the fast and continuous growth of cold-worked or SRA zirco- nium alloys (Figure 21 (d)). Here the growth created by the point-defect flux to the grain boundaries is relatively unimportant because they are not dominat- ing sinks. Irradiation growth under such circumstances is thus not sensitive to the grain size or shape.177 It has also been discussed by several authors, espe- cially by Holt,162 that due to the polycrystalline nature of the material, the growth strain of the indi- vidual grains can induce strain incompatibilities between adjacent grains that exhibit different orien- tations. Intergranular stresses can then result from these strain incompatibilities, leading to a local irradiation creep of individual grains even without D2 D1 5 > 1 MeV) � 10–25 6 7 8 9 e component cations Many component dislocations Fuel assembly guide tubes in calvert cliffs-1 (508–583 K) –580 K, showing accelerating growth at 4� 1025 nm�2 , 22–42. 24 Radiation Effects in Zirconium Alloys macroscopic applied stress on the material. This phe- nomenon can also affect the growth behavior of the polycrystalline material. It has also been shown that the intergranular stresses resulting from a deformation prior to irradiation can lead to a complex transient growth behavior at the beginning of the irradiation due to intergranular stress relaxation.162,178 4.01.3.2 Irradiation Creep 4.01.3.2.1 Irradiation creep: Macroscopic behavior Under neutron irradiation, metals exhibit a high creep rate, much higher than the out-of-reactor ‘ther- mal’ creep rate, the creep rate increasing as the neu- tron flux increases. The behavior under irradiation of zirconium alloys, and particularly the creep behav- ior, has been studied extensively as pointed out by Franklin et al.134 and Fidleris,150 because of the major importance of the prediction of the in-reactor defor- mation of the fuel assembly in the case of PWR and boiling-water reactor (BWR)169 or in-reactor struc- ture especially in the case of the CANDU reactor.163,179 It is usually assumed, for practical considerations, that the in-pile deformation consists of the sum of (i) the growth, (ii) the classical thermally activated out-of-pile creep, or so-called thermal creep, and (iii) the irradiation creep, strictly speaking.100,150,163,180 The ‘pure’ irradiation creep, subtracted from the two other components of the deformation, is the result of mechanisms which differ from the thermal creep and the growth. Nevertheless, these mechanisms are certainly coupled since they all imply disloca- tion loops, slip and climb of line dislocations, and point-defect bulk diffusion toward these defects. But very few authors have studied these potential couplings.134,181 The creep deformation under irradiation results, in fact, from two antagonistic phenomena. Indeed, while new deformation processes are activated, caus- ing the creep rate to increase, the thermal creep rate is strongly reduced by irradiation due to the irradiation-induced hardening. Indeed, it has been shown150 that a preirradiation reduces the thermal creep component of the deformation under irradia- tion. The effect of preirradiation on the reduction of the irradiation creep rate is particularly noticeable for RXA alloys. However, the hardening effect saturates at fluence of about 4� 1024 nm�2 and is followed by a steady-state creep rate. Concerning cold-worked materials, the effect of the preirradia- tion is much lower, according to Fidleris.150 As reported by several authors,134,150,153,182 the metallurgical state of the zirconium alloy has a sig- nificant effect on the in-reactor creep resistance. Indeed, while cold working may improve the thermal creep resistance of Zircaloy in certain test directions and stress range, it increases the in-reactor creep rate appreciably.150,153 Nevertheless, the creep sensitivity to the initial dislocation density is significantly lower than the growth sensitivity to the initial dislocation density.171 On the other hand, the grain size does not seem to have a significant effect on the creep strength in the range from 1 to 70 mm. The in-reactor creep rate is very sensitive to irra- diation as well as loading conditions. The effects of flux, as well as the effect of stress, are usually described by a power correlation. The effect of tem- perature is usually described by an Arrhenius equa- tion.134 However, since it is in general very complex to distinguish between the ‘pure’ irradiation creep and the thermal creep, the authors usually use an overall creep constitutive law (eqns [1] and [2])163,180 and only growth is taken into account as a separate defor- mation component. _e ¼ _ethermal�creep þ _eirradiation�creep þ _egrowth ¼ _ecreep þ _egrowth ½1� with _ecreep ¼ Ksnfpexp Q RT � � ½2� where _e is the strain rate in s�1; s is the effective stress for thermal creep inMPa; n is the stress exponent;T is the temperature in K; Q is the activation energy in J; R is the gas constant, 8.31 J K�1 mol�1; � is the fast neutron flux in nm�2 s�1 (E> 1 MeV); p is the flux exponent; and K is a constant for thermal creep in s�1 (MPa)�n(nm�2 s�1)�p. According to various authors,134,150 the flux exponent (p) has been assigned values ranging from 0.25 to 1. A flux exponent of p¼ 1 is commonly obtained for CANDU pressure tube deformation.163,183 For uniaxial creep tests per- formed at 280 �C on cold-worked Zy-2, Tinti184 has obtained a flux exponent increasing from 0.6 to 1.0 with increasing instant flux. A stress exponent of n¼ 1 is obtained at 300 �C for low applied stress (s 100MPa). As the stress increases, the stress exponent increases, reaching values up to n¼ 25 for 450 MPa applied stress for cold-worked Zr–2.5% Nb.183 Temperature (�C) C re ep r at e (h –1 ) 400 10–4 10–5 10–6 10–7 10–8 1.4 1.5 1.6 1.7 1.8 1/T � 103 (K) 1.9 2.0 2.1 350 300 250 200 + + + + + + + 207 MPa 138 MPa 207 MPa138 MPa Laboratory tests, t > 6000 h In-reactor tests, t > 6000 h Flux = 9�1016n m–2s–1, E > 1 MeV Figure 23 Temperature dependence of laboratory and in-reactor creep rates of cold-worked Zircaloy-2. Adapted from Fidleris, V. J. Nucl. Mater. 1988, 159, 22–42. Radiation Effects in Zirconium Alloys 25 The effect of temperature on the creep rate can be rationalized by plotting the creep rate in an Arrhenius plot (logarithm of the creep rate vs. inverse temperature). The activation energy is then the slope obtained in this plot. It can be seen in Figure 23 that for low temperatures, the creep activa- tion energy Q /R is very low, between 2000 and 5000 K.150,163 The irradiation creep at low temperature is therefore nearly athermal. At higher temperatures, the dependence increases rapidly toward values of Q /R of 25 000–30 000 K. These last values are close to the activation energy measured for thermal creep. These observations tend to prove that for low-temperature regime, mainly ‘pure’ irradiation creep mechanisms are activated. As the temperature increases, the ther- mal creep mechanisms become activated, yielding to activation energy close to the thermal creep values. It has also been shown by several authors that while the thermal creep of zirconium alloys is aniso- tropic, the irradiation creep remains strongly anisotropic.150 According to Holt,171 the anisotropy of irradiation creep is nevertheless slightly lower than that of thermal creep. 4.01.3.2.2 Irradiation creep: Mechanisms Various mechanisms for irradiation creep have been proposed in the literature as reviewed by Franklin et al.,134 Holt,163,171 Matthews and Finnis,181 and Was.9 A nice history of the proposed mechanisms for both zirconium alloys and stainless steels is given by Franklin et al.134 These mechanisms can fall mainly into two large categories: 1. The mechanisms based on stress-induced prefer- ential absorption (SIPA) of point defects by line dislocations arising from different fundamental phenomenon. These mechanisms lead to the climb of edge dislocations under applied stress, yielding a creep deformation. 2. The mechanisms based on climb-enhanced disloca- tion glide mechanisms, which are essentially a com- bination of climb of dislocations due the absorption of point defects under irradiation and glide resulting in a creep deformation. For this categoryofmechan- isms, the strain is essentially produced by glide but the strain rate is controlled by the climb. Other mechanisms involving irradiation-induced loops have also to be added to these two categories of deformation mechanisms involving line disloca- tions. Indeed, the stress-induced preferential nucle- ation (SIPN) of loops or the stress-induced preferential growth of loops due to SIPA can lead to an additional creep strain. The SIPA mechanism is based on the fact that under an applied stress, the bias of the dislocation becomes dependent on the orientation of the Burgers vector with respect to the direction of the stress.105,134,181 Indeed, as described previously, due to a higher relaxation volume, the sink strength of an edge dislocation toward SIAs is higher than toward vacancies. This difference in sink strength is the bias of the edge dislocation. It can be shown that a dislo- cation with a Burgers vector parallel to the applied stress exhibits a higher bias toward SIAs than a dislo- cation with a Burgers vector perpendicular to the applied stress. Therefore, under irradiation, the net flux of SIAs (SIA flux minus vacancy flux) toward dislocations, with Burgers vector parallel to the applied stress, is higher than the net flux of SIAs toward dislocations with Burgers vector perpendicular to the applied stress. This difference in the absorption of point defects by different types of dislocations leads to dislocation climb, resulting in a creep strain. The SIPA creep rate is insensitive to the grain size but is sensitive to the dislocation network. 26 Radiation Effects in Zirconium Alloys However, it has been seen that for growth, the anisotropic diffusion of SIAs is believed to play an important role in the deformation mechanism. There- fore, any irradiation creep model proposed for zir- conium should also include anisotropic diffusion. The SIPA model that includes anisotropic diffusion is called the SIPA-AD model and has been reviewed by Matthews and Finnis.181 In the case of RXA zirconium alloys, the irradia- tion creep mechanisms are not clearly identified yet. Indeed, since the initial dislocation density is very low, another deformation mechanism has to be acti- vated. The creep strain could be partly due to the preferred nucleation and/or growth of the hai type loops in the prismatic planes. Indeed, according to the SIPN or SIPA mechanism, the nucleation or growth of interstitial hai loops can be favored in the prismatic planes perpendicular to the applied stress. For the same reason, the nucleation or growth of vacancy hai loops can be favored in the prismatic planes parallel to the applied stress, leading to a resulting creep strain. According to Faulkner and McElroy,185 an applied stress increases the mean diameter of hai loops without affecting the density, proving that the SIPA mechanism is efficient in their experiment. However, the growth of hai loops under an applied stress can explain the measured creep strain only for low strain levels. Indeed, this creep strain should remain limited since the hai loop den- sity and mean loop diameter saturate at relatively low doses. Since the initial dislocation density is very low in RXA zirconium alloys, creep mechanisms involv- ing climb of dislocations due to the SIPA mechanism or climb-plus-glide of dislocations require the gener- ation of a dislocation network. It is possible that hai loops coalescence occurs, resulting in the creation of a dislocation network that is able to climb and glide under stress.181,186 However, this network is clearly observed only at 400 �C.67 Other types of dislocation sources, such as Frank–Read or Bardeen–Herring sources,147 can also be activated under both irradia- tion and applied stress, leading to the creation of a dislocation network that undergoes a SIPA or climb- enhanced glide mechanism. It should also be pointed out that in order to explain the observed creep rate, some mechanisms must be activated that allow the dislocations to over- come the high density of dislocation loops during their climb and glide motion, even for low applied stress. It is possible, as pointed out by MacEwen and Fidleris187 in the case of Zr single crystal, that the gliding dislocations are able to clear the loops during in-pile deformation, leading to the dislocation chan- neling mechanism. All these mechanisms probably occur in series, as proposed by Nichols,188 explaining the evolution of the stress dependency as the stress increases. Indeed, according to this author, for zero applied stress, growth of zirconium occurs, and then as the stress increases, hai loop alignment occurs (SIPA on loops). For higher stress, the climb of line dislocations via SIPA takes place, and then the dislo- cation climb and glide processes occur at even higher stress. For very high stress, close to the YS, disloca- tion channeling occurs. For cold-worked zirconium alloys, such as SRA Zircaloy or cold-worked Zr–2.5Nb alloy,163 the SIPA mechanism on the initial dislocations is a likely mech- anism for irradiation creep. However, according to Holt,171 the creep anisotropy of cold-worked zirco- nium alloys computed from the SIPA mechanism assuming only hai type dislocations is not in agree- ment with the experimental anisotropy. The anisot- ropy computed from the climb-plus-glide mechanism assuming 80% prism slip and 20% basal slip is in good agreement with the experimental anisotropy, demon- strating that climb-plus-glide mechanism is probably the effective mechanism. It should also be pointed out that, since dislocations climb toward grain boundaries or toward other dislocations, recovery of the initial dislocation network occurs. In order to maintain a steady-state creep rate, multiplication of dislocations should also occur either via loop coalescence or via dislocation sources, as discussed previously. It should also be pointed out that, as there is a coupling between swelling and irradiation creep in stainless steel,181 we could assume a coupling between growth and irradiation creep to occur in zirconium alloys due to the effect of the stress on the partitioning of point defects.134,162 Nevertheless, the simple assumption of two separable deformation components has proved to hold correctly for the results given in the literature.163,180 4.01.3.3 Outlook Concerning damage creation and point-defect clus- ter formation, improvement in the knowledge of anisotropic diffusion of SIAs as well as better under- standing of the microstructure of vacancy and inter- stitial hai loops and basal hci vacancy loops (origin of the loop alignment, origin of the corduroy contrast Radiation Effects in Zirconium Alloys 27 for instance) has to be aimed at. Multiscale modeling approaches coupled with fine experimental analyses of the irradiation microstructure (high-resolution TEM, synchrotron radiation analyses, tomography atom probe, etc.) should bring new insight concerning the previous points mentioned but also elements in order to propose modeling of the microstructure evolution during irradiation: for instance, origin of the alignments of Nb precipitates, stability of b-Nb pre- cipitates, etc. Concerning the mechanical behavior of Zr alloys after irradiation, multiscale modeling of the postirra- diation deformation with a better understanding of the dislocation channeling mechanism and under- standing of its effects on the postirradiation mechan- ical behavior are needed. Moreover, better understanding of the postirradi- ation creep deformation mechanisms is also needed using multiscale modeling. The last point concerns the deformation mechan- isms under irradiation. 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All rights reserved. 4.02.1 Introduction 34 4.02.2 Basic Damage Processes 35 4.02.2.1 Atomic Displacements 35 4.02.2.2 Transmutation 37 4.02.3 Differences in Neutron Spectra 37 4.02.4 Transmutation Issues for Stainless Steels 40 4.02.5 Evolution of Radiation-Induced Microchemistry and Microstructure 44 4.02.6 A Cross-Over Issue Involving Radiation-Induced Microstructural Evolution and Transmutation 49 4.02.7 Radiation-Induced Changes in Mechanical Properties 50 4.02.8 Radiation-Induced Changes in Dimension 61 4.02.8.1 Precipitation-Related Strains 62 4.02.8.2 Void Swelling and Bubble Swelling 65 4.02.8.3 Parametric Dependencies of Void Swelling 67 4.02.8.3.1 Stress state 67 4.02.8.3.2 Elemental composition 68 4.02.8.3.3 Alloy starting state 69 4.02.8.3.4 Irradiation temperature 69 4.02.8.3.5 Influence of dpa rate on swelling 70 4.02.9 Irradiation Creep 74 4.02.9.1 Introduction 74 4.02.9.2 Stages of Irradiation Creep 78 4.02.9.3 Examples of Creep Behavior 79 4.02.9.4 Creep Disappearance 79 4.02.9.5 Recent Revisions in Understanding of Irradiation Creep 83 4.02.9.5.1 Dependence of irradiation creep on dpa rate 83 4.02.9.5.2 Dependence of creep and creep relaxation on neutron spectra 84 4.02.9.5.3 Dependence of creep modulus on hydrostatic stress 85 4.02.9.6 Stress Relaxation by Irradiation Creep 86 4.02.9.7 Stress Rupture 88 4.02.9.8 Fatigue 89 4.02.10 Conclusions 90 References 91 Abbreviations ATR Advanced Test Reactor in Idaho Falls, Idaho BN-350 Russian acronym for Fast Neutron at 350 MW in Actau, Kazakhstan BN-600 Russian acronym for Fast Neutron at 600 MW in Zarechney, Russia BOR-60 Russian acronym for Fast Experimental Reactor at 60 MW in Dimitrovgrad, Russia BR-2 Belgium Research Reactor-II in Mol, Belgium BR-10 Russian acronym for Fast Reactor at 10 MW in Obninsk, Russia BWR Boiling water reactor CAGR Commercial Advanced Gas Reactor CANDU Registered trademark for Canadian Deuterium Uranium Reactor DFR Dounreay Fast Reactor in Dounreay, Scotland 33 34 Radiation Damage in Austenitic Steels DMTR Dounreay Materials Test Reactor in Dounreay, Scotland EBR-II Experimental Breeder Reactor-II in Idaho Falls, Idaho FFTF Fast Flux Test Facility, fast reactor in Richland, WA HFIR High Flux Isotope Reactor at Oak Ridge National Laboratory HFR High Flux Reactor in Petten, Netherlands IASCC Irradiation-assisted stress corrosion cracking IGSCC Intergranular stress corrosion cracking JMTR Japan Material Testing Reactor in Oarai, Japan NRU National Research Universal Reactor in Chalk River, Canada ORNL Oak Ridge National Laboratory: ORR Oak Ridge Research Reactor in Oak Ridge, Tennessee PWR Pressurized water reactor T/F Thermal-to-fast neutron ratio VVER Russian acronym for water-cooled, water moderated energetic reactor 4.02.1 Introduction Austenitic stainless steels are widely used as struc- tural components in nuclear service in addition to being employed in many other nonnuclear engineering and technological applications. The description of these steels and their as-fabricated properties is covered in Chapter 2.09, Properties of Austenitic Steels for Nuclear Reactor Applica- tions. This chapter describes the evolution of both microstructure and macroscopic property changes that occur when these steels are subjected not only to prolonged strenuous environments but also to the punishing effects of radiation. While various nuclear environments involve mixtures of charged particles, high-energy photons and neutrons, it is the latter that usually exerts the strongest influence on the evolution of structural steels and thereby determines the lifetime and continued functionality of structural components. To describe the response of austenitic stainless steels in all neutron environments is a challenging assignment, especially given the wide range of neutron spectra characteristic of various neutron devices. This review of neutron-induced changes in properties and dimensions of austenitic stainless steels in all spectral environments has therefore been compiled from a series of other, more focused reviews directed toward particular reactor types1–8 and then augmented with material from a recently published textbook9 and journal articles. It should be noted, however, that many of the behavioral char- acteristics of iron-based stainless steels following neutron irradiation are also observed in nickel- based alloys. Whenever appropriate, the similarities between the two face-centered-cubic alloy systems will be highlighted. A more comprehensive treat- ment of radiation effects in nickel-base alloys is provided in Chapter 4.04, Radiation Effects in Nickel-Based Alloys. This review is confined to the effects of neutron exposure only on the response of irradiated steels and does not address the influence of charged particle irradiation. While most of the phenomena induced by neutrons and charged particles are identical, there are additional processes occurring in charged par- ticle studies that can strongly influence the results. Examples of processes characteristic of charged par- ticle simulations are the injected interstitial effect,10,11 strong surface effects,12,13 dose gradients,14,15 and atypical stress states.16,17 Chapter 1.07, Radiation Damage Using Ion Beams addresses the use of charged particles for irradiation. Austenitic stainless steels used as fuel cladding or structural components in various reactor types must often withstand an exceptionally strenuous and chal- lenging environment, even in the absence of neutron irradiation. Depending on the particular reactor type, the inlet temperature during reactor operation can range from �50 to �370 �C. The maximum temper- ature can range from as high as 650 to 700 �C for structural components in some reactor types, although most nonfueled stainless steel components reach maximum temperatures in the range of 400–550 �C. During operation, the steel must also withstand the corrosive action of fission products on some surfaces and flowing coolant on other surfaces. The coolant especially may be corrosive to the steel under operating conditions. Some of these environmental phenomena are synergized or enhanced by the effect of neutron irradiation. Dependent on the nature of the component and the length of its exposure, there may also be sig- nificant levels of stress acting on the component. Stress not only influences cracking and corrosion (see Chapter 5.08, Irradiation Assisted Stress Cor- rosion Cracking) but can also impact the dimen- sional stability of stainless steel, primarily due to Radiation Damage in Austenitic Steels 35 thermal creep and irradiation creep, and also from the influence of stress on precipitation, phase stabil- ity, and void growth, some of which will be discussed later. However, it will be shown that neutron irradia- tion can strongly affect both the microstructure and microchemistry of stainless steels and high-nickel alloys, with strong consequences on physical proper- ties, mechanical properties, dimensional stability, and structural integrity. Stainless steels are currently being used or have been used as structural materials in a variety of nuclear environments, most particularly in sodium- cooled fast reactors, water-cooled and water-moderated test reactors, water-cooled and water-moderated power reactors, with the latter subdivided into light water and heavy water types. Additionally, there are reactor types involving the use of other coolants (helium, lithium, NaK, lead, lead–bismuth eutectic, mercury, molten salt, organic liquids, etc.) and other moderators such as graphite or beryllium. The preceding reactor types are based on the fission of uranium and/or plutonium, producing neutron energy distributions peaking at �2MeV prior to moderation and leakage effects that produce the operating spectrum. However, there are more energetic sources of neutrons in fusion-derived spectra, with the source peaking at �14MeV and especially from spallation events occurring at ener- gies of hundreds of MeV, although most spallation spectra are mixtures of high-energy protons and neutrons. It is important to note that in each of these various reactors, there are not only significant differences in neutron flux-spectra but also signifi- cant differences in neutron fluence experienced by structural components. These differences in fluence arise not only from differences in neutron flux characteristic of the different reactor types but also the location of the steel relative to the core. For instance, boiling water reactors and pressurized water reactors have similar in-core spectra, but stainless steels in boiling water reactors are located much farther from the core, resulting in a factor of reduction of �20 in both neutron dose rate and accumulated dose compared to steels in pressurized water reactors. 4.02.2 Basic Damage Processes 4.02.2.1 Atomic Displacements What are the nature and origins of neutron-induced phenomena in metals? The major underlying driving force arises primarily from neutron collisions with atoms in a crystalline metal matrix. When exposed to displacive irradiation by energetic neutrons, the atoms in a metal experience a transfer of energy, which if larger than several tens of eV, can lead to displacement of the atom from its crystalline position. The displace- ments can be in the form of single displacements resulting from a low-energy neutron collision with a single atom or a glancing collisionwith a higher energy neutron. More frequently, however, the ‘primary knock-on’ collision involves a larger energy transfer and there occurs a localized ‘cascade’ of defects that result from subsequent atom-to-atom collisions. There are several other contributions to displace- ment of atoms from their lattice site, but these are usually of second-order importance. The first of these processes involve production of energetic elec- trons produced by high-energy photons via the photoelectric effect, Compton Effect, or pair produc- tion.18 These electrons can then cause atomic displa- cements, but at a much lower efficiency than that associated with neutron-scattering events. The sec- ond type of process involves neutron absorption by an atom, its subsequent transmutation or excitation, followed by gamma emission. The emission-induced recoil of the resulting isotope often is sufficient to displace one or several atoms. In general, however, such recoils add a maximum of only several percent to the displacement process and only then in highly thermalized neutron spectra.4 One very significant exception to this generalization involving nickel will be presented later. For structural components of various types of nuclear reactors, it is the convention to express the accumulated damage exposure in terms of the calcu- lated number of times, on the average, that each atom has been displaced from its lattice site. Thus, 10 dpa (displacements per atom) means that each atom has been displaced an average of 10 times. Doses in the order of 100–200 dpa can be accumulated over the lifetimes of some reactor components in various high-flux reactor types. The dpa concept is very useful in that it divorces the damage process from the details of the neutron spectrum, allowing comparison of data generated in various spectra, providing that the damage mechanism arises primar- ily from displacements and not from transmutation. The use of the dpa concept also relieves research- ers from the use of relatively artificial and sometimes confusing threshold energies frequently used to describe the damage-causing portion of the neutron spectrum. Neutrons with ‘energies greater than 36 Radiation Damage in Austenitic Steels X MeV,’ where X is most frequently 0.0, 0.1, 0.5, or 1.0MeV, have been used for different reactor con- cepts at different times in history. The threshold energy of 0.1MeV is currently the most widely used value and is most applicable to fast reactors where large fractions of the spectra lay below 0.5 and 1.0MeV. Many older studies employed the total neutron flux (E> 0.0) but this is the least useful threshold for most correlation efforts. Caution should be exercised when compiling data from many older studies where the neutron flux was not adequately identified in terms of the threshold energy employed. There are rough conversion factors for ‘displace- ment effectiveness’ for 300 series austenitic steels that are useful for estimating dpa from>0.1MeV fluences for both in-core or near-core spectra in most fission spectra. Examples are �7 dpa per 1022 n cm�2 (E > 0.1) for most in-core light water spectra with lower in-core values of�5 dpa per 1022 n cm�2 (E> 0.1) for metal fueled fast reactors and�4 dpa per 1022 n cm�2 (E > 0.1) for oxide-fueled fast reactors.4 Such con- version factors should not be trusted within more than (10–15%), primarily due to spatial variations across the core resulting from neutron leakage. For fast reactor spectra, E > 1.0 conversion factors are completely unreliable. When E > 1.0 fluxes are employed in light water reactor studies, the conversion factor increases from �7 dpa per 1022 n cm�2 (E > 0.1) to �14 dpa per 1022 n cm�2 (E > 1.0). In Russia, a threshold energy of >0.5 MeV is popular for light water Neutron fluence, E > 0.1 MeV 1017 0 50 100 150 200 250 300 1018 1019 1020 Y ie ld s tr es s ch an ge (M P a) LASREF, 40 �C RTNS-II, 90 �C OWR, 90 �C Figure 1 Radiation-induced yield stress changes of 316 stainl and (right) displacements per atom. Reproduced from Heinisch, J. Nucl. Mater. 1992, 191–194, 1177, as modified by Greenwoo reactors with �9 dpa per 1022 n cm�2 (E > 0.5). All of these conversion factors assume that within several percent pure iron is a good surrogate for 300 series alloys. Note that other metals such as Cu, Al, W, etc. will have different conversion values arising from different displacement threshold energies and some- times different displacement contributions. A standard procedure for calculating dpa has been published,19 although other definitions of dpa were used prior to international acceptance of the ‘NRT model’ where the letters represent the first letter of the three author’s last name (see Garner1 for details on earlier models). Caution must be exercised when compiling doses from older studies where displacement doses were calculated using other mod- els (Kinchin-Pease, Half-Nelson, French dpa, etc.) sometimes without clearly identifying the model employed. Conversion factors between the NRT model and various older models of dpa are provided in Garner,1 but all models agree within �23%. While sometimes controversial with respect to how far the dpa concept can be stretched to cover the full range of spectral differences for neutron and especially for charged particle environments, it appears that the dpa concept is very efficient to stretch over light water, heavy water, fusion, and spallation spectra, providing that all energy deposition and displacement processes are included. Note in Figure 1 how well the dpa concept collapses the data on neutron-induced strengthening of stainless steel into one response function for three very different spectra (light water fission, pure D–T fusion and ‘beam-stop’ spallation).20 10-3 10-2 0 50 dpa 100 150 200 250 300 Y ie ld s tr es s ch an ge (M P a) LASREF, 40 �C RTNS-II, 90 �C OWR, 90 �C ess steel versus (left) neutron fluence (n cm�2 E> 0.1 MeV), H. L.; Hamilton, M. L.; Sommer, W. F.; Ferguson, P. d, L. R. J. Nucl. Mater. 1994, 216, 29–44. Radiation Damage in Austenitic Steels 37 4.02.2.2 Transmutation It is important to note that material modification by radiation arises from two primary spectral-related processes . In addition to the neutron-induced dis- placement of atoms there can be a chemical and/or isotopic alteration of the steel via transmutation. With the exception of helium production, transmuta- tion in general has been ignored as being a significant contributor to property changes of stainless steels and nickel-base alloys. In this chapter, transmutation is shown to be sometimes much more important than previously assumed. Both the displacement and transmutation pro- cesses are sensitive to the details of the neutron flux-spectra, and under some conditions each can synergistically and strongly impact the properties of the steel during irradiation. In addition to the brief summary presented below on flux-spectra issues rel- evant to stainless steels, the reader is referred to various papers on transmutation and its consequences in different reactor spectra.5–8,18,21–23 Transmutation may be subdivided into four cate- gories of transmutants. Three of these are relevant to fission-derived or fusion-derived spectra, and the fourth is associated with spallation-derived spectra. The first three are solid transmutants, gaseous trans- mutants, and ‘isotope shifts,’ the latter involving pro- duction of other isotopes of the same element. While the latter does not change the chemical composition of stainless steels, it is an underappreciated effect that is particularly relevant to nickel-containing alloys such as stainless steels and nickel-base alloys when irradiated in highly thermalized neutron spectra. Whereas the first three categories arise from dis- crete nuclear reactions to produce discrete isotopes of specific elements, the spallation-induced transmuta- tion arising in accelerator-driven devices involves a continuous distribution of every conceivable fragment of the spalled atom, producing every element below that of the target atom across a wide range of isotopes for each element. While individual solid transmutants in spallation spectra are usually produced at levels that do not change the alloy composition significantly, the very wide range of elements produced allows the possibility that deleterious impurities not normally found in the original steel may impact its continued viability. This possibility has not received sufficient attention and should be examined further if spallation devices continue to be developed. Another consequence of spallation-relevant trans- mutation is that the induced radioactivity per unit mass is correspondingly much higher than that pro- duced per dpa in other spectra. The majority of the spalled fragments and their daughters/granddaugh- ters are radioactive with relatively short half-lives, leading to materials that are often much more diffi- cult to examine than materials irradiated in fission spectra. Most importantly, there is a very strong produc- tion of hydrogen and helium in spallation spectra at levels that are one or two orders of magnitude greater than produced in most fission or fusion spectra.5,6,21 While there is a tendency to view displacement and transmutation processes as separate processes, it will be shown later that under some circum- stances the two processes are strongly linked and therefore inseparable in their action to change alloy behavior. 4.02.3 Differences in Neutron Spectra There are significant differences in neutron spectra for water-cooled, sodium-cooled, and other types of fission-based reactors. It should be noted that there is a conventional but slightly misleading practice to differentiate between ‘fast’ and ‘thermal’ reactors. Thermal reactors have a significant portion of their spectra composed of thermal neutrons. Thermalized neutrons have suffered enough collisions with the moderator material that they are in thermal equilib- rium with the vibrations of the surrounding atoms. Efficient thermalization requires low-Z materials such as H, D, and C in the form of water, graphite, or hydrocarbons. At room temperature the mean energy of thermalized neutrons is 0.023 eV. The designation ‘fast’ reactor, as compared to ‘thermal’ reactor, refers to the portion of the neutron spectrum used to control the kinetics of ascent to full power for each type of reactor. As shown later, this practice incorrectly implies to many that fast reactors have ‘harder’ neutron spectra than do ‘softer’ thermal reactors. Actually, the opposite is true. Examples of typical flux-spectral differences in fission-based reactors are shown in Figures 2–5. The local spectrum at any position is determined primarily by the fuel (U, Pu) and fuel type (metal, oxide, carbide, etc.), the coolant identity and density, the local balance of fuel/coolant/metal as well as the proximity to control rods, water traps, or core bound- aries. Additionally, it is possible to modify the neutron spectra in a given irradiation capsule by including in it Neutron energy (MeV) 1.E - 9 1.E + 08 1.E + 10 Fl ux /le th ar gy 1.E + 12 1.E + 14 1.E + 16 1.E - 7 1.E - 5 1.E - 3 1.E - 1 1.E + 1 HFIR-RB* HFIR-PTP ATR-ITV FFTF EBRII Figure 4 Comparison of flux-spectra in various test reactors. Note that FFTF is softer in spectrum compared to EBR-II due to the use of oxide fuel rather than metal fuel. Neither fast reactor has measurable fluxes of thermal neutrons. In the PTP position of HFIR a water trap strongly contributes to a high thermal-to-fast ratio, while in the RB* (removable beryllium) position the predominance of Be over water reduces the thermal population. In the ATR position where the ITV assembly was located, the use of strong absorber sleeves strongly depressed the thermal flux. 10-8 1012 Neutron energy (MeV) EBR II ORR HFIR Fl ux p er u ni t le th ar gy 1013 1014 1015 10-7 10-6 10-5 10-4 10-3 10-2 10-1 100 101 102 Figure 2 Difference in neutron flux-spectra of two water-cooled test reactors (high-flux HFIR and lower-flux ORR) and one high-flux sodium-cooled fast reactor (EBR-II). Upper core plate Baffle bolt Top of bolt head T/F ~0.15 Neutron energy (MeV) 109 10-8 10-7 10-6 10-5 10-4 10-3 10-2 10-1 100 101 102 1010 Fl ux p er u ni t le th ar gy 1011 1012 1013 1014 Figure 3 Typical neutron flux-spectra of internal components of a pressurized water reactor, having a thermal-to-fast neutron ratio smaller by factors of 10–20 than that of typical light water test reactors. Reproduced from Garner, F. A.; Greenwood, L. R. In 11th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors; 2003; pp 887–909. 38 Radiation Damage in Austenitic Steels or enclosing it with a moderator or absorber. Metal hydrides are used in fast reactors to soften the spec- trum, while in mixed-spectrum reactors the thermal- to-fast ratio can be strongly reduced by incorporating elements such as B, Hf, Gd, and Eu. The most pronounced influence on neutron spectra in fission reactors arises from the choices of coolant and moderator, which are often the same material (e.g., water). Moving from heavy liquid metals such as lead or lead–bismuth to lighter metals such as sodium leads to less energetic or ‘softer’ spectra. Use of light water for cooling serves as a much more effective moderator. Counterintuitively, however, this leads to both more energetic and less energetic spectra at the same time, producing a two- peaked ‘fast’ and ‘thermal’ distribution separated by a wide energy gulf at lower fluxes. Such two-peaked spectra are frequently called ‘mixed spectra.’ The ratio of the thermal and fast neutron fluxes in and near such reactors can vary significantly with position and also with time.4 Using heavy water, we obtain a somewhat less efficient moderator that does not absorb neutrons as easily as light water, but one that produces an even more pro- nounced two-peak spectral distribution where the thermal-to-fast neutron ratio can be very large. These spectral differences lead to strong varia- tions between various reactors in the neutron’s ability to displace atoms and to cause transmutation. Depend- ing on the reactor size and its construction details there can also be significant variations in neutron spectra and ‘displacement effectiveness’ within a given reactor and its environs, especially where more energetic neutrons can leak out of the core. Examples of these variations of displacement effec- tiveness for fast reactors are shown in Figures 6 and 7. Compared to fission-derived spectra, there are even larger spectral differences in various fusion or spallation neutron devices. The reader should note the emphasis placed here on flux-spectra rather than simply spectra. If we focus only on light water-cooled reactors for example, there are in general three regimes of neutron flux of relevance to this review. First, there are the relatively low fluxes typical of many experimental reactors that Axial position (cm) Row 4 Row 2 -20 -10 0 4.0 4.5 5.0 5.5 dpa 1022 (E > 0.1) 10 20 Figure 6 Displacement effectiveness values of dpa per 1022 n cm�2 (E > 0.1MeV) across the small core (30 cm tall and �30 cm diameter) of the EBR-II fast reactor, showing effects of neutron leakage to soften the spectrum near the core axial boundaries. Near core center (Row 2) the spectrum and displacement effectiveness are dictated primarily by the use of metal fuel, producing a maximum of �5.2 dpa per 1022 n cm�2 (E> 0.1 MeV). In mid-core Row 4 the radial leakage is just becoming significant. Radial distance from core center (cm) 0 1013 1014 1015 2 5 2 N eu tr on flu x (n eu tr on s p er c m 2 s– 1 ) 5 1016 2 5 4 H 2O a nn ul us H 2O o ut er a nn ul us H 2O 30 0- g P u ta rg et In ne r fu el a nn ul us O ut er fu el a nn ul us C on tr ol r eg io n R em ov ab le b er yl liu m P er m an en t b er yl liu m 8 12 16 20 24 28 32 36 40 44 48 52 56 60 Thermal flux, clean core Thermal flux, 21-day core Total nonthermal flux >0.821 MeV >0.111 MeV Figure 5 Variation in fast and thermal fluxes in HFIR as a function of radial position at mid-core at 85MW, also showing change in thermal population with burn-up (Source: ORNL website). Radiation Damage in Austenitic Steels 39 can produce doses of 10 dpa or less over a decade. Second, there are moderate flux reactors that are used to produce power that can introduce doses as high as 60–100 dpa maximum over a 30–40 year life- time and finally, some high-flux thermal reactors that can produce 10–15 dpa year�1 in stainless steels. Most importantly, fast reactors also operate in the high-flux regime, producing 10–40 dpa year�1. Therefore, the largest amount of published high- dpa data on stainless steels has been generated in fast reactors. Some phenomena observed at high exposure, such as void swelling, have been found to be exceptionally sensitive to the dpa rate, while others are less sensitive (change in yield strength) or essentially insensitive (irradiation creep). These sensitivities will be covered in later sections. For light water-cooled reactors, the various flux regimes need not necessarily involve large differ- ences in neutron spectra, but only in flux. However, the very large dpa rates characteristic of fast reactors are associated with a significant difference in spec- trum. This difference is a direct consequence of the fact that fast reactors were originally designed to breed the fissionable isotope 239Pu from the relatively nonfissile isotope 238U, which comprises 99.3% of natural uranium. In order to maximize the breeding of 239Pu, it is necessary to minimize the unproductive capture of neutrons by elements other than uranium. One strategy used to accomplish this goal is to avoid thermalization of the reactor neutrons, which requires that no low atomic weight materials such as H2O, D2O, Be, or graphite be used as coolants or as moderators. For this purpose, sodium is an excellent coolant with a moderate atomic weight. The use of sodium results BC 6.0 5.0 dpa 4.0 3.0 −100 −75 −50 −25 0 Distance from core midplane (cm) 25 50 75 100 125 150 1022 (E > 0.1) 1 2 3 4 5 FFTF cycles 2 and 3 FFTF core FFTF cycle 10 6 7 Above core 8 Figure 7 Values of dpa per 1022 n cm�2 (E> 0.1MeV) across the much larger core of FFTF for two different fuel/experiment loadings, showing a lesser effect of neutron leakage in larger cores. Note, however, that the in-core values are less than the in-core values of EBR-II, reflecting the softer spectra arising from the use of oxide fuel. Far from the core the displacement effectiveness values are lower, determined primarily by the absence of fuel and the balance of sodium and steel. 40 Radiation Damage in Austenitic Steels in a neutron spectrum that is nominally single-peaked rather than the typical double-peaked (thermal and fast) neutron spectrum found in light water or heavy water reactors. The single-peaked fast reactor spec- trum is significantly less energetic or softer, however, than that found in the fast peak of light water reactors. Depending on the fuel type (metal vs. oxide) the mean energy of fast reactor spectra varies from �0.8 to �0.5–0.4MeV while light water-cooled reactors have a fast neutron peak near �1.2MeV. One consequence of attaining successful breeding conditions is that the spectrum-averaged cross- section for fission is reduced by a factor of 300–400 relative to that found in light water spectra. To reach a power density comparable to that of a light water power-producing reactor, the fast reactor utilizes two concurrent strategies: increases in fissile enrichment to levels in the order of 20% or more, and most importantly, an increase in neutron flux by one or two orders of magnitude. Thus, for a given power density, the fast reactor will subject its structural materials to the punishing effects of neutron bombardment at a rate that is several orders of magnitude greater than that in light water reactors. At the same time, however, the softer ‘fast’ spectrum without thermalized neutrons leads to a significant reduction in transmutation compared to typical light water spectra, at least for stainless steels and nickel-base steels. 4.02.4 Transmutation Issues for Stainless Steels For most, but not all fission-derived spectra, stainless steels are relatively immune to transmutation, espe- cially when compared to other elements such as aluminum, copper, silver, gold, vanadium, tungsten, and rhenium,5,21,24–27 each of which can rapidly become two or three component alloys via transmu- tation arising from thermal or epithermal neutrons. Whereas the properties of these metals are particu- larly sensitive to formation of solid transmutation products, stainless steels in general do not change their composition by significant amounts compared to preexisting levels of impurities, but significant amounts of helium and hydrogen can be produced in fission-derived spectra, however. In stainless steels the primary transmutant changes that arise in various fission and fusion reactor spectra involve the loss of manganese to form iron, loss of chromium to form vanadium, conversion of boron to lithium and helium, and formation of helium and hydrogen gas.4,28 While each of these changes in solid or gaseous elements are produced at relatively small concentrations, they can impact the evolution of alloy properties and behavior. For instance, vanadium is not a starting compo- nent of most 300 series stainless steels, but when included it participates in the formation of carbide 0.14 0.12 0.10 Ni C ro ss -s ec tio n (b ar ns ) 0.08 0.06 0.04 0.02 1 2 4 6 8 10 Fe Ti Cr 20 Energy (MeV) Figure 8 Cross-sections for (n, a) reactions as a function of neutron energy for common elements used in stainless steels. Reproduced from Mansur, L. K.; Grossbeck, M. L. J. Nucl. Mater. 1988, 155–157, 130–147. Nickel dominates the production of helium at higher neutron energies. Radiation Damage in Austenitic Steels 41 precipitates that change the distribution and chemi- cal activity of carbon in the alloy matrix. Carbon plays a number of important roles in the evolution of microstructure1 and especially in grain boundary composition. The latter consideration is very impor- tant in determining the grain boundary cracking behavior, designated irradiation-assisted stress corro- sion cracking (IASCC), especially with respect to the sensitization process.29 The strong loss of manganese in highly therma- lized neutron spectra has been suggested to degrade the stability of insoluble MnS precipitates that tie up S, Cl, and F, all of which are elements implicated in grain boundary cracking.30 Late-term radiation- induced release of these impurities to grain bound- aries may participate in cracking, but this possibility has not yet been conclusively demonstrated. In some high-manganese alloys such as XM-19 manganese serves to enhance the solubility of nitrogen which serves as a very efficient matrix strengthener. In highly thermalized spectra the loss of manganese via transmutation has been proposed to possibly lead to a decrease in the strength of the alloy and perhaps to induce a release of nitrogen from solution to form bubbles.31 The overwhelming majority of published trans- mutation studies for stainless steels and high-nickel alloys steels have addressed the effects of He/dpa ratio on mechanical properties and dimensional instabilities. Much less attention has been paid to the effect of H/dpa ratio based on the long-standing perception that hydrogen is very mobile in metals and therefore is not easily retained in steels at reactor-relevant temperatures. As presented later, this perception is now known to be incorrect, espe- cially for water-cooled reactors. The focus of most published studies concerned the much higher helium generation rates anticipated in fusion spectra (�3–10 appm He/dpa) compared to the lower rates found in fast reactors (�0.1–0.3 appm He/dpa).32 It was later realized that in some highly thermalized test reactors, such as HFIR, very large generation rates could be reached (�100 appm He/dpa), and even in pressurized water reactors the rate could be very high (�15 appm He/dpa).33 In heavy water reactors the rate can be much larger, especially in out-of-core regions.34,35 While some helium arises from (n, a) reactions with thermal and epithermal neutrons interacting with the small amounts of boron found in most stainless steels, the major contribution comes initially from high-energy threshold-type (n, a) reactions with the major alloy components. This type of reac- tion occurs only above high neutron threshold ener- gies (>6MeV). Figure 8 shows that nickel is the major contributor to helium production by (n, a) reactions,36 and thus the helium generation rate scales almost directly with nickel content for a large number of commercial steels. A similar behavior occurs for production of hydrogen by transmutation via high-energy neutrons, where nickel is also the major source of hydrogen compared to other elements in the steel.4,7 In this case, the threshold energy is around 1MeV with 58Ni being the major contributor. This generality concerning nickel as the major source of He and H is preserved in more energetic fusion-derived spectra, although the He/dpa and H/dpa generation rates in fusion spectra are much larger than those of fast reactor spectra. When moving to very energetic spallation-derived neutron and proton spectra, however, the observation that nickel accounts for most of the helium and hydrogen is no longer correct. Iron, nickel, chromium, cobalt, and copper produce essentially the same amounts of helium and hydrogen for energies above �100MeV as shown in Figure 9.6 Another very important helium-generation pro- cess also involves nickel. Helium is produced via the two-step 58Ni(n, g)59Ni(n, a)56Fe reaction sequence.37,38 This sequence operates very strongly in mixed-spectrum reactors. 59Ni is not a naturally occurring isotope and is produced from 58Ni. Thus, this helium contribution involves a delay relative to Natural nickel 58Ni 67.85% 60Ni 26.2% 60Ni 58Ni 59Ni 61Ni 62Ni 64Ni 1.6 1.2 0.8 R at io t o in iti al v al ue 0.4 Thermal fluence (n cm-2) 0.0 1021 1022 1023 1024 6.1% total Figure 10 Transmutation-induced evolution of three nickel isotopes during irradiation in thermalized neutron spectra. Reproduced from Garner, F. A.; Greenwood, L. R. In 11th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors; 2003; pp 887–909. Reproduced from Garner, F. A.; Griffiths, M.; Greenwood, L. R.; Gilbert, E. R. In Proceedings of the 14th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors; American Nuclear Society, 2010; pp 1344–1354. 2500 Inconel 304L 316L 9Cr–1Mo Fe Co Ni Cu 2000 1500 1000H e (a p p m ) 500 0 0 5 10 dpa 15 Figure 9 Measured amount of helium in alloys and pure metals that were irradiated by a mixed spectrum of high energy neutrons and protons produced by 800MeV proton irradiation of tungsten rods. There is some significant uncertainty in the dpa assignment for Inconel 718 at the highest dose. Otherwise the He/dpa ratio appears to be independent of composition. Reproduced fromGarner, F. A.; Oliver,B.M.;Greenwood,L.R.; James,M.R.; Ferguson,P.D.; Maloy, S. A.; Sommer, W. F. J. Nucl. Mater. 2001, 296, 66–82. 42 Radiation Damage in Austenitic Steels that of single-step threshold (n, a) reactions. Since both steps of the sequence involve cross-sections that increase with decreasing energy and the second step exhibits a resonance at 203 eV, the generation rate per dpa in fast reactors increases near the core boundaries and out-of-core areas. It is in thermalized neutron spectra characteristic of light and heavy water-cooled reactors, however, where the 59Ni(n, a) reaction can produce He/dpa generation rates that are significantly larger than those characteristic of fusion-derived spectra. Nickel has five naturally occurring stable isotopes with 58Ni comprising 67.8% natural abundance, 60Ni comprising 26.2%, and�6.1% total of 61Ni, 62Ni, and 64Ni. There is no natural 59Ni or 63Ni at the beginning of radiation. During irradiation in a highly therma- lized neutron spectrum, all nickel isotopes are strongly transmuted, primarily to the next higher isotopic number of nickel. 59Ni has a half-life of 76 000 years and is progressively transmuted to 60Ni while 58Ni is continuously reduced in concentration. Therefore, the 59Ni concentration rises to a peak at a thermal neutron fluence of 4� 1022 n cm�2 where the 59/58 ratio peaks at�0.04 and then declines, as shown in Figure 10. This transmutation sequence in nickel is an exam- ple of the isotopic shift category of transmutation defined earlier. For other elements used to make stain- less steels, there are no consequences to such a shift since the total amount of the element is unchanged and isotope shifts induce no significant consequences. However, in the case of nickel there is an intimate linkage between the displacement and transmutation processes that arises from the isotope shift. The recoil of the 59Ni upon emission of the gamma ray produces only about five displacements per event, and usually is not a significant addition to the displacement dose. However, the isotope 59Ni undergoes three strong reactions with thermal and resonance (�0.2 keV) neutrons, two of which are exceptionally exothermic and can significantly add to the dpa level. These reactions, in order of highest-to-lowest thermal cross-section, are (n, g) to produce 60Ni, followed by (n, a) and (n, p) to produce helium and hydrogen, respectively. Even at relatively low thermal-to-fast neutron ratios, the reaction sequence can produce significant amounts of helium. For example, He/dpa ratios in the order of �3–8 appm dpa�1 can be experienced along the length of a 316 stainless baffle bolt in the baffle-former assembly of a pressurized water Pure nickel in HFIR-PTP 340 keV 1701 dpa 4.8 MeV 62 dpa 0 20 40 P er ce nt ag e in cr ea se 60 80 100 20 40 60 80 Displacements (dpa) neglecting 59Ni (n, a) 56Fe reaction 100 120 140 160 56Fe 4He Figure 11 Increase in dpa arising from the effect of 59Ni to produce helium when pure nickel is irradiated in the HFIR test reactor in the peripheral target position (PTP) where the thermal-to-fast ratio is 2.0. Reproduced from Garner, F. A.; Greenwood, L. R. In 11th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors; 2003; pp 887–909. The rate of dpa acceleration will be increased �3% further if the 59Ni(n, p) and (n, g) reactions are taken into account. Reproduced fromGarner, F. A.; Griffiths, M.; Greenwood, L. R.; Gilbert, E. R. In Proceedings of the 14th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors; American Nuclear Society, 2010; pp 1344–1354. Radiation Damage in Austenitic Steels 43 reactor4,33,39 while comparable rates in fast reactors are in the order of 0.1–0.2 appm dpa�1. In therma- lized spectra the latter two reactions can quickly overwhelm the gas production produced by nickel at high neutron energies. As mentioned previously, the thermal neutron reactions of 59Ni are quite exothermic in nature and release large amounts of energy, thereby causing increases in the rate of atomic displacements, and con- comitant increases in nuclear heating rates. Nuclear heating by elastic collisions with high-energy neu- trons is usually too small to be of much significance. The 59Ni(n, a) reaction releases 5.1MeV, produc- ing a 4.8MeV alpha particle which loses most of its energy by electronic losses, depositing significant thermal energy but producing only �62 atomic dis- placements per each event. However, the recoiling 56Fe carries 340 keV, which is very large compared to most primary knock-on energies, and produces an astounding �1701 displacements per event. The thermal (n, p) reaction of 59Ni produces about one proton per six helium atoms, reflecting the difference in thermal neutron cross-sections of 2.0 and 12.3 barns, and is somewhat less energetic (1.85MeV), producing a total of �222 displacements per event.7,40 In addition, approximately five dis- placed atoms are created by each emission-induced recoil of 60Ni. This reaction occurs at six times higher rate compared to the 59Ni(n, a) reaction, resulting from a thermal neutron cross-section of 77.7 barns. In effect, the dpa rate increases during irradiation due to the three 59Ni reactions even though the neutron flux-spectrum may not change. The major point here is that use of standardized computer codes to calculate dpa does not track shifts in isotopic distribution and therefore will underpredict the dpa level when 59Ni production is an important consideration. A strong example of this time-dependent increase in dpa rate in highly thermalized light water spectra is shown for pure nickel in Figure 11 for a thermal- to-fast ratio of 2.0. Note that the calculated increase in this figure addresses only the 59Ni(n, a) reaction. Additional increases occur as a result of the 59Ni(n, p) and 59Ni(n, g) reactions, resulting in almost doubling of dpa by the three 59Ni reactions before a calculated dose of �40 dpa is attained. Recently, however, an even stronger example of the linkage of the 59Ni transmutation effect and the displacement process has been observed.34,35 In-core thermal-to-fast ratios in heavy water-moderated reactors such as CANDUs are in the order of �10, but far from the core the ratio can be near �1000. Compression-loaded springs constructed of high- nickel alloy X-750 were examined after 18.5 years of operation far from the core and were found to be 44 Radiation Damage in Austenitic Steels completely relaxed. Calculating the 59Ni contribu- tion, it was deduced that full relaxation occurred in �3–4 years rather than the 650–700 years one would predict based on dpa calculated without taking into account the 59Ni contribution. Therefore, in this case 59Ni contributed �95% of the dpa. Additionally, 1100 appm of helium was cal- culated to have been produced at the mid-section of the spring in �3 years, with �20 000 appm helium having been produced when the spring was examined after 18.5 years of exposure. There is another consequence of the 59Ni sequence that causes the temperature to increase during irradia- tion. At the peak 59Ni level reached at 4� 1022 n cm�2, the nuclear heating rates from the energetic (n, a) and (n, p) reactions are 0.377 and 0.023Wg�1 of nickel, significantly larger than the neutron heating level of �0.03Wg�1 of natural nickel. Thus, an increase in nuclear heating of �0.4Wg�1 of nickel must be added to the gamma heating rate at the peak 59Ni level. Fractions of the peak heating rates that are pro- portional to the current 59Ni level should be added at nonpeak conditions. Depending on the nickel level of the steel and the level of gamma heating, which is the primary cause of temperature increases in the interior of thick plates, this additional heating contribution may or may not be significant. Gamma heating is also a strong function of the thermal-to-fast (T/F) neutron ratio and the neutron flux, being �54Wg�1 in the center of the HFIR test reactor where the T/F ratio is �2.0. In pressurized water reactors at the austenitic near-core internals, however, the T/F ratios are lower by a factor of 2–10, depending on location, and the gamma heating rates in the baffle-former assembly are�1–3Wg�1. In this case, an additional 0.4Wg�1 of nuclear heating can be a significant but time-dependent addition to total heating, especially for high-nickel alloys. It should be noted that thermal neutron populations can vary during an irradiation campaign with conse- quences not onlyon 59Ni production but also on gamma heating levels. In PWRs boric acid is added to the water as a burnable poison at the beginning of each cycle. As the 10B burns out the thermal neutron population increases, leading to an increase in gamma heating and transmutation.3,4 Over successive cycles there is a saw- tooth variation of gamma heating rate in the baffle- former assembly and therefore in DT, with the latter reaching values as large as�20 �C in the worst case. Additionally, another concern may arise in that small radiation-induced nickel-rich phases such as g0, Ni-phosphides, and G-phase may become less stable. This concern arises due to cascade-induced dissolution as the 56Fe from the 59Ni(n, a) reaction recoils within the precipitates, thereby altering the phase evolution in thermalized neutron spectra com- pared to nonthermalized spectra typical of fast reac- tors. These precipitates are known to form as a direct result of irradiation and contribute to hardening, swelling, and irradiation creep processes.1 The size of these precipitates at PWR-relevant temperatures (290–400 �C) is often comparable to or smaller than the �80 nm range of the recoiling 56Fe atom. Finally, another significant source of helium can arise from the implantation of energetic helium resulting from collisions with neutrons into the sur- face layers of helium gas-pressurized or gas-cooled components, often involving hundreds and often thousands of appm of injected helium. In gas-cooled reactors helium injection has been investigated as a possible degradation mechanism of alloy surfaces.41 In fast reactor fuel cladding helium was found to be injected into the inner surface, coming from two major sources, ternary fission events (two heavy fis- sion fragments plus an alpha particle) in the fuel and from helium recoiling from the pins’ helium cover gas as a result of collisions with neutrons.42 The injection rates from these two sources of injected helium are slowly reduced during irradia- tion, however, as heavy fission gases build up in the space between the fuel pellet and the cladding. These gases slow down the energetic helium atoms, reducing their energy sufficiently to prevent most of them from reaching the cladding. Helium injection at high levels was also found on the inner surface of helium-pressurized creep tubes.42 Although helium injection tends to saturate in fuel pin cladding with increasing dose, it does not saturate in pressurized tubes due to the lack of increasing fission gases to reduce the range of helium knock-ons in the gas phase. Some studies have cited this early source of helium as contributing to the embrittlement of fuel pin clad- ding and its poor performance during transient heating tests,43 although more recent studies have linked the major mechanism to delayed grain boundary attack by the fission products cesium and tellurium.44,45 4.02.5 Evolution of Radiation- Induced Microchemistry and Microstructure When metals are subjected to displacive irradiation, especially at elevated temperatures, an intricate and coordinated coevolution of microstructure and Radiation Damage in Austenitic Steels 45 microchemistry commences that is dependent pri- marily on the alloy starting state, the dpa rate, and the temperature, and secondarily dependent on vari- ables such as He/dpa rate and applied or internally generated stresses. In general, the starting microstructure and micro- chemistry of the alloy determine only the path taken to the radiation-defined quasi-equilibrium state, and not the final state itself. If an alloy experiences enough displacements, it effectively forgets its start- ing state and arrives at a destination determined only by irradiation temperature and dpa rate. This quasi- equilibrium or dynamic-equilibrium state consists of microstructural components existing at relatively fixed densities and size distributions, but individual dislocations, loops, precipitates, or cavities at any one moment may be growing, shrinking, or even disap- pearing by shrinkage or annihilation. The displacement process produces two types of crystalline point defects, vacant crystalline posi- tions (vacancies) and displaced atoms in interstitial crystalline positions (interstitials). These two defect types are both mobile, but move with different dif- fusional modes and at vastly different velocities, with interstitials diffusing much faster than vacan- cies. Therefore it is obvious that all diffusion-driven processes will be strongly affected by radiation. Both defect types have the ability to recombine with the opposite type (annihilation) or to form agglomerations of various types and geometries. These agglomerations and their subsequent evolution alter both the microstructure and elemental distribu- tion of the alloy. It is important to note that interstitial agglomera- tions are constrained to be two-dimensional, while vacancies can agglomerate in both two-dimensional and three-dimensional forms. This dimensional dis- parity is the root cause of the void swelling phenom- enon covered in a later section. The developing ensemble of various defect agglomerations with increasing dose induces signifi- cant time-dependent and dose-dependent changes in physical and mechanical properties, as well as resulting in significant dimensional distortion. Most importantly, under high displacement rates stainless steels and other alloys are driven far from equilibrium conditions as defined in phase diagrams, affecting not only phase stability but also all physical, mechanical, and distortion processes that involve phase changes in their initiation or evolution. During irradiation, the phase evolution can be significantly altered, both in its kinetics and in the identity and balance of phases that form.46,47 Phases can be altered in their composition from that found in the absence of irradiation, and new phases can form that are not found on the equilibrium phase diagram of a given class of steels. In 300 series stainless steels these new or altered phases have been classified as radiation-induced phases, radiation-modified phases, and radiation-enhanced phases.48–51 These classifica- tions are equally applicable to phases formed in other classes of steel. Radiation-induced alterations of microstructure and microchemistry occur because newdriving forces arise that do not occur in purely thermal environ- ments. The first of these new driving forces is the presence of very large supersaturations of point defects, especially at relatively low irradiation tem- peratures (250–550 �C). Not only are vacancies present in uncharacteristically high levels, thereby accelerating normal vacancy-related diffusional pro- cesses, but interstitials are also abundant. Solutes that can bind with either type of point defect tend to flow down any microstructurally induced gradient of that defect, providing a new mechanism of solute segre- gation referred to as solute drag.52 This mechanism has been proposed to be particularly important for binding of smaller solute atoms such as P and Si, and sometimes Ni, with interstitials. A second new driving force is the inverse Kirkendall effect 53 whereby differences in elemental diffusivity via vacancy exchange lead to segregation of the slowest diffusing species at the bottom of sink- induced vacancy gradients. This mechanism is par- ticularly effective in segregating nickel in austenitic Fe–Cr–Ni alloys at all sinks which absorb vacancies, leading to nickel-rich shells or atmospheres on grain boundaries and other preexisting or radiation- produced microstructural sinks. This type of segre- gation arises because the elemental diffusivities of Fe–Cr–Ni alloys are significantly different, with DCr>DFe>DNi at all nickel levels. 54–57 A third newdriving force results from the action of the other two driving forces when operating on microstructural sinks that are produced only in irra- diation environments. These are Frank interstitial loops, helium bubbles, and voids that may have devel- oped from helium bubbles. Precipitates are often observed to form and to co-evolve on the surface of such radiation-induced sinks. Examples of typical radiation-induced microstructures in stainless steels are shown in Figures 12–15. These microstructural sinks have been implicated as participating in the evolutionary path taken by the precipitates and thereby influencing the microchemical evolution of the matrix.1,58–60 (a) (c) (b)CW 316 SS, thimble tube 70 dpa, 315 ºC 50 nm 50 nm 50 nm CW 316 SS, thimble tube 33 dpa, 290 ºC CW 316 SS, thimble tube 33 dpa, 290 ºC Figure 12 Frank loops observed in a 316 stainless flux thimble from a PWR power reactor (a) 70 dpa, 315 �C and (b) 33dpa, 290 �C imaged edge-on on one set of the four (111) planes using the dark-field relrod technique. Reproduced from Edwards, D. J.; Garner, F. A.; Bruemmer, S. M.; Efsing, P. G. J. Nucl. Mater. 2009, 384, 249–255. The image in (c) is from Frank loops that are slightly inclined to the beam direction imaged using a relrod in the diffraction pattern. G-phase 50 nm Figure 13 Electron micrograph of radiation-induced voids in annealed ‘PCA’ stainless steel irradiated in the ORR water-cooled test reactor at 500 �C to 11dpa. The largest voids have radiation-induced G-phase particles attached to them that are rich in Ni, Si, and Ti. Reproduced from Maziasz, P. J. J. Nucl. Mater. 1989, 169, 95–115. 46 Radiation Damage in Austenitic Steels Minor solute elements such as Si and P have much higher diffusivities than those of Fe, Ni, and Cr and also participate in the segregation process. Addition- ally, these elements increase the diffusivities of the major elements Fe, Ni, and Cr.54 When the solute drag mechanism, operating between interstitials and smaller size Si and P atoms, combines with nickel segregation via the inverse Kirkendall mechanism, phases that are rich in nickel, silicon, or phosphorus often form (g0, G-phase and Ni2P for example), although in 300 series stainless steels these phases do not form thermally. Other phases that are normally stable in the absence of radiation (carbides, intermetallics) can be forced dur- ing irradiation to become enriched in these elements.1 The removal of nickel, silicon, and phosphorus from the matrix by radiation-induced precipitation exerts a large effect on the effective vacancy diffusiv- ity.57,61 On a per atom basis, phosphorus has been Figure 15 Reverse contrast image showing void and line dislocation microstructure in Fe–10Cr–30Mn model alloy irradiated in FFTF fast reactor to 15 dpa at 520 �C. Average void sizes are � 40 nm. Reproduced from Brager, H. R.; Garner, F. A.; Gelles, D. S.; Hamilton, M. L. J. Nucl. Mater. 1985, 133–134, 907–911. Frank loops have unfaulted to produce a line dislocation network whose segments end either on void surfaces or on upper and lower surfaces of the thin microscopy specimen. The voids are coated with ferrite phase due to Mn depletion from their surfaces via the Inverse Kirkendall effect. 50 nm Figure 14 Void swelling (�1%) and M23C6 carbide precipitation produced in annealed 304 stainless steel after irradiation in the reflector region of the sodium-cooled EBR-II fast reactor at 380 �C to 21.7dpa at a dpa rate of 0.84� 10�7 dpa s�1. Reproduced from Garner, F. A.; Edwards, D. J.; Bruemmer, S. M.; et al. In Proceedings, Fontevraud 5, Contribution of Materials Investigation to the Resolution of Problems Encountered in Pressurized Water Reactors; 2002; paper #22. Dislocations and dislocation loops are present but are not in contrast. Radiation Damage in Austenitic Steels 47 shown to exert an even larger effect on the effective vacancy diffusivity57 and its removal into Ni2P and other precipitates has a strong influence on matrix diffusion. Silicon is the next most effective element on a per atom basis. As the effective vacancy diffusion coefficient falls with decreasing matrix levels of Ni, Si, and P, conditions for void nucleation become more favorable. The radiation-induced evolution of diffusional properties has been strongly implicated in determin- ing the transient duration before void swelling accel- erates.1 This evolution often does not necessarily proceed by only one path but occurs in several inter- active stages. Some phases such as nickel phosphides and TiC, especially when precipitated on a very fine scale, are thought to be beneficial in resisting the evolution of nickel silicide type phases.59,62,63 It has been shown, however, that continued radiation- induced segregation eventually overwhelms these phases by removing critical elements such as Ni and Si from solution, causing their dissolution and replacement with nickel-rich and silicon-rich phases that coincide with accelerated swelling.63–65 In high-nickel alloys that normally form the g0 and g00 ordered phases, irradiation-induced segregation processes do not significantly change the identity or composition of the phases, but can strongly change their distribution, dissolving the original distribution but plating these phases out on voids, dislocations, and grain boundaries, with the latter often leading to severe grain boundary embrittlement.66,67 The original dislocation microstructure quickly responds to mobile displacement-generated point defects, increasing their mobility and leading to reductions in dislocation density and distribution in the cold-worked steels most frequently used for fuel cladding and structural components.1 These dislocations are quickly replaced by new micro- structural components, often at very high densities, with two-dimensional interstitial Frank loops first dominating the microstructure, then generating new line dislocations via unfaulting and interaction of loops. In well-annealed alloys there are very few pre- existing dislocations but the same radiation-induced loop and dislocation processes occur, eventually reaching the same quasi-equilibrium microstructure reached by cold-worked alloys. At lower temperatures found in water-cooled test reactors especially, the microstructural features appear to be three-dimensional vacancy clusters or stacking fault tetrahedra and two-dimensional vacancy or interstitial platelets, which are probably also small dislocation loops. These ‘defect clusters’ at temperatures below �300�C are usually too small to be easily resolved via conventional transmission Figure 16 (top) Spiral distortion of 316-clad fuel pins induced by swelling and irradiation creep in an FFTF fuel assembly where the wire wrap swells less than the cladding. Reproduced from Makenas, B. J.; Chastain, S. A.; Gneiting, B. C. In Proceedings of LMR: A Decade of LMR Progress and Promise; ANS: La Grange Park, IL, 1990; pp 176–183; (middle) Swelling-induced changes in length of fuel pins of the same assembly in response to gradients in dose rate, temperature, and production lot variations as observed at the top of the fuel pin bundle. Reproduced from Makenas, B. J.; Chastain, S. A.; Gneiting, B. C. In Proceedings of LMR: A Decade of LMR Progress and Promise; ANS: La Grange Park, IL, 1990; pp 176–183; (bottom) swelling-induced distortion of a BN-600 fuel assembly and an individual pin where the wire swells more than the cladding. Reproduced from Astashov, S. E.; Kozmanov, E. A.; Ogorodov, A. N.; Roslyakov, V. F.; Chuev, V. V.; Sheinkman, A. G. In Studies of the Structural Materials in the Core Components of Fast Sodium Reactors; Russian Academy of Science: Urals Branch, Ekaterinburg, 1984; pp 48–84, in Russian. 48 Radiation Damage in Austenitic Steels electron microscopy and are often characterized as either ‘black dots’ or ‘black spots.’ These dots are generally thought to be very small Frank interstitial loops. The cluster and dislocation loop evolution is fre- quently concurrent with or followed by the loss or redistribution of preexisting precipitates. Most importantly, new radiation-stabilized precipitates at high density often appear with crystal structure and composition that are not found on an equilibrium phase diagram for austenitic steels. As a consequence of these various processes the microstructure at higher doses often develops very high densities of crystallographically faceted, vacuum- filled ‘cavities’ called voids, thought to nucleate on helium clusters formed by transmutation, although residual gases in the steel often help nucleate voids at lower concentrations. Voids have frequently been observed in charged particle irradiations where no helium was introduced. The void phenomenon is not a volume- conservative process and the metal begins to ‘swell’ as the microscopic voids in aggregate contribute to macroscopic changes in dimension, sometimes increasing the metal volume by levels of many tens of percent. Concurrently, the dislocation microstructure responds to the local stress state, moving mass via a volume-conservative process designated irradiation creep. In general, irradiation creep is not a directly damaging process but it can lead to component failures resulting from distortion that causes local blockage of coolant flow or strong postirradiation withdrawal forces. Both swelling and irradiation creep are interrelated and are interactive processes that can produce significant distortions in component dimensions. Figure 16 shows some pronounced examples of such distortion.68,69 Eventually, the microstructural/microchemical ensemble approaches a quasi-equilibrium condition Matrix bubbles 1.6 � 1023 m−3 CW 316 SS, thimble tube 70 dpa, 330 �C CW 316 SS, thimble tube 70 dpa, 330 �C Bubbles on grain boundary 20 nm 20 nm−256 nm UF Figure 17 High densities of nanocavities observed using highly under-focus conditions in a PWR flux thimble tube constructed from cold-worked 316 stainless steel. Reproduced from Edwards, D. J.; Garner, F. A.; Bruemmer, S. M.; Efsing, P. G. J. Nucl. Mater. 2009, 384, 249–255. The irradiation conditions were�70 dpa and 330 �C, producing�600 appm He and 2500 appm H. Note the high density of cavities on the grain boundary. Radiation Damage in Austenitic Steels 49 or ‘saturation’ state, usually at less than 10 dpa for mechanical properties but at higher doses for swelling. As a consequence, the mechanical properties tend to stabilize at levels depending primarily on tempera- ture and to a lesser extent on dpa rate. The two major deformation processes, swelling and irradiation creep, do not saturate but reach steady-state defor- mation rates when quasi-equilibrium microstructures are attained. This coupling of saturation microstruc- ture with steady-state behavior has been character- ized as ‘persistence.’70 Interestingly, the saturation states of each property change are almost always independent of the starting thermal–mechanical state of the material.1,70,71 If irradiation continues long enough, the memory of the starting microstructural state and the associated mechanical properties is almost completely lost. The only deformation-induced microstructural component that succeeds in resisting this erasure process is that of preexisting, deformation-induced twin boundaries. If this quasi-equilibrium is maintained to higher neutron exposure no further change occurs in the steel’s mechanical properties. However, some slowly developing second-order processes are nonsaturable and are often nonlinear. Eventually, these processes force the system to jump toward a new quasi- equilibrium. These new states usually arise from either the microstructural or microchemical evolu- tion, with voids dominating the former and the latter involving continued segregation, continued transmutation, or a combination of these factors.70–72 A number of such late-stage changes in quasi- equilibrium state are discussed later in this paper. 4.02.6 A Cross-Over Issue Involving Radiation-Induced Microstructural Evolution and Transmutation Recently, it been discovered that significant levels of hydrogen can be stored in bubbles and voids in both stainless steels and pure nickel when the hydrogen is cogenerated with helium, especially in light water spectra where there are also environmental sources of hydrogen.73–75 It was shown in these studies that this phenomenon is a direct result of the 59Ni nuclear reactions. Previously, it was a long-standing percep- tion that such storage could not occur at reactor- relevant temperatures. The retained hydrogen levels are in significant excess of the levels predicted by Sievert’s Law and appear to be increasing with both cavity volume and neutron fluence. Since these gases are known to assist in nucleation and stabilization of cavities, it is expected that the nonlinear 59Ni reactions discussed earlier may lead to a rapidly developing, nonlinear, cavity-dominated microstructure in stainless steels irradiated at temperatures characteristic of pressur- ized water reactors. Figure 17 presents such a microstructure observed in a PWR flux thimble tube (cold-worked 316 stain- less steel) at �70 dpa and 330 �C.76 There is a very high density (>1017 cm�3) of nanocavities with dia- meters 50 Radiation Damage in Austenitic Steels under well-focused imaging conditions and can only be imaged using very large levels of under-focus. This implies that previous studies on similar mate- rials may have overlooked such cavity-dominated structures. When this specimen and near-identical specimens were subjected to slow strain rate testing after irradi- ation, the fracture surface was indicative of �100% intergranular stress corrosion cracking (IGSCC), with lower doses and gas levels producing propor- tionally less IGSCC.77 As hydrogen is known to be a contributor to grain boundary cracking, it appears plausible that hydrogen storage may accelerate the cracking process and that higher exposure will lead to an increasing susceptibility to cracking. This issue may therefore become increasingly important as PWRs previously licensed for 40 years are being considered for life extension to 60 and possibly 80 years. E E ng . s tr es s (M P a) (c) 1000 10.7 dpa 3.64 dpa 0. 0 1.36 dpa 2.53 dpa 800 600 400 200 0 0.0 0.2 0.4 Eng 1000 316 SS tested at RT 800 E ng . s tr es s (M P a) 600 400 200 (a) 0 0.0 0.2 0.4 0.6 Eng. strain 0.8 Unirradiated 0.0001 dpa 0.001 dpa 0.01 dpa 0.1 dpa 0.78 dpa 1.0 Figure 18 Engineering stress–strain curves for irradiated auste mixed spectrum reactor at 60–100 �C and tested at 25 �C, (b) an 288 �C, and (c) annealed EC316LN irradiated in the LANSCE spa at 25 �C. Reproduced from Kim, J. W.; Byun, T. S. J. Nucl. Mate 4.02.7 Radiation-Induced Changes in Mechanical Properties Long before the onset of significant phase evolution or void swelling is observed, the first manifestation of the radiation-induced microstructural/microchemical evolution appears in changes of the mechanical prop- erties. As shown in Figure 18 the stress–strain dia- grams of stainless steels begin to change significantly even at very low dpa levels. The strength of the alloy increases, the elongation decreases, and there is a progressive decrease in work-hardening. This behav- ior is dependent somewhat on test temperature but is not very sensitive to neutron spectrum. Movement of dislocations in metals during deformation following irradiation is impeded by the microstructural components produced by radiation (dislocations, dislocation loops, voids, bubbles, preci- pitates) and therefore the strength of annealed steel C316LN tested at RT 86 dpa .45 dpa 0.0 dpa . strain 0.6 0.8 1.0 316 SS tested at 288 �C (b) Unirradiated 0.01 dpa 0.1 dpa 1000 800 E ng . s tr es s (M P a) 600 400 200 0 0.0 0.2 0.4 0.6 Eng. strain 0.8 1.0 nitic stainless steels: (a) annealed 316 SS irradiated in HFIR nealed 316 SS irradiated in HFIR at 350 �C and tested at llation neutron and proton spectrum at 60–100 �C and tested r. 2010, 396, 10–19. Radiation Damage in Austenitic Steels 51 increases. The strength increase usually saturates at relatively low exposure levels ( 1100 1000 900 800 700 600 Y ie ld s tr en gt h (M P a) 427 �C 371 �C 483 �C 538 �C 593 �C 649 �C 704 �C 760 �C 816 �C 500 400 300 200 100 0 0 1 2 3 4 Neutron fluence (n cm–2) (E > 0.1 MeV) 5 6 7 8 9 10 � 1022 Figure 21 Evolution of yield strength in 20% cold-worked 316 stainless steel irradiated in EBR-II over a wide range of temperatures. Reproduced from Garner, F. A.; Hamilton, M. L.; Panayotou, N. F.; Johnson, G. D. J. Nucl. Mater. 1981, 103 and 104, 803–808. 600 400 200 0 600 400 200 0 800 600 400 200 0 0 1 2 3 4 5 Neutron fluence (1022n cm–2) Y ie ld s tr en gt h (M P a) 20% Cold-worked 20% Cold-worked 20% Cold-worked Annealed Annealed Annealed 650 �C 538 �C 427 �C 6 7 8 9 10 Figure 22 Influence of temperature and neutron exposure on evolution of yield strength in both annealed and 20% cold-worked AISI 316 irradiated in EBR-II, showing that the saturation strength level is independent of starting condition, converging at doses of 5–15 dpa. Reproduced from Garner, F. A.; Hamilton, M. L.; Panayotou, N. F.; Johnson, G. D. J. Nucl. Mater. 1981, 103 and 104, 803–808. 52 Radiation Damage in Austenitic Steels (27–58 appmdpa�1) and 395 and 450 �C in EBR-II at very low He/dpa ratios (0.7–1.2 appmdpa�1).85 Note that there are very significant differences in hardening observed between the two experiments and that the differences arose primarily from a very large difference in cavity density, a difference that was too large to be explained in terms of helium content alone. It was later shown that the ORR experiment suffered a very large number (237 over 2 years) of unrecognized negative temperature set- backs of 1–2 h, with decreases varying from 50 to 500 �C.86 Even though the total dpa accumulated during these setbacks was only �1% of the total dose, the frequent bloom of high densities of small Frank loops at lower temperatures provided a very large periodic increase in nucleation sites for helium bubbles on the new Frank loops that significantly strengthened the matrix. The loops could subse- quently dissolve but the bubbles could not. In addition to temperature, the most prominent irradiation variable is the dpa rate and it is known that the microstructural densities, especially Frank loops and voids, are known to increase in concentra- tion as the dpa rate increases. Various radiation-stable phases such as w0 are also known to be flux-sensitive, while other phases such as carbides and intermetallics are more time-sensitive.1 Thus, it is not surprising that some sensitivity to dpa rate might be observed in strength properties, as 1200 Test temperature = Irradiation temperature 40 20 10 5 2 1 0.5 1000 800 600 400 Y ie ld s tr en gt h (M P a) U ni fo rm e lo ng at io n (% ) 200 0 1 2 3 4 585–625 �C 525–565 �C 525–565 �C 450–500 �C 215–245 �C 400 �C 400 �C 5 ´ 1026 0 1 2 3 4 5 ´ 1026 215–245 �C Neutron fluence (n m-2) E > 0.1 MeV Neutron fluence (n m-2) E > 0.1 MeV Figure 23 Neutron-induced changes in tensile properties of annealed 1.4988 stainless steel irradiated in the DFR fast reactor. Reproduced from Ehrlich, K. J. Nucl. Mater. 1985, 133–134, 119–126. Ductility declines as strength increases. 2 4 0 0 10 20 30 40 50 400 800 12006 8 100 1200 20% CW 316 25% CW PCA 20% CW 316 25% CW PCA800 Yield stress (MPa) Y ie ld s tr es s (M P a) dpa 400 SA 316L U ni fo rm e lo ng at io n (% ) SA PCA SA 316L SA PCA 0 Figure 24 Strengthening and ductility loss observed in two stainless steels irradiated in the HFIR, HFR, and R2 mixed spectrum reactors at 250 �C at He/dpa ratios ranging from 10 to 35appmdpa�1. Note that both annealed and cold-worked (CW) steels quickly converge to the same elongation levels, while convergence of strength is not developing as quickly. Reproduced from Elen, J. D.; Fenici, P. J. Nucl. Mater. 1992, 191–194, 766–770. Radiation Damage in Austenitic Steels 53 suggested by the behavior shown in Figure 28 where both the transient rate of strength rise and saturation strength appear to increase with increasing dpa rate. Unfortunately, this figure does not represent a single variable comparison, and by itself is not sufficiently convincing evidence of flux sensitivity. The data shown in Figure 29 is much closer to a single variable comparison, indicating that the transient rise may or not be somewhat flux-sensitive, depending on the details of the microstructural evolution of each alloy. The authors of this study used microscopy to confirm the microstructural origins of the observed differ- ences of behavior as a function of dpa rate. More recently, Chatani and coworkers showed that at relatively low irradiation temperatures char- acteristic of boiling water reactors, the radiation- induced increments in strength of 304 stainless steel increased by the 1/4 power of the increase in dpa rate.87 It was demonstrated that the black-spot micro- structure dominated the strengthening. It was also shown that the concentration of black spots varied with the square root of the flux as expected, and it is known that hardening varies with the square root of the loop density, thereby producing a fourth-root dependence. Thus, in the absence of any significant microchemical or phase stability contributions, it 25 Ni 45 Ni 25 Ni + 0.04 P 25 Ni + 0.04 P 45 Ni Annealed dpa With 59Ni Without ~0.5 and ~15 appm He per dpa Annealed To ta l e lo ng at io n (% ) Cold-worked Cold-worked 25 NiY ie ld s tr en gt h (M P a) 365 �C 1000 800 600 400 200 0 40 30 20 10 0 0 10 20 0 10 20 0 10 20 30 Figure 25 Influence of starting state, composition of isotopically doped alloys and He/dpa ratio on changes in mechanical properties produced during isothermal irradiation at 365 �C in FFTF. Reproduced from Garner, F. A.; Hamilton, M. L.; Greenwood, L. R.; Stubbins, J. F.; Oliver, B. M. In Proceedings of 16th ASTM International Symposium on Effects of Radiation on Materials; ASTM STP 1175; 1992, pp 921–939. 54 Radiation Damage in Austenitic Steels appears that radiation-induced strengthening is affected by dpa rate but not very strongly. The loss of ductility proceeds in several stages, first involving convergence of the yield and ultimate strengths as shown in Figures 29 and 30, such that a loss of work-hardening occurs and very little uniform elongation is attained. As the irradiation pro- ceeds, there is a progressive tendency toward flow localization followed by necking. As seen in Figure 31 the failure surface shows this evolution with increas- ing dose. The flat faces observed at highest exposure in Figure 31 are often referred to as ‘channel fracture’ but they are not cleavage faces. They are the result of intense flow localization, resulting from the first moving dislocations clearing a path of radiation- produced obstacles, especially Frank loops, and thereby softening the alloy along that path. It is not possible to remove the voids by channeling but the distorted voids provide a microstructural record of the flow localization as shown in Figure 32. Linkage of the elongated voids is thought to contribute to the failure. Such a failure surface might best be characterized as ‘quasi-embrittlement’, which is a suppression of uniform deformation, differentiating it from true embrittlement, which involves the complete sup- pression of the steel’s ability for plastic deformation. This distinction is made because under some con- ditions quasi-embrittlement can evolve into true embrittlement. The tendency toward quasi-embrittlement grows with increasing swelling but the alloy is actually softening with increasing swelling rather than hard- ening. As shown in Figure 33 brittle fracture (defined as strength reduction with zero plasticity) of a Fe–18Cr–10Ni–Ti stainless wrapper in BOR-60 at 72 dpa maximum was observed at positions where peak swelling occurs.88 Some decrease of strength is 25 Ni To ta l e lo ng at io n (% ) Y ie ld s tr en gt h (M P a) 25 Ni 45 Ni 45 Ni 495 �C 800 600 Original series Original series Isothermal repeat series Isothermal repeat series 400 200 0 40 30 20 10 0 0 20 40 20 dpa 400 20 40 600 25 Ni + 0.04 P 25 Ni + 0.04 P ~0.5 and ~5.0 appm He per dpa With 59Ni Without Figure 26 Comparison of isothermal and nonisothermal behavior on convergence behavior. The original target temperature of 495 �C was maintained for some time but thereafter there was a large, relatively brief over-temperature event, followed by a prolonged and significant under-temperature event. Reproduced from Garner, F. A.; Hamilton, M. L.; Greenwood, L. R.; Stubbins, J. F.; Oliver, B. M. In Proceedings of 16th ASTM International Symposium on Effects of Radiation on Materials; ASTM STP 1175; 1992, pp 921–939. When the target temperature was reestablished in the second and third irradiation segments the mechanical properties returned to the isothermal destination. Radiation Damage in Austenitic Steels 55 observed with increasing irradiation temperature, but the primary strength reduction for specimens tested at the irradiation temperature arises from the magnitude of swelling. Testing at temperatures below the irradia- tion temperature (e.g., 20 �C) demonstrates the same dependence on swelling and irradiation temperature, but the strength and plasticity values are higher. As expected, the strengths for tests conducted at 800 �C are uniformly much lower than that observed at lower temperatures, but there is an absence of any relation- ship between strength and swelling at this temperature. As shown in Figure 34 failure surfaces at high swelling levels exhibit transgranular cup-cone mor- phology where failure proceeded by micropore coa- lescence arising from stress concentration between deforming voids.88 Similar fracture morphology has been observed in studies on other stainless steels.1 Although voids and bubbles initially serve to harden the microstructure,78 large swelling levels allow previously second-order void effects to become dominant.1,88,89 One of these second-order effects is the strong decrease of elastic moduli at high swelling levels. All three elastic moduli decrease initially at �2% per each percent of void swelling.90–93 At >10% swelling this leads to significant reduction in strength. As a consequence, the slope of the elastic region (Young’s modulus) of the stress–strain curve decreases, and more importantly, the barrier strengths of all sinks decrease as the shear modulus likewise decreases. Therefore, the yield and ultimate strengths decrease with increasing swelling, even though the elongation strongly decreases. Similar behavior has also been observed in pure copper.94 MFE-4 experiment in ORR AD-1 experiment in EBR-II 450 �C 43 23 25 23 nm 40 35 9.4 Ni 7 15 20 Cr 3525 45 1023 1022 1021 1020 1019 20 30 40 50 Nickel (wt%) C av ity d en si ty (m -3 ) 1024 5.6 5.5 3.4 1.9 1.9 nm 2.92.42.32.1 500 �C 400 �C 395 �C MFE-4 AD-1 300 800 700 600 500 400 300 200 100 0 400 Temperature (�C) A ve ra ge D Y S (M P a) 500 600 Figure 27 Comparison of hardening of Fe-YCr-XNi ternary alloys observed in the MFE-4 experiment in ORR at �13 dpa and the AD-1 experiment in EBR-II at �10 dpa. Reproduced from Hamilton, M. L.; Okada, A.; Garner, F. A. J. Nucl. Mater. 1991, 179–181, 558–562; Garner, F. A.; Sekimura, N.; Grossbeck, M. L.; et al. J. Nucl. Mater. 1993, 205, 206–218. Higher levels of hardening in ORR arise from a refinement and elevation of cavity density arising from frequent negative temperature excursions at high He/dpa rates. Mean cavity sizes are shown next to each data point. Fluence (dpa) Phénix Rapsodie 0 200 400 600 800 10 20 30 40 0. 2% p ro of s tr es s (M P a) 50 Figure 28 Differences in strength change exhibited by annealed 316 stainless steel after irradiation at 390 �C in the PHENIX and RAPSODIE fast reactors. Dupouy, J. M.; Erler, J.; Huillery, R. In Proceedings International Conference on Radiation Effects in Breeder Reactor Structural Materials, Scottsdale; The Metallurgical Society of AIME: New York, 1977, pp 83–93. Phénix operated at a displacement rate that was �three times higher than that of RAPSODIE. 56 Radiation Damage in Austenitic Steels The nature of the void-related failure changes from quasi-embrittlement to true embrittlement for tests at or near room temperature, demonstrating another example of a late-term second-order process growing to first-order importance at higher swelling levels. Hamilton and coworkers observed that above �10% swelling the previously established saturation strength level of 316 stainless steel suddenly increased very strongly in room temperature tensile tests.95 Similar results were observed in Russian steels.96,97 As shown in Figures 35 and 36 the failure surfaces in such tests had rotated from the expected 45� (relative to the stress axis) to 90� as swelling approached 10%, indi- cating complete brittle failure, as also indicated by the fully transgranular nature of the failure surface. Con- currently, the ductility vanished and the tearing mod- ulus plunged to zero, indicating no resistance to crack propagation. Once a crack has initiated it then propa- gates completely and instantly through the specimen. Neustroev and coworkers observed such failures in Russian steels that are subject to greater amounts of precipitation and determined that the critical micro- structural condition was not defined solely by the level of swelling, but by the obstacle-to-obstacle distance of the void-precipitate ensemble, indicating that stress concentration between obstacles was one contributing factor.96 However, it was the progressive segregation of nickel to increasing amounts of void surface and the concurrent rejection of chromium from the sur- faces that precipitated the rather abrupt change in failure behavior.1,95 This late-term void-induced micro- chemical evolution induces a martensite instability in the matrix, as evidenced by the failure surface being completely coated with alpha-martensite.95 Exposure (dpa) 0.1 (a) 100 300 500 S tr en gt h (M P a) 700 Yield Yield Ultimate Ultimate AISI 304 AISI 316 1 10 100 0.75 E, MeV Ti, �C odf, dpa s–1 7.9 ´ 10-7 3.9 ´ 10-7 1.8 ´ 10-7 0.8 ´ 10-7 0.6 ´ 10-7 1.5 ´ 10-7 0.53 0.29 0.29 0.19 0.17 392 376 373 426 371 371 Exposure (dpa) 0.1 (b) 100 300 500 S tr en gt h (M P a) 700 1 10 100 0.76 E, MeV Ti, ºC odf, dpa s–1 8.4 ´ 10-7 5.1 ´ 10-7 2.3 ´ 10-7 1.0 ´ 10-7 0.6 ´ 10-7 1.9 ´ 10-7 0.63 0.38 0.35 0.21 0.17 399 378 374 424 372 371 Figure 29 Strength changes observed in annealed 304 and 316 stainless steels irradiated in EBR-II at 371–426 �C and tested at 385 �C. Reproduced from Brager, H. R. Blackburn, L. D.; Greenslade, D. L. J. Nucl. Mater. 1984, 122–123, 332–337. Microscopy showed that the dependence of microstructure on displacement rate was consistent with themacroscopic behavior exhibited by each alloy. In AISI 316, the flux dependence of precipitation canceled the opposite dependence of other microstructural components. Neutron fluence (E > 0.1 MeV, n cm–2) Ultimate strength 370 �C Yield strength S tr en gt h (M P a) 0 0 200 400 600 800 1000 1200 2 4 6 8 10 ´ 1022 Figure 30 Convergence of ultimate and yield strengths of annealed 304 stainless steel irradiated in EBR-II and tested at 370 �C. Reproduced from Holmes, J. J.; Straalsund, J. L. In Proceedings of International Conference: Radiation Effects in Breeder Reactor Structural Materials; 1977; pp 53–63. Radiation Damage in Austenitic Steels 57 The abrupt jump in strength just before failure observed by Hamilton and coworkers is the result of a stress-induced blossoming of a high density of small, thin, epsilon-martensite platelets, as seen in Figure 37. These platelets are essentially stacking faults that form under stress as a result of the influ- ence of both falling nickel level and low deforma- tion temperature to decrease the stacking fault energy of the matrix.1 When encountered by the advancing crack tip, the epsilon-martensite is con- verted to alpha-martensite in the strain field ahead of the crack, providing a very brittle path for further cracking. The correlation between void swelling and both quasi-embrittlement and true embrittlement is observed not only in slow tensile tests (Figures 36, 38, and 39) but also in Charpy impact tests as shown in Figure 39. Figures 40–44 present examples of swelling-induced failures in components experien- cing a wide range of physical insults. The example of Porollo et al. in Figure 44 (top) is particularly noteworthy in that it results from significant swelling at �335 �C, a temperature earlier thought not to produce significant amounts of swelling. If there are no physical insults experienced by the component during irradiation, the continued segregation of nickel to void surfaces and the con- current rejection of chromium can lead to strong changes in composition in the matrix during irradia- tion, pushing the matrix toward ferrite rather than martensite at higher temperatures, especially for steels with nickel content of 100 80 60 40 20 0 0 1 2 3 4 5 Fluence, ´1022 n cm–2 (E > 0.1 MeV) Uniform elongation Uniform elongation Proportional elastic limit Proportional elastic limit Unirradiated plastic dimpling mechanism Intermediate Channel fracture2.8 ´ 1022 n cm–2 10.7 ´ 1022 n cm–2 Yield strength Yield strength Symbols Open-5C3 CRT Crossed-5A3 CRT Closed-3A1 SRT EBR-II 304 SS Irradiated at 700 �F Tested at 700 �FSt re ng th (k si ) U ni fo rm e lo ng at io n (% ) 6 7 8 9 10 11 0 1 2 3 Figure 31 Increase in strength, loss of ductility, and change in failure mode observed during tensile testing in annealed 304 safety and control rod thimbles (SRT and CRT) after irradiation at �370 �C in EBR-II. Reproduced from Fish, R. L.; Straalsund, J. L.; Hunter, C. W.; Holmes, J. J. In Effects of Radiation on Substructure andMechanical Properties of Metals and Alloys; ASTM STP 529; 1973; pp 149–164. Figure 32 Intense flow localization manifested as shearing of voids below a ‘channeled’ failure surface in a 304 steel tensile specimen at 40 dpa and �400 �C when tested at 370 �C. There is 100–200% strain in the 0.05mm wide deformation band. Reproduced from Fish, R. L.; Straalsund, J. L.; Hunter, C. W.; Holmes, J. J. In Effects of Radiation on Substructure and Mechanical Properties of Metals and Alloys; ASTM STP 529; 1973; pp 149–164. The swelling was �5% in this specimen. 1000 800 600 10, 6,5 Ttest= 20 �C Ttest= 800 �C Ttest= Tirr. 18, 18, 21, 17, 15, 25, 23, 6,2 0,8 0,6 2,7 400 200 0 11, 22, 26, 7,8 1,8 0 50 100 Position from core central plane (mm) U lti m at e te ns ile s tr en gt h (M P a) X-strength without elongation 150 200 250 300−50 Figure 33 Ultimate tensile strength of Fe–18Cr–10Ni–Ti stainless steel wrapper specimens irradiated in the BOR-60 fast reactor to a maximum dose of 72 dpa. Reproduced from Neustroev, V. S.; Garner, F. A. J. Nucl. Mater. 2009, 386–388, 157–160. Three tensile test temperatures are shown: closed circle, 20 �C; open circle, 450–550 �C; triangle; 800 �C. Swelling values in % are given near the points. 58 Radiation Damage in Austenitic Steels regained not because the steel has softened, but because it becomes exceptionally strong and hard- ened during deformation. As a consequence, the steel has lost the ability to neck. 10 mm Figure 34 Fracture surface of Fe–18Cr–10Ni–Ti stainless steel specimen at a swelling level of 26%. Reproduced from Neustroev, V. S.; Garner, F. A. J. Nucl. Mater. 2009, 386–388, 157–160. Micrograph corresponds to open circle at 70mm position in Figure 33. Ti = 460 �C Tt = 20 ft = 15.5 � 1022n cm–2 Ti = 385 �C ft = 12.8 � 1022n cm–2 Ti = 460 �C ft = 15.5 � 1022n cm–2 et = 1.9 Tt = 20 et = 5.2 Tt = 460 et = 7.2 100 mm Figure 35 Fractographs of failure surfaces of 20% cold-work Reproduced from Hamilton, M. L.; Huang, F. H.; Yang, W. J. S.; International Symposium (Part II) Influence of Radiation on Mate change of fracture mode from channel fracture when tested at 2 Radiation Damage in Austenitic Steels 59 Instead of necking, a moving wave of deformation is initiated at the first attempted necking point. The wave front then travels nearly the full length of the gage section. Initially, there is a local deforma- tion in the order of 40–45%, but as the wave moves forward it leaves a relatively uniform local defor- mation in its wake. Everywhere behind the wave front there is measured 30–35% volume percent of martensite, as shown in Figures 45 and 46. The martensite is not only a product of the wave, but also the cause of the wave. Deformation-induced martensite resists further necking and forces the deformation to be displaced to the adjacent lesser deformed material. The mechanisms that cause the late-term onset of martensite instability have not yet been determined. A property of important engineering interest is the fracture toughness Jc. While the fracture toughness of various unirradiated stainless steels can be quite �C % 5 �C % �C % 10 mm ed 316 specimens cut from an FFTF duct at high exposure. Garner, F. A. In Effects of Radiation on Materials: 13th rial Properties; ASTM STP 956; 1987; pp 245–270. Note 05 and 460 �C to brittle fracture when tested at 20 �C. A ng le o f f ra ct ur e Test at 20 �C Test at irradiation temperature 1.0 0.8 s 0 .2 s U TS 0.6 90 60 30 3020 Swelling (%) 100 Test at 20 �C Test at 20 �C Test at irradiation temperature Test at irradiation temperature 800 15 10 5 U E (% ) 0 600 400 200 U TS (M P a) 0 0 10 20 Nil ductility 30 40 Swelling (%) 0 10 20 30 40 Swelling (%) , Figure 36 Influence of swelling on fracture properties during tensile testing of an annealed Fe-18Cr-10Ni-Ti steel irradiated in BOR-60 at 400–500 �C. Neustroev, V. S.; Shamardin, V. K. Atomnaya Energiya 1990, 71(4), 345–348, in Russian. Note that softening and rotation of fracture surface by voids is observed at both room and elevated temperatures. 60 Radiation Damage in Austenitic Steels different, it appears that all austenitic steels studied undergo the same general evolution in toughness during irradiation. Mills has shown that three regimes of evolution occur.103,104 The first regime involves a low-dose threshold exposure range ( 422 (a) (b)16049 0.1µm 0.1µm 16185 200 Figure 37 (a) Bright field image of voids and deformation bands observed in a highly embrittled 20% cold-worked 316 hexagonal duct. (b) Dark field image showing a high density of thin stacking fault platelets of epsilon-martensite on one of the four sets of close-packed planes. Reproduced from Hamilton, M. L.; Huang, F. H.; Yang, W. J. S.; Garner, F. A. In Effects of Radiation on Materials: 13th International Symposium (Part II) Influence of Radiation on Material Properties; ASTM STP 956; 1987; pp 245–270. 1000 100 OKh16N15M3B Some ductility retained Nil ductility 10 U TS (M P a) 1000 100 Kh18N10T 10 0.01 0.1 1.0 Swelling (%) 10 1 Figure 38 Influence of swelling on ultimate tensile strength of 0Khl6N15M3B cladding and Kh18N10T hexagonal ducts irradiated in BOR-60. Tests were performed at the irradiation temperature, using specimens cut from regions of maximum swelling. Reproduced from Neustroev, V. S.; Shamardin, V. K. Phys. Met. Metallogr. 1997, 83(5), 555–560. The two steels develop loss of strength with swelling differently, probably reflecting the very different precipitate structures of the two steels. Radiation Damage in Austenitic Steels 61 4.02.8 Radiation-Induced Changes in Dimension One of the most challenging engineering conse- quences of neutron irradiation is the development of dimensional instability, whereby a structural com- ponent can shrink or grow in volume and where it can be distorted in shape, often with both processes occurring at the same time. There are two major categories of such changes: conservative of volume and nonconservative of volume. A distinction can also be made between processes that distribute the resulting strains isotropically or anisotropically. Additionally, a further distinction can be made con- cerning whether the process to the first-order is stress-driven or not, or whether it is stress-sensitive to the second-order. Depending on the crystal structure there are a variety of such distortion processes, some more prominent than others in a given crystal system. For austenitic stainless steels the phenomenon of radiation-induced growth (volume-conservative, an- isotropic distribution of strains in the absence of stress) is not an issue, whereas for hexagonal close packed alloys based on zirconium and rhenium growth is often a dominant process.9,106 Austenitic steels also Test temperature = 180 �C Test temperature = 180 �C 30 20 Total elongation 80 60 40 Wrapper #1–low dose Wrapper #2–high dose 20 0 10 0 0 5 Swelling (%) Uniform elongation E lo ng at io n (% ) To ta l a b so rb ed e ne rg y (J c m −2 ) 10 0 5 Swelling (%) 10 Figure 39 (Left) Correlation between ductility loss and swelling in several heats of irradiated Ti-modified steels in PHENIX. At�5% swelling the total and uniform elongations converge and by�10% no ductility remains. (Right) Correlation of swelling and embrittlement in Charpy impact tests of cold-worked Ti-modified 316 steel irradiated in PHENIX. Reproduced from Fissolo, A.; Cauvin, R.; Hugot, J. P. Levy, V. In Effects of Radiation on Materials: 14th International Symposium; STP 1046; 1990; Vol. II; pp 700–713. Figure 40 Failure during mounting in a vise of severely void-embrittled 316 stainless steel creep tube irradiated in the EBR-II fast reactor to 130 dpa at �400 �C with a hoop stress of 276MPa. Reproduced from Porter, D. L.; Garner, F. A. J. Nucl. Mater. 1988, 159, 114–121. Swelling at the initial failure point was �14%. Figure 41 Void-induced embrittlement of an annealed 304 steel EBR-II assembly duct after �54 dpa at �400 �C. Reproduced from Flinn, J. E.; Krajcinovic, D.; Phipps, R. D.; Franklin, D. G.; Miller, S. C. Evaluation of Ex-Reactor Loading Event on High-fluence EBR-II Control-rod Thimble 5E3, ANL/EBR-068, February 1973. The duct broke during routine handling in the hot cell. 62 Radiation Damage in Austenitic Steels are not very prone to significant transmutation- induced changes in lattice parameter as sometimes observed in alloys based on rhenium and vana- dium.106,107 See also Chapter 4.01, Radiation Effects in Zirconium Alloys. Stainless steels experience three general cate- gories of radiation-induced strain processes. These are precipitation-related strains, void swelling, and irradiation creep. In general, these three processes are not fully independent but are interrelated and often synergistic. 4.02.8.1 Precipitation-Related Strains Stainless steels undergo an evolution of phase struc- ture at reactor-relevant temperatures, even in the absence of radiation. These changes involve the for- mation of various carbides, later followed by various intermetallic phases.1,108 This evolution is accompa- nied by net changes in average lattice parameter arising from differences in partial molar volume of elements when passing from one phase to another. 53 dpa 52 dpa 34 dpa max 27.8% 29.8% 14% swelling By-97 By-92 U-796 Figure 42 Severe embrittlement and failure in three BOR- 60 reflector assembly ducts. The ducts were made of annealed X18H10T, the Russian equivalent of 321 steel. Reproduced from Neustroev, V. S.; Ostrovsky, Z. E.; Teykovtsev, A. A.; Shamardin, V. K.; Yakolev, V. V. In Proceedings of 6th Russian Conference on Reactor Materials Science; 11–15 September 2000, Dimitrovgrad, Russia, in Russian. The maximum swelling values (from left to right) were 27.8, 29.8, and 14%. Failure was the result of high withdrawal loads arising from both swelling and bending, the latter a consequence of radial dpa gradients in the reflector. Figure 43 Failure of 20% cold-worked D9 (Ti-modified 316) cladding during routine handling. Failure occurred where 90 dpa was attained at �460 �C in FFTF, producing �32% swelling. Reproduced from Makenas, B. J.; Chastain, S. A.; Gneiting, B. C. , ‘‘Dimensional Changes in FFTF Austenitic Cladding and Ducts, Westinghouse Hanford Company Report WHC-SA-0933VA, Richland WA, 1990. Fuel was lost from the open section. 50 50 100 Annealed tubes 200 0 50 100 Cold-worked tubes 200 200 MPa 200 MPa50 50 100 Annealed tubes 200 200 M Figure 44 EI-847 pressurized tubes irradiated to 75 dpa, 330–342 �C (top) and 82 dpa, 365 �C, (bottom) in BN-350. Reproduced from Porollo, S. I.; Vorobjev, A. N.; Konobeev, Yu. V.; Dvoraishin, A. M.; Krigan, V. M.; Budylkin, N. I.; Mironova, E. G.; Garner, F.A. J. Nucl. Mater. 1998, 258–263, 1613–1617. All tubes lost pressure, either by cracking or by completely failing during removal from their canister. Before breaking, the tubes were also bent by irradiation creep due to swelling-induced interaction with the top of the canister. Swelling of 6.2%wasmeasured in the zero stress, annealed tube and 11.2% in the cold-worked zero stress tube. ~30–35% of martensite Figure 45 Deformation at room temperature of theRussian analog of AISI 321 following irradiation in BN-600 to 55 dpa at 310 �C. Distortion of painted circular dots shows where the deformation wave has passed, moving toward the left. The specimen was cut from a hexagonal duct of a fuel assembly. Reproduced from Gusev, M. N.; Maksimkin, O. P.; Garner, F. A. J. Nucl. Mater. 2010, 403, 121–125. Radiation Damage in Austenitic Steels 63 The resulting macroscopic strains are sometimes very counterintuitive, however, especially with re- spect to their sign. For example, formation of the less dense carbide phases leads to macroscopic densification of the alloy and shrinkage of volume,109 while the formation of denser intermetallic phases (Chi, Sigma, Laves) usu- ally leads to an increase in volume, a form of nonvoid swelling.110,111 This counterintuitive behavior is the result of the different partial molar volumes of criti- cal elements (C and Mo primarily) between the new precipitates and the alloy matrix in which they form. Both the carbide and intermetallic phase evolution appear to be accelerated and sometimes altered under irradiation. Other radiation-produced phases (w0, G-phase) also appear to induce changes in lattice parameter Test portion length (mm) Vo lu m e fr ac tio n of m ar te ns ite (% ) 35 30 25 20 15 10 5 0 0 2 4 6 8 10 12 14 16 Figure 46 Deformation-inducedmartensite (vol.%) produced at 20 �C in the Russian analog of AISI 321 following irradiation in BN-600 to 26 dpa at 423 �C. Reproduced from Gusev, M. N.; Maksimkin, O. P.; Garner, F. A. J. Nucl. Mater. 2010, 403, 121–125. The leading edge of the wave was moving from right to left and was at 2–3mm when the test was interrupted. 600 500 400 316 SS 304 SS Inconel 600 Inconel 800 CF8 SS as-cast 308 SS weld Annealed Irradiated temperature = 400–427 �C Test temperature = 427 �C 1070 kJ m–2 J c (k J m –2 ) 300 200 100 0 0 5 10 Neutron exposure (dpa) Weld Base metal 15 20 25 Figure 47 Irradiation-induced evolution of fracture toughness Jc in various austenitic steels and welds. Reproduced from Mills, W. J. ‘‘Irradiation Effects on the Fracture Toughness of Austenitic Fe-Cr-Ni Alloys,’’ Hanford Engineering Development Laboratory Report HEDL-TME-82–17, Richland, WA, 1982; Mills, W. J. Nucl. Technol. 1987, 82, 290–303. 64 Radiation Damage in Austenitic Steels but these have not been well characterized, primarily because these phases develop concurrently with void swelling that masks their contribution.1 Garner1 provides a reviewof precipitation-induced strains. For the current purpose it is sufficient to note that carbide-induced densification increases with carbon content and with increasing irradiation temperature. Such volume changes for the most common carbon levels range from 0.1% to 0.4% decrease in volume. The resulting strains may or may not be isotropically distributed, depending on whether there is a pronounced starting dislocation texture on which the carbides nucleate. This process is most pro- nounced for titanium carbides in Ti-stabilized steels. Carbide-induced strains usually develop quickly enough to be measurable before swelling strains become dominant and therefore are relatively easy to identify compared to those of slower forming phases. The formation of intermetallic phases can gener- ate strains in the order of 1–3%. There is insufficient evidence to support anisotropy of resulting strains, but there exists some evidence that tensile stress states may accelerate the formation of these phases.110 Additionally, there is a decrease in density and a concurrent increase in volume when ferrite is formed from austenite as a result of radiation-induced segre- gation of nickel. Formation of ferrite from austenite can lead to volume increases as large as 3%, but there are no available data on potential anisotropy or stress dependence. As opposed to carbide-induced strains that develop relatively quickly, ferrite and interme- tallic strains develop rather slowly, and therefore are usually unrecognized, especially when other strain contributions arising from swelling and creep are present. Such precipitation-induced strains are important in that while they usually saturate in magnitude, they can be a significant portion of the total net strain at low dpa levels, thereby complicating the analysis and extrapolation of void swelling and irradiation creep data. Such strains can also affect the stress distribu- tion and level in a structural component. For instance, a preloaded tie-rod or bolt will initially increase in load as a result of carbide-induced shrink- age even while irradiation creep proceeds to relax the load. 100 11.0–11.3 ´ 1022n cm–2 (E > 0.1 MeV) 80 377 �C 388 �C Transgranular fracture Intergranular fracture Fatigue precracked zone Fatigue precracked zone Test temperature = 538 �C Test temperature = 649 �C 382 �C 400 �C 60 40 20 0 200 300 400 500 600 700 Irradiation temperature Test temperature (�C) K IC (M P a �m ) Figure 48 Dependence on test temperature of fracture toughness and fracture mode of highly irradiated 20% cold-worked 316. Reproduced from Huang, F. H.; Wire G. L. In Proceedings of Conference on Dimensional Stability and Mechanical Behavior of Irradiated Metals and Alloys; British Nuclear Energy Society: London, 1983; pp 135–138. Radiation Damage in Austenitic Steels 65 It should be noted that radiation-induced segre- gation can lead to overall changes in average lattice parameter without actually culminating in observable precipitation. Although there is no convincing evi- dence that segregation to void and grain boundaries produces measurable strains, it has been shown that radiation-induced spinodal-like decomposition in Fe–35Ni and Fe–Cr–35Ni alloys produces periodic oscillations in composition that are accompanied by densification in the order of �1%.112,113 Oscillations in nickel level are almost exactly offset by out-of- phase oscillations in chromium. This demonstrates that in a single phase system the lattice parameter of a given element is not constant but is influenced by its local concentration and its association with other elements. 4.02.8.2 Void Swelling and Bubble Swelling The progressive accumulation of high ‘cavity’ densi- ties (1012–1017 cm�3) leads to a macroscopic increase in volume of the steel. The concentration of these cavities tends to increase with decreasing tempera- ture or with increasing He, H, and residual gases such as O and N. ‘Cavity’ is a generic distinction for a hole in the matrix. Identifying a specific cavity as being either a bubble or void is not as simple as might be imagined, however. In general, bubbles are rela- tively small, gas-pressurized, existing sometimes at equilibrium pressures, although not necessarily at lower temperatures where they can be significantly over-pressurized. One defining feature is that bubbles tend to grow slowly by gas accumulation while voids are either totally or partially vacuum-filled, but which are free to grow rapidly via vacancy accumu- lation without further gas addition. It is well known that bubbles can serve as nuclei for voids, accounting for the known tendencyof helium especially to accelerate the onset of void swelling and to increase the cavity density. In some strongly helium-generating environments, there can also develop a late-term surge of tiny bubbles forming at very high densities in the interstices between earlier nucleated voids at much lower densities. This is a consequence of the 59Ni two-step transmutation sequence that accelerates helium production after voids are already nucleated and growing.114 As dis- cussed earlier in Section 4.02.6 these ‘helium-filled’ bubbles are probably pressurizedwith stored hydrogen as well as helium. Interestingly, the onset of this late- term bubble evolution does not change the steady-state swelling rate even though the cavity density increased by several orders of magnitude once helium generation accelerated strongly with the 59Ni sequence. For most engineering applications in nuclear sys- tems it is void swelling that is the most important contributor to dimensional instability. In the absence of physical restraint or applied stress field void swell- ing distributes its strains isotropically with the most 66 Radiation Damage in Austenitic Steels famous published example shown in Figure 49.115 When restrained in any direction, however, the swelling-induced stresses activate irradiation creep (to be discussed later), which then redistributes the strain in the unrestrained directions, as shown earlier in Figure 16 where fuel pins locally restrained by a spirally wrapped wire evolved into spiral fuel pins. At any given altitude on the fuel pin the interaction between wire and cladding the cross-section becomes oval in shape and the resulting deformation is called ‘ovality.’ It is important to note that, contrary to popular opinion, swelling and irradiation creep are not sepa- rate processes, but are ‘two sides of the same coin.’ These phenomena are two manifestations of the radiation-enhanced dislocation motion required to move the material previously located at the void positions to the outer boundaries of the grains. This process is operating even in the absence of stress to produce swelling, but responds selectively to shear stresses generated either by externally applied or internally generated forces. While swelling attempts to be isotropic, irradiation creep redirects mass flow anisotropically. As will be shown later irradiation creep can operate before the onset of swelling but is accelerated when swelling begins. Unirradiated control 20% CW 316 Fluences beyond FFTF goal 1 cm Figure 49 Macroscopic swelling (�10% linear as measured by length change, �33% volumetric, as measured by density change) observed in unfueled 20% cold-worked AISI 316 open cladding tube at 1.5�1023 n cm�2 (E > 0.1MeV) or �75 dpa at 510 �C in EBR-II. Note that in the absence of physical restraints all relative proportions were preserved. Reproduced from Straalsund, J. L.; Powell, R. W.; Chin, B. A. J. Nucl. Mater. 1982, 108–109, 299–305. Void swelling is probably the most heavily researched and published radiation-induced phe- nomenon, although pressure vessel embrittlement has also received a similar amount of attention. A comprehensive review on void swelling and irradi- ation creep was written in 1994 1 and is now being revised116 not only to incorporate new insights devel- oped over the past decade and a half, but also to revise some earlier perceptions that have not sur- vived more recent examination. A brief summary of current knowledge relevant to the purpose of this review is provided in the following sections. In some crystal systems, especially simple body- centered cubic (bcc) metals, the void swelling process is inherently self-limiting, usually saturating at some value below 5%.9 Such saturation is usually accompa- nied by a process referred to as ‘self-organization’ whereby voids arrange themselves in three- dimensional arrays that exhibit the same crystalline orientation as that of the crystal structure. Unfortu- nately, for most face-centered cubic (fcc) metals, espe- cially stainless steels, self-organization and saturation of void swelling do not operate under most reactor- relevant conditions, and as a result swelling in austen- itic stainless steels is an inherently unsaturable process. Void swelling normally exhibits a transient or incu- bation regime where either no swelling or very slow swelling occurs before swelling moves to a steady-state rate. Tens of percent swelling have been reached in a number of reactor-relevant irradiation histories, and values of 80–90% swelling without hint of impending saturation have been attained in both model and com- mercial alloys during neutron irradiation.1,117,118 Swelling in excess of 200% was reached during proton irradiation of 316 stainless steel and saturation was eventually observed at �260% swelling.119 An example of apparently nonsaturable void swelling in AISI 316 is presented in Figure 50.117 Note that the onset of rapid swelling, defined by termination of a ‘transient’ regime, is dependent on both irradiation temperature and dpa rate. The dpa rate dependence of the transient is not easily dis- cerned in Figure 50 where each irradiation temper- ature in this experiment is coupledwith a specific dpa rate, with the range of dpa rates increasing �65% from lowest to highest. It will be shown later that dpa rate is a very strong determinant of void swelling. The transient regime is terminated when the condi- tions for both void nucleation and especially rapid void growth have been attained. The conditions for void nucleation must be favor- able to end the transient. This usually requires 80 1% per dpa 510 �C 538 �C 593 �C 427 �C 454 �C 650 �C 0.2% per dpa S w el lin g (% ) 400 �C 482 �C60 40 20 0 0 Neutron fluence (n cm-2) (E > 0.1 MeV) 10 20 30 ´ 1022 Figure 50 Swelling determined by density change as a function of irradiation temperature and dose, as observed in 20% cold-worked AISI 316 irradiated in the EBR-II fast reactor. Reproduced from Garner, F. A.; Gelles, D. S. In Proceedings of Symposium on Effects of Radiation on Materials: 14th International Symposium; ASTM STP 1046; 1990; Vol. II, pp 673–683. All measurements at a given temperature were made on the same specimen after multiple exposures with subsequent reinsertion into the reactor. This procedure minimized specimen-to-specimen data scatter and assisted in a clear visualization of the posttransient swelling rate. Radiation Damage in Austenitic Steels 67 attainment of a dislocation network to the quasi- equilibrium value of �3� 1010 cm�2, either by reduction of higher cold-worked densities or build up from lower densities characteristic of annealed alloys.1 It also requires that the temperatures be low enough to guarantee sufficient supersaturation of vacancies or that elements (P, Si, Ni) that strongly increase the effective vacancy diffusion coefficient, and thereby depress void nucleation, be low enough or have been reduced via precipitation. Helium and other gases influence void nucleation and under some situations where nucleation is difficult can serve to shorten the transient duration. Rapid void growth after sufficient nucleation of voids requires not only the attainment of the quasi-equilibrium dislocation density, but also that dislocation network be a ‘glissle’ network capable of moving mass quickly. Voids previously nucleated but still embedded in a ‘sessile’ microstructure composed primarily of Frank loops can grow but not quickly. Therefore, significant unfaulting of Frank loops is a prerequisite for termination of the transient and the onset of the high swelling rate. As also shown in Figure 50, the terminal post- transient swelling rate of AISI 316 is typical of all austenitic stainless steel at �1% per dpa, essentially independent of all irradiation or material vari- ables.1,120 This terminal rate also appears to be char- acteristic of Fe–Cr–Mn, Fe–Cr–Mn–Ni, and simple Ni-base alloys, although for the nickel-base w0/w00 stabilized alloys the transients are generally much longer and insufficient amounts of swelling were attained in most studies to allow confirmation of the full generality of this statement of a universal steady- state swelling rate for all fcc alloys.1,121,122 In Fe–Cr–Mn and Fe–Cr–Mn–Ni alloys removal of highly diffusing Mn from voids and grain bound- aries via the inverse Kirkendall effect leads to these sinks becoming coated with lower-swelling ferrite phase, thereby producing a late-term decrease in the average swelling rate.121,122 4.02.8.3 Parametric Dependencies of Void Swelling The duration of the transient regime of swelling in austenitic and high-nickel steels is known to be exceptionally sensitive to irradiation parameters but also to be very sensitive to fine details of composition, heat treatment, and mechanical processing. It would require a very long article to review all of the para- metric sensitivities of the transient duration to such a wide array of variables, so only a brief summary will be presented here. The reader is referred to Garner1,116 for a more detailed description. 4.02.8.3.1 Stress state The dependence of void swelling on stress state is an example of a second-order sensitivity mentioned at the beginning of this section. If a material swells rather easily, stress has only a small or unnoticeable effect on swelling. If the transient regime is large, however, stress can shorten the transient significantly. The effect of stress during irradiation is almost always to increase swelling. One significant exception arises if an annealed steel is subjected to a load above its yield stress during the rise to power. This often leads to a decrease in swelling relative to that pro- duced at a stress below yield. In effect, the steel is plastically deformed and warm-worked during the rise to power, raising the dislocation density. Figure 51 Top of a fuel assembly from the BN-600 fast reactor showing larger swelling-induced elongation of annealed EI-847 steel in pins with lower (0.09 vs. 0.20%) silicon content, with both heats having concentrations below the specified maximum of 0.4%. Reproduced from Porollo, S. I.; Shulepin, S. V.; Konobeev, Yu. V.; Garner, F. A. J. Nucl. Mater. 2008, 378, 17–24. 68 Radiation Damage in Austenitic Steels Applied stresses have been shown to participate in the evolution of Frank loop and dislocation evolution and to produce the anisotropy of Burger’s vector distribution that is important to the operation of irradiation creep.123 Since shear stresses also assist in the unfaulting of Frank loops and in the evolution toward quasi-equilibrium network densities, it is not surprising that applied stress accelerates the onset of swelling. Although most previously reported experi- ments involved only tensile stress states, some experi- ments suggested that both tensile and compressive stress states shortened the transient regime.1 Two recent studies have convincingly shown that the hydro- static component of stress is relatively unimportant and that it is the deviatoric component or shear stress that accelerates swelling.124,125 This is especially evi- dent for loads applied to springs where there is a pure shear stress state without a hydrostatic component. In this case stress-enhanced swelling is also observed.124 Until recently it was not known if the stress- enhanced increment of swelling during constantly applied stress was distributed isotropically or not. A recent publication by Gilbert and Garner showed that both the stress-free and stress-enhanced incre- ments of swelling were distributed isotropically.126 The history of the stress state is as important as its magnitude and relative contribution of shear and hydrostatic components. In fuel pins, for instance, the stress is initially low and builds up slowly. In this case, swelling is usually in progress long before stress can participate. In pressurized tubes, however, creep starts long before swelling begins. The loop and dislocation microstructures of the swelling-before- creep and creep-before-swelling scenarios are differ- ent and therefore the swelling and creep behaviors are also somewhat different.1 Stress can also leave a memory in a component after the stress is removed and irradiation continues.123,127 Garner and coworkers recently showed that when stress was removed from previously stressed tubes they continued for a short time to distribute mass in the directions dictated by irradiation creep in response to the stress state characteristic of a pressur- ized tube, although the memory faded as irradiation continued.127 The memory is thought to reside in the stress-induced anisotropic distribution of Burger’s vec- tors, which was eventually replaced with an isotropic distribution. 4.02.8.3.2 Elemental composition The duration of the transient regime of austenitic and nickel-base alloys depends to the first-order on major element composition, primarily on the Fe, Cr, and Ni content.1,57,120 Increases in chromium content decrease the effective vacancy diffusion coefficient and thereby increase the vacancy supersaturation, increasing void nucleation, and decreasing the tran- sient duration. Increases in nickel initially increase the effective vacancy diffusion and thereby the tran- sient, but behavior reverses at some mid-nickel level (40–60%), reflecting the nonmonotonic dependence of both the effective vacancy diffusion coefficient and the dislocation bias on nickel content.55,128,129 With respect to minor solutes, the most important elements influencing swelling are P and Si.57,130 On a per atom basis phosphorus has the most pronounced effect on the transient duration, followed by silicon. Additions of small amounts of silicon and phosphorus initially increase swelling, but then strongly decrease it at higher content, producing a nonmonotonic swelling behavior. This response reflects the two competing roles of these elements on solute– interstitial binding at low concentration and their much stronger enhancement of vacancy diffusion at higher content. Very small differences in silicon between two otherwise identical heats of steel can produce quite different transient duration and there- fore swelling, as shown in Figure 51.130 Looking back at the FFTF fuel assembly in Figure 16, it can be seen that there are three clusters Double-aged to precipitate all carbon from solution Solution annealed 20% cold-worked Distance along fuel pin DD D (% ) Figure 52 Schematic illustration of swelling-induced changes in pin diameter observed in EBR-II for one heat of AISI 316 stainless steel irradiated in various starting conditions. Reproduced from Garner, F. A. In Materials Science and Technology: A Comprehensive Treatment; VCH: New York,1994; Vol. 10A, pp 419–543. Radiation Damage in Austenitic Steels 69 of pins that also extend above their neighbors. The pins in these clusters were made from a nominally similar heat with differences in phosphorus level, 0.002 versus 0.009wt%, both below the maximum specification of 0.04 wt%. In both the silicon and phosphorus examples shown here, the compositions fell under the specified maximum value, indicating the necessity to specify both the upper and lower limits of active elements when attempting to control swelling.131 Other common solute additions such as boron, carbon, manganese, molybdenum, niobium, vana- dium, and others have some impact on diffusion, but appear to exert their greatest influence on the formation of various precipitates that remove the more active elements from solution. 4.02.8.3.3 Alloy starting state To the first-order most researchers concentrate on the cold-work level as the primary way to delay void swelling, although it is known that increasing cold work beyond a certain level specific to each alloy yields diminishing returns, with the optimum level usually chosen to be 20–25% for austenitic alloys. Larger levels are often counter-productive in that the additional stored energy at higher cold-work levels sometimes induces recrystallization during irradia- tion, often resulting in higher swelling.1 Additionally, in some alloys and metals it is diffi- cult to nucleate voids under some combinations of temperature and dpa rate due to the difficulty to establish a stable dislocation network. Cold working in some cases can actually shorten the transient by providing a stable glissile dislocation network and thereby accelerate swelling, as observed in model Fe–Cr–Ni alloys and simple metals such as nickel and iron.132–134 The starting thermal–mechanical condition of the alloy plays an important role in determining the transient duration via its influence on the starting dislocation density, but more importantly in deter- mining the distribution or chemical activity of the active elements. For instance, aging of an alloy at intermediate temperatures that encourage carbide precipitation, for instance, is the most effective way to produce the shortest transient and the highest swelling.1 There are many other examples. For instance, the chemical activity of an element like phosphorus is very sensitive to the inter-pass annealing tempera- ture range employed in producing cold-worked tub- ing by multiple drawings. It is speculated that phosphorus can be either in solution or existing as small invisible precipitates of lesser chemical activity depending on the inter-pass annealing temperature or tube feed rate through the furnace.1,135 As carbon plays a role in both carbide and inter- metallic phase evolution, and its chemical activity can be strongly affected by thermal and mechanical his- tory, it exerts a strong and often complex effect on the transient duration. One aspect of this complexity is the often-observed two-peak swelling behavior ver- sus temperature that strongly varies with thermal– mechanical treatment.1 This effect is so strong that the swelling valley between the two peaks often occurs at the peak flux position. Cold-working tends to suppress the low temperature peak more than the high temperature peak due to its effect to delay and homogenize carbide formation. Removing almost all carbon into precipitates by aging erases the double peak behavior and usually produces the largest amount of swelling, as shown in Figure 52. 4.02.8.3.4 Irradiation temperature With respect to the irradiation environment there are four major variables that determine the duration of the transient. The first three are related to each other: irradiation temperature, temperature gradi- ents, and temperature history. The fourth is strongly synergistic with temperature and is the dpa rate, which will be covered in the following section. Some temperature histories, especially when gradually falling from one temperature to a lower temperature, produce a shorter transient compared to that of either the starting or final temperature, 70 Radiation Damage in Austenitic Steels primarily because such histories tend to accelerate the radiation-induced formation of nickel and silicon-rich phases, especially that of the g0 phase.1,136 Formation of these phases usually precedes swelling.1 Strong gradients in temperature across thin fuel cladding have also been shown to accelerate the onset of swelling, producing more swelling than what iso- thermal irradiation would produce at either the upper or lower cladding temperature.137,138 The exact cause is unknown but it was speculated that the stress gradients associated with strong tempera- ture gradients might be a contributing factor. For isothermal irradiation the temperature is an important determinant of the transient duration, not only because it directly impacts diffusion and void nucleation, but also because of its influence on phase stability and phase evolution. However, over the wide range of temperatures experienced in fast reactors, temperature has no effect on the posttransient steady-state swelling rate of 300 series stainless steels at �1% per dpa. However, it is frequently assumed that at constant dpa rate there is a peak swelling temperature or peak swelling rate as a function of temperature for swelling of austenitic steels. This persistent misperception is a consequence of the historical use of fast reactors. All of the earlydata on swellingwas derived from small fast reactor cores such as EBR-II and DFR, which have strongly peaked dpa rate profiles, both axially and radially. Later studies conducted in larger cores such as that of FFTF showed that assuming a temperature- dependent steady-state swelling rate was incorrect. More careful analyses of other data from these smaller cores also support this point of view. 4.02.8.3.5 Influence of dpa rate on swelling Historically, the influence of differences in dpa rate across small cores was perceived as an effect of tem- perature on swelling rate rather than a flux effect, primarily because it was difficult to separate the influence of dpa, dpa rate, and temperature in limited data fields from small cores. While it was recognized for many years that there was some effect of dpa rate to determine the transient duration, until rather recently the full strength of the rate effect was underappreciated. The new appreciation for the strong influence of dpa rate arises from two categories of studies con- ducted over the past decade. The first type involved direct single variable comparisons of the effect of dpa rate on swelling. The second category involved the examination of components irradiated at very low dpa rates and often at temperatures below the previ- ously perceived lower limit of swelling. 4.02.8.3.5.1 Category I of dpa rate effects Three examples of the first category of dpa rate studies are presented here. The first experiment by Garner and coworkers involved the examination (density change and microscopy) of five unfueled hexagonal subassemblies constructed of a single heat of annealed AISI 304 stainless steel irradiated for many years in the reflector rows 8, 9, 10, and blanket row 14 of the EBR-II fast reactor.139,140 These com- ponents were chosen because they were made of the same steel used to construct the baffle-former-barrel assembly of PWR internals and the hexagonal sub- assemblies spanned the full range of dpa rates and temperatures found in the most swelling-vulnerable parts of the PWR baffle-former assembly. The EBR-II experiment isolated the effect of dpa rate by concentrating on a limited range of tempera- tures (373–388 �C), but covering a very large range of dpa rates (0.06–3.8� 10�7 dpa s�1), with no sig- nificant difference in He/dpa ratio. The data in Figure 52 clearly shows that the transient regime of swelling is progressively shortened as the dpa rate decreases, such that only 10 dpa is required to reach 1% swelling in row 14. In a previous publication it was shown that 30–50 dpa were required to exceed 1% swelling when data were collected at these tempera- tures from rows 2 to 4 inside the EBR-II core at higher dpa rates.141 In this experiment the swelling rates at the highest doses reached are still far from the 1% per dpa known to be a characteristic of this alloy (Figure 53). Voids and carbide precipitates were found in all examined specimens with swelling ranging as high as 2.8%. Examples of the void microstructure and its sensitivity to dpa rate are shown in Figure 54.142 Universally, it was found that lower dpa rates at a given temperature increased the swelling. The second series of experiments were reported by Okita and coworkers and involved simple model alloys, ternary Fe15Cr16Ni and quaternary Fe15Cr16Ni–0.25Ti, with very low levels of other solutes.143–145 These alloys have no possibility to be involved in segregation-induced precipitation of Ni-rich phases, so any dependence on dpa rate must involve the evolution only of voids, loops, and dislocations. These simple austenitic alloys were irradiated in the FFTF fast reactor with actively controlled tem- peratures near 400 �C at seven different dpa rates. Measurement techniques used were density change 3.0 2.5 1.5 S w el lin g (% ) 1.0 0.5 0.0 0 5 10 15 20 dpa(a) (b) 25 30 35 0 5 10 15 20 dpa 25 30 35 U9807 Row 8 1.25–3.60 ´ 10-7 dpa/sec U8972 Row 9 1.00–2.05 ´ 10-7 U9009 Row 10 0.38–0.96 ´ 10-7 U9009 0.38–0.96 ´ 10-7 dpa/sec U9007 0.44–1.12 ´ 10-7 dpa/sec U1603 Row 14 0.062–0.156 ´ 10-7 2.0 3.0 2.5 1.5 S w el lin g (% ) 1.0 0.5 0.0 2.0 U1603 U8972U9009 U9807 U9009 U9007 Figure 53 Swelling of annealed 304 stainless steel in the range 373–388 �C measured by density changes in the lower halves of EBR-II reflector subassemblies, designated by identification numbers such as U9807. Note that each data set spans a range of dpa rates. (a) Comparison of four subassemblies in different rows of the reactor. (b) Comparison of two subassemblies in Row 10 but on opposite sides of the reactor, with dpa rates varying only �16%, showing that lower dpa rates lead to an earlier acceleration of void swelling. Reproduced from Garner, F. A.; Makenas, B. J. In Proceedings of Fontevraud-6 Symposium on Contribution of Materials Investigations to Improve the Safety and Performance of LWRs; 2006; pp 625–636. Radiation Damage in Austenitic Steels 71 and microscopy. Multiple specimens were irradiated side-by-side and the measured swelling was remark- ably reproducible. Figure 55 shows swelling for five of the seven dpa rates where there was a progressive shortening of the transient regime as the dpa rate decreased. At the lower two dpa rates (not shown here) the transient regime had decreased to less than 1 dpa. Most importantly, the steady-state swelling rate appeared to be approach- ing or to have reached 1% per dpa at all seven dpa rates. The most illuminating observation came from the microscopy, however, showing that the microstructural feature most prominently associated with attaining the steady-state swelling rate was the loss of all Frank loops and the establishment of a glissile rather than sessile dislocation structure. In a companion experiment the ternary Fe15Cr16Ni alloy was irradiated over a range of temperatures using nickel ions at three much higher dpa rates; it was shown that while voids can nucleate in a highly sessile microstructure, they cannot grow at a high rate.146 Most importantly, it was confirmed that increases in dpa rate led to a progressive decrease in swelling even in sessile networks. 100 nm 100 nm 10 dpa 0.15 ´ 10-7dpa s–1 1.2% swelling 14.3 dpa 1.8 ´ 10-7dpa s–1 0.42% swelling Figure 54 Void microstructures observed in annealed AISI 304 reflector ducts from EBR-II showing variation of swelling in response to differences in dpa rate at 379 �C. Reproduced from Bond, G. M.; Sencer, B. H.; Garner, F. A.; Hamilton, M. L.; Allen, T. R.; Porter, D. L. In Proceedings of 9th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors; 1999, pp 1045–1050. Less swelling per dpa was produced at the higher dpa rate. Small dark features are M23C6 precipitates that form concurrently, producing a densification of �0.2%. 30 25 20 15 10 5 0 0 10 20 30 40 50 60 Cumulative Fe–15Cr–16Ni 0.54 0.78 1.70 ´ 10-6dpa s–1 0.31 0.09 S w el lin g (% ) Figure 55 Swelling of simple ternary and quaternary model au decrease in the transient duration as the dpa rate decreases. Re Greenwood, L. R.; Wolfer, W. G.; Isobe, Y. In Proceedings of 10t Materials in Nuclear Power Systems – Water Reactors; 2001. Not terminal swelling rate of �1% per dpa (see dotted line). 72 Radiation Damage in Austenitic Steels Whereas most void swelling models focus on the rate dependence of void nucleation and growth, Okita showed by microscopy that the effect of dpa rate was most strongly manifested in the Frank loop population. High dpa rates produced a high density of loops of smaller size, while low dpa rates produced fewer loops at larger size. The latter ensemble is more prone to unfaulting, the first step toward pro- ducing a glissile microstructure. Denser ensembles at smaller sizes resisted unfaulting for a longer period. Thus the dependence of transient duration on dpa rate arose primarily from its influence on the stability against loop unfaulting. In the third series of experiments, Budylkin prepared two experimental alloy series to be irra- diated in very similar neutron spectra in both the BOR-60 and BN-350 fast reactors at nearly identical temperatures and dpa levels.147 The first four-alloy series was Fe–16Cr–15Ni–3Mo–0.6Nb–0.6Mn–0.06C– 0.008P but varying in silicon content from 0.4 to 1.2 wt%. The second three-alloy series contained the 0.63% silicon variant from the first series and two other alloys where 0.15% titanium either was added to or replaced the 0.6% Nb. The irradiations proceeded at 5.06� 10�7 dpa s�1 and 480 �C in BOR-60 and at 1.58� 10�6 dpa s�1 and 490 �C in BN-350. Thus there was approximately a factor of three difference in dpa rate. As shown in Figure 56, significantly higher swelling was uni- formly observed in the lower flux irradiation in BOR-60, regardless of alloy composition. 0 dose (dpa) Fe–15Cr–16Ni–0.25Ti 10 20 30 40 50 60 70 stenitic alloys at �400 �C in FFTF, showing a progressive produced from Okita, T.; Sekimura, N.; Garner, F. A.; h International Conference on Environmental Degradation of e that all swelling curves have reached or are approaching a Avg. size = 8.6 nm Swelling = 0.20% Density = 0.61 ´ 1022m-3 2200 nmnm 1100 nmnm Figure 57 Voids observed in Tihange baffle-former bolt made with cold-worked 316 stainless steel after irradiation at �345 �C to 12 dpa. Reproduced from Edwards, D. J.; Simonen, E. P.; Garner, F. A.; Greenwood, L. R.; Oliver, B. A.; Bruemmer, S. M. J. Nucl. Mater. 2003, 317, 32–45. 25 20 15 10 5 0 25 20 15 10 5 0 0 0.5 Nb Nb+Ti Ti Solute addition 1 Silicon (wt%) S w el lin g (% ) S w el lin g (% ) 1.5 BOR-60 BN-350 Figure 56 Comparison of swelling measured by density change for two experimental alloy series based on Fe16Cr15Ni3MoNbB that were irradiated in BOR-60 (480 �C, 52 dpa, 5�10�7 dpa s�1) and BN-350 (490 �C, 53 dpa, 15.6�10�7 dpa s�1), showing that swelling was always higher at the lower dpa rate. Reproduced from Budylkin, N. I.; Bulanova, T. M.; Mironova, E. G.; Mitrofanova, N. M.; Porollo, S. I.; Chernov, V. M.; Shamardin, V. K.; Garner, F. A. J. Nucl. Mater. 2004, 329–333, 621–624. Radiation Damage in Austenitic Steels 73 4.02.8.3.5.2 Category 2 of dpa rate effects For many years it was assumed that void swelling would not be an issue for the 304 and 316 stainless components comprising the internals of power- producing light water-cooled reactors. Such a conclu- sion was easily accepted for boiling water reactors since steels used in the shroud assembly are separated from the core by a substantial water gap and therefore experience less than 5 dpa over a 40-year lifetime. For pressurized water reactors, however, the steel is much closer to the core and some regions can reach 80–100 dpa over 40 years. Swelling was still not thought to be a problem because swelling was per- ceived to inhabit a temperature range that did not extend down to the 280–290 �C inlet temperatures of PWRs, and based on most fast reactor irradiations, swelling was thought to vanish below 345 �C, themax- imumwater temperature in PWRs. It was also thought that the lower dpa rates characteristic of PWRswould reduce vacancy supersaturations and would therefore inhibit void nucleation. Unfortunately, gamma and nuclear heating of thick plates can raise the internal temperatures in some areas of the baffle-former plates to temperatures above 400 �C, known to be prime territory for void swelling. Also, as seen in the previous section, void nucleation does not dominate the swelling response to decreasing dpa rate. The shortening of the tran- sient regime at lower dpa rates raised the strong possibility that void swelling might indeed occur in PWR internals. Theoretical considerations based on void nucleation also suggested that the temperature regime of swelling might move to lower temperatures with decreasing dpa rates. Therefore an effort was made to find stainless steels irradiated at lower dpa rates and/or lower temperatures. The first clear example of void swelling in PWRs was found in a cold-worked 316 baffle bolt removed from the Tihange PWR reactor located in Belgium.39 The bolt was removed in response to an ultrasonic indication of cracking under the bolt head. Although the bolt shown in Figure 57 was con- structed from cold-worked 316 austenitic stainless steel known to be more resistant to the onset of swelling than the annealed AISI 304 plate in which it was embedded, well-faceted voids of easily resolv- able size were clearly observed in three sections removed along the bolt axis. The doses in the bolt were relatively low and the calculated temperatures were also relatively low compared to typical fast reactor observations, but the swelling exceeded expectations based on fast reactor experience. As cold-worked 316 is known to always swell less than 74 Radiation Damage in Austenitic Steels annealed 304 at the same temperature and dpa rate the worrisome inference is that the 304 plate sur- rounding the bolt might be swelling at higher levels. Significantly, hydrogen was also found to be stored in the microstructure at unexpectedly high levels that increased as void swelling increased along the bolt length. Subsequently, voids were observed in other AISI 316 bolts from this same reactor by other research- ers148 often at even lower doses and temperatures, producing lesser but measurable amounts of swelling. An example is shown in Figure 58, but it should be noted that there appear to be two populations of cavities, a few that are recognizable as voids and an exceptionally high population of nanometer-sized cavities that are only visible using a large level of defocusing, similar to the behavior shown earlier in Figure 17. Figure 58 (top) Voids at very low density (see arrows) and (bottom) an exceptionally high density of subvisible cavities or ‘nano-bubbles’ observed in another Tihange baffle-former bolt designated 2K1R1 after 8.5 dpa at �299 �C. Micrographs supplied courtesy of L. E. Thomas of Pacific Northwest National Laboratory. The smaller cavities can only be seen with significant under-focusing. Black bars are 50nm in length. Voids have been sometimes but not always observed in bolts of various steels removed from US PWRs.149,150 These studies were conducted before the need for defocusing was recognized, however. Small cavities that could be either voids or bubbles have also been observed in thin-walled flux thimble tubes removed from various PWRs.76,151–153 Neus- troev and coworkers also found voids in a thimble tube removed from a VVERoperating in the Ukraine, noting that voids were observed at unexpectedly low temperatures and dpa levels.154 The potential for void swelling at PWR-relevant dpa rates and temperatures is best demonstrated in more comprehensive studies conducted in four USSR sodium-cooled fast reactors located in Russia and Kazakhstan. Whereas the inlet temperature of most Western or Asian fast reactors was of the order of 365–375 �C, the Soviet BOR-60 and BN-350 fast reactors had inlet temperatures of the order of 270–280 �C. Components from regions below the core or in the reflector region have been extracted for study at dpa rates and temperatures that were comparable to those of PWRs.155–161 A summary paper containing an overview of these studies shows that in all studies conducted on compo- nents removed from low flux positions in Soviet fast reactors, certain recurrent trends were observed.155 First, whenever the dpa rate was significantly lower at any investigated temperature, swelling was observed at surprisingly very low dpa levels. An excellent exam- ple is shown in Figure 59 where significant void swelling was observed at only 0.64 dpa at 350 �C.156 Second, whenever a comparison could be made within one reactor at a given temperature, the transient dura- tion decreased with lower dpa rate.157–159 Most impor- tantly, whenever temperatures approaching 280 �C could be reached, swelling was observed not only at these low temperatures but also at surprisingly low dpa levels.160,161 Other examples are shown in Figures 60 and 61. 4.02.9 Irradiation Creep 4.02.9.1 Introduction While the deleterious impact of thermal creep at higher temperatures has long been known to be of engineering concern, the discovery of orders of mag- nitude increase in creep rate at relatively low tempera- tures was as unexpected and worrisome as was the discovery of void swelling. The first observations of creep indeed occurred in systems and at doses where Radiation Damage in Austenitic Steels 75 void swelling had not yet happened. Thus, it was natural to assume that the creep and swelling phenomena were independent processes. One early example of radiation-accelerated creep at 454 �C is shown in Figure 62.162 The posttransient creep rate 50 nm 50 nm Figure 59 Voids observed in annealed 12X18H9 steel at 350 �C in the BR-10 fast reactor at only 0.64 dpa produced at 1.9� 10�9 dpa s�1. Reproduced from Porollo, S. I.; Dvoriashin, A. M.; Konobeev, Yu. V.; Ivanov, A. A.; Shulepin, S. V.; Garner, F. A. J. Nucl. Mater. 2006, 359, 41–49. This steel is analogous to AISI 321. 3 dpa 6.5 dp 50 nm Figure 60 Microstructure of annealed 12X18H9 specimens irra 320–330 �C for 27 years. Lower dpa levels were reached at lowe Konobeev, Yu. V.; Neustroev, V. S.; Maksimkin, O. P. In Procee Materials Investigations to Improve the Safety and Performance in this experiment was later calculated by Foster and coworkers to be 0.95� 10�6 (MPa dpa)�1.163 In this case, the specimen is increasing in length and decreasing in cross-section. Precipitation of car- bides leads to a small densification and shrinkage of the specimen as shown in the thermal creep behavior. A similar densification process occurs during irradia- tion but its strains are overwhelmed by the irradiation creep strains. The accumulated damage is relatively small ( 0.65 dpa, 281 �C 7.7 dpa, 285 �C 8.8 dpa, 430 �C12.6 dpa, 380 �C12.3 dpa, 363 �C Figure 61 Void microstructure observed in a wrapper duct constructed from annealed 12X18H9 stainless steel and irradiated in the BN-350 reactor at various axial distances from the midplane. Reproduced fromMaksimkin, O. P.; Tsai, K. V.; Turubarova, L. G.; Doronina, T. A.; Garner, F. A. J. Nucl. Mater. 2007, 367–370, 990–994. Lowest temperatures correspond to the bottom of the duct. 15 ´ 10-4 10 5Te ns ile s tr ai n 0 0 500 1000 Thermally induced densification and creep Irradiation creep 138 MPa 454 �C Time (h) 1500 2000 Figure 62 Comparison of thermal and irradiation creep strains observed in 20% cold-worked 316 stainless steel in uniaxial creep tests during neutron irradiation in the EBR-II fast reactor or during ex-reactor thermal aging. Reproduced from Gilbert, E. R.; Kaulitz, D. C.; Holmes, J. J.; Claudsen, T. T. In Proceedings Conference on Irradiation Embrittlement and Creep in Fuel Cladding and Core Components; British Nuclear Energy Society: London, 1972; pp 239–251. 0.1 M id w al l h oo p s tr ai n (% ) 0.01 0.001 1.0 1.1 1.2 1000/T (K-1) 1.3 1.4 1.5 1.6 1 10 1300 1200 1100 1000 Temperature �F 900 20% CW 316 SS Dose–15 dpa Time–4138 h Thermal creep In-reactor creep Hoop stress ksi/MPa 30/207 20/138 5/34 10/69 30 20 10 5 800 700 Figure 63 Early comparison of thermal creep and irradiation creep in EBR-II of 20% cold-worked 316 pressurized tubes at 15 dpa and various stress levels. Reproduced fromGilbert, E. R.; Straalsund, J. L.; Wire, G. L. J. Nucl. Mater. 1977, 65, 277–294. 76 Radiation Damage in Austenitic Steels anisotropy, it has been shown that during charged particle simulation experiments on void swelling that 100% of the mass flow is directed perpendicular to the surface normal. During irradiation the swelling region experiences only a small fraction of the yield stress and does not produce plastic deformation.16 In effect, ion-induced swelling experiments might be better characterized as irradiation creep experiments 300 mm Figure 65 Micrographic section through a short length of Nb-stabilized Fe-20Cr-20Ni stainless steel after irradiation in a Commercial Advanced Gas Reactor, showing collapse of the cladding at �650 �C into a large inter-pellet gap caused by thermal ratcheting of pellets during power level changes. The pressure difference driving the collapse is only �4MPa but produces �55% strain without failure. The square features are cross-sections of a spiral rib on the cladding used to improve thermal performance. Reproduced from Crossland, I. G.; Roberts, G.; Jones, K. W.; Bradshaw, K. A. In Nuclear Fuel Performance; British Nuclear Energy Society: London, 1985; Vol. 1; pp 37–41. 15 dpa Hoop stress = 70 MPa Temperature (�C) Temperature (�F) 1.0 1400 700 600 500 400 1200 1000 800 0.1 M id w al l h oo p s tr ai n (% ) 0.01 0.001 1.0 1.1 1.2 1000 / T (K-1) 1.3 1.4 1.5 1.6 8 dpa 2 dpa 4138 h 2400 h 600 h Figure 64 Another comparison of thermal creep and irradiation creep in EBR-II of 20% cold-worked 316 pressurized tubes at various dpa levels at 70MPa hoop stress. Reproduced from Gilbert, E. R.; Bates, J. F. J. Nucl. Mater. 1977, 65, 204–209. Radiation Damage in Austenitic Steels 77 in that two-thirds of the mass flow has been redir- ected along the third coordinate. One distinguishing characteristic of irradiation creep as opposed to purely thermal creep is that it is inherently nondamaging on the microstructural level. Its operation does not result in triple-point cracks or grain boundary voids and does not produce plastic deformation or martensite. In that sense irra- diation creep is beneficial in that its action mitigates both microscopic and macroscopic problems arising from both externally loaded and internally generated stresses. Irradiation creep is especially effective to avoid stress concentrations anywhere in the micro- structure, reducing the possibility of failure initiation. There is some evidence to support the proposal that radiation creep is not simply additive to thermal creep but that it modifies thermal creep, making it less dam- aging. Figure 65 shows an excellent example of the power of irradiation-modified thermal creep to pro- duce a large amount of distortion without failure in CAGR cladding. Irradiation creep is a not completely understood phenomenon whose microstructural origins and complexity continue to evolve in our understanding, requiring additional evaluation as new data and insight are collected. As noted earlier, a compre- hensive review on both void swelling and irradia- tion creep was written in 19941 and is now being revised,116 not only to incorporate new insights devel- oped over the past decade and a half, but also to revise some earlier perceptions that have not survived more recent examination. Most of the revised percep- tions concern irradiation creep rather than swelling. On the microscopic level, however, there is still debate about what modes of dislocation movement and what balance of diffusion versus dislocation con- tributions combine to produce the observed strains under various irradiation conditions. Additional background information on this subject can be found in two other papers.166,167 While there are a very large number of creep mechanisms that have been proposed, most fall in sev- eral broad categories with respect to the microstruc- tural components and the point defects involved. Matthews and Finnis have presented the most exten- sive review of these mechanisms and ranked them in terms of their relative plausibility.166Garner andGelles 78 Radiation Damage in Austenitic Steels demonstrated that irradiation creep often leavesmicro- structural evidence in the alloymatrix that allows some mechanisms to be identified as strong contributors.123 A brief summary of current state of knowledge relevant to the purpose of this review is provided in the following sections. 1.0 0.5 0.2 0.1 S tr es s re d uc tio n ra tio 0 2 4 dpa 0 0.5 1.0 1.5 300 �C 60 �C Fluence (n m-2) (E > 1.0 MeV) 172 < s0< 262 MPa 2.0 6 8 ´ 1024 Figure 66 Irradiation-induced stress relaxation of X-750 bent beams in the NRU reactor at two temperatures, showing a greater relaxation at 60 �C due to an increased creep rate compared to that at 300 �C. Reproduced from Causey, A. R., Carpenter, C. K. C.; MacEwen, S. R. J. Nucl. Mater. 1980, 90, 216–223. Similar behavior in this study was observed in pure nickel and to a lesser extent in 304 stainless steel. 4.02.9.2 Stages of Irradiation Creep Irradiation creep of austenitic steels can be envisaged as having four stages. These are the transient regime, the creep regime in the absence of swelling, swelling- enhanced creep, and creep disappearance. The first three contributions have traditionally been described using the following equation.1 _2 s ¼ A½1� expð�dpa=tÞ� þ B0 þDS : : The equivalent strain per unit equivalent stress, sometimes called the creep modulus B, is the sum of a transient contribution that saturates usually at a dpa or less, the creep compliance B0 in the absence of swelling, and stress-enhanced creep where the enhanced creep rate is proportional to the void swelling rate. Bubble swelling also accelerates irradi- ation creep, but the influence is expressed primarily in the early stages of creep.167 The coefficients A and t are empirical, experimen- tally determined constants that are very material- specific (composition, thermal–mechanical treatment and texture) and sometimes stress-state dependent. For many high-exposure applications, the transient can be ignored. Even more importantly, the transient term can be obscured if significant precipitation- related strains are developing concurrently. In gen- eral, the magnitude of the transient-induced strain increases with increasing stress, but the duration in time or dpa usually does not increase with stress. The creep transient involves both the dose needed to establish equilibrium densities of point defects, but most importantly, it involves the dose required for establishment of the quasi-equilibrium dislocation density. The transient is most pronounced for cold- worked steels that start at higher dislocation density. As recombination–annihilation mechanisms reduce the dislocation density the instantaneous creep rate drops until the quasi-equilibrium dislocation density is reached and the steady-state creep rate B0 is estab- lished. For any austenitic steel it can be safely assumed that B0 is �1� 10�6 (MPa dpa)�1 and is effectively a ‘crystal constant’ similar to the 1% per dpa swelling rate of austenitic steels. The creep compliance B0 is not only independent of composition but also starting state, dpa rate, and temperature over the range of reactor-relevant condi- tions. There appears to be one exception, however, in that the creep rate at temperatures somewhere below �100 �C can increase significantly above B0, as shown at 60 �C by Grossbeck and Mansur for various austenitic steels.168,169 At very low temperatures, vacancies are relatively immobile and cannot cancel the climb of dislocations produced by more mobile interstitials. Examples of such behavior have been seen in other studies.170,171 A good example of acceler- ated creep at low temperatures is shown in Figure 66. The swelling–creep coupling coefficient D was originally assumed also to be a crystal constant at �0.6� 10�2MPa,�1 but as discussed later, it was found that D declines with increasing swelling to approximately one-third of this value or even to zero, depending on the stress state, stress history, and swelling history. This decline is an expression of the fourth stage of irradiation creep, variously designated as creep cessation, creep disappearance, or creep damping. One feature of irradiation creep that distinguishes it as different from thermal creep is that it varies with stress to power 1.0 rather than a higher power typical of thermal creep. This linearity is shown in Figure 67. The creep equation presented above also predicts Radiation Damage in Austenitic Steels 79 that the creep rate is proportional to dpa both before and after swelling begins. As discussed later, some important characteristics of creep have been redefined in the past two decades, especially for the creep compliance. B0 is known to be generally independent of alloy composition, thermal–mechanical treatment, irradiation tempera- ture, and dpa rate, but swelling is known to be very sensitive to all of these variables. This means that the irradiation creep modulus B quickly assumes all of the parametric sensitivities of void swelling. When the swelling rate reaches only 0.017% per dpa the swelling-enhanced contribution equals the B0 contri- bution, effectively doubling the creep rate. There are a number of consequences of the cou- pling between swelling, creep, and precipitation- related strains. 1. The onset of swelling can be detected by a jump in creep modulus B long before measurable swelling- induced changes in dimension can be detected, and often before microscopy confirms the pres- ence of voids. 2. Attempts to measure B0 in the presence of low and sometimes undetectable levels of voids or bubbles will lead to misleading values, usually higher than �1� 10�6 (MPa dpa)�1. 12 PCA 384–406 �C 405 �C, 24.4 ´ 1022 354 �C, 18.4 ´ 1022 390 �C, 15.1 ´ 1022 316 �C, 11.6 ´ 1022 401 �C, 6.7 ´ 1022 406 �C, 4.6 ´ 1022 n cm- 3 (E > 0.1 MeV)10 8 6 4 2 0 0 50 100 150 Hoop stress (MPa) 200 250 300 (% ) DD D 0 Figure 67 Linear stress dependence of total diametral strain (creep and swelling) for 20% cold-worked PCA (Ti-modified 316 stainless) pressurized tubes irradiated in FFTF at 400 �C. Reproduced from Garner, F. A.; Puigh, R. J. J. Nucl. Mater. 1991, 179–181, 577–580. Stress-free swelling is approximately three times the Y-intercept value with the largest swelling at �8%. 3. Any local stress gradient generated by a swelling gradient will be reduced to a very low level by a local gradient of creep exactly matched to that of swelling. 4. Attempts to measure B0 in the presence of precipitate-related strains will lead to mislead- ingly different values, either too large, too small, and even negative values. 5. Whenever the stress state is generated solely by swelling, the coupling between creep and swelling guarantees that the system cannot operate at a stress level higher than D�1 or �160MPa.1,16 4.02.9.3 Examples of Creep Behavior Various aspects of behavior resulting in irradiation creep can be illustrated with some examples pre- sented in Figures 68–75. 4.02.9.4 Creep Disappearance The previous figures demonstrate the swelling–creep correlation at its simplest when swelling is either zero or just beginning, but not yet provoking the next shift in quasi-equilibrium. When looking across a wider 1.2 D /D e 1.0 0.8 0.6 0.4 0.2 0.0 0 1 2 3 ´ 1024 Fluence (n m–2) (E > 1 MeV) 43.3 MPa 26.6 MPa FV548 52.2 MPa PE1656.0 MPa 57.0 MPa Figure 68 Creep-induced deflections of helical springs constructed from two steels with different composition that were irradiated in DMTR at 100 �C, normalized to the elastic deflection, showing that both the transient and steady-state creep rate B0 are proportional to the stress level. While the transients are different in the two steels, the posttransient creep rate is independent of composition. Reproduced from Lewthwaite, G. W.; Proctor, K. J. J. Nucl. Mater. 1973, 46, 9–22. The maximum dose is �0.5 dpa. 0.0 7 6 5 4 3 2 1 0 0.1 0.2 0.3 0.4 Dose (dpa) S tr ai n/ st re ss (0 .0 00 1 % /M P a) 0.5 0.6 0.7 0.8 175–200 �C Annealed 304 20% CW 316 at 370 �C 300–370 �C 2 ´ 10-8 dpa s–1 Figure 69 Irradiation creep of austenitic steels during uniaxial testing in the K Reactor, showing independence of creep of annealed 304 on temperature in the range 175–370 �C. Reproduced from Foster, J. P.; Gilbert, E. R.; Bunde, K.; Porter, D. L. J. Nucl. Mater. 1998, 252, 89–97; Gilbert, E. R. Reactor Technol. 1971, 14, 258–285. B0 is 0.92� 10�6 (MPa dpa)�1. The larger transient of cold-worked 316 is due to its much higher dislocation density. 370 �C, 130 MPa 380 �C, 50 MPa Le ng th c ha ng e (% ) 420 �C, 130 MPa 390 �C, 100 MPa 320 �C, 100 MPa 350 �C, 100 MPa 316 PCA AMCR 0033 0.06 (a) (b) (c) 0.04 0.02 −0.02 0.00 0.15 0.05 −0.05 0.0 0.15 0.10 0.05 0 0 1 2 dpa 3 4 0.1 AMCR 0033 20% Cold-worked AMCR 0033 As received Aged 400 �C, 1 h Aged 600 �C, 1 h 20% Cold-worked Figure 70 Length changes observed in HFR during uniaxial creep tests of (a) three different cold-worked steels at 370 �C; (b) 20% cold-worked AMCR 0033 at different irradiation temperatures; (c) 20% cold-worked AMCR 0033 in different starting conditions. Reproduced from Hausen, H. Schüle, W.; Cundy, M. R. Fusion Technol. 1988, 88, 905–909. Note that precipitate-induced strains can be positive or negative, and vary with composition, starting condition, and irradiation temperature. The posttransient creep rate is not sensitive to these variables, however. 80 Radiation Damage in Austenitic Steels � 10−6MPa−1dpa−1 C = 0.046 wt% C = 0.006 wt% Fluence (n m−2) (E > 0.1 MeV) 12 � 1026840 0 2 4 C re ep c oe ffi ci en t B 6 Figure 71 Acceleration of irradiation creep in two carbon variants of a stainless steel by a low rate of swelling at 350 and 420 �C. Reproduced from Neustroev, V. S.; Shamardin, V. K. In Effects of Radiation on Materials: 16th International Symposium; 1993; pp 816–823. The lower carbon steel has a longer transient regime of swelling. The height of the plateau is determined by the swelling rate. B0 was determined to be �1� 10�6 (MPa dpa)�1 and D to be 0.6�10�2MPa�1. Swelling Irradiation creep 400–420 �C 400–420 �C 400–420 �C 450–470 �C 450–470 �C 450–460 �C 316Ti 316Ti+P 170 MPa C re ep s tr ai n (% ) 0 2 4 DV V 0 6 8 6 4 2 0 90 MPa 0 20 40 60 dpa 80 100 120 (% ) Figure 73 Swelling and creep strains observed in two French steels irradiated as pressurized tubes in PHENIX, showing strong correlation between the two types of strain as the swelling rate increases. Reproduced from Dubuisson, P.; Maillard, A.; Delalande, C.; Gilbon, D. D.; Seran, J. L. In Effects of Radiation on Materials: 15th International Symposium; STP 1125; 1992; pp 995–1014. 20% CW 316 25% CW PCA 13.3 dpa MPa−1dpa−1 BO= 2.8 � 10 −6 MPa−1dpa−1 BO= 3.2 � 10 −6 300 �C 400 �C 500 �C 600 �C 13.1 dpa 12.0 dpa 330 �C 400 �C 500 �C 600 �C 12.1 dpa 0 0.0 1.0 2.0 1.0 2.0 0.0 100 200 Effective stress (MPa) 300 400 500 e (% ) e (% ) Figure 72 Temperature-independent creep strains observed in 20% cold-worked 316 and 25% cold-worked PCA during irradiation in the ORR test reactor at a high He/ dpa ratio. Reproduced from Grossbeck, M. L.; Horak, J. A. J. Nucl. Mater. 1988, 155–157, 1001–1005. Note that the two steels have very similar values of creep modulus B and are independent of irradiation temperature over a wide range. The creep modulus B is about three times that of B0¼ 1�10�6 (MPa dpa)�1, however, probably arising from observed high densities of helium bubbles to produce bubble-enhanced creep. Radiation Damage in Austenitic Steels 81 range of swelling behavior some unusual behaviors are often observed. An example is shown in Figure 76 where the two-peaked swelling behavior frequently observed in 300 series steels is mirrored in the creep strains, but the relative proportions of the two strains are distorted.172 This is one manifestation of the creep disappearance phenomenon in which the attainment of significant swelling causes irradiation creep to strongly drop in rate or even to disappear under some conditions as seen in Figures 77 and 78. In early fuel pin studies it was often observed that irradiation creep strains would increase and then abruptly decrease and sometimes stop entirely, even though fission gas pressures continued to increase.173,174 These results were interpreted as evi- dence of fuel swelling very quickly to meet and thereby put stress on the cladding but later the onset of swelling in the clad caused it to out-swell the fuel and break contact. Actually, the driving force 2.0 1.0 0.5 0 0 50 dpa dpa M id w al l c re ep s tr ai n/ ho op s tr es s (% p er M P a) S w el lin g- in d uc ed d ia m et ra l s tr ai n (% ) 100 150 1.5 83508, T-420 �C K280, T-395 �C A095, T-415 �C 2.5 0.05 0 MPa 60 MPa 83508 K280 A095 Failed in next cycle Failed in next cycle 100 MPa 140 MPa 200 MPa 300 MPa 0.04 0.03 0.02 0.01 0 0 50 100 150 A094, T-415 �C C42, T-415 �C C38, T-390 �C C39, T-390 �C C40, T-390 �C C44, T-390 �C Figure 74 (left) Diametral strains resulting from void swelling at 400 �C in neutron-irradiated stress-free tubes constructed from nine titanium-modified 316 stainless steels, (right) Stress-normalized midwall creep strains observed in three of these steels, showing a strong correlation of swelling and irradiation creep rates in each steel. Reproduced from Toloczko, M. B.; Garner, F. A.; Eiholzer, C. R. J. Nucl. Mater. 1992, 191–194, 803–807. 10 20 Ni-equivalent (%) 30 6.3 dpa 10 dpa 56 dpa 40 500 0 B E Q � 10 -6 (M P a d p a) -1 2 4 6 8 10 12 14 16 Figure 75 Creep modulus measured for six austenitic steels irradiated in BOR-60 fast reactor at 420 �C, showing an enhancement of creep versus Ni-equivalent. Reproduced from Neustroev, V. S.; Shamardin, V. K. J. Nucl. Mater. 2002, 307–311, 343–346. This behavior corresponds to the known effect of nickel on void swelling, indicating swelling-enhanced creep. 82 Radiation Damage in Austenitic Steels was primarily increasing levels of fission gas but irradiation creep had disappeared by ~7% burn-up. Several features of creep disappearance are noteworthy. 1. The combined creep and swelling strain rate in a fuel pin or pressurized tube cannot exceed 0.33% per dpa or one-third of the eventual steady-state swelling rate. 2. As swelling approaches 1% per dpa the creep rate backs down proportionately to maintain this max- imum rate as shown in Figures 78–80. 3. The limit of 0.33% per dpa is reached before swelling gets to a significant fraction of 1% per dpa, as shown in Figure 80. Some tubes had already reached the maximum strain rate limit, but then lost their gas pressure and continued to swell at less than 1% per dpa. 4. As the creep cessation process gets underway the creep strain loses its responsiveness to the magnitude of the stress. Note in Figures 79 and 80 that doubling the hoop stress did not double the strain rate in the tube. 5. The coupling coefficient D tends to fall to zero rather quickly when swelling-before-creep occurs but falls more slowly in creep-before-swelling scenarios (fuel pins vs. pressurized tubes).175 A consensus explanation of the creep disappearance phenomena has not yet been reached. Various models have been proposed involving voids acting to erase the anisotropy of dislocation Burgers vector176,177 and the involvement of precipitate sinks to serve as strong sinks that compete with dislocations.175 8 4 2 6´10-6 0 4 2 0 400 500 600 Temperature (ºC) 700 Pin 47 Pin 47 Pin 1 Pin 1 Pin 32 Pin 32 Pin 31 Pin 31 Bottom Fuel column length Top 6 4 2 0 1.2 0.8 C re ep s tr ai ns (% ) C re ep m od ul us (M P a d p a F) -1 S w el lin g (% ) S w el lin g (% ) 0.4 0 Figure 76 Swelling and creep behavior observed along the length of AISI 316 fuel pins irradiated in the RAPSODIE fast reactor; (left) solution annealed and (right) 20% cold-worked. Reproduced from Boutard, J. L.; Carteret, Y.; Cauvin, R.; Guerin, Y.; Maillard, A. In Proceedings Conference on Dimensional Stability and Mechanical Behavior of Irradiated Metals and Alloys; British Nuclear Energy Society: London, 1983; pp 109–112. 30 20 10 0 0 10 1030 20 30 dpa Instantaneous creep coefficient Swelling-enhanced creep psi n cm-2 Swelling in the absence of creep Onset of creep disappearance 40 50 60 Figure 77 Instantaneous creep coefficient B derived from strain measurements on pressurized tubes constructed from a double-aged higher-swelling condition of 316 stainless steel irradiated in EBRII at 550 �C. Reproduced from Porter, D. L.; Garner, F. A. J. Nucl. Mater. 1988, 159, 114–121. Radiation Damage in Austenitic Steels 83 4.02.9.5 Recent Revisions in Understanding of Irradiation Creep 4.02.9.5.1 Dependence of irradiation creep on dpa rate As mentioned earlier, once swelling begins, irradia- tion creep quickly assumes all the parametric depen- dencies of void swelling. However, for many years it was assumed that the B0 component of creep was also strongly dependent on dpa rate, increasing as the dpa rate fell, as shown in Figure 81. The original research that established this percep- tion was performed by Lewthwaite and Mosedale on various cold-worked steels in the Dounreay Fast Reactor at temperatures in the 270–350 �C range.178 6 4 487-543 �C Swelling deformation Plastic deformation Total 2 0 0 2 4 6 8 Local atomic burnup (%) 10 12 DD D (% ) Figure 78 Creep and swelling strains observed in annealed 347 stainless clad fuel pins irradiated in EBR-II, showing the disappearance of further creep strain as irradiation continues. Reproduced from Appleby, W. K.; Hilbert, R. F.; Bailey, R. W. In Proceedings Conference on Irradiation Embrittlement and Creep in Fuel Cladding and Core Components; British Nuclear Energy Society: London, 1972, pp 209–216. These data were originally explained in terms of fuel-clad interaction acting as the major source of stress in the cladding, with fuel contact and stress-driven creep eventually terminated by the onset of clad swelling to move the clad away from the fuel. Continually increasing gas loading was actually the primary loading on the cladding, not the fuel. 10 8 6 4 2 -2 -2 0 8 6 4 2 0 0 20 40 60 0.33% per dpa 0.33% per dpa 30 ksi Hoop stress = 0 ksi 30 ksi 15 ksi (a) (b) Irradiation creep Stress- free swelling Stress-affected swelling at 30 ksi 80 100 dpa D ia m et er c ha ng e (% ) Figure 79 (a) Deformation observed in pressurized tubes of 20% cold-worked AISI 316 irradiated in EBR-II at 550 �C. Reproduced from Porter, D. L.; Garner, F. A. J. Nucl. Mater. 1988, 159, 114–121; Porter, D. L.; Garner, F. A. In Effects of Radiation on Materials: 13 International Symposium (Part II) Influence of Radiation on Material Properties; ASTM STP 956; 1987; pp 11–21. Note that doubling the hoop stress from (from 15 to 30 ksi: 103 to 206MPa) does not double the deformation rate, which never exceeds 0.33% per dpa. (b) Density measurements on the 30 ksi (206MPa) tube show that stress accelerates the rate of swelling, but also causes the creep rate to approach zero at high swelling levels. 84 Radiation Damage in Austenitic Steels The explanation advanced for such a dependence was the decreasing amount of annihilation of point defects by recombination at lower dpa rates, where such an effect is expected to be more pronounced at the lower irradiation temperatures characteristic of this experiment. An earlier review article was published where this and other data sets were assessed to determine the appropriate rate dependence.1 Some data sets avail- able at that time supported a flux dependence and other data sets supported an independence of dpa rate. On balance it appeared that a strong depen- dence of irradiation creep rate on dpa rate was the more defendable conclusion. With hindsight and additional published data sup- porting the opposite conclusion, it was later realized that apparent dependence of creep rate on dpa rate was an artifact of the analysis procedure used by Mosedale and Lewthwaite. The authors had not prop- erly separated the transient and post-transient strains, and all of the lower flux data were in the higher-rate transient regime. When the DFR creep data were reanalyzed by Garner and Toloczko, the creep com- pliance B0 was found to be independent of dpa rate. 179 4.02.9.5.2 Dependence of creep and creep relaxation on neutron spectra It is sometimes assumed that thermalized neutron spectra can produce more effectively surviving point defects since gamma-recoil events do not pro- duce cascades and therefore there is less in-cascade annihilation. Thus, a larger fraction of thermally pro- duced defects are postulated to survive to contribute to creep or embrittlement.180,181 12 550 �C Core 1 575 �C Core 1 575 �C Core 4 550 �C Core 4 8 0.33% per dpa 4 0 12 8 4 0 0 20 40 60 0 20 dpa 40 60 80 0 MPa 35 MPa 70 MPa 117 MPa 163 MPa 233 MPa 0 MPa 16 MPa 31 MPa 63 MPa 104 MPa 146 MPa DD D (% ) Figure 80 Diametral strains observed in two related heats of 20% cold-worked AISI 36 irradiated in FFTF as pressurized tubes. Reproduced from Garner, F. A.; Toloczko, M. B.; Puigh, R. J. In Effects of Radiation on Materials: 19th International Symposium; ASTM STP 1366; 2000; pp 667–678. Note that many of the tubes have reached the limiting deformation rate of 0.33% per dpa. Those tubes which subsequently fail show that swelling had not yet reached its limiting rate of 1% per dpa. Radiation Damage in Austenitic Steels 85 Foster and coworkers have published three papers over the past several decades where it appeared that irradiation creep indeed occurred at a higher rate in thermal reactors than in fast reactors.182–184 In the last of these papers it was noted that, as proposed by Garner 34 the previously unsuspected 59Ni con- tributions to dpa might account for the apparent but possibly misleading increase in creep rate. The T/F ratio in the experimental test reactors cited by Foster was rather high compared to that in PWRs. An additional reason for such enhancement of creep probably lies in the large amounts of trans- muted helium and stored hydrogen in thermalized spectra that results from the 59Ni sequence and the stored hydrogen concept, producing bubbles and voids that accelerate the creep rate. Therefore, it does not appear necessary to invoke an enhanced survivability or displacement effectiveness of gamma recoil events to explain the apparently higher creep rates in thermal reactors. 4.02.9.5.3 Dependence of creep modulus on hydrostatic stress Although it is well known that it is the deviatoric component of any stress state that drives creep, there were previously very little data to show whether the creep coefficient is identical in both dilational and compressive stress states. Recent papers by Hall,185,186 Neustroev,187 and Garzarolli188 show that creep coefficients are unchanged by the sign of the hydrostatic stress. As shown in the next section, additional confirma- tion of the independence of creep compliance on the sign of the hydrostatic stress component can be found in some stress relaxation experiments. 4.0 3.0 2.0 1.0 0.6 0 1 2 3 Displacement rate (dpa s–1) N or m al iz ed c re ep r at e 4 5 6 ´ 107 4.8 J EN58 347 S.S. 240–360 �C Annealed M316 T < 304 �C Mk 1 helices H6 helices (1) (2) (3) FV548 E B Figure 81 Dependence of irradiation creep rate of springs made from various austenitic steels on dpa rate in and below the DFR core, normalized to the highest displacement rate studied. Reproduced from Lewthwaite, G. W.; Mosedale, D. J. Nucl. Mater. 1980, 90, 205–215. 0.5 s0= 216 MPa 0 1.0 0.9 0.7 0.5 0.3 0 1 2 3 Neutron dose (dpa) 4 5 S tr es s re d uc tio n ra tio S tr es s re d uc tio n ra tio 0 0.5 1.0 1.5 dpa Preload 23.6 N 36.5 N 2.0 1.0 Figure 82 (top) Stress relaxation experiment conducted on X-750 in the NRU heavy-water reactor at 300 �C using constant curvature bent beams. Reproduced from Causey, A. R., Carpenter, C. K. C.; MacEwen, S. R. J. Nucl. Mater. 1980, 90, 216–223; (bottom) stress relaxation of compressed springs in EBR-II at 375–415 �C. Reproduced from Walters, L. C.; Reuther, W. E. J. Nucl. Mater. 1977, 68, 324–333. 86 Radiation Damage in Austenitic Steels 4.02.9.6 Stress Relaxation by Irradiation Creep There are situations where the applied load is initi- ally fixed and then declines during irradiation. There is usually a transient followed by an instantaneous creep rate defined by B0, but the load is constantly falling, leading to an exponentially declining load. Two examples of in-reactor creep relaxation experi- ments are shown in Figure 82, both conducted on a high-nickel alloy Inconel X-750. Foster and coworkers have very convincingly demonstrated that creep coefficients derived from creep experiments could be used to successfully pre- dict stress relaxation for the same steel in similar neutron spectra.163 Note that the creep coefficient derived for X-750 from the EBR-II experiment is 1.6� 10�6 (MPa dpa)�1, just slightly larger than B0 and probably enhanced by low levels of voids or bubbles in this high-nickel alloy. In NRU, however, the creep relax- ation proceeded much faster, partially due to a larger transient, but also because the steady-state creep rate is larger. In this experiment the thermal-to-fast ratio was �10, so there was significant 59Ni enhancement of dpa rate and probably also bubble formation to enhance the creep rate. The greater scatter at very low residual stresses in the EBR-II experiment is mostly due to frictional variations on the compressed Radiation Damage in Austenitic Steels 87 springs and grain-to-grain interactions that come into play at low stress levels. Stress relaxation experiments can be conducted using a wide variety of specimen types and usually yield similar results, although the transient regimes often vary with specimen geometry, preparation, and texture versus stress field relationship, as shown in Figure 83. 0 0 1 2 Dose (dpa) Bend C ring 304 Irradiated temperature: - 561 K (A2= 1.8 ´ 10-8MPa dpa–1) (A2= 1.2 ´ 10-6MPa dpa–1) S tr es s ra tio (a /d ) 3 4 5 1 0.8 0.6 0.4 0.2 Figure 83 Stress relaxation experiments conducted on 304 stainless steel at 288 �C in water-cooled JMTR at 0.82–1.7�10�7 dpa s�1, showing creep coefficients close to B0, and also demonstrating different transient behavior in different test geometries. Reproduced from Ishiyama, Y.; Nakata, K.; Obata, M.; et al. In Proceedings of 11th International Conference on Environmental Degradation of Materials in Nuclear Systems; 2003; pp 920–929. 250 3W-H200 Before irradiation After 3–6 dpa irradiation 150 100 50 0 0 10 20 30 40 50 Distance from surface, 4 mm Distance from left edge (mm) s y ( M P a) 6 -50 -100 -150 Figure 84 Residual stresses in SA 304 associated with a one-p with depth from the surface. Reproduced from Obata, M.; Ishiya J. ASTM Int. 2006, 3, 15–31. Residual stresses before and after B0 was determined to be �1�10�6 (MPa dpa)�1 and is indepen Creep relaxation by irradiation is important in that it can reduce the opportunity for irradiation- assisted stress corrosion cracking. It does so by decreasing internal or surface stresses produced by deliberate or inadvertent damage, as well as by reducing internal stresses arising from welding, abrupt cooling, etc. Figure 84 demonstrates the radiation-induced relaxation that occurs in a weld that proceeds with a creep compliance of B0 that is independent of the sign of the hydrostatic stress.189 Therefore, it appears that the creep compliance B0 can be confidently applied to any stress state. As a rule of thumb one can anticipate that by 10 dpa, >90% of any preload will be relaxed even in the absence of a transient. The fractional unload- ing is not dependent on the magnitude of the preload as long as the bolt or component was not loaded beyond the yield point. Stress relaxation in structural components of operating reactors is not always operating in isolation. Frequently, a component experiences time-dependent stresses that develop with time as a result of the growth or movement of adjacent components. In pressurized water reactors there are bolts that join baffle plates to former plates. These bolts are usually cold-worked 316 but the plates they join are annealed 304 stainless, a higher swelling steel. Initially, the bolt will start to relax its preload but if the plates are swelling faster than the bolts, then differential swelling will begin to reload the bolt. Additionally, if a bolt is replaced with a fresh bolt, the reloading can be even stronger due to larger amount of 0 1 0.8 0.6 0.4 0.2 0 0 1 Tensile (6 mm from surface) Irradiation temperature: 561 K Compressive (4 mm from surface) 2 Dose (dpa) S tr es s re la xa tio n s/ s 0 3 4 sy ass weld with mechanical constraint. Stress reversals occur ma, Y.; Nakata, K.; Sakamoto, H.; Anzai, H.; Asano, K. irradiation were measured by neutron diffraction. Note that dent of the sign of the hydrostatic stress. 350 300 250 200 150 A xi al s tr es s (M P a) 100 50 0 0 10 20 30 40 Time (year) 50 60 70 10 years 40 years Figure 85 Calculated bolt relaxation and reloading is shown for two conditions of bolt replacement in a 304 stainless baffle-former assembly. Reproduced from Simonen, E. P.; Garner, F. A.; Klymyshyn, N. A.; Toloczko, M. B. In Proceedings of 12th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors; 2005; pp 449–456. The cold-worked 316 bolt is replaced and reloaded at either 10 or 40 years. Note that differential swelling does not reverse the loading until almost 10 dpa as the bolt approaches full relaxation. 200 To rq ue (n m ) 180 160 140 120 100 80 60 0 10 20 Dose (dpa) 30 40 A B C Figure 86 Torques measured during removal of bolts from French PWRs of the CPO series. Only bolts showing no indication of cracking are included. The results are in agreement with predicted creep relaxation when applied to upper or lower preload values, but the predictions do not include any reloading. A, B, and C denote measurements from three different CPO plants. Reproduced fromMassoud, J. P.; Dubuisson, P.; Scott, P.; Ligneau, N.; Lemaire, E. In Proceedings of Fontevraud; 2002; Vol. 5; paper 62, 417. 88 Radiation Damage in Austenitic Steels differential swelling. Figure 85 shows several calcu- lated histories of bolt loading for PWR-relevant tem- peratures and dpa rates.190 While bolts are generally preloaded to a specified level, there is always some range of attained loads. It is difficult to measure the stress level in a bolt while it is still in place, but a rough measure of the remaining load can be made from the torque required to remove the bolt. While this is not an exact measurement with friction, corrosion, irradiation-induced self-welding, and other complications possibly participating to define the torque, Figure 86 shows that the measured torques are in reasonable agreement with predictions of creep equations based on experiments conducted in BOR-60 fast reactor. The fact that most of the data lie above the predictions may indicate that many of the bolts are indeed being reloaded by differential swelling to some degree. 4.02.9.7 Stress Rupture While irradiation creep is relatively well understood the effect of radiation on thermal creep and thereby creep rupture is not as well defined. In general it appears that creep rupture properties are not improved by irradiation and are adversely affected as shown in the example of Figure 87.191,192 As shown in Figure 88 Ukai and coworkers have compared the reduction in rupture life in air, sodium, and after irradiation in FFTF, demonstrating that the largest influence is due to irradiation.193 There is some evidence that irradiation in neutron spectra that pro- duce high He/dpa ratios will decrease rupture life, especially at higher temperatures, compared to irradi- ation in fast reactors due to the accumulation of helium bubbles on grain boundaries and triple points.191,192 It is possible to improve the in-reactor stress rupture properties of a given steel by additions of selected trace elements such as P and B, both of which are known to affect the distribution and stability of carbide phases. An example is shown in Figure 89.194 Fortuitously, such additions also add to the swelling resistance of such steels. Radiation Damage in Austenitic Steels 89 4.02.9.8 Fatigue Fatigue loading can be very detrimental for situations involving cyclic loading, especially when associated with thermal cycling such as might occur in the first wall of a fusion device. As shown in preceding sections, radiation changes the microstructure and affects the phase stability of steels as well as generating deleterious gases such as helium and hydrogen. 500 300 100 H oo p s tr es s (M P a) 80 60 14 14.5 15.515 1 LMP = T (14.04 Irradiation effect Figure 88 Creep rupture behavior of 20% cold-worked modif irradiation to reduce failure lifetimes. Reproduced from Ukai, S.; 320–327. 102 101 103 700 �C Cold-worked Irradiated in BR-2 1.3 1.4 S tr es s (M P a) 1.5 1.6 T(13.5 + log tR) (K) 1.7 1.8 Thermal aged Annealed Figure 87 Effect of starting condition and irradiation in the BR-2 reactor on stress rupture behavior of DIN 1.4970 at 700 �C. Reproduced from Wassilew, C.; Ehrlich, K.; Bergmann, H. J. In Influence of Radiation on Material Properties: 13th International Symposium; ASTM STP 956; 1987; pp 30–53; Grossbeck, M. L.; Ehrlich, K.; Wassilew, C. J. Nucl. Mater. 1990, 174, 264–281. Data are plotted versus the Larson Miller Parameter (LMP). The effect of radiation is stronger than the effect of cold-working. Therefore it is not unexpected that fatigue life will be adversely affected by irradiation as shown in Figure 90.192 Fatigue tests are by necessity conducted out-of- reactor and therefore are not fully representative of in-reactor conditions, especially not being subject to the mitigating influence of radiation creep to reduce local stress concentrations. In this sense out-of- reactor results may be conservative. The tests can be conducted in a variety of ways, however, generally using either strain-controlled or load-controlled methods, with the former being more relevant to low cycle fatigue arising from thermal cycling. Guidance on the application of fatigue data is provided by Tavassoli.195 Figure 90 presents the usual engineering curves of total strain versus the number of cycles to failure. In this representation the lifetimes of irradiated and unirradiated materials are not really so dissimilar. The observed difference is the result of competing influences, degradation due to irradiation, and improvement due to hardening. As pointed out by Boutard,196 it is better to isolate the irradiation effect on the lifetime in which the controlling parameter is the plastic strain range. As shown inFigure 91, there is a significant effect of radiation on the lifetime at a given plastic strain.196,197 The lower the plastic strain, the greater the decrease in lifetime. Under conditions where the crack initiation phase controls the lifetime of the unirradiatedmaterial, irradiation will result in much earlier crack formation 16.56 + log tR) / 1000 17 18 Sodium effect In-air In-sodium In-reactor 17.5 ied 316 stainless steel, showing effect of sodium and Mizuta, S.; Kaito, T.; Okada, H. J. Nucl. Mater. 2000, 278, 90 Radiation Damage in Austenitic Steels and much earlier failure. Other researchers have reached the same conclusion.198 In general it appears that most researchers agree that helium is a contributing but not primary cause of the radiation-induced degradation in lifetime.195–199 4.02.10 Conclusions In general there are no beneficial aspects of radia- tion on austenitic steels when exposed to neutron 10.0 1.0 To ta l s tr ai n ra ng e, D 1 (% ) 0.1 Cycles 102 103 104 1 ' Figure 90 Fatigue life of 20% cold-worked AISI 316 stainless 900appm He. Reproduced from Grossbeck, M. L.; Ehrlich, K.; W 1000 100 H oo p s tr es s (M P a) 10 12 14 D9I LMP, T (13.5 + l D9 575 �C 575 �C 605 630 695 775 670 750 Figure 89 Improvement of in-reactor [FFTF fast reactor] stres additions of B and P. Reproduced from Hamilton, M. L.; Johnson on Residual Elements in Steel; ASTM STP 1042; 1989, pp 124–1 irradiation. Structural components used in various nuclear reactors may have been constructed from alloys with carefully tailored and optimized proper- ties, but there is an inevitable degradation of almost all engineering properties of interest as irradiation proceeds. Even more importantly, having labored to build a device with well-defined dimensions, separations, and tolerances, it must be recognized that these dimensional attributes can also change dramatically, requiring that the design anticipate such changes in order to maximize safe and efficient operation for the longest possible lifetime. - Unirradiated - Unirradiated, aged 115 days at 430 �C - f1 = 0.7-2 ´ 1026n m–2 to failure 05 106 107 108 steel irradiated in HFIR to a maximum dose of 15dpa and assilew, C. J. Nucl. Mater. 1990, 174, 264–281. 16 18 D9I og tR) ´ 10-3 (K, h) D9 20 s rupture properties of D9 stainless steel by controlled , G. D.; Puigh, R. J.; et al. In Proceedings ASTM Symposium 49. Number of cycles to failure 103 10-1 100 10-2 104 2 5 2 5 EC - 316L : 430 °C Nonirrad. Irrad.: 10 dpa 105 106 P la st ic s tr ai n ra ng e (% ) Figure 91 Plastic strain versus number of cycles to failure of annealed EC-316L irradiated to 10dpa at �430 �C in BR2. Reproduced from Grossbeck, M. 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All rights reserved. 4.03.1 Introduction 97 4.03.2 Basic Metallurgy of Ferritic–Martensitic Steels 98 4.03.3 Radiation Damage of Core Components in Fast Reactors 101 4.03.4 Development of Ferritic Steels for Fast Reactor Core 102 4.03.4.1 Influence of Composition and Microstructure on Properties of Ferritic Steels 103 4.03.4.2 Void Swelling Resistance 105 4.03.4.3 Irradiation Hardening in Ferritic Steels 106 4.03.4.4 Irradiation Creep Resistance of Ferritic Steels 108 4.03.4.5 Irradiation Embrittlement in Ferritic Steels 110 4.03.4.5.1 GBE to reduce embrittlement in ferritic steels 112 4.03.5 Development of Advanced ODS Ferritic Steels 114 4.03.6 Ferritic Steels for Out-of-Core Applications: Improvements in Joining 116 4.03.7 Summary 119 References 119 Abbreviations bcc Body-centered cubic CSL Coincident site lattice DBTT Ductile to brittle transition temperature DICTRA Diffusion-controlled transformations dpa Displacements per atom EBR Experimental breeder reactor EBSD Electron back scattered diffraction fcc Face-centered cubic FFTF Fast flux test facility GBCD Grain boundary character distribution GBE Grain boundary engineering HAADF High angle annular dark field HAZ Heat-affected zone HFIR High flux isotope reactor ITER International Thermonuclear Experimental Reactor ODS steel Oxide dispersion strengthened steel PAGS Prior austenite grain size PFR Power fast reactor PWHT Postweld heat treatment RIS Radiation-induced segregation SIPA Stress-induced preferential absorption SIPN Stress-induced preferential nucleation TEM Transmission electron microscopy ▽DBTT Change in DBTT 4.03.1 Introduction The widespread acceptance of nuclear energy depends1 on the improved economics, better safety, sustainability, proliferation resistance, and waste man- agement. Innovative technological solutions are being arrived at, in order to achieve the above goals. The anticipated sustainability, rapid growth rate, and eco- nomic viability can be ensured by the judicious choice of fast reactor technology with a closed fuel cycle option. The fast reactor technology has attained (http://www.world-nuclear.org/info/inf98.html) a high level of maturity in the last three decades, with 390 years of successful operation. The emerging inter- national collaborative projects (http://www.iaea.org/ INPRO/; http://www.gen4.org/) have, therefore, cho- sen fast reactors as one of the important constituents of the nuclear energy in the twenty-first century. The nuclear community has been constantly striv- ing for improving the economic prospects of the technology. The short-term strategies include the development of radiation-resistant materials and extension of the lifetime of the components. The achievement of materials scientists in this field is remarkable. Three generations of materials have been developed,2 increasing the burn-up of the fuel from 45 dpa for 316 austenitic stainless steel to above 180 dpa for ferritic steels. Presently, efforts are in 97 http://www.world-nuclear.org/info/inf98.html http://www.iaea.org/INPRO/ http://www.iaea.org/INPRO/ http://www.gen4.org/ 98 Ferritic Steels and Advanced Ferritic–Martensitic Steels progress to achieve a target burn-up of 250 dpa, using advanced ferritic steels. The attempts by nuclear technologists to enhance the thermal effi- ciency have posed the challenge of improving the high temperature capability of ferritic steels. Addi- tionally, there is an inherent disadvantage in ferritic steels, that is, their susceptibility to undergo embrit- tlement, which is more severe under irradiation. It is necessary to arrive at innovative solutions to overcome these problems in ferritic steels. In the long time horizon, advanced metallic fuels and cool- ants for fast reactors are being considered for increasing the sustainability and thermal efficiency respectively. Fusion technology, which is ushering (http://www.iter.org/proj) in a new era of opti- mism with construction of the International Ther- monuclear Experimental Reactor (ITER) in France, envisages the use of radiation-resistant advanced ferritic steels. Thus, the newly emerging scenario in nuclear energy imposes the necessity to reevalu- ate the materials technology of today for future applications. The genesis of the development of ferritic steels is, indeed, in the thermal power industry. The develop- ment of creep-resistant, low alloy steels for boilers and steam generators has been one of the major activities in the last century. Today, the attempt to develop ultra super critical steels is at an advanced stage. Extensive research of the last century is responsible for identifying certain guidelines to address the concerns in the ferritic steels. The merit of ferritic steels for the fast reactor industry was established3 in the 1970s and since then, extensive R&D has been carried out4 on the application of ferritic steels for nuclear core component. A series of commercial ferritic alloys have been developed, which show excellent void swelling resis- tance. The basic understanding of the superior resistance of the ferrite lattice to void swelling, the nature of dislocations and their interaction with point defects generated during irradiation have been well understood. The strengthening and deformation mechanisms of ferrite, influence of various alloying elements, microstructural stability, and response of the ferrite lattice to irradiation temperature and stress have been extensively investigated. Themechanism of irradiation hardening, embrittlement and methods to overcome the same are studied in detail. Of the dif- ferent steels evaluated, 9–12%Cr ferritic–martensitic steels are the immediate future solution for fast reac- tor core material, with best void swelling resistance and minimum propensity for embrittlement. The high temperature capability of the ferritic steels has been improved from 773 to 973K, by launching the next generation ferritic steels, which are currently under evaluation for nuclear applica- tions, namely the oxide dispersion strengthened (ODS) ferritic steels (see Chapter 4.08, Oxide Dis- persion Strengthened Steels). Conceptually, this series of steels combines the merits of swelling resis- tance of the ferrite matrix and the creep resistance offered by inert, nanometer sized, yttria dispersions to enhance the high temperature limit of the ODS steels to temperature beyond 823K. The concerns of this family of materials include optimization of the chemistry of the host lattice, cost effective fabrication procedure, and stability of the dispersions under irra- diation, which will be discussed in this article. The present review begins with a brief introduc- tion to the basic metallurgy of ferritic steels, summar- izing the influence of chemistry on stability of phases, decomposition modes of austenite, different types of steels and structure–property correlations. The main thrust is on the development of commercial ferritic steels for core components of fast reactors, based on their chemistry and microstructure. Hence, the next part of the review introduces the operating conditions and radiation damage mechanisms of core compo- nents in fast reactors. The irradiation response of ferritic steels with respect to swelling resistance, irra- diation hardening, and irradiation creep are high- lighted. The in-depth understanding of the damage mechanisms is explained. The main concerns of fer- ritic steels such as the inferior high temperature irra- diation creep and severe embrittlement are addressed. The current attempts to overcome the problems are discussed. Finally, the development of advanced creep-resistant ferritic steels like the ODS steels, for fission and fusion applications are presented. The application of ferritic steels for steam generator cir- cuits and the main concerns in the weldments of ferritic steels are discussed briefly. The future trends in the application of ferritic steels in fast reactor technology are finally summarized. 4.03.2 Basic Metallurgy of Ferritic–Martensitic Steels The advanced ferritic and ferritic–martensitic steels of current interest have evolved5 from their prede- cessors, the creep-resistant ferritic steels, over nearly a century. The first of the series was the carbon and C–Mn steels with a limited application to about http://www.iter.org/proj 1500 Liquid + a + g Liquid 1400 1300 1200 1100 Te m p er at ur e (� C ) 1000 900 800 700 600 0 5 (a) 10 15 Chromium (%) 20 25 Liquid 9Cr steel a g a a + g a + g + (CrFe)7C3 a + (CrFe)7C3 a + g + (FeCr)3C a + (FeCr)3C a + (CrFe)4C a + (FeCr)3C + (CrFe)7C3 a + (CrFe)7C3 + (CrFe)4C Ae3 Martensite (b) Bainite Widmanstatten ferrite Ferrite Pearlite Upper bainite Lower bainite Log {time} Te m p er at ur e (� C ) Ws Ms Bs Figure 1 (a) Pseudobinary phase diagram for a Fe–Cr–C steel with 0.01% C. Reprinted, with permission, from High chromium ferritic and martensitic steels for nuclear applications, copyright ASTM International, 100 Barr Harbor Drive, West Conshohocken, PA 19428. (b) Decomposition modes of high-temperature austenite during cooling. Ferritic Steels and Advanced Ferritic–Martensitic Steels 99 523K. Subsequent developments through different levels of chromium, molybdenum have increased the high temperature limit to 873, leading to the current ferritic and ferritic–martensitic steels, that is, the 9–12% Cr–Mo steels. In addition to being economi- cally attractive, easy control of microstructure using simple heat treatments is possible in this family of steels, resulting in desired mechanical properties. The propensity to retain different forms of bcc ferrite, that is, ferrite or martensite or a mixture at room temperature in Cr–Mo steels, depends crucially on the alloying elements. Extent of the phase field traversed by an alloy on heating also depends on the amount of chromium, silicon, molybdenum, vana- dium, and carbon in the steel. The combined effect of all the elements can be represented by the net chromium equivalent, based on the effect of the aus- tenite and ferrite stabilizing elements. A typical pseu- dobinary phase diagram6 is shown in Figure 1(a). Increase in chromium equivalent by addition of ferrite stabilizers or V or Nb would shift the Fe–9Cr alloy into the duplex phase field at the normalizing temperature. The phase field at the normalizing tem- perature and the decomposition mode7–9 of high temperature austenite (Figure 1(b)) dictate the result- ingmicrostructure at room temperature and hence, the type of steel. Accordingly, the 9CrMo family of steels can either be martensitic (9Cr–1Mo (EM10) or stabi- lized 9Cr–1MoVNb (T91)), ferritic (12Cr–1MoVW (HT9)) or ferritic–martensitic (9Cr–2Mo–V–Nb (EM12)) steel. The stabilized variety of 9–12 CrMo steels could result10 in improved strength and delayed grain coarsening due to the uniform distribution of fine niobium or vanadium carbides or carbonitrides. The transformation temperatures and the kinetics of phase transformations depend strongly on the composition of the steels. Sixteen different 9Cr steels have been studied11,12 and the results, which provide the required thermodynamic database are shown in Figure 2, with respect to the dependence of melting point, Ms temperature and the continuous heating transformation diagrams. The constitution and the kinetics of transformations dictate microstructure and the properties. In the early stages, the oxidation resistance and creep strength were of prime importance, since the Cr–Mo steels were developed4 for thermal power stations. In addition to the major constituent phases discussed above, the minor carbides which form at temperatures less than 1100K, dictate the long term industrial performance of the steels. Evaluation of tensile and creep properties of Cr–Mo steels exposed to elevated temperature for prolonged durations have been extensively studied.5,13,14 The following trends were established: The optimized initial alloy compo- sition considered was 9Cr, W–2Mo¼ 3, Si¼ 0.5, with C, B, V, Nb, and Ta in small amounts. Higher chro- mium content has two effects: it increases the hard- enability leading to the formation of martensite and also promotes the formation of d-ferrite thereby reducing the toughness. A reduction in the chromium 750 1820 1815 1810 1805 1800 1795 1790 1785 1 2 3 4 5 6 7 8 9 10 11 Ms/K = 904 - 474 (C + 0.46(N - 0.15Nb ) - 0.046Ta) -{17Cr + 33Mn + 21Mo + 20Ni + 39V + 5W) -45Mn2- 25Ni2- 100V2+ 10Co } - 44.5Ta 9Cr–ferritic martensitic steels 725 700 675 650 625 600 600 (a) (c) (b) M s, e xp er im en ta l ( K ) 625 650 1000 950 900 850 Time (s) 9Cr–ferritic steel Continuous heating transformation (CHT) diagram Te m p er at ur e (� C ) Te m p er at ur e (K ) 800 100 101 102 1098 1148 1198 1248 103 50% transformed Ferrite + carbide Austenite 99 60 40 Ferrite+ austenite+ carbide Ac3 Ac1 20 15 10 5 1 Ms, empirical estimate (K) Steel designation M el tin g p oi nt (K ) 675 700 725 750 Experimental Estimated P la in 9 C r 1M o 0. 24 S i a d d ed 9 C r 1M o 0. 42 S i a d d ed 9 C r 1M o 0. 6 S i a d d ed 9 C r 1M o M od . 9 C r 1M o ( M n+ N i)/ 2. 32 M od . 9 C r 1M o (M n+ N i)/ 1. 85 M od . 9 C r 1M o (M n+ N i)/ 1. 7 M od . 9 C r 1M o (M n+ N i)/ 13 5 M od . 9 C r 1M o 1W -0 .2 3V -0 .0 5T a 9C r 1M o M od . 9 C r 1M o: b as e m od el Figure 2 Influence of chemistry on transformation temperatures (Ms and melting point) and kinetics of transformation of g ! a þ carbide, in various ferritic steels. 100 Ferritic Steels and Advanced Ferritic–Martensitic Steels content lowers the oxidation resistance. If W þ Mo concentration is kept Ferritic Steels and Advanced Ferritic–Martensitic Steels 101 This would ensure a satisfactory solidification process with a fully austenitic structure. Additionally, this enables easier hot workability during primary proces- sing and tubemaking, without losing high tempera- ture creep resistance. The formation of d-ferrite reduces toughness due to the notch sensitivity, pro- motes solidification cracking and embrittlement due to sigma-phase precipitation and reduces the creep ductility at elevated temperatures of operation. Other problems relate to solidification cracking, hydrogen cracking, and reheat cracking, which have been exten- sively studied.18 The Type IV cracking in ferritic steel weldments and the brittle layer formation in the dis- similar welds are discussed in detail later. 4.03.3 Radiation Damage of Core Components in Fast Reactors The core components in fast reactors include the following: clad (cylindrical tubes which house the fuel pellets) for the fuel and wrapper (a container which houses fuel elements, in between which the coolant flows) for fuel subassemblies. Figure 3 shows a schematic of clad and wrapper in a typical fuel subassembly. The necessity to develop robust tech- nology for core component materials arises from the fact that the ‘burn-up’ (energy production from unit Head Shield pins, pellet stack Section−MM Section−E E Section−HH Section−BB Fuel 217 fuel pins Adaptor B&C shielding Steel shielding Top plannet Fuel Bottom plannet Coolant entry tube Discriminator Section–XX Pin Figure 3 Schematic of a typical fuel subassembly. quantity of the fuel) of the fuel depends on the performance of the clad materials. The higher burn- up of the fuel increases the ‘residence time’ of the subassembly in the core, eventually lowering the cost. The core component materials in fast breeder reactors are exposed to severe environmental service conditions. The differences in the exposure condi- tions of the clad and wrapper in a fast reactor core are listed in Table 1. Under such exposure condi- tions, materials in the fast reactor fuel assemblies exhibit many phenomena (Figure 4), specific to fast reactor core: Void swelling, irradiation growth, irra- diation hardening, irradiation creep, irradiation, and helium embrittlement. Another selection criterion, namely the compati- bility of the core component materials with the cool- ant, the liquid sodium, has already been established. Presently, methods are known to avoid interaction of the clad material with the coolant. Detailed books and reviews19,20,21,22,23 are avail- able on all the degradation mechanisms mentioned above, which are related to the production, diffu- sion, and interaction of point defects in the specific lattice of the material. Hence, a brief introduction is presented below (see also Chapter 1.03, Radiation- Induced Effects on Microstructure; Chapter 1.11, PrimaryRadiationDamageFormation; andChapter 1.04, Effect of Radiation on Strength and Ductility of Metals and Alloys). Void swelling in a fast reactor core can change a cube of nickel to increase (20%) its side from 1 cm to 1.06 cm, after an exposure to irradiation of 1022 n cm�2. Void swelling is caused by the conden- sation of ‘excess vacancies’ left behind in the lattice after ‘recombination’ of point defects produced dur- ing irradiation. Void swelling is measured using the change in volume (▽V/V) of bulk components of the reactor or image analysis of voids observed using transmission electron microscope (TEM). The ‘irradiation growth’ (fluence �1020 n cm�2) can increase the length of a cylindrical rod of uranium three times and reduce its diameter by 50%, retaining the same volume. This occurs mainly in anisotropic crystals, introducing severe distortion in core compo- nents. It is caused by the preferential condensation of interstitials as dislocation loops on prism planes of type (110) of hcp structures and vacancies as loops on the basal planes (0001), which is equivalent to transfer of atoms from the basal planes to prism planes, via irradiation-induced point defects. Irradiation hardening refers to the increase in the yield strength of the material with a Table 1 Comparison of exposure conditions of clad and wrapper of fast reactor core Criterion Clad tube Wrapper tube Exposure conditions (only trends; exact values depend on core design) Maximum temperature: 923–973K Lower temperature range than clad: 823K Steeper temperature gradient Lower temperature gradient Higher stresses from fission gas pressure Moderate stresses from coolant pressure Chemical attack from fuel Flowing sodium environment Average neutron energy: 100 keV Neutron environment similar Neutron flux: 4–7� 1011 nm�2 s Neutron fluence: 2–4�1019 nm�2 Major damage mechanisms Void swelling Void swelling Irradiation creep at higher temperatures Irradiation creep Irradiation embrittlement Irradiation embrittlement Interactions with fuel and fission products Interaction with sodium Selection criteria: mechanical properties Tensile strength Tensile strength Tensile ductility Tensile ductility Creep strength Creep ductility Corrosion criteria Compatibility with sodium Compatibility with sodium Compatibility with fuel Compatibility with fission products General common selection criteria Good workability International neutron irradiation experience as driver or experimental fuel subassembly Availability 102 Ferritic Steels and Advanced Ferritic–Martensitic Steels concomitant reduction in ductility, under irradiation at temperatures Linear swelling regime Threshold dose (a) Irradiation dose (dpa) S w el lin g (Ñ V /V ) Transient swelling regime c axis (b) a axis Neutron fluence (E>1 MeV) 1024 n m–2 3 2 1 0 G ro w th s tr ai n (1 0- 4 ) -1 -2 -3 Irradiated Unirradiated Strain(c) S tr es s Irradiated Unirradiated Time (h)(d) S tr ai n (% ) Irradiated Unirradiated Test temperature(e) A b so rb ed e ne rg y Figure 4 Schematic representation of major damage mechanisms in the core component materials of fast reactors: (a) The different stages of void swelling, (b) irradiation growth, (c) increase in strength with a concomitant reduction in ductility during irradiation hardening, (d) increase in creep strain and reduction in creep life after irradiation caused by irradiation creep, and (e) increase in ductile to brittle transition temperature and reduction in upper shelf energy after irradiation caused by irradiation embrittlement. Ferritic Steels and Advanced Ferritic–Martensitic Steels 103 creep-resistant ferritic steels. The microstructural instability during service exposure is briefly pre- sented. The superior swelling performance of ferritic steels is understood based on mechanisms of void swelling suppression. Following this, the irradiation- induced/-enhanced segregation/precipitation causing irradiation hardening is discussed. The irradiation creep and embrittlement, their mechanisms and meth- ods to combat the problems are highlighted. The R&D efforts of today to reduce the severity of embrittle- ment in ferritic steels, using modeling methods, are outlined. Finally, typical problems in the weldments of ferritic steels, when used for out of core applica- tions, are presented, emphasizing the advantage of modeling in predicting the materials’ behavior. 4.03.4.1 Influence of Composition and Microstructure on Properties of Ferritic Steels Rapid strides have been made the world over, in the design and development of advanced creep-resistant ferritic or ferritic–martensitic steels. The low alloy steels can be used as either 100% ferrite–martensite 104 Ferritic Steels and Advanced Ferritic–Martensitic Steels or a mixture of both. It is possible to choose the required structure by the appropriate choice of either the chemistry or the heat treatment. For example, a completely ferrite matrix, yielding high toughness, can be obtained in steels with chromium content higher than 12%, with carbon reduced to less than 0.03%. The same steel can be used to provide higher strength by choosing the 100% martensite structure, if carbon content is increased to about �0.1%. The 9Cr steels have always been used in the 100% mar- tensite state. Extensive studies have been carried out on phase stabilities of these steels, with changes in chemistry and heat treatment. The creep resistance of the plain Cr–Mo steels has, further, been increased by the addition of carbide stabilizers like Ti or V or Nb, leading to the modified variety of 9–12Cr–Mo steels. These Table 2 Optimizing the constitution in the development of ferritic steels Element Function Cr Basic alloying element, corrosion resistance, hardenability Mo, W, Re, Co Solid solution strengthening V, Nb, Ti, Ta Strengthening by formation of MX-carbonitride C, N Austenite stabilizer, solid solution strengthening, carbonitride formers B Grain boundary strengthening, stabilization of carbide Ni, Cu, Co Austenite former, inhibits d-ferrite formation Table 3 Beneficial and harmful effects of different elements Element Beneficial Carbon Strength Mo Creep strength Ni Mn S scavenger Si Void swelling, spheroidization Ti, V, Nb, Ta, and W Precipitation strengthening S, P, As, Sb, Sn, and Bi Cu Co N, O B Delayed coarsening elements led24 to copious, uniform precipitation of Monte Carlo (MC) type of monocarbides, which are very fine and semicoherent. Such precipitates are very efficient in pinning the mobile dislocations, lead- ing to improved creep behavior at higher tem- peratures. These carbides are stable at temperatures higher than even 1273K and hence, do not cause deterioration of long-term mechanical properties during service exposure. The development of high creep-rupture strength 9–12% steels with various combinations of N, Mo, W, V, Nb, Co, Cu, and Ta is based on optimizing the constitution (Table 2.) and d-ferrite content, increasing the stability of the martensite, dislocation structure and maximizing the solid solution and pre- cipitation hardening. The concentration of each ele- ment in ferritic steels has been optimized based on an in-depth understanding of the influence of the specific element on the behavior of the steel. The extensive studies related to optimization of chemistry are summarized in Table 3. Based on the strong scientific insights, large number of commercial steels have been developed (Table 4) in the later half of the last century. Most of this family of ferritic–martensitic steels is used in the normalized and tempered condition or fully annealed condition to achieve the desirable phase. The type of structure that is deliberately favored in a given steel depends on the end application. The microstructure of the steels in normalized and tempered conditions consists24 (Figure 5) of (a) martensite laths containing dislocations with a Burgers vector 1/2a0 with a density of approximately 1� 1014m�2 (b) coarse M23C6 particles located at during design of creep-resistant ferritic steels Detrimental Optimum (wt%) Weldability 0.1 Hardenability d-ferrite formation 1 Laves phase Intermetallic Ni3P 0.1 Induced radioactivity 0.5 d-ferrite formation 0.2–0.4 Silicide formation Undissolved carbides, low hardness of martensite Table 4 List25 of commercial ferritic steels, their chemistry, and properties Commercial name Chemistry 105 h creep strength at 873KMPa�1 T22 2.25Cr1Mo 35 Stab. T22 2.25Cr1MoV 60–80 HCM2S 2.25Cr1MoWNb 100 T9 9Cr1Mo 35 EM12 9Cr2MoVNb 60–80 F9 9Cr1MoVNb 60–80 T91 9Cr1MoVNb (optimized) 100 T92 9Cr(MoW)VNb 120 Eurofer 9CrWTiV �120 HT91 12Cr1MoV 60–80 HT9 12Cr1MoWV 60–80 HCM12A 12CrMoWVNbCu 120 SAVE12 12CrWVNbCo 180 Ferritic Steels and Advanced Ferritic–Martensitic Steels 105 prior austenite and ferrite grain boundaries with finer precipitateswithin the laths and at martensite lath and subgrain boundaries. M2X precipitates rich in Cr are isomorphous with (CrMoWV)2CN. The initial microstructure of the normalized and tempered steels described above does not remain stable during service in a nuclear reactor. Pro- longed exposure at high temperature causes changes in the initial microstructure, which has been studied extensively. The M2X precipitates in the normal- ized and tempered stabilized 9Cr–1Mo steels are gradually replaced (Figure 6) by MX, intermetallic, and Laves phases during prolonged aging at high temperature. The high temperature and the irradiation over prolonged time of exposure introduce microstructural instabilities. These instabilities are caused mainly by the point defects caused by irradiation and complex coupling of these defects with atoms in the host lattice, their diffusion or segregation and finally the precipitation. There is a recovery of the defect structure since the irradiation-induced vacancies alter the dislocation dynamics. There are three types of processes with respect to evolution of secondary phases: irradiation-induced precipitation, irradiation- enhanced transformations, and the irradiation modi- fied phases. It is seen that the evolution of these phases depends on the composition and structure of the steel and the irradiation parameters like the temperature, dose rate, and the dose. Evolution of irradiation- induced phases and their influence on hardening and embrittlement is discussed later. 4.03.4.2 Void Swelling Resistance Extensive experimental investigations found3 that the ferritic steels, whose high temperature mechanical properties are far inferior to austenitic stainless steels, displayed excellent radiation resistance. The ferritic– martensitic steels (9–12% Cr) have, therefore, been chosen for clad and wrapper applications, in order to achieve the high burn-up of the fuel. This is based26–29 (Table 5) on their inherent low swelling behavior. The 9Cr–1Mo steel, modified 9Cr–1Mo (Grade 91), 9Cr–2Mo, and 12Cr–1MoVW (HT9) have low swelling rates at doses as high as 200 dpa. For example, HT9 shows 1% swelling at 693 K for 200 dpa. The threshold dose for swelling in ferritic steels is as high as nearly 200 dpa in contrast to 80 dpa for the present generation D9 austenitic stainless steel. It is established that the void swelling depends crucially on the structure of the matrix lattice, in which irradiation produces the excess defects. Extensive basic studies have identified19,30–33 the following reasons as the origin of superior swelling resistance in ferritic steels: 1. The relaxation volume for interstitials, that is, the volumeof thematrix inwhichdistortionis introduced bycreatingan interstitial, inbcc ferrite is larger19 than fcc austenite. For every interstitial introduced, the lattice distortion is high and hence the strain energy of the lattice. Hence, the bias toward attracting or accommodating interstitials in the bcc lattice is less. This leaves higher density of ‘free’ interstitials in the bcc lattice than fcc lattice. As a result, recombina- tionprobabilitywithvacancies increases significantly and supersaturation of vacancy reduces. Conse- quently, the void nucleation and swelling is less. 2. The migration energy of vacancies in bcc iron is only 0.55 eV, against a high value in fcc austenite, 1.4 eV. Vacancies are more mobile in bcc than fcc, increasing the recombination probabilities in bcc ferrite. Another factor is the high binding energy between carbon and vacancy in bcc iron (0.85 eV), while it is only 0.36–0.41 eV in austenite. This leads19 to enhanced point defect recombina- tion in bcc than fcc, due to more trapping of vacancies by carbon or nitrogen. 3. In bcc iron, it is known30 that there is a strong inter- action between dislocations and interstitials solutes, forming atmospheres of solutes around dislocations. The formation of ‘atmospheres’ around dislocations makes them more effective sinks for vacancies than interstitials, resulting in suppression of void growth, A B (a) 0.5 mm (001) (b) 2.00 (c) 4.00 6.00 8.00 10.00 12.00 14.00 16.00 1 NbKaFeKa V Kb V Ka NbLa 2.00 (d) V Ka NbLa NbKa 4.00 6.00 8.00 10.00 12.00 14.00 16.00 Figure 5 Initial structure24 of normalized and tempered modified 9Cr–1Mo steel: (a) Monocarbides (MC) and M23C6 along lath boundaries in a carbon extraction replica of the sample and (b) Microdiffraction of fine particle marked B, confirming the crystal structure of MC. Energy dispersive analysis of X-rays (EDAX) identifying theMCparticles (B) to be rich in (c) V and (d) Nb. 106 Ferritic Steels and Advanced Ferritic–Martensitic Steels provided the following conditions are satisfied: ‘atmo- spheres’ comprise of either oversized substitutional atoms or interstitials, dislocations have high binding energy with solutes, and concentration of solute atoms at the core of the dislocation exceeds a critical value. On the other hand, if ‘atmosphere’ is made up of undersized atoms like Si or P, the voids can grow. The ‘atmosphere’ of interstitials reduces the dislocation bias for additional capture and inhibits dislocation climb, thus converting them to saturable sinks. Such a scenario would increase the recombination prob- abilities, suppressing the void growth. These fundamental differences in the behavior of solutes and point defects in bcc lattice make ferritic steel far superior to austenitic steels, with respect to radiation damage. The challenging task for materials scientists to use ferritic steels directly in fast reactor fuel assembly was with respect to enhancing the high temperature mechanical properties of the ferritic steels, especially high temperature creep life and irradiation creep resistance. 4.03.4.3 Irradiation Hardening in Ferritic Steels The initial microstructure of the steels evolves dur- ing service, due to high temperature and irradiation for prolonged times, leading to modification of defect structure and secondary phases. These changes harden the steel, leading to concomitant embrittle- ment, which is discussed below. Ferritic Steels and Advanced Ferritic–Martensitic Steels 107 It is reported that carbon content in 12% chro- mium steel is maintained high in order to use the steel as martensitic steels. The high amount of car- bon in 12% chromium steel leads to copious precip- itation of carbides, that is, twice as much in 9Cr steels. Both the steels have predominantly M23C6 carbides with a small fraction of monocarbides, Table 5 Void swelling resistance26–29 of some commercial f Commercial name Chemistry and country of origin R w FV448 12Cr–MoVNb, UK P EM10 9Cr–1Mo, France P 1.4914 12CrMoVNb, Germany P EP450 12Cr–MoVNb, Russia B EP450 – B (a) 1 mm Laves phase (1 21 3) 2.00 4.00 6.00 8.00 10.00 12.00 14.00 16.00 18.00 MoKa CrKa MoLa FeKa (b) SiKa Figure 6 Effect24 of prolonged exposure (823K per 10 000h) of modified 9Cr–1Mo steel. Transmission electron micrograph showing (a) formation of detrimental Fe2Mo Laves intermetallic phase around the M23C6. The insets show the microdiffraction pattern and magnified view of the nucleation of Laves phase (b) EDAX spectrum confirming the enrichment of iron and molybdenum. eventually leading34 to deterioration of their resis- tance to brittle failure. The critical stress to propa- gate a crack is inversely proportional to the crack length. If it is assumed that fracture initiates at an M23C6 precipitate and the crack length at initiation equals the diameter of a carbide particle then the fracture stress will decrease with increasing precipi- tate size. The precipitates coarsen during irradiation in the range of 673–773 K, thus causing a decrease in fracture stress and an increase in DBTT even in the absence of further hardening. Additionally, Cr rich, bcc a0 precipitates formed35 in the higher chromium steel during thermal expo- sure and irradiation lead to hardening and embrittle- ment of the steel. The d-ferrite, into which there is a repartitioning of chromium, is harmful, since it pro- motes formation of a0 . The presence of very fine coherent particles of the w (Fe2Mo) phase has also been reported in theT91 andHT9 steels. The w phase was observed to form more rapidly in the 9Cr–2Mo type of steels, both in the d-ferrite and martensite phases. This is possibly due to the higher amount of Mo in the EM12 type of steels. The w phase is enriched in Fe, Si, and Ni and contains significant amount of Mo and P. The G phase (Mn7Ni16Si17) has been found to form very occasionally in the mod- ified 9Cr–1Mo and HT9 (12% Cr) variety of steels. The s phase (Fe–Cr phase, enriched in Si, Ni, and P) has been observed to form around the M23C6 particles in 9–13% Cr martensitic steels after irradia- tion at 420–460 �C in Dounray Fast Reactor. In addition Cr3P needles and MP (M¼Fe, Cr, and Mo) particles have also been detected in the 12 and 13Cr steels in the range of 420–615 �C. The formation of these phases during irradiation may be understood in terms of the strong radiation-induced segregation (RIS) of alloying and impurity elements to point defect sinks in the steels (see Chapter 1.18, Radiation- Induced Segregation). The RIS of alloying/impurity elements could lead36 to either enrichment or deple- tion near the sinks, depending on the size of the atom and its binding energy with iron self-interstitials. erritic steels eactor in which irradiation as carried out Burn-up achieved (dpa) FR 132 henix 142 henix 115 N-350 45 N-600 144 Lath boundary Lath boundary 108 Ferritic Steels and Advanced Ferritic–Martensitic Steels Generally, a large number of alloying elements, W, Nb, Mo, Ta, V, or Ti are dissolved into thematrix of ferritic steels, some of them being larger than the iron atom. This could lead to the expansion of the unit cell of ferrite, making an element say, chromium undersized, with a positive binding energy with iron self-intersti- tial. Such a situation would lead to enrichment of chromium near the sink-like grain boundary. The reverse could happen if the size of the alloying ele- ments happen to be smaller than iron. The w, G and s phases are all enriched in Si and Ni – elements which are known to segregate to in- terfaces during irradiation. With the exception of G phase, all the other phases and the a0 phase are rich in Cr. In those ferritic steels, where Cr is depleted near voids and at other interfaces which act as point defect sinks, it follows that in steels containing higher than 11 or 12% Cr, the chromium enrichment within the matrix may lead to local con- centrations exceeding those (�14%) at which a0 forms thermally. Further, enrichment of Cr may also result from the partial dissolution of chromium rich precipitates such as M23C6 during irradiation. In addition, RIS of phosphorus can also lead to the formation of phosphides in some of the steels. The irradiation-induced point defect clusters and loops may also facilitate and enhance nucleation of these phases. Although the relatively soft d-ferrite im- proves the ductility and toughness of the 12Cr steel, the fracture could be initiated at the M23C6 pre- cipitates on the d-ferrite–martensite interface. The presence of d-ferrite, extensive precipitation and radiation-induced growth of M23C6 precipitates and formation of the embrittling intermetallic phases in the 12Cr–1MoVW steel in the temperature range 573–773 K are together responsible37 for the relative change in impact behavior of 9Cr–1MoVNb and 12Cr–1MoVW between 323 and 673K. Irradiation-induced microstructural changes are the factors that govern the creep and embrittlement behavior, which therefore, has to be minimized using appropriate chemistry and structure. (a) (b) MX M23C6 Figure 7 The schematic39 of most undesirable (a) and desirable (b) microstructures for design of creep-resistant steels. 4.03.4.4 Irradiation Creep Resistance of Ferritic Steels An essential prerequisite for maximizing the ‘irra- diation creep resistance’ is to ensure38 the best combination of thermal creep behavior and long- term microstructural stability at high temperature. Hence, the present section would discuss irradiation creep in the same sequence as mentioned above. The design principles of development of creep- resistant steels are as follows: � Introduce high dislocation density by either trans- formations or cold work to increase the strength of the basic lattice; � Strengthen the host lattice by either solid solution strengtheners or defects; � Stabilize the boundaries created by phase transfor- mations by precipitating carbides along the boundaries; � Arrest dislocation glide and climb by appropriate selection of crystal structure, solid solution, inter- faces, dislocation interactions, and crystal with low diffusivity; � Resist sliding of grain boundaries by introducing special type of boundaries and anchoring the boundaries with precipitates; � Ensure long-term stability of the microstructure, especially against recovery and coarsening of the fine second phase particles; In the case of 9–12 Cr steels, the martensitic lath structure (Figure 7) decorated with only MX which should39 be stable over long-term service life is the most desired structure. Thermo-Calc evalua- tions show39 that MX can be stabilized at the expense of M23C6 only by reducing carbon to as low a value as 0.02% in 9 Cr–1Mo steel. This value is too low to ensure acceptable high temperature mechani- cal behavior of the steels. In the context of fast reactor core components, the high chromium 9–12% ferritic–martensitic steels assume relevance. Hence, an extensive database25 for a large number of commercial ferritic steels has been generated and P91 P92E9119Cr1Mo(a) 0 50 10 5 h ru p tu re s tr es s (M P a) 100 + 0.04N + 0.2V + 0.08Nb +1 W -0.5Mo + 0.8 W 150 60 0 �C 65 0 �C 5 4 3 2 1 0 0 20 (b) 40 60 80 Hoop stress (MPa) 316SS 540 �C 1 � 1023 n cm–2 DD /D (% ) D9 HT-9 D21 D68 100 120 140 Figure 8 (a) Thermal creep40 of 9Cr1Mo ferritic steel. (b) irradiation creep41 of ferritics in comparison to austenitics. Reprinted, with permission, from J. ASTM Int., copyright ASTM International, 100 Barr Harbor Drive, West Conshohocken, PA 19428 (a) (b) stensile stensile s T 0 0 0 0 0 0 0 0 0 0 0 0 T T T T Figure 9 The mechanisms of stress-induced preferential absorption (a) and stress-induced preferential nucleation (b) during irradiation creep. Ferritic Steels and Advanced Ferritic–Martensitic Steels 109 Figure 8(a) shows40 the continuous improvement achieved by careful modification of alloying elements, in the thermal creep behavior of successive grades of different commercial ferritic steels. While understanding thermal creep is essential to narrow down the choice of ferritic steels for use in a fast reactor, ‘benchmarking’ the steels developed under irradiation is an essential stage before actually using the radiation-resistant steels in the reactor. The irradiation creep depends on the stress level, the temperature, and the dose. Figure 8(b) shows41 the comparison of irradiation creep of ferritic steels with competing materials like the austenitics and nickel-based alloys. It is clear that the point defects generated during irradiation act against the design principles of devel- oping creep-resistant materials, listed earlier. The point defects accelerate the kinetics of dislocation climb, coarsen the precipitates, and generally enhance the diffusivity. In addition, the excess point defects precipitate into either interstitial or vacancy loops, but not randomly. The interaction between point defects and stress leads to the precipitation of intersti- tial loops parallel to the applied stress, while vacancy loops form in planes perpendicular to the stress. This process (Figure 9(a)) called the stress-induced pref- erential nucleation (SIPN) results in additional creep strain solely due to irradiation. The excess point defects under temperature migrate randomly. But in the presence of an additional factor, that is, stress, the vacancies migrate preferentially to grain boundaries perpendicular to the applied stress, while the intersti- tials toward boundaries parallel to the stress. This is equivalent to removing material from planes parallel to the stress to those which are perpendicular to the applied stress, introducing additional creep strain. This process is called the stress-induced preferential absorption (SIPA) (Figure 9(b)). The radiation-induced defects also evolve from isolated point defect to loops and voids, which have different types of influence on irradiation creep. Most often, irradiation creep occurs19,42 simultaneouslywith swelling and sometimes, swelling influences irradia- tion creep. At very small dose levels, swelling enhances creep rates. Beyond a certain dose levels, the creep component reduces and at high dose levels, creep dis- appears, while swelling continues.Figure 10 shows the variation in creep coefficient at various dose levels, and the regimes where swelling has an influence. The dynamics of point defects during irradiation continu- ously evolve with change in structure of dislocation network and loops. At small dose levels, there is a uniform distribution of very fine voids, which act as effective pinning centers for mobile dislocations. Thus the creep rate increases. With increase in dose levels, voids growandmultiply. The chance of interstitials and Onset of disappearance of creep Swelling enhanced creep Swelling without creep Dose (dpa) In st an ta ne ou s cr ee p c oe ffi ci en t Figure 10 Schematic of variation of instantaneous creep coefficient with dose, showing the interplay between irradiation creep and void swelling. 110 Ferritic Steels and Advanced Ferritic–Martensitic Steels vacancies impinging on the void surface becomes more than their reaching dislocations. The number of inter- stitials reaching a dislocation reduces. Additionally, the defect clusters, that is, the dislocation loops also undergo ‘faulting’ contributing to the density of dis- locations in the matrix. Hence, creep rate reduces, due to two factors: increased dislocation density of the matrix due to unfaulting of dislocation loops and reduced availability of interstitials to dislocations. The above process continues until complete cessation of creep, with swelling continue to take place. At very high temperatures, the point defect migra- tion along the grain boundaries in preferential routes causes the grain boundary aided creep. This high temperature limit of ferritic–martensitic steels restricts the application of these steels to at best, wrappers of present generation fast reactors based on oxide fuel. It is necessary to develop materi- als with better high temperature irradiation creep properties and void swelling for clad applications. The future scenario, which envisages the development of metallic fuel to ensure sustainability by breeding, could make use of ferritic steels for both clad and wrapper. This advantage arises due to the lower value of the anticipated clad temperatureswith metallic fuels (see Chapter 3.14, Uranium Intermetallic Fuels (U-Al, U-Si, U-Mo)), whose choice is mainly to ensure sustainability using high breeding ratio. 4.03.4.5 Irradiation Embrittlement in Ferritic Steels The stabilized ferritic steels in the normalized and tempered condition have a tempered martensitic structure with a preponderance of monocarbides that impart the necessary creep strength, while the prior austenite grain and lath boundaries are deco- rated with Cr rich M23C6 precipitates which increase the thermal stability of the steel. It is reported that thermal aging at temperatures above 773K causes gradual but continuous degradation in upper shelf properties in addition to increase in the DBTT. The nature of embrittlement varies for different compo- nents of the reactor. For removable components such as clad, which are subjected to high temperature and pressure, with a residence time of a few years, creep embrittlement is the issue which decides their design and performance, while for permanent sup- port structures increase in hardening and loss in fracture toughness on irradiation are major issues. The origin of embrittlement is two-fold: segrega- tion of tramp elements to prior austenite grain boundaries which make the grain boundaries deco- hesive and evolution of carbides and intermetallic phases. The latter causes progressive changes in the tempered martensitic microstructure, which deterio- rate the fracture properties of the steel, by introdu- cing irradiation hardening effects. The increase in the ductile to brittle transition temperature, DDBTT, is known to be related to irradiation hardening, which is generally observed to saturate with fluence. Evidence for a possible maxi- mum in DBTTwas observed for the 12Cr steel irra- diated in the range of 35–100 dpa in fast flux test facility (FFTF). Based on observed data in a number of cases it appears that a high fluence and/or high tempera- ture are required before a maximum is observed. This implies that the strength and impact properties are a balance between the point defect production and irradiation-induced precipitation. The precipitation during irradiation hardens the steel and irradiation accelerated recovery and aging soften the steel. The latter process ismore important at high fluences and/or higher irradiation temperatures. Hence, hardening in most of these Cr–Mo steels is more than compensated for by the recovery and aging processes, leading to saturation in irradiation hardening above 723K. For body centered cubic materials such as ferritic martensitic steels, radiation hardening at low tempera- tures ( Ferritic Steels and Advanced Ferritic–Martensitic Steels 111 with fuel and coolant. Void swelling is not expected to be significant in F/M steels up to damage levels of about 200 dpa. Extensive evaluation14,15,43–58 of the embrittle- ment behavior of the ferritic steels for different chemistry is shown in Figure 11. The merit in focus- ing on chemistry around 9% chromium is very clear based on the observation of minimum shift in DBTT around this composition, under irradiation. However, higher chromium improves corrosion resistance and ease of reprocessing. Hence, chromium content has to be selected balancing these requirements. It is known44–48 that addition of phosphorous, copper, vanadium, aluminum, and silicon would increase the DBTT while sulfur reduces the upper shelf energy (USE). The 12Cr steels, HT9, show a larger shift (125K) in DBTT as compared to modified 9Cr–1Mo steel (�54K). Hence, the balance is always between nearly nil swelling resistant 12Cr steels and 9Cr steel which is less prone to embrittlement than 12Cr steels. Microstructural parameters, like the prior aus- tenite grain size, lath/packet size, carbides, and their distribution influence49,50 the embrittlement behav- ior. Studies on the effects of heat treatment and microstructure on the irradiation embrittlement in 9Cr–1MoVNb and HT9 steels are summarized below: � Prior austenite grain size (PAGS) influences51 the DBTT for the 9Cr–1MoVNb steel, but not in 250 200 150 2.25Cr–2W 10 dpa: 365 �C 10 dpa: 365 �C 36 36 dpa: 410 �C 7 dpa: 365 �C 7 dpa: 365 �C 2.25Cr–1WV R. Klueh R. Klueh et al. 2Cr–1.5V JLF-4 7.5 2.25CrV 2.25Cr–2W 5Cr–2WV 100 50 0 0 2 4 6 Chromium D B TT s hi ft (C ) Figure 11 Variation43 of shift in ductile to brittle transition tem to different dose levels at around 673K. The ferritic steel with 9C 12Cr–MoVW steel. This is attributed to the precipitates in the microstructure controlling the fracture behavior rather than the PAGS, in the 12Cr steel. � The size of martensitic lath and packet, which is sensitive52 to austenitization temperature, can also affect51 the fracture behavior. Examination of the fracture surface revealed cleavage and regions of ductile tearing along prior austenite grain and lath packet boundaries. Subsurface microcracks and sec- ondary surface cracks were found associated with large boundary carbides. It was suggested that cleav- age fracture initiated in HT9 by propagation of a microcrack from a coarse carbide into the matrix. Propagation was inhibited by the intercepted boundaries, lath or grain and ductile tearing was required53 to continue propagation. The amount of tearing increased with increasing austenitization temperature. � Tempering for the two normalization tempera- tures had very small effect on the DBTT, for the two steels. � Irradiation of the two steels at 638 and 693K resulted37 in an increase in DBTT and a decrease in USE for all conditions with the shift in DBTT for the 12 Cr steel being almost twice that for 9Cr steel. � Although the 12Cr steel with the smallest grain size had55 the lowest DBTT after 20 dpa, the effect of tempering was different. In the case dpa: 410 �C et al. A. Kohyama et al. JLF-3 JLF-6 12Cr–6Mn–1W 12Cr–6Mn–1V JLF-1 Cr–1V 9Cr–1W Cr–2W 9Cr–2WVTa F82H 9Cr–2WV 12Cr–2WV content (wt%) 8 10 12 perature (DBTT) for various Cr–Mo steels with irradiation r–1Mo has the least variation in DBTT. 112 Ferritic Steels and Advanced Ferritic–Martensitic Steels of 12Cr steel, the higher tempering temperature causes coarsening of precipitates thus accelerating fracture. � The saturation of shift in DBTT with fluence is independent54 of tempering conditions for the 9Cr steel, while for the 12Cr steel, a maximum is observed, probably due to faster growth of preci- pitates during irradiation. The generation of helium through (n,a) reaction in elements of structural materials is known to cause severe damage to the embrittlement behavior of core component materials. Table 6 lists the shift in DBTT, for 9 and 12CrMo steels, under reactor irra- diation, with and without helium, which demon- strates56 the harmful effect of helium. These results become more pertinent in the case of fusion reactors, where the operating conditions include the genera- tion of helium up to about �100 appmyear�1. The increase in the DBTT due to irradiation is a cause of serious concern for use of ferritic steels, since it makes the postirradiation operations very difficult. Several methods have been attempted57,58 to address this problem, which includes modification of the steel through alloying additions, control of tramp elements by using pure raw materials and improved melting practices, and grain boundary engineering (GBE). However, the propensity of the problem is less if the clad thickness is low, which normally is the case to ensure best heat transfer properties. For low thickness components, the triaxial stress necessary for the embrittlement does not develop, which reduces the intensity of this otherwise serious problem of embrittlement in ferritic steels. An approach to reduce shift in DBTT is an immediate concern in ferritic steels for core com- ponent applications and efforts to overcome this problem by selection of high purity metals, adoption of double or triple vacuum melting for steel making, strict control of tramp and volatile elements, and development of special processing methods, which would improve the nature of grain boundaries (GBE) are in progress. Table 6 Comparison56 of embrittlement behavior of 9 and 1 Irradiation conditions Reactor Temperature (K) Dose (dpa) EBR II 663 13 EBR II 663 26 HFIR 673 40 4.03.4.5.1 GBE to reduce embrittlement in ferritic steels GBE is an emerging field, which promises methods to improve the performance of materials, whose deg- radation in service is caused by the presence of high angle boundaries. The concept, first proposed59 by Prof. T. Watanabe in the early 1980s, envisages improvement of properties of materials by controlling the grain boundary character distribution (GBCD). Many processes like diffusion, precipitation, segrega- tion, sliding, cavitation, and corrosion are kinetically faster along high angle grain boundaries. Hence, it is possible to decelerate these detrimental pro- cesses by replacing the random boundaries with low energy ones, coincident site lattice (CSL) boundaries (denoted by the ‘sigma number,’ S, which is defined as the reciprocal of the fraction of lattice points in the boundaries that coincide between the two adjoining grains on the basis of CSL model). Another prerequi- site for GBE is to completely destroy the interconnec- tivity of random grain boundary network. The insight in the field of GBE was achieved with the advent of computer assisted EBSD (electron back scatter diffraction) technique developed during the 1980s. The embrittlement in ferritic steels is known to be caused by segregation phenomena. The kinetics of segregation can be controlled by suitable selection of the nature of grain boundaries. GBE has been applied60–63 to combat embrittlement problems in ferritic steels. The task of carrying out GBE using experimental methods is time consuming. Hence, it is prudent to resort to computational methods, which need to be validated using selected experiments. A 3D Poisson–Voronoi grain structure, simulated using MC technique was employed to study60 (Figure 12(a)) intergranular crack percolation using percolation theory. The percolation threshold was estimated to be 80%. To apply this model to specific alloys like ferritic steel, system specific characteris- tics need to be incorporated61 in the model. One such attempt is to define the propensity of the grain boundaries for propagation of cracks based on rela- tive values of the grain boundary energy and the 2Cr steels, with and without helium Shift in DBTT (K) 9Cr1Mo(VNb) 12Cr1Mo(VW) 50 125 50 �150 200 (30 appm He) 250 (110appm He) (a) 0.9 0.8 0.7 0.6 0.5 0.4 0.3 0.2 0.1 0 0 100 (b) 200 300 400 500 Critical crack length (mm) Fa ilu re p ro b ab ili ty 600 700 800 900 Fine grains (12 mm) 30 50 60 70 80 % crack- resistant boundaries Coarse grains (25 mm) 1000 1 300 250 200 150 100 50 Two-step normalization and tempering treatment Conventional N&T treatment 67.8 J lower bound criteria 0 -80 -70 -60 -50 -40 -30 -20 -30 �C-45 �C -10 0 Temperature (�C) C ha rp y ab so rb ed e ne rg y (J ) 10 20 30 (c) 0.185 0.180 0.175 0.170 0.165 0.160 0.155 0.150 0.145 0.140 0.135 0.130 0.125 0.185 0.180 0.175 0.170 0.165 0.160 0.155 0.150 0.145 0.140 0.135 0.130 0.125 0 5 10 S p ec im en B S p ec im en A 15 20 25 Charpy impact energy (J) 30 35 40 45 0 Fr ac ta l d im en si on in cr em en t 5 (d) 10 15 20 25 30 35 40 45 Average Crack initial stage Figure 12 Modeling and electron back scattered diffraction studies in grain boundary engineering of ferritic steels: (a) Percolation of a crack in the 3D P–V model of grain structure generated60 using Monte Carlo methods. (b) Percolation probability for two different grain size (c) experimental confirmation63 of reduction in ductile to brittle transition temperature (DBTT) with grain size and (d) fractal analysis63 of the fracture surface revealing the tortuous path being responsible for the improvement of DBTT in fine grain size. N and T refers to normalized and tempered. Ferritic Steels and Advanced Ferritic–Martensitic Steels 113 10 nm Figure 13 Z-contrast in the high angle annular dark field (HAADF) micrograph of dispersoids in oxide dispersion strengthened (ODS)-9Cr–1M0 ferritic steels, which are responsible for the superior high temperature creep behavior. 114 Ferritic Steels and Advanced Ferritic–Martensitic Steels energy required for propagation of cracks. These calculations were carried out (Figure 12(b)) for two different grain sizes. The prediction of finer grain size being favorable to reduce embrittlement was con- firmed (Figure 12(c)) experimentally. The GBCD, that is, the distribution of various grain boundary types has been evaluated62 in modified 9Cr–1Mo fer- ritic steel using EBSD technique. The experimental observations confirmed the reduction of DBTT by 20K with reduction in grain size. The fractal analysis of the fracture surface demonstrated (Figure 12(d)) that the tortuous path which cracks need to follow in fine grain sample is responsible63 for the observed reduction in the propensity for embrittlement. It is shown clearly that the low energy boundaries can be introduced in engineering materials in three different methods: preferential nucleation of low angle boundaries around twins or controlled recov- ery or orientation relations during phase transforma- tion, if some of the variants happen to result in CSL boundaries. Significant improvements in properties using GBE have been achieved64 in many austenitic stainless steels, in contrast to ferritic steels. The major challenges in the application of GBE to Cr–Mo fer- ritic steels arise from the following factors: lower twinning probability, higher stacking fault energy, and limited variants with CSL boundaries during g! a transformation during cooling. 4.03.5 Development of Advanced ODS Ferritic Steels In recent years, an attempt to increase the high tem- perature creep life of ferritics to 973K and target burn-up of the fuel to 250 dpa, has enabled a ‘revisit’ to the concept of strengthening the steel using 5 nm particles of yttria (see Chapter 4.08, Oxide Disper- sion Strengthened Steels), leading to the ODS fer- ritic steels. ODS ferritic steels are prospective candidate materials for sodium cooled fast reactors with peak burn-up of 250 dpa as well as GenIV and fusion reactors. Earliest developments of ODS steels can be traced to the efforts65 of Belgium in 1960s, followed by Japan66 since 1987, and France67 in the last decade. The ODS steels for fast and fusion reac- tors68,69 are in the R&D stage. The design of ODS steels for fast and fusion reactor applications is based on Fe–Cr–W–Ti– Y2O3, either the martensitic 9 or 12Cr or the ferritic 12Cr steels. The dispersoids which confer the high temperature creep life to the ferrite matrix are70 (Figure 13) in the size range of around 5 nm with a volume fraction around 0.3%. The yttria dissolves in it some amount of titanium, leading to the formation of mixed, complex oxide, namely TiO2�Y2O3. The rationale for the choice of the matrix compo- sition is as follows: Chromium: Choice of 9% Cr and 0.1% C ensures 100%martensite, during normalization of the steel. It is possible to ensure 100% martensite in 12% chro- mium steel by ensuring the carbon content to be above 0.1%. Ferritic ODS steels can be obtained in 12% chromium steels by lowering the carbon content to be less than 0.03%. Higher chromium provides the corrosion and decarburization resistance in sodium at 973K, with acceptable oxidation resistance. Carbon: Addition of 0.1% carbon ensures 100% martensite in 9% Cr steels, thus ensuring absence of anisotropy during g! a transformation. Higher amount of carbon would promote precipitation of M23C6, thus reducing the toughness. On the other hand, M23C6 along the lath boundaries offers the long-term microstructural stability of the lath structure. Nitrogen: The solubility of nitrogen in ferrite is very low. This is useful in non-ODS ferritic steels like T91, due to enhanced creep resistance by forma- tion of V or Nb carbides/carbonitrides. But, in ODS steels, Ti is used for refining yttria. Hence, nitrogen content is restricted to 0.01%, preventing the forma- tion of deleterious TiN compound. Table 7 List of few commercial ODS ferritic steels and their chemistry Commercial name Chemistry MA956 Fe–20Cr–4.5Al–0.5Y2O3 MA957 Fe–14Cr–0.3Mo–Ti–0.27Y2O3 M11 Fe–9Cr–Mo–0.37Y2O3 M92 Fe–9Cr–Mo–0.30Y2O3 PM2000 Fe–20CrAlTi–0.5Y2O3 Ferritic Steels and Advanced Ferritic–Martensitic Steels 115 Tungsten: Tungsten is a more effective solid solution strengthener than Mo, but at the cost of ductility. Tungsten stabilizes d-ferrite and accelerates formation of Laves phase, both of which cause reduc- tion in toughness. Hence, it is optimized to 2.0%. Yttria: The most important constituent of ODS steels is the yttria, which enhances high temperature creep strength by pinning mobile dislocations and delays void swelling by acting as sinks for point defects produced during irradiation. The strength increase is accompanied by a concomitant loss of ductility and saturates around 0.4% yttria. Hence, it is optimized to 0.35%. Titanium: The major role of titanium in ODS steels is to refine the yttria particles (20 nm after mechanical alloying) to ultra-fine (2–3 nm) particles. The complex Y–Ti–O particle imparts the necessary high temperature creep strength. The beneficial effect of titanium saturates around 0.2%. Further increase introduces manufacturing problems of the tubes and hence titanium is chosen as 0.2%. Excess oxygen: Oxygen is present during processing of the ODS steels. The oxygen present in excess of the amount required for formation of required amount of Y–Ti–O complex leads to increase in tensile and creep strength. The Y–Ti–O complex oxides requires about 0.07þ 0.01% excess oxygen. Argon : A strict control of argon ( 116 Ferritic Steels and Advanced Ferritic–Martensitic Steels beyond 250 dpa, at temperatures exceeding 973K. Additionally, ODS ferritic steels are also being consid- ered for fusion reactor applications. The rich experi- ence in the development of fast reactor materials would enable launching the advanced ferritic steels for fusion technology, in a shorter time span. 4.03.6 Ferritic Steels for Out-of-Core Applications: Improvements in Joining An ambitious target of increasing the temperature and pressure of steam in many power plants has provided a high impetus for the development of steels with better high temperature properties. Very often, the weld joints play a crucial life limiting role in these components. One of the recurrent problems is the frequent failure of weldments due to Type IV cracking (see below), in weldments of ferritic steels subjected to creep loading. Another problem encoun- tered during service exposure of joints of dissimilar ferritic steels is the failure due to the formation of hard brittle zone at the heat-affected zone (HAZ). Both these issues are discussed below. The modified 9Cr–1Mo steel fusion weld joint (Figure 14) consisting of base metal, deposited weld metal, and the HAZ produces a complex heteroge- neous microstructure due to thermal cycle. The base metal and weld metal consist of a tempered martens- ite structure, with columnar grains in the weld metal. Weld metal Interface 20 mm 9Cr–1Mo weld metal Weld metal HAZ CGHAZ FGHAZ C G H A Z Interface Figure 14 Schematic showing different zones in a ferritic stee metal, base metal, interface, coarse grained heat-affected zone and intercritical zone (ICZ) of the weldment. The HAZ comprises coarse prior-austenitic grain martensite, fine prior-austenitic grain martensite and an intercritical structure, as one traverses from the weld fusion interface toward the unaffected base metal. This is dictated by the peak temperatures experienced by the base metal during weld thermal cycle and the phase transformation characteristics of the steel. It has been established that the localized microstructural degradation in the intercritical region of HAZ is mainly responsible for the prema- ture creep-rupture strength of Cr–Mo weld joint and can be overcome if residual stresses of the weld are adequately relieved by PWHT. The lower creep-rupture strength of weld joint than the base metal is due38,77 to the different types of cracking developed during creep exposure. Four types of cracking have been identified (Figure 15) in Cr–Mo steel weld joint. They have been categor- ized as Type I, Type II, Type III, and Type IV. The Type I and Type II cracks originate in the weld metal, propagate either through the weld metal itself (Type I) or cross over in the HAZ (Type II). The Type III cracking occurs in the coarse grain region of HAZ and can be avoided by refining the grain size. Type IV cracking nucleates and propagates in the intercritical/fine grain region of HAZ. Type IV fail- ure occurs at longer creep exposure and higher test temperature, by coalescence of fine cavities leading to microcracks (Figure 16(a)) and their eventual propagation to the surface. Base metal 20 mm 20 mm20 mm 20 mm ICZ Base metal FG H A Z IC Z 21/4Cr–1Mo Base metal l weldment and optical micrographs obtained from weld (CGHAZ), fine grained heat-affected zone (FGHAZ), (a) Cavity Cavity associated with precipitate 10 mm 1 mm Z-phase [111] Z phase M23C6 (b) Figure 16 Type IV cracking in same sample as in Figure 15. (a) cavities in the intercritical region and (b) Z-phase77 in creep-tested 9Cr–1Mo steel. The inset shows the microchemistry of the Z-phase. Base metal Base metalHAZ HAZ IV II III I Weld metal Base metal HAZ HAZ (a) (b) Weld metal 5 Cm Figure 15 Locations77 of different types of failure in weld geometry of the ferritic steels: (a) schematic representation and (b) experimental observation in creep tested weldment of 9Cr–1Mo steels. Ferritic Steels and Advanced Ferritic–Martensitic Steels 117 The type IV cracking susceptibility, defined as the reduction in creep-rupture strength of weld joint compared to its base metal, depends on the type of ferritic steel. 2.25Cr–1Mo steel is most susceptible to type IV cracking; whereas the plain 9Cr–1Mo steel is the least susceptible. At higher test temperature, the type IV cracking susceptibility is higher in modified 9Cr–1Mo steel than the plain steel. This is related77 to the precipitation of Z-phase (Figure 16(b)), a com- plex Cr (V, Nb) N particle, in the modified steel. The Z-phase grows rapidly at elevated temperatures dur- ing long term exposure, by dissolving the beneficial MX types of precipitates. This promotes the recovery of the substructure with associated decrease in strength in the intercritical region of HAZ. Although it is difficult to completely eliminate Type IV cracking, several methods are being adopted to improve type IV cracking resistance. It is more severe in thick sections due to the imposed geometri- cal constraint. A design modification can be adopted to decrease the variation in tensile stresses across the welded section of the component or avoid joints in critical regions having high system stresses and relo- cate them in the less critical region. Strength homo- geneity across the weld joint can also be improved by a suitable PWHT. An increase in width of the HAZ can reduce the stress triaxiality such that the soft intercritical region deforms with less constraint with the consequence of reduced creep cavitation, to minimize type IV cracking tendency. The width of the HAZ can be increased both by changing pre- heat and heat-input during welding. Another con- trasting approach to overcome type IV cracking is to avoid or minimize the width of the HAZ, to elimi- nate the intercritical zone. This is being attempted by employing advanced welding techniques such as laser welding. The resistance against intercritical softening can also be improved by increasing the base strength of the steel with the addition of solid solution hard- ening elements such as W, Re, and Co and also by microalloying the steel with boron. Microalloying with boron retards the coarsening rate of M23C6 by replacing some of its carbon. The boron content needs to be optimized with the nitrogen content to avoid BN formation. Addition of Cu is also found to be beneficial. Copper is almost completely insoluble in the iron matrix and when added in small amounts, precipitates as nanosize particles to impart creep resistance. A suitable adjustment of the chemical composition of steel within the specifi- cation range also reduces the large difference in creep strength between the softened HAZ, the base metal, and the coarse grain HAZ of the joint. A weld joint of modified 9Cr–1Mo steel with low carbon, nitrogen, and niobium has been reported to possess creep strength comparable to that of the base steel. It is expected that a judicious combination of changes in chemistry and process variables would reduce the failures due to type IV cracking in weld- ments of ferritic steels subjected to creep loading. Another frequent problem78–81 is the formation of ‘hard brittle zone’ during service exposure of dissim- ilar joints between ferritic steels, leading to failures. The formation (Figure 17(a)) of microscopic layer of 118 Ferritic Steels and Advanced Ferritic–Martensitic Steels hard, brittle zone along the HAZ in dissimilar weld- ments of steels is known to be responsible for the cold cracking, stress corrosion cracking, and higher fre- quency of failures of the weldments. This is one of the cases where modeling has enabled an in-depth understanding of the problem, in addition to providing an industrial solution to prevent the for- mation of brittle zone. The brittle layer at the interface between 9Cr–1Mo weld and 2.25Cr–1Mo base metal is shown77 to be a manifestation of a number of syner- gistic factors: (a) microstructural changes in regions close to the heat source during welding (b) migration of carbon during PWHT, driven by the gradient in its activity and (c) formation (inset in Figure 17(a)) (a) A A B 9C r– 1M o 2. 25 C r– 1M o p p t zo ne 100 m S of t zo ne (b) M 1C M23C5 500 nm ( 0.5 0.4 0.3 0.2 C m ax im um in h ar d z on e (w t% ) 0.1 0.00 (c) 0.02 0.04 0.06 d (mm) 0.08 0 Ni Co Cu Figure 17 Modeling in preventing formation of hard zone in d formation of hard zone in the heat-affected zone between 9Cr–1 1023K for 15 h. The hard zone is marked as A and the soft zone a individual carbides in the hard zone. (b) Finite difference method the same weld geometry. (c) Variation79 of amount of carbon in introduced between the 9Cr–1Mo and the 2.25Cr–1Mo to preve atoms in a bcc iron lattice calculated using molecular dynamics of series of fine carbides when there is a local supersaturation of carbon. It has been possible to use modeling methods like Finite Difference Meth- ods to predict the carbon profile across the weld region of 9Cr–1Mo and the base metal of 2.25Cr– 1Mo (Figure 17(b)), which were in good agreement with the profiles obtained using electron probe microanalysis. These calculations could be refined using Thermo-Calc and diffusion-controlled trans- formations (DICTRA) to take into account the simultaneous precipitation of carbides and diffusion of carbon. These computational methods were instru- mental in predicting the methods to prevent the formation of hard zone in dissimilar joints of ferritic steels. Three elements which would repel carbon 0.50 0.45 0.40 0.35 0.30 0.25 0.20 0.15 0.10 0.05 b) 0.00 -0.1 0.0 ‘x’ (mm) C ar b on c on ce nt ra tio n (w t% ) 9Cr–1Mo HZ SZ Temp: 1023 K 0.1 1 h 15 h 21/4Cr–1Mo .10 Z 7 6.5 6 5.5 5 4.5 4 6 5 4 3 4.5 5.0, 5.5, 5.5 5 6 6.5 7 5.5Y(d) X issimilar joints of ferritic steels: (a) Optical micrograph78 of Mo weld and the 2.25Cr–1Mo base metal, exposed to s B. The inset shows the transmission electronmicrograph of calculation78 predicting the diffusion profile of carbon in the hard zone versus thickness of the ‘diffusion barrier’ nt the formation of hard zone and (d) positions81 of carbon . Ferritic Steels and Advanced Ferritic–Martensitic Steels 119 atoms, that is, with the positive interaction energy were chosen for this purpose. Figure 17(c) shows78 the comparison of three different metals, Ni, Cu, and Co in preventing the formation of hard zone. Experimental confirmation was obtained79 using inter- layer between the two dissimilar ferritic steels. Fur- ther insight could also be arrived81 at in the diffusion behavior (Figure 17(d)) of carbon interstitial in the lattices of bcc iron and fcc nickel using molecular dynamics. These calculations could demonstrate that the activation energy for diffusion of carbon in a fcc nickel lattice is higher than bcc iron. This sluggish diffusion kinetics is due to the repulsive potential of nickel toward carbon, which is the main reason for the choice of nickel as the most effective diffusion barrier between the two ferritic steels. Thus, an indus- trial solution to prevent the formation of brittle zone in joints of dissimilar ferritic steels after service exposure could be arrived at, based on an in-depth understanding of the interaction between the lattice potentials of atoms. It has been demonstrated in the above studies that modeling methods could be used most effectively to reduce the experimental time required for overcom- ing an industrial problem. Experimental benchmark- ing was required only for final confirmation of the predictions. These trends are becoming more common in almost all problems in materials technol- ogy, in recent years, be it atomistic mechanisms or fabrication technologies or prediction of life of com- ponents. It is hoped that this approach of knowledge- based design of materials would gradually replace the time consuming empirical methods of today. 4.03.7 Summary Future trends in the global fast reactor industry are toward higher operating temperatures, higher burn- up (250GWd t�1), higher breeding ratios (�1.4) and longer lifetime for reactor (60–100 years). These goals require several developments in materials sci- ence and technology across all components of nuclear plants, especially for core component materials. Ferritic steels have a much better void swelling resistance compared to currently used austenitic stainless steels and are capable of enhancing the burn-up of the fuel up to about �200GWd t�1. Ferritic–martensitic steels based on 9–12% Cr com- positions exhibit the highest swelling resistance and a number of commercial swelling resistant materials have been marketed. The principles behind the design of swelling resistant ferritic steels for core components of fast reactors have been discussed. However, their use is rendered difficult due to their poorer creep strengths at temperatures higher than�873K. Improvement of higher temperature ten- sile and creep strengths in these alloys will enable us to achieve higher temperatures, in addition to higher burn-up, thus improving the economics of nuclear power production. Presently, the reduced creep strength of 9–12Cr ferritic steels at temperatures above 798K, has restricted their use to certain low stressed components such as subassembly wrappers. Another crucial problem is ‘embrittlement’ in ferritic steels. The mechanisms and methods which are being attempted to overcome embrittlement problems are discussed. Alloy development programmes are in progress to explore ferritic–martensitic oxide dispersion strength variants, for higher target burn-up of 250 dpa, with enhanced high temperature (�973K) capability, by improving mechanical properties. Conventional alloy melting routes will have to be abandoned in favor of powder metallurgy techniques of ball-milling, hot isostatic pressing, and hot extrusion for the synthesis of these ODS steels. Process optimization for the development of 9Cr-based ferritic–martensitic steels strengthened by a fine dispersion of yttria nanoparticles has been completed. The major con- cerns in this family of ferritic or ferritic–martensitic steels are the anisotropy of properties in ferritic 12Cr steels or oxidation resistance in 9Cr steels, fabrication procedure, microstructural stability under irradia- tion, and dissolution during back-end technologies. 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All rights reserved. 4.04.1 Introduction 123 4.04.2 Void Swelling 124 4.04.2.1 Compositional Dependence of Void Swelling 124 4.04.2.2 Void-Swelling Models 129 4.04.2.3 Swelling Behavior of Neutron-Irradiated Nimonic PE16 133 4.04.3 Irradiation Creep 136 4.04.4 Microstructural Stability 138 4.04.4.1 Dislocation Structures 138 4.04.4.2 Precipitate Stability 139 4.04.5 Irradiation Embrittlement 140 4.04.5.1 Fast Neutron Irradiation Experiments 140 4.04.5.2 Helium Implantation Experiments 145 4.04.6 Concluding Remarks 147 References 148 Abbreviations AGR Advanced gas-cooled reactor DFR Dounreay Fast Reactor EBR-II Experimental Breeder Reactor-II EDX Energy dispersive X-ray HVEM High-voltage electron microscope N/2 dpa calculated according to half-Nelson model NRT dpa calculated according to Norget, Robinson, and Torrens model OA Overaged PFR Prototype Fast Reactor PS Proof stress SIPA Stress-induced preferred absorption ST Solution treated STA Solution treated and aged TEM Transmission electron microscope UTS Ultimate tensile strength VEC Variable energy cyclotron 4.04.1 Introduction Research into the effects of irradiation on nickel- based alloys peaked during the fast reactor develop- ment programs carried out in the 1970s and 1980s. Interest in these materials focused on their high resis- tance to radiation-induced void swelling compared to austenitic steels, though a perceived susceptibility to irradiation embrittlement limited their application to some extent. Nevertheless, the Nimonic alloy PE16 was successfully used for fuel element cladding and subassembly wrappers in the United Kingdom, and Inconel 706 was utilized for cladding in France. Both of these materials are precipitation hardened and consequently have high creep strength, and much research and development of alternative alloys was directed toward maintaining swelling resistance and creep strength while aiming to alleviate, or at least understand, irradiation embrittlement effects. There has been some revival of interest in nickel-based alloys for nuclear applications, and various aspects of radiation damage in suchmaterials have recently been reviewed by Rowcliffe et al.1 in the context of Gener- ation IV reactors, and by Angeliu et al.2 in consider- ation of their use for the pressure vessel of the Prometheus space reactor. Nickel-based alloys are also candidate structural materials for molten salt reactors, for which resistance to corrosion by molten fluoride salts and high-temperature creep strength are prime requirements, though intergranular attack by the fission product tellurium and irradiation embrittlement due to helium production are poten- tially limiting factors for this application.3 This chapter focuses on the void swelling behav- ior, irradiation creep, microstructural stability, and irradiation embrittlement of precipitation-hardened nickel-based alloys. Fundamental to all of these effects are the basic processes of damage production 123 124 Radiation Effects in Nickel-Based Alloys by the creation of vacancies and interstitial atoms in displacement cascades, and the ways in which these point defects migrate and interact with, causing the redistribution of, solute atoms. Detailed discus- sions of damage processes and radiation-induced segregation are beyond the scope of this chapter but these topics will be introduced where necessary, particularly in relation to void swelling models. More detailed reviews are given in Chapter 1.01, Fundamental Properties of Defects in Metals; Chapter 1.03, Radiation-Induced Effects on Microstructure; Chapter 1.11, Primary Radiation Damage Formation; Chapter 1.12, Atomic-Level Level Dislocation Dynamics in Irradiated Metals and Chapter 1.18, Radiation-Induced Segregation. Typical compositions of nickel-based alloys and some precipitation-hardened steels, which are con- sidered in this chapter, are shown in Table 1. Alloy compositions are generally given in weight percent throughout this chapter unless stated otherwise. Precipitation-hardened alloys may be utilized in a number of different heat-treated conditions, which are generally abbreviated here as ST (solution trea- ted), STA (solution treated and aged), and OA (over- aged). Further information on the material properties of nickel alloys is given in Chapter 2.08, Nickel Alloys: Properties and Characteristics. Neutron fluences are generally given for E> 0.1MeV unless indicated otherwise. Atomic dis- placement doses (dpa) are generally given in NRT (Fe) units, although the half-Nelson (N/2) model was Table 1 Nominal compositions (wt%) of commercial and de Alloy Ni Cr Mo Ti Al Nimonic PE16 43 16.5 1.1 1.2 1.2 Inconel 706 41.5 16 – 1.8 0.2 Inconel 718 52.5 19 3.0 0.9 0.5 Inconel 600 75 16 – 0.3 0.2 Inconel 625 61 22 9.0 0.3 0.3 Incoloy 800 34 20.5 – 0.4 0.4 Hastelloy X 48 21 9.0 – – D21a 25 8.4 1.0 3.3 1.7 D25a 30 10.5 3.7 1.8 1.3 D66a 45 12 3.0 2.5 2.5 D66b 40 11 2.0 3.0 1.5 D68a 45 12 – 1.8 0.4 D68b 34 12.5 – 1.6 0.25 PE16 matrix 36 20 4.0 – – Incoloy DS 39 18 – 0.04 0.02 Alloy 7817 40 15 3.2 2.0 0.9 Alloy 7818 40 15 3.0 0.3 – aComposition indicated by Yang et al.4 bComposition indicated by Toloczko et al.5 widely used particularly in the United Kingdom in the 1970s6. The exact relationship between these units will vary depending on the neutron spectrum (which may differ, not only from one reactor to another, but also depending on location within a reactor), but approximate conversion factors for fast reactor core irradiations are 1026nm�2 E > 0:1MeVð Þ ¼ 5dpa NRT Feð Þ ¼ 6:25dpa N=2ð Þ 4.04.2 Void Swelling 4.04.2.1 Compositional Dependence of Void Swelling Nimonic PE16 was first identified as a low-swelling alloy in the early 1970s. Void swelling data derived from densitymeasurements on fuel pin claddingmate- rials from the Dounreay Fast Reactor (DFR) were reported by Bramman et al.7 and were complemented by electron microscope examinations described by Cawthorne et al.8 Swelling in STA PE16 was found to be lower than in heat-treated austenitic steels and comparable to cold-worked steels. Comparison of data for PE16 and FV548 (a Nb-stabilized austenitic steel) irradiated under identical conditions in DFR to a peak neutron fluence of �6� 1026 nm�2 indicated that the lower swelling of PE16 was due to smaller void concentrations at irradiation temperatures up to �550 �C and reduced void sizes at higher velopmental nickel-based alloys Nb Mn Si C Other Fe – 0.1 0.2 0.05 Bal. 2.9 0.2 0.2 0.03 Bal. 5.2 0.2 0.2 0.04 Bal. – 0.2 0.2 0.08 0.3Cu Bal. 3.5 0.2 0.2 0.05 Bal. – 0.9 0.5 0.07 0.5Cu Bal. – 0.5 0.5 0.10 0.5W, 2.0Co Bal. – 1.0 1.0 0.04 Bal. – 1.0 1.0 0.04 Bal. – – 0.5 0.03 Bal. – 0.2 0.5 0.04 Bal. 3.6 0.3 0.4 0.03 Bal. 2.8 0.2 0.4 0.02 Bal. – 0.1 0.2 0.07 Bal. – 1.0 2.0 0.08 Bal. – 0.2 0.5 0.02 Bal. 3.0 0.2 0.5 0.02 Bal. 8 9 7 6 4 53 2 1 13121110 0 10 0 10 20 30 40 50 60 70 80 90 20 30 40 50 60 Nickel (wt%) S w el lin g (% ) Commercial alloys Fe–15Cr–Ni alloys Commercial alloys: 13 Inconel 702 1 Type 304 2 Type 321 3 Type 316 4 Type 318 5 12R72HV 6 A286 7 Incoloy 800 8 Inconel 706 9 Nimonic PE16 10 Hastelloy X 11 Inconel 625 12 Inconel 600 Figure 1 Swelling versus nickel content of commercial alloys and ternary Fe–15Cr–Ni alloys bombarded with Ni2þ ions to a damage level of 116dpa at 625 �C. Reproduced from Johnston, W. G.; Rosolowski, J. H.; Turkalo, A. M.; Lauritzen, T. J. Nucl. Mater.1974, 54, 24–40. Radiation Effects in Nickel-Based Alloys 125 temperatures. At around the same time, Hudson et al.9 compared the swelling behavior of PE16, type 316 steel, and pure nickel, using 20MeV C2+ ion irradia- tions in the Harwell VEC (variable energy cyclotron). The materials were implanted with 10 appm (atomic parts per million) of helium prior to ion bombardment to peak displacement doses>200 dpa (N/2) at 525 �C. Void swelling in 316 steel and nickel exceeded 10% at the highest doses examined, compared to �0.5% in PE16. Void nucleation appeared to occur earlier in nickel (at �0.1 dpa) than in PE16 or type 316 (�2 dpa), but the peak void concentration was higher by a factor of about 10 in the austenitic steel than in nickel or PE16. Hudson et al.9 originally attributed the swelling resistance of PE16 to the presence of the coherent, ordered face-centered cubic, Ni3(Al,Ti) g0 precipi- tates, which were thought either to trap vacancies and interstitials at their surface, thereby enhancing point- defect recombination, or to inhibit the climb of dis- locations, thereby preventing them from acting as preferential sinks for interstitial atoms. In support of the first of these two suggested mechanisms, Bullough and Perrin10 argued that the surface of a coherent precipitate would be a more effective trapping site than an incoherent one where the identity of the point defects would immediately be lost (and where, as a consequence, void nucleation was likely to occur). The efficiency of point defect trapping would be expected to be greater the higher the total surface area of the g0 precipitates, that is, to be inversely proportional to the precipitate size at constant volume fraction. On the other hand, the second mechanism proposed by Hudson et al. should be most effective at an intermedi- ate particle size where dislocation pinning is strongest. Support for the latter processwas provided byWilliams and Fisher11 from HVEM (high-voltage electron microscope) irradiations of PE16 at a damage rate of about 10�2 dpa s�1 at 600 �C, where the swelling rate was higher at small (3 nm) and large (70 nm) g0 particle diameters than at intermediate sizes of about 20 nm. However, it is now considered that any effect that the g0 precipitates may have on the swelling resistance of Nimonic PE16 is secondary to that of the matrix composition. The generally low-swelling behavior of Ni-based alloys compared to austenitic steels was shown by Johnston et al.12 following bom- bardment with 5MeVNi2+ ions at 625 �C. The dam- age rate in these experiments was 10�2 dpa s�1 and the displacement dose was originally estimated as 140 dpa but this was subsequently revised by Bates and Johnston13 to 116 dpa (based on displacement energy Ed¼ 40 eV). In addition to precipitation- hardened alloys, including PE16 and Inconel 706, this experiment included nonhardenable high-Ni alloys, such as Inconel 600 and Hastelloy X, a range of commercial steels, and Fe–Cr–Ni ternary alloys containing 15% Cr and 15–35% Ni. The alloys were preimplanted with 15 appm helium prior to ion bom- bardment, and the irradiation temperature was chosen as being close to the peak swelling temperature for ion- irradiated austenitic steels. The extent of void swelling was determined by electron microscope examinations in low-swelling alloys, but was estimated from step- height measurements (comparing the surfaces of irra- diated and nonirradiated regions) in high-swelling materials. As illustrated in Figure 1, the results showed negligible swelling ( 126 Radiation Effects in Nickel-Based Alloys an increase in overall swelling from 0.1MeV) at tem- peratures from 454 to 593 �C. Swelling rates during Ni ion irradiations at 675 �C were higher by a factor of about five in reactor-conditioned material than in a nonconditioned sample. The increased swelling rate was attributed to changes in the matrix composition resulting from an increased volume fraction of g0 in the reactor-conditioned material. Early attempts to account for the effects of matrix composition on void swelling focused on the stability of the austenite phase. Harries16 suggested that the swelling behavior of austenitic steels and nickel- based alloys could be rationalized in terms of their Ni and Cr equivalent contents (i.e., the relative austenite and ferrite stabilizing effects of their con- stituent elements), with the composition of high- swelling alloys then falling into the (gþs) phase field in the Fe–Cr–Ni ternary phase diagram. Harries postulated that the partitioning of solute elements into the sigma phase would have a detrimental effect on the swelling resistance of austenite. Watkin17 took a similar approach, but found that an improved correlation could be obtained using the concept of electron vacancy numbers rather than Ni and Cr equivalents. The average electron vacancy number, Nv, of the matrix is calculated from the atomic frac- tions of its constituents, with allowance being made for the precipitation of carbides and g0 (or g00, etc.), and has been widely used to predict the susceptibility of nickel-based alloys to the formation of intermetal- lic phases.18Nv was calculated from: Nv ¼ 0:66Niþ 1:70Coþ 2:66Feþ 3:66Mn þ 4:66 CrþMoð Þ Watkin found that void swelling in a range of alloys with Ni contents up to �43%, which were irradiated in DFR to a peak dose of 30 dpa at 600 �C, remained low for Nv below about 2.5 (corresponding to low susceptibility to s phase formation), but increased approximately linearly at higher Nv. However, as was clearly argued by Bates and Johnston,13 correlations based on sigma-forming tendency could not account for the minimum in swelling observed at about 45% Ni, since higher Ni contents should continue to be beneficial. Radiation Effects in Nickel-Based Alloys 127 A better understanding of the swelling behavior of Fe- and Ni-based alloys resulted from a series of fast neutron irradiation experiments which were carried out in EBR-II in the early 1980s. Irradiation tem- peratures in these experiments ranged from about 400 to 650 �C. Initial data for a range of commercial alloys, including ferritic and austenitic steels, as well as nickel-based alloys, were reported by Bates and Powell19 and Powell et al.,20 with higher dose data (up to a peak fluence (E> 0.1MeV) of �25� 1026 nm�2, corresponding to �125 dpa) being reported by Gelles21 and Garner and Gelles.22 Swelling data for Fe–Cr–Ni ternary alloys, irradiated in EBR-II to a peak fluence of 22� 1026 nm�2 (�110 dpa), were presented by Garner and Brager.23 The extent of void swelling in these experiments was determined by density change measurements. In general, alloys with nickel contents in the range of 40–50% exhib- ited the lowest swelling. Swelling in commercial nickel-based alloys was generally lower in ST than in aged conditions, this being attributed to the bene- ficial (though temporary) effect of minor elements remaining in solution and being able to interact with point defects19; subsequent precipitation during irra- diation would be expected to reduce this benefit and the resulting densification, though small, would also effectively reduce the measured swelling. Swelling data for a number of STalloys, which were irradiated in the AA-1 rig in EBR-II, are shown in Figure 2; data are shown for two withdrawals, at peak fluences of 14.7� 1026 nm�2 and 25.3� 1026 nm�2, with mea- surements for Inconel 600 and Inconel 625 reported at both fluence levels, data for Nimonic PE16 and Inconel 706 at the lower level, and data for Incoloy 800 and Hastelloy X at the higher level. The nickel contents of the alloys range from about 34% in Inco- loy 800 to 75% in Inconel 600. Swelling remained relatively low in the three Inconel alloys and in PE16. However, both Incoloy 800 and Hastelloy X exhibited high swelling at some temperatures, with swelling in the latter reaching �80% at 593 �C. The reason for such high swelling in neutron-irradiated Hastelloy X (nickel content �48%) is unclear, but it was noted that densification up to 3% occurred at the lower irradiation temperatures – indicating microstructural instability and possibly signaling changes in the com- position of the matrix which may have affected the swelling behavior. (Note that Hastelloy X was identi- fied as a low-swelling alloy in the Ni2+ ion irradiation experiments described by Johnston et al.12) Some data for different heat-treated conditions of PE16 at the higher fluence level were reported by Garner and Gelles,22 and are compared for irradia- tions at 538 �C (more or less corresponding to the peak swelling temperature for PE16 in the AA-1 experiment) with lower fluence data from Bates and Powell19 in Figure 3. The heat-treated condi- tions indicated in Figure 3 are ST (ST 4 h at 1080 �C), A1 (ST and aged 16 h at 705 �C), A2 (ST and aged 1 h at 890 �C plus 8 h at 750 �C), and OA (ST and aged 24 h at 840 �C). Note that the silicon content of the PE16 used in these experiments was much lower at 0.01% than the level of �0.2% typi- cally found in UK heats of the alloy. Overall, the data appear to show little effect of initial heat treatment on the swelling of PE16, except that the OA condi- tion exhibited the most swelling (5.2%) at the higher fluence. Although it is clear that the swelling behavior of austenitic alloys is largely dependent on nickel con- tent, there is ample evidence to show that minor solute additions can have significant effects. Much of the work on minor solutes has focused on steels similar to type 316, but some data are available for higher nickel alloys. For example, Mazey and Hanks24 used 46.5MeV Ni6+ ion irradiations to examine the effects of Si, Ti, and Al additions on the swelling response of model alloys with base compositions approximating that of the matrix phase in PE16. Solute additions of �0.25% Si or 1.2% Ti reduced swelling, but the addition of �1.2% Al (in the absence of Si or Ti) markedly increased it. The beneficial effect of Si was believed to arise from its high diffusivity in solution (this is discussed further in Section 4.04.2.2), whereas that of Ti appeared to be related to the formation of Z phase (hexagonal-structured Ni3Ti). The addition of Al resulted in an increase in the concentration of voids, the surfaces of which were coated in a thin layer of the g0 phase (Ni3Al). A beneficial effect of Si on the swelling response of modified Incoloy DS alloys under Ni6+ ion irradiation was also reported byMazey et al.25 However, it should be noted that high Si con- tents can give rise to the formation of radiation- induced phases which are enriched with Ni and Si, such as the Ni3Si form of g0 and the silicide G-phase (M6Ni16Si7, where M is usually Ti, Nb, or Mn). G-phase particles are generally found in association with large voids and their formation may therefore give rise to an increase in the swelling rate.26,27 Swelling data derived from density measurements for neutron irradiated, modified Incoloy DS alloys, with Si contents ranging from 0.19 to 2.05% (com- pared to a specified level of 1.9–2.6% in the commer- cial alloy), are compared with data for a ‘PE16 matrix −0.5 0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5 4.0 4.5 5.0 Temperature (°C) S w el lin g (% ) 0 5 10 15 Fl ue nc e (1 02 6 n m −2 , E > 0. 1 M eV ) Fl ue nc e (1 02 6 n m −2 , E > 0. 1 M eV ) In 600 In 625 PE16 In 706 Fluence 350 400 450 500 550 600 650 700 350 400 450 500 550 600 650 700 −10 0 10 20 30 40 50 60 70 80 90 100 Temperature (°C) S w el lin g (% ) 0 5 10 15 20 25 In 600 In 625 In 800 Hast X Fluence Figure 2 Void swelling of nickel-based alloys irradiated in AA-1 rig in Experimental Breeder Reactor-II. Based on data from Bates, J. F.; Powell, R. W. J. Nucl. Mater.1981, 102, 200–213; Garner, F. A.; Gelles, D. S. In Effects of Radiation on Materials: 14th International Symposium; Packan, N. H., Stoller, R. E., Kumar, A. S., Eds.; American Society for Testing and Materials: Philadelphia, PA, 1990; Vol. II, pp 673–683, ASTM STP 1046. 128 Radiation Effects in Nickel-Based Alloys alloy’ and Nimonic PE16 in Figure 4. The materials were all in ST condition apart from PE16 which was in an STA condition (aged 4 h at 750 �C). The alloys were irradiated in the UK-1 rig in EBR-II to fluences in the range of 9–16� 1026 nm�2 (E> 0.1MeV) at temperatures of �390–640 �C. These data are pre- viously unpublished except those for STA PE16 (heat DAA 766) which were reported by Boothby.28 Swelling in the modified Incoloy DS alloys generally decreased with increasing Si content. The 0.19% Si alloy exhibited high swelling at all temperatures with indications of swelling peaks at about 440 and 640 �C. Increased Si levels tended to suppress the high tem- perature swelling peak and reduce the magnitude of swelling at lower temperatures. The PE16 matrix alloy containing 0.24% Si exhibited a high tempera- ture swelling peak but moderate swelling below �550 �C, suggesting a beneficial effect of Mo (this being the main compositional difference between the PE16 matrix alloy and the modified Incoloy DS alloys) 0.0 2.0 4.0 6.0 8.0 10.0 12.0 Temperature (°C) 300 350 400 450 500 550 600 650 S w el lin g (% ) 0 5 10 15 Fl ue nc e (1 02 6 n m –2 , E > 0. 1 M eV ) PE16 STA PE16 Matrix DS (0.19 Si) DS (0.56 Si) DS (1.05 Si) DS (2.05 Si) Figure 4 Void swelling data derived from density measurements for Nimonic PE16, a PE16 matrix alloy, and modified Incoloy DS alloys, irradiated in the UK-1 rig in Experimental Breeder Reactor-II. Unpublished data from Boothby, R. M.; Cattle, G. C. Void Swelling in EBR-2 Irradiated Nimonic PE16 and Incoloy DS; FPSG/P(90)10, with permission from AEA Technology Plc. 0.0 ST A1 A2 OA 1.0 2.0 3.0 4.0 5.0 6.0 S w el lin g (% ) 68 dpa 116 dpa Figure 3 Effect of heat treatment on void swelling of Nimonic PE16 irradiated in Experimental Breeder Reactor-II at 538 �C. Adapted from Bates, J. F.; Powell, R. W. J. Nucl. Mater.1981, 102, 200–213; Garner, F. A.; Gelles, D. S. In Effects of Radiation on Materials: 14th International Symposium; Packan, N. H., Stoller, R. E., Kumar, A. S., Eds.; American Society for Testing and Materials: Philadelphia, PA, 1990; Vol. II, pp 673–683, ASTM STP 1046. Radiation Effects in Nickel-Based Alloys 129 at lower temperatures. However, swelling in the PE16 matrix alloy remained significantly higher at all tem- peratures than in STA Nimonic PE16 (containing 0.15% Si), indicating a significant benefit of the g0 forming elements Al and Ti. 4.04.2.2 Void-Swelling Models Point defects created by atomic displacements are lost either through mutual recombination or by migration to sinks. Void swelling requires a mobile population of excess vacancies and can only occur over a limited temperature range, typically �350– 700 �C in neutron-irradiated steels and nickel-based alloys. Rapid diffusion at higher temperatures reduces the concentration of radiation-induced vacancies to near thermal equilibrium levels. Recombination dom- inates at lower temperatures, where reduced vacancy mobility prevents the formation of voids as the neces- sary counter-migration of matrix atoms cannot occur. In the swelling regime, an increased bias for intersti- tials over vacancies at dislocation sinks gives rise to the surplus vacancies which agglomerate to form voids. The flux of point defects to sinks, including void surfaces, dislocations, and grain boundaries, results in the segregation of particular solute atoms at the sinks and the depletion of others. In austenitic steels and nickel-based alloys, it is generally found that nickel segregates at the point defect sinks. This is generally attributed to the inverse Kirkendall effect described by Marwick,29 whereby faster diffusing solutes such as Cr move in the opposite direction to the vacancy flux and are depleted at the sink, and slower diffusing solutes such as Ni are enriched. One of the earliest observations of nickel segregation at void surfaces due to the inverse Kirkendall effect was made by Marwick et al.30 in an alloy with a composition cor- responding to that of the matrix phase in Nimonic PE16. (For more detailed discussions on radiation- induced segregation effects, see the reviews of Wiedersich and Lam,31 and Rehn and Okamoto.32) Venker and Ehrlich33 recognized that differences in the partial diffusion coefficients of alloy constitu- ents might account for the effects of composition on swelling. Any effect of this kind would generally be expected to be more significant the larger are the differences between the partial diffusion coefficients of the alloy components. Garner and Wolfer34 exam- ined Venker and Ehrlich’s conjecture and concluded that the addition of even small amounts of a fast- diffusing solute such as silicon to austenitic alloys would greatly increase the effective vacancy diffusion 130 Radiation Effects in Nickel-Based Alloys coefficient (i.e., would enhance the diffusion rate for all matrix elements). The overall effect is analogous to an increase in temperature – resulting in an effective decrease in the vacancy supersaturation and hence a reduction in the void nucleation rate. This mecha- nism is generally accepted as the explanation for the beneficial effect of silicon in reducing swelling in austenitic steels and nickel-based alloys. Although this relies on the diffusion of silicon via vacancy ex- change, silicon is also generally observed to segregate to point defect sinks and since it is an undersized solute, this is believed to occur by the migration of interstitial–solute complexes. There is, however, no reason to suppose that both diffusion mechanisms cannot operate simultaneously. Garner and Wolfer34 originally considered that since nickel diffuses relatively slowly in austenitic alloys, an increase in nickel content would have the opposite effect to silicon. However, a later assessment made by Esmailzadeh and Kumar,35 based on diffu- sion data reported by Rothman et al.,36 indicated that the void nucleation rate in Fe–15Cr–Ni alloys would decrease with an increase in nickel content from 20 to 45%. This result is obtained because, although nickel remains the slowest diffusing species, the effective vacancy diffusion coefficient of the system is calcu- lated to increase at the higher nickel content. Esmail- zadeh and Kumar’s calculations also confirmed the beneficial effect of silicon, with the addition of 1% Si predicted to be as effective in suppressing void nucleation as increasing the nickel content from 20 to 45%. Effects at nickel contents above 45% could not be examined due to a lack of appro- priate diffusion data. As well as affecting the nucleation of voids, differ- ences in the diffusion rates of the various solutes might also be expected to influence void growth. Simplistically, this can be thought of as being partly due to the segregation of slower diffusing solutes reducing the rate of vacancy migration in the vicinity of the voids. However, a further consequence of such nonequilibrium solute segregation was identified by Marwick,29 who realized that it would give rise to an additional vacancy flux which would oppose the radiation-induced flux to the sink. As discussed by Marwick, this additional flux (the Kirkendall flux) may itself be an important factor in limiting void growth, since it will reduce the probability of vacancy annihilation at sinks and increase the likelihood of point defect recombination. The effect of nickel content on void swelling was considered further in a model developed by Wolfer and coworkers.37,38 The model examined the compo- sitional dependence of the void bias and focused on the effects of nickel segregation at void surfaces. Wolfer’s model indicated that the compositional gra- dients produced by radiation-induced segregation give rise to additional drift forces which affect the point defect fluxes and thereby modify the bias terms. These additional drift forces arise from the effects of composition on point defect formation and migration energies, on the lattice parameter and the elastic mod- uli, and from the Kirkendall flux. Wolfer’s calculations for binary Fe–Ni alloys indicated that the effect of the Kirkendall flux is small for interstitials but significant for vacancies. Nevertheless, it was considered that the overall effect of compositional gradients on the bias terms is likely to be greater for interstitials than for vacancies due to other factors, particularly the effect of variations in the elastic moduli. As noted by Garner and Wolfer,39 an increase in the shear modulus in the segregated regions around voids would reduce the bias for interstitials and therefore help to stabilize voids. It is difficult to predict the significance of this effect in complex alloys, however, since depletion of Cr in the segregated region will tend to reduce the shear modu- lus, whereas enrichment of Ni in high-Ni alloys will tend to increase it.38 A more significant result of the model with regard to the effect of nickel on swelling is that there is a reversal in the sign of the Kirkendall force for vacancies in Fe–Ni alloys at�35%Ni. Below this level, vacancies are predicted to be attracted into regions of higher Ni concentration, but above it, the opposite occurs. Wolfer et al. considered that this effect may account for the dependence of swelling on Ni content in austenitic alloys containing less than 35% Ni. A generalized description of the swelling behavior of austenitic alloys, which was consistent with the model developed by Wolfer et al., was put forward by Garner40 (see also Chapter 4.02, Radiation Damage in Austenitic Steels). Garner’s ideas were largely based on the results of the EBR-II irradiation experiments and the earlier ion bombardment work of Johnston et al., both of which showed a strong dependence of swelling on nickel content. It was considered that swelling was characterized by a tran- sient period followed by a regime in which the swelling rate became constant. In neutron-irradiated alloys, the swelling rate in the posttransient regime was generally found to be �1% per dpa. In swelling- resistant alloys, however, it was argued that such high swelling rates might not be observed owing to ex- tended transient periods. The duration of the Radiation Effects in Nickel-Based Alloys 131 transient regime was shown to be dependent on alloy composition and could extend for many tens of dpa in low-swelling materials. The duration of the transient regime was implicitly linked to the completion of void nucleation but, at the time these ideas were put forward, relatively few measurements of void concentrations were available, as swelling data were mainly derived from dimensional or density changes. Factors that were proposed to account for the influence of nickel on the void nucleation rate in- cluded the effect on vacancy diffusivity described by Esmailzadeh and Kumar35; a possible correlation with the development of fine scale compositional fluctua- tions by a spinodal-like decomposition process (observed by Dodd et al.41 in ion-irradiated ternary Fe–Cr–Ni alloys); and an effect of nickel on the minimum critical radius for the formation of stable voids.42 Voids are unstable below a critical size, and will generally shrink unless stabilized by gas atoms; the minimum stable void radius is dependent on a number of factors, including temperature and defect bias, and Coghlan and Garner suggested that the compositional dependence of the vacancy diffusivity would also affect this critical size. In other words, it was considered that the transition from gas bubble to void would require a larger bubble size in high- nickel alloys, particularly at relatively high tempera- tures in the swelling regime where void nucleation becomes increasingly difficult. Hoyt and Garner43 subsequently argued that the minimum critical void radius concept might account for the minimum in swelling found at the intermediate nickel contents, provided that a compositional-dependent bias factor for dislocations was also incorporated into the model. The compositional dependence of the bias factor arises from solute segregation, which reduces the strain energy of dislocations and decreases the ratio of the bias for interstitials compared to vacancies. It is of interest that early evidence for the opera- tion of the bubble to void transition was obtained by Mazey and Nelson,44 who implanted Nimonic PE16 (STA condition) and a PE16 matrix alloy (ST condi- tion) with 1000 appm He to produce a high density of gas bubbles before subsequent irradiation with 46.5MeV Ni6+ ions. The PE16 matrix alloy used in this particular experiment was a low Si variant (0.1MeV) (see Figure 6). Alloys containing 19% and 30% Ni exhib- ited high swelling rates at higher fluences, but swelling remained low in higher nickel alloys. Similar effects were found in the ion-bombarded samples, where, for example, it was shown that there was no significant change in the void concentration in Fe–15Cr–45Ni at doses above 50 dpa in irradiations at 675 �C, yet a marked increase in swelling rate occurred above 120 dpa. Thus, contrary to earlier ideas, these investi- gations clearly demonstrated that the onset of a high swelling rate was not related to the cessation of void nucleation. It follows that the transition to a high rate of swelling must be due to an increase in the growth rate of existing voids. Muroga et al.45,46 observed that the total disloca- tion density in the irradiated Fe–15Cr–Ni alloys was only weakly dependent on nickel content. This sug- gested that at the intermediate nickel levels, where the void concentration was low, dislocations were weak sinks (for both vacancies and interstitials) relative to voids. In addition, it was observed that 20 16 12 8 4 0 0 0 10 20 30 40 Cavity diameter, d (nm) 50 60 70 80 90 100 110 120 130 4 8 12 16 0 4 8 12 16 N um b er o f c av iti es in in te rv al N c � 10 –1 9 m –3 30 dpa 60 dpa 0.2 dpa Bubbles 0 300 200 100 300 200 100 0 0 5 10 15 20 25 30 35 40 45 50 55 60 0 5 10 15 0 5 10 15 0 100 200 300 400 Cavity diameter d nm(a) (b) Bubbles Bubbles Voids Voids N um b er o f b ub b le s in in te rv al N b � 10 –1 9 m –3 N um b er o f v oi d s in in te rv al N v � 10 –1 9 m –3 30 dpa 60 dpa 0.2 dpa Figure 5 Histograms showing size distributions of bubbles/voids in (a) solution treated and aged Nimonic PE16 and (b) solution treated PE16 matrix alloy, irradiated with Ni6þ ions at 625 �C to damage levels of 30 and 60dpa following implantation with 1000appm. He (producing �0.2 dpa) at the same temperature. Reproduced from Mazey, D. J.; Nelson, R. S. J. Nucl. Mater.1979, 85–86, 671–675. 132 Radiation Effects in Nickel-Based Alloys dislocation loops persisted to higher doses at the intermediate nickel contents, indicating a lower growth rate for the loops – again implying an effect of nickel on dislocation sink strength. Based on these observations, Muroga et al. suggested that a reduced dislocation bias for interstitials at the intermedi- ate nickel contents might explain the influence of nickel on the early stages of void development. An additional factor was required to account for the eventual transition to a high swelling rate. Microche- mical data presented by Muroga et al.46 suggested that this transition was related to the depletion of nickel in the matrix owing to its enrichment at void surfaces. A complete description which incorporates all of the composition-dependent factors which affect the nucleation and growth of voids is lacking. However, there is a general consensus that the major influence of alloy composition arises through its effects on the effective vacancy diffusivity and on segregation aris- ing from the inverse Kirkendall effect. A correlation between the magnitude of void swelling and radia- tion-induced segregation was shown for Fe–Cr–Ni ternary alloys by Allen et al.48 The compositional dependence of radiation-induced segregation was determined using a model based on the earlier work of Marwick,29 which incorporates both the vacancy flux to the voids and the back-diffusion of vacancies due to the solute gradients set up by the inverse Kirkendall effect. Vacancy diffusivities for various alloy compositions were determined by the measure- ments of grain boundary segregation in proton- irradiated samples. Swelling data for ion and neutron-irradiated alloys were then compared with the expected swelling propensity defined by the ratio of the forward to back diffusion terms calculated at the appropriate irradiation temperature. The materials for which vacancy diffusivity data were determined included Fe-based alloys containing 16–24% Cr and 9–24% Ni, and Ni-based alloys containing 18% Cr and either zero or 9% Fe. This work did not specifi- cally examine 40–50%Ni alloys corresponding to the highest swelling resistance, though the results indi- cated that swelling generally decreased with increas- ing levels of nickel enrichment and chromium depletion at void surfaces. 0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5 4.0 300 350 400 450 500 550 600 650 Temperature (°C) S w el lin g (% ) DFR ~ 17 dpa DFR ~ 80 dpa DFR ~ 17 dpa DFR ~ 80 dpa 1018 1019 1020 1021 1022 300 350 400 450 500 550 600 650 Temperature (°C) V oi d c on ce nt ra tio n (m −3 ) DFR ~ 17 dpa DFR ~ 80 dpa 0 10 20 30 40 50 60 70 80 90 100 300 350 400 450 500 550 600 650 Temperature (°C) V oi d d ia m et er (n m ) Figure 7 Swelling data, void concentrations, and void diameters for Nimonic PE16 Dounreay Fast Reactor fuel pin cladding.Unpublished data fromSharpe, R.M. Void Swelling in Fast Reactor Irradiated Commercial High Nickel Alloy; DFMC/P(82)27, with permission from AEA Technology Plc. 1020 1021 0 10 20 S w el lin g (% ) Vo id d en si ty (m −3 ) 30 Immersion density Garner and Kumar47 19Ni 75Ni Fe–15Cr–XNi 510 �C 19Ni 75Ni 30Ni 35Ni 45Ni 120 2 4 6 Fluence (1026n m-2) 8 10 35Ni 45Ni 30Ni Figure 6 Fluence dependence of swelling and void density of Fe–15Cr–Ni alloys irradiated in Experimental Breeder Reactor-II at 510 �C. Swelling data obtained by immersion density measurements by Garner and Kumar47 are also shown. Reproduced fromMuroga, T.; Garner, F. A.; Ohnuki, S. J. Nucl. Mater.1991, 179–181, 546–549. Radiation Effects in Nickel-Based Alloys 133 4.04.2.3 Swelling Behavior of Neutron- Irradiated Nimonic PE16 Brown et al.49 compared the swelling behavior of STA Nimonic PE16 and two cold-worked austenitic steels (M316 and Nb-stabilized FV548) which were irradiated in DFR as fuel pin cladding. Two PE16 clad pins were examined, which were irradiated to burn-ups of 6.1% and 21.6% of heavy atoms, corresponding to peak damage levels of about 17 and 80 dpa, respectively. Void concentrations and swelling were lower in PE16 than in the austenitic steels. Swelling data, void concentrations, and void diameters for the two PE16 pins examined by Brown et al. are shown in Figure 7. Note that Brown et al.49 only showed trend lines for void concentration and void size in the less highly irradiated pin and com- pared the swelling tendencies of the two pins; the individual data points were not plotted and those shown in Figure 7 are previously unpublished data obtained by Sharpe. Brown et al. stated that the void concentration in PE16 decreased with increasing irradiation temperature but did not alter greatly with an increasing dose above �17 dpa. It should be noted, however, that swelling measurements for the higher burn-up pin were restricted to temperatures below 525 �C, so that a direct comparison of void concentrations in the two pins cannot be made at higher temperatures. Although there were fewer voids in PE16 than in the two steels, the voids 0.0 2.0 4.0 6.0 8.0 10.0 300 350 400 450 500 550 600 650 Temperature (°C) S w el lin g (% ) 0 5 10 15 Fl ue nc e (1 02 6 n m -2 , E > 0. 1 M eV )STA DAA766 OA DAA766 STA Z260D STA Z184 STA DAA766 OA DAA766 STA Z260D STA Z184 STA DAA766 OA DAA766 STA Z260D STA Z184 Fluence 1022 1021 1020 1019 1018 300 350 400 450 500 550 600 650 Temperature (°C) V oi d c on ce nt ra tio n (m –3 ) 0 10 20 30 40 50 60 70 80 90 100 300 350 400 450 500 550 600 650 Temperature (°C) V oi d d ia m et er (n m ) Figure 8 Swelling data, void concentrations, and void diameters for Nimonic PE16 samples irradiated in UK-1 rig in Experimental Breeder Reactor-II. Adapted from Boothby, R. M. J. Nucl. Mater.1996, 230, 148–157; Unpublished data for Boothby, R. M. The Microstructure of EBR-II Irradiated Nimonic PE16; AEA TRS 2002 (FPSG/P(90)23), with permission from AEA Technology Plc. 134 Radiation Effects in Nickel-Based Alloys appeared to be homogeneously distributed and to have developed during the early stages of irradiation; once nucleated, the growth rate of voids in PE16 remained low. These observations are clearly contrary to early models which suggested that low swelling rates result from incomplete void nucleation and extended tran- sient regimes. Rather, in agreement with the more recent observations of Muroga et al.,45,46 it appears that the swelling resistance of PE16 is due to a combi- nation of a comparatively low saturation void concen- tration, which is reached at a relatively low displacement dose, and a low void growth rate. There does not appear to be any evidence of an accelerated swelling rate in PE16 once void nucleation is complete. Additional data on void concentrations in neutron- irradiated PE16 are available from Cawthorne et al.,8 Sklad et al.,50 and Boothby.28 The results presented by Cawthorne et al. for PE16 fuel pin cladding irra- diated in DFR to a peak fluence of 5.6� 1026 nm�2 (�28 dpa) differ from those shown in Figure 7 in that, although void number densities are similar for irradiations at �380–520 �C, void concentrations are about an order of magnitude higher at 350 �C and 600–630 �C. Such discrepancies might arise from uncertainty and/or variability in irradiation tempera- tures. Another possibility is that void nucleation was incomplete at the higher irradiation temperatures in the lower burn-up pin examined by Brown et al. Data from Sklad et al. show an increase in void numbers in unstressed PE16 specimens irradiated in EBR-II at 500 �C from an average (for two differently heat trea- ted conditions) of about 4� 1019 to 1.2� 1020m�3 with increasing fluence from 1.2� 1026 to 4.0� 1026 nm�2 (E> 0.1MeV), that is, from �6 to 20 dpa. In this case, the void concentration and overall swelling of �0.2% at�20 dpa remain below the levels shown in Figure 7 for the DFR-irradiated pin at �17 dpa; this may reflect the effect of stress on swelling for fuel pin cladding. Void swelling data determined from Transmission electron microscope (TEM) examinations of three heats of PE16 which were irradiated in the UK-1 rig in EBR-II are shown in Figure 8, which includes previously unpublished results for the low boron (4 ppm) heat Z184 as well as data for heats DAA766 and Z260D (with 18 and 70 ppm boron, respectively) which were reported by Boothby.28 Data are shown for all three heats in the STA condition (ST 1020 �C and aged 4 h at 750 �C) and for DAA766 in the OA condition (a multistage heat treatment that included aging at 900 �C, slow cooling to 750 �C, and then aging for 16 h at that temperature, resulting in the precipitation of TiC and an overaged g0 structure). Swelling data derived from the density measure- ments of STA PE16 heat DAA 766 from the same experiment are shown in Figure 4. An example of the 200 nm Figure 9 Void structure in PE16 (OA condition) irradiated in Experimental Breeder Reactor-II to 58dpa at 513 �C. Reproduced from Boothby, R. M. J. Nucl. Mater.1996, 230, 148–157. Radiation Effects in Nickel-Based Alloys 135 void distribution in the OA condition is shown in Figure 9. Note that the voids in neutron-irradiated PE16 tend to be cuboidal and that enhanced growth of voids attached to TiC precipitates (located at the site of a prior grain boundary) has occurred. Neutron fluences and irradiation temperatures in the UK-1 experiment were similar to those for the first withdrawal of the AA-1 rig for which data is shown in Figure 2. Void concentrations for heats DAA766 and Z260D shown in Figure 8 appear to be less temperature-dependent than for the fuel pin clad- ding data shown in Figure 7. Void numbers are gener- ally lower than in the cladding at temperatures up to �550 �C, but are intermediate between the results of Brown et al.49 and Cawthorne et al.8 for irra- diations at �600 �C. Void concentrations for PE16 irradiated to fast neutron fluences (E>0.1MeV) of 9.4–12.3� 1026 nm�2 at 477–513 �C in the UK-1 ex- periment were very similar to those determined by Sklad et al.50 for 4.0� 1026 nm�2 at 500 �C. The low boron heat Z184 showed atypical behavior, with a very high concentration of small voids and low swelling at 438 �C, but high swelling owing to increased void sizes at normal void concentrations at temperatures above 513 �C. It is probable that the effect of boron on swelling is related to the formation of boron–vacancy complexes, which can give rise to the nonequilibrium segregation of boron in the presence of quenched-in thermal vacan- cies as well as to radiation-induced effects.51 Some variability in the swelling response of Nimonic PE16 in PFR (Prototype Fast Reactor) components was reported by Brown and Linekar.52 Increased swelling in PE16 subassembly and guide tube wrappers in PFR compared to expectations based on the performance of DFR pin cladding appeared to be related to temperature fluctuations, particularly at temperatures below 400 �C during the early operation of PFR. Void concentrations were reported to be higher in the PFR components, and it was suggested (by Cawthorne, unpublished data) that this may have been due to the release of vacan- cies from vacancy loops which had formed during lower temperature excursions. In fact, the void con- centration reported by Cawthorne et al.8 for DFR pin cladding irradiated at 350 �C was higher than the highest value reported for the PFR components by a factor of about 3, but this comparison was not made by Brown and Linekar. There were also indications of heat-to-heat variability and effects of the fabrication route on the swelling of PE16 wrappers in PFR. Nevertheless, swelling of PE16 wrappers, although higher than expected, remained low in absolute terms and did not give rise to any operational problems. Although PE16 was originally selected as the ref- erence wrapper material for PFR and as an alterna- tive to cold-worked M316 steel for fuel pin cladding, PE16 was favored as a cladding material with 12%Cr ferritic–martensitic steel wrappers in subsequent subassembly designs.53 The 12%Cr steel was chosen as a wrapper material because of its superior swelling resistance, but its use was limited to relatively low temperatures owing to inadequate strength at the higher operating temperatures experienced by pin cladding. Design calculations for PE16 fuel pin clad- ding made by Cole54 indicated that cladding hoop stresses, which arise from the internal pressure from the gaseous fission products released from the fuel, were much lower than the yield stress of the material and were generally expected to remain below about 70MPa. In addition, the void swelling and irradiation creep behavior of PE16 were considered to be well matched to the fuel swelling, so that fuel–clad inter- action stresses also remain low. Fuel pins with PE16 cladding successfully attained high burn-ups in PFR, with some 3500 pins exceeding dose levels of 100 dpa and 265 pins reaching maximum doses of 155 dpa.55 Very few failures of PE16 clad pins were recorded – three failures occurred in pins which had reached burn-ups over 17 at.%, with one failure at 11.3 at.% burn-up which was believed to have resulted from a fabrication defect.56 In addition to the four PE16 cladding failures in PFR, Plitz et al.57 recorded 14 failures in austenitic steel cladding, all at lower burn- ups than in PE16. The failures in PE16 cladding were 136 Radiation Effects in Nickel-Based Alloys regarded as benign and permitted continued opera- tion, with no significant loss of fuel into the primary circuit coolant. A peak burn-up of 23.2 at.%, corresponding to a peak dose in the PE16 cladding of 144 dpa, was achieved in PFR in an experimental fuel cluster. Postirradiation examinations of pins from this cluster and a high burn-up subassembly (18.9 at.%, with a peak cladding dose of 148 dpa) were carried out by Naganuma et al.58 Maximum diametral strains of less than 1% were measured, attributable to the combined effects of void swelling, creep deformation arising from internal gas pressure in the pins, and small contributions from mechanical interactions between the fuel and cladding in the lower part of the pins. 4.04.3 Irradiation Creep A detailed discussion of irradiation creep mechan- isms is beyond the scope of this chapter, which will instead concentrate on experimental data which enable comparisons to be made between nickel- based alloys and austenitic steels. However, some insight into irradiation creep mechanisms is given in Section 4.04.4.1, where the effect of stress on the evolution of dislocation structures is described. Irra- diation creep mechanisms are discussed more fully in Chapter 1.04, Effect of Radiation on Strength and Ductility of Metals and Alloys. Several reviews of irradiation creep data are available in the literature, for example, by Harries,59 Ehrlich,60 and Garner,61 and although these have tended to focus on austenitic steels, the behavior of nickel-based alloys generally appears to be similar. Different types of test specimen, including pres- surized tubes and helical springs, have been used to measure irradiation creep strains. The data are therefore generally converted to effective strain �e and effective stress �s values, using the Soderberg formalism60: �e=�s ¼ e=s ¼ g=3t ¼ 4eH=3sH where e, g, and eH are tensile, surface shear and hoop strains; and s, t, and sH are tensile, surface shear and hoop stresses, respectively. Irradiation creep experiments carried out in DFR used helical spring specimens, which were loaded in tension and periodically removed for measurements. DFR data for austenitic steels and Nimonic PE16 were reviewed by Mosedale et al.62 and Harries,59 and results for PE16 were reported in full by Lewthwaite and Mosedale.63 Average irradiation temperatures for PE16 specimens ranged from about 280 to 340 �C, with displacement doses up to a maximum of �13 dpa (N/2). For austenitic steels, the irradiation creep strain was found to be linearly dependent on the applied stress and the displacement dose, comprising transient and steady-state compo- nents as follows: g ¼ Atþ Bdt where d is the displacement dose and A and B are material-dependent creep coefficients. For PE16 in a STA condition (1 h at 1080 �C plus 16 h at 700 �C), creep at dose rates of �5� 10�7 dpa (N/2) s�1 was characterized by an initial period of low strain and an increased creep rate at higher displacement doses. Mosedale et al.62 described the g=t versus dpa creep curve for STA PE16 as parabolic, though the maxi- mum observed creep rate was similar to that in aus- tenitic steels and Harries59 represented the creep strain above a threshold dose of 8 dpa (N/2) by g ¼ 4:3� 10�6tðd � 8Þ where t is in MPa; converting to effective strain/ stress values and to NRT units of displacement dose (assuming 1 dpa (N/2)¼ 0.8 dpa (NRT-Fe)) would reduce the creep coefficient by a factor of 2.4. Data presented by Lewthwaite and Mosedale63 showed that ST PE16 behaved similarly to the STA condi- tion, though OA conditions exhibited higher creep strains due to a combination of increased creep rates and low threshold doses (around 1 dpa). An apparent dose-rate dependency was observed, with steady- state creep coefficients for STA and OA PE16 increased by factors of �2 at lower damage rates of �0.5–1.5� 10�7 dpa (N/2) s�1 and threshold doses reduced to �0.5 dpa or less. A similar effect of dose rate on the creep strain per dpa was also reported for austenitic steels.64 Steady-state creep coefficients (MPa�1 dpa�1) and creep strain rates (MPa�1 s�1) for PE16 as a function of dose rate are compared with data for cold-worked steels M316 and FV548 in Figure 10. The data plotted in Figure 10 are derived from the results of Lewthwaite andMose- dale63,64 but are converted to effective strain/stress values and NRT(Fe) dpa units to enable comparison with other published data. It is evident that the irra- diation creep behavior of STA and OA (24 h at 800 �C) PE16 is similar to that of the austenitic steels. Creep rates at higher dose rates are generally lower than would be indicated from the linear extrapolation of low dose rate data. Lewthwaite and Mosedale63 0.0 1.0 2.0 3.0 4.0 5.0 6.0 0.0 1.0 2.0 3.0 4.0 5.0 Displacement rate (10–7dpa s–1) C re ep c oe ffi ci en t, B (1 0– 6 M P a– 1 d p a– 1 ) 0.0 1.0 2.0 3.0 4.0 5.0 6.0 7.0 8.0 9.0 10.0 0.0 1.0 2.0 3.0 4.0 5.0 Displacement rate (10–7dpa s–1) C re ep s tr ai n ra te (1 0– 13 M P a– 1 s– 1 ) M316 FV548 PE16 OA PE16 STA M316 FV548 PE16 OA PE16 STA Figure 10 Steady-state creep coefficients and creep strain rates for Nimonic PE16 and austenitic steels, derived from the measurements of Lewthwaite and Mosedale. Adapted from Lewthwaite, G. W.; Mosedale, D. In Proceedings of International Conference on Irradiation Behaviour of Metallic Materials for Fast Reactor Core Components, Ajaccio, Corsica, June 4–8, 1979; Poirier, J., Dupouy, J. M., Eds.; Le Commissariat a l’Energie Atomique (CEA): Saclay, France, 1979; pp 399–405; Lewthwaite, G. W.; Mosedale, D. J. Nucl. Mater. 1980, 90, 205–215. Radiation Effects in Nickel-Based Alloys 137 considered that the measured irradiation creep rates for PE16 at low dose rates were in close agreement with the expected rates for SIPA-(stress-induced pre- ferred absorption of interstitials at dislocations) con- trolled creep. It was suggested by Mosedale et al.62 that reduced creep rates at higher dose rates might be attributable to increased recombination rates for vacancies and interstitial atoms, although a more detailed assessment of this effect by Lewthwaite and Mosedale64 proved inconclusive and a dose-rate de- pendency has not generally been observed in other experiments.60 Garner and coworkers65,66 considered that the higher creep rates measured by Lewthwaite and Mosedale at lower displacement rates were an aberration due to transient effects at low dpa levels. Nevertheless, this does not alter the finding that the irradiation creep behavior of PE16 is comparable to that of austenitic steels. Paxton et al.67 examined the in-reactor creep behavior of a number of alloys, including Nimonic PE16, Inconel 706, and Inconel 718, as well as aus- tenitic and ferritic steels, in pressurized tube experi- ments carried out in EBR-II at 540 �C to fluences up to 4� 1026 nm�2 (E>0.1MeV). Diametral strains measured in pressurized tubes (with hoop stresses in the approximate range of 25–175MPa) were cor- rected for void swelling and/or densification observed in unstressed specimens (though this does not allow for any effects of stress on swelling or precipitation processes). Precipitation-hardened alloys exhibited lower creep strains than solid solution strengthened steels, with the Inconel alloys superior to PE16 at 540 �C. The creep resistance of the precipitation- hardened materials was also dependent on heat treat- ment, with ST conditions generally superior to aged conditions. However, it was noted that ST conditions also exhibited greater densification – giving rise to the possibility of increased fuel–clad interactions in fuel elements. In-reactor creep strains were discussed in terms of a widely used model which includes a term for creep enhancement due to swelling. The total effective creep strain �e is given by �e ¼ B0ft�sþ DS�s where B0 is the creep compliance, ft is the neutron fluence, D is the creep–swelling coupling coefficient, and S is the fractional swelling. A contribution from thermal creep may be expected at 540 �C, but data to correct for this component were not available and hence the creep coefficients could not be determined precisely. The stress dependence of the measured creep strain was approximately linear in the low swelling precipitation-hardened alloys, though non- linearity attributable to the effects of stress on swelling was observed in the solid solution alloys. An approximate value of B0 of 1.5� 10�28MPa�1 (n cm�2)�1, which is equivalent to 3� 10�7MPa�1 138 Radiation Effects in Nickel-Based Alloys dpa�1, was derived by Paxton et al. for the Inconel alloys. Ehrlich60 subsequently made estimates of B0 for the other materials included in this study, which ranged from �1.4� 10�6MPa�1 dpa�1 for ST PE16 to �10�5MPa�1 dpa�1 for cold-worked 316 steel. Paxton et al. noted that values of the creep–swelling coefficient D appeared to be much larger for the solid solution strengthened steels than for the precipitation- hardened alloys, with the higher values being attrib- utable to increased thermal creep components and/or the effects of stress on swelling. Gilbert and Chin68 examined the nonisothermal creep behavior of EBR-II-irradiated PE16 and Inconel 706. Both materials were in ST conditions. Pressurized tubes, with nominal hoop stresses of 100MPa for PE16 and 200MPa for Inconel 706, were irradiated at 425, 540, and 590 �C, both isother- mally and with temperature steps. Diametral strains for isothermally irradiated PE16 increased with increasing fluence and temperature as expected. Fol- lowing temperature changes from 540 to 590 �C or 425 �C, the creep rate for PE16 adjusted to the iso- thermal rate at the new temperature. For Inconel 706, however, the isothermal creep rate was highest at 425 �C, and an upward step to 540 �C resulted in a reduced creep rate; a downward step from 540 to 425 �C gave rise to an increased creep rate that exceeded the isothermal rate at 425 �C; and an upward step from 540 to 590 �C reduced the creep rate, even though the isothermal creep rate was higher at 590 than 540 �C. The complex in-reactor creep behavior of Inconel 706 appeared to be related to the stability of the ordered body-centered tetrago- nal, Ni3Nb g00 phase and its effect on thermal creep resistance. Gilbert and Chin considered that the in- reactor deformation of Inconel 706 was primarily controlled by thermal rather than irradiation creep processes, since similar creep rates were reported to occur in thermal control tests. Microstructural examinations made by Thomas69 indicated that g00 precipitated during irradiation above �500 �C but dissolved at lower temperatures, thereby reducing the creep strength of the material. Gelles70 subse- quently reported that the dissolution of g00 at low irradiation temperatures appeared to be promoted by the application of stress since more of this phase was retained in unstressed material. Toloczko et al.5 investigated the swelling and creep behavior of five austenitic alloys which were irradiated in PFR in a pressurized tube experiment at �420 �C. The materials examined included the solid solution strengthened steels 316 and D9, and the higher-Ni precipitation-hardened alloys D21, D68, and D66. Dose rate variations were examined by positioning specimens at different axial loca- tions within the reactor core. The tubes were removed periodically for diameter measurements, with peak doses of �50 dpa being attained at the highest flux level. Hoop stresses ranged from 0 to 150MPa, and swelling as a function of dose was estimated from measurements on unstressed tubes assuming that densification effects were completed during the first irradiation cycle. There was some scatter in the results but the creep coefficient B0 was found to be relatively independent of alloy composition and dose rate, with typical values of �1.0–1.4� 10�6MPa�1 dpa�1 (though higher values were determined for type 316 steel). The creep–swelling coupling coefficient D was also independent of dose rate but appeared to be mate- rial dependent (with values in the approximate range of 0.4–1.6� 10�2MPa�1), though this varia- bility could not be associated with any particular compositional factor. Similar results for two pre- cipitation-hardened high-nickel alloys (with simi- lar compositions to Nimonic PE16, but with additions of �0.5% Nb), which were irradiated in a pressurized tube experiment in the Russian fast reactor BN-350 to �90 dpa at 400 �C, were also reported by Porollo et al.71 4.04.4 Microstructural Stability 4.04.4.1 Dislocation Structures Dislocation structures in irradiated pressurized tube samples were examined by Gelles et al.72 The materi- als which were examined included stressed and unstressed samples of ST PE16, and stressed samples of ST and STA Inconel 706. A subsequent paper by Gelles73 extended these investigations to the stressed samples of PE16 in STA and OA conditions. Further details of this work were also provided by Garner and Gelles74, and by Gelles.70 Examination of ST PE16, which was irradiated at 550 �C to 2� 1026 nm�2 (E>0.1MeV) at hoop stres- ses of 0 and 167MPa, revealed that the distribution of Frank dislocation loops was similar on all the four {111} planes in the unstressed sample but was aniso- tropic in the stressed material. In the stressed sample, the loop density on any particular {111} plane increased with increasing magnitude of the normal stress component on that plane. A near-linear rela- tionship between the loop density and the normal Radiation Effects in Nickel-Based Alloys 139 component of the deviatoric stress tensor, sDN (¼ sN � sH, where sN is the normal component of the applied stress on a particular plane and sH is the hydrostatic stress), was found for PE16. This result is in line with the SIPA loop growth model described by Garner et al.75 No such correlation was found in the similarly irradiated and stressed Inconel 706 samples, however, possibly because the low creep rate of this material at 550 �C did not allow the relaxation of internal stresses. Unfaulting of Frank dislocation loops with a/3 {111} Burgers vectors proceeds via interaction with a/6{112} Shockley partials to produce perfect a/2 {110} line dislocations. Gelles70 described how this occurs via a two-step process, with the necessary partial dislocations (two per interstitial loop) first being nucleated by an interaction of the faulted loop with a suitable perfect dislocation and then sweeping across the loop to reestablish the perfect dislocation. Gelles73 examined the distribution of Burgers vectors among the six possible a/2{110} perfect dislocation types in irradiated pressur- ized tube samples of PE16. The samples examined included the stressed ST condition irradiated at 550 �C, and STA and OA conditions which were both irradiated at 480 �C to a fluence of 8� 1026 nm�2 at a hoop stress of 331MPa. The results showed highly anisotropic distributions in the Burgers vectors of perfect dislocations in all the three heat-treated conditions, with dislocation den- sities of the various types differing by factors of up to 10–40 in each sample. The level of anisotropy pro- duced in the population of perfect dislocations was significantly greater than in the dispersion of Frank loops. This is a feasible outcome since, in principle, all loops may be unfaulted by just two variants of the six a/2{110} perfect dislocation types. In effect, the development of anisotropic dislocation structures is a response of the material to produce the strain which is required to accommodate the applied stress. Furthermore, it was found that the perfect disloca- tions in the irradiation creep samples of PE16 were primarily of edge type lying on {100} planes rather than {111} slip planes, indicating that they could only contribute to the creep strain via climb (i.e., by the SIPA mechanism) and not by processes involving dislocation glide. 4.04.4.2 Precipitate Stability Early models of precipitate stability under irradiation were based on the ideas of Nelson et al.,76 who suggested that precipitates would evolve to an equi- librium size determined by competing processes affecting their growth, via the radiation enhanced and/or thermal diffusion of solutes, and their simul- taneous dissolution due to damage arising in collision cascades. Two dissolution mechanisms were sug- gested: recoil dissolution due to the displacement of atoms from the precipitate into the matrix, and dis- ordering dissolution of ordered phases such as g0, with the latter predicted to be the more effective. The model predicted that fine precipitates would continue to grow to some equilibrium size (depen- dent on temperature, dose rate, and solute levels), but that precipitates greater than this size would shrink. Experimental evidence for the dissolution of large preexisting Ni3Al g0 precipitates in heavy-ion- irradiated Ni–Al alloys was shown by Nelson et al.76 These ideas were developed further and applied to g0 precipitates in ion and neutron-irradiated alloys by Baron et al.77 The model developed by Baron et al. indicated that, at a given particle size, a higher solute supersaturation was required under irradiation than in a purely thermal environment. The model appeared to be consistent with the observed coarsen- ing behavior of g0 precipitates during irradiation, though no evidence for the shrinkage of large particles was presented. For example, data for PE16 irradiated at fluences up to 7.5� 1026 nm�2 at 560 �C, which were reported by Chang and Baron,78 only examined the growth of g0 particles up to a maximum radius of �15 nm under conditions where the predicted maximum equilibrium radius was �35 nm. However, detailed examinations of g0 structures in neutron-irradiated Nimonic PE16 which were made by Gelles79 found no evidence to indicate that irra- diation-induced dissolution mechanisms limited the particle size. Microstructural examination of PE16, originally in ST, STA, and OA conditions, irradiated in EBR-II to�27 dpa (5.4� 1026 nm�2, E>0.1MeV) at 600 �C, revealed that preexisting g0 dispersions in aged material were maintained but continued to coarsen even in the OA condition, and that a fine dispersion formed in ST material. Coarsening of the g0 particles in the OA material was accompanied by the formation of fine background precipitates in some regions. Further in-reactor precipitation of g0 also occurred at point defect sinks, including void surfaces and dislocations, in all the heat-treated con- ditions. Additional examinations by Gelles80 of ST PE16, irradiated to �30–50 dpa at temperatures in the range of 430–650 �C, indicated that g0 coarsening was controlled by radiation-enhanced diffusion 140 Radiation Effects in Nickel-Based Alloys below 600 �C with an activation energy that (in agreement with theoretical predictions for a process governed by point defect recombination) was approx- imately a quarter of that for thermal diffusion. As described in Section 4.04.5.1 in relation to irradiation embrittlement effects, Yang81 examined an identically irradiated set of ST PE16 samples as Gelles, focusing on the precipitation of g0 at grain boundaries. Similar g0 structures to those described by Gelles and Yang were also observed by Boothby28 in the aged conditions of EBR-II-irradiated PE16, though at higher irradiation temperatures (�540 �C for the STA condition, and �600 �C for the OA condition), where doses were in the range 66–74 dpa, the spherical g0 precipitates which formed during thermal aging were almost entirely replaced by ‘skel- etal’ forms nucleated at point defect sinks. Figure 11 shows an example of the g0 distribution, imaged in dark field, in STA PE16 irradiated to 69 dpa at 570 �C; although small spherical precipitates were retained in a narrow region adjacent to the grain boundary, a much coarser dispersion is evident at the boundary itself and within the bulk of the grain. 4.04.5 Irradiation Embrittlement The effects of fast neutron irradiation on the tensile properties of several precipitation-hardened nickel- based alloys were investigated in the 1970s and 1980s. 200 nm Figure 11 Dark field, transmission electron micrograph, illustrating the distribution of g0 precipitates in solution treated and aged Nimonic PE16 irradiated in Experimental Breeder Reactor-II to 69 dpa at 570 �C. Unpublished data from Boothby, R. M. The Microstructure of EBR-II Irradiated Nimonic PE16; AEA TRS 2002 (FPSG/P(90)23), with permission from AEA Technology Plc. The materials examined included a number of g0/g00- hardened alloys, such as the Inconel alloys 706 and 718 and the developmental alloys D68 and 7818, as well as g0-hardened alloys similar to Nimonic PE16. Earlier work by Broomfield et al.82 on thermal reactor irradiated materials indicated that PE16 was more susceptible to irradiation embrittlement at elevated test temperatures than austenitic steels. Broomfield83 found that thermal neutron irradiated PE16 was most severely embrittled in low strain tests at �550–650 �C, and attributed this to an increased tendency for intergranular failure arising from the effects of helium generated from the 10B(n,a)7Li reaction. Boron itself is considered to have a benefi- cial effect on (unirradiated) creep rupture life, as it segregates to grain boundaries and inhibits intergran- ular cracking, and additions of a few 10s of ppm are therefore, generally made to nickel-based alloys, including PE16.84 Nickel is also a major source of helium in neutron-irradiated alloys, with the two- stage 58Ni(n,g)59Ni(n,a)56Fe reaction becoming the dominant source at high thermal neutron fluences, and nickel threshold reactions accounting for the greater part of helium production in fast neutron spectra.85 For example, the rate of helium generation in fast reactor irradiated PE16 was estimated by Boothby28 to be �1.2 appm per dpa, with about 85% of the helium being generated from nickel threshold reactions (see also Chapter 1.06, The Effects of Helium in Irradiated Structural Alloys). Nevertheless, other factors, including irradiation- induced strengthening and grain boundary segrega- tion and precipitation effects, have been implicated in the embrittlement of fast neutron irradiated nickel-based alloys. 4.04.5.1 Fast Neutron Irradiation Experiments Rowcliffe and Horak86 investigated the tensile prop- erties of Inconel 706 (in a multistep ‘fully aged’ condition) and Inconel 718 (ST condition) following irradiation in EBR-II to fluences of 4–5� 1026 nm�2 (E> 0.1MeV). Irradiation temperatures (Ti) ranged from 450 to 735 �C, with tensile tests being per- formed at a strain rate of 4� 10�4 s�1 at temperatures corresponding to Ti and to Tiþ 110 �C. Yield stresses and total elongation data for Inconel 706 are shown in Figure 12 and for Inconel 718 in Figure 13. Data for Inconel 706 showed very high (>1000MPa) yield stresses and ultimate tensile strengths (UTS) in Radiation Effects in Nickel-Based Alloys 141 specimens irradiated at temperatures up to and including 500 �C. This high tensile strength was maintained in a specimen irradiated at 500 �C but tested at 610 �C. Although there was some reduction in strength in specimens irradiated at 560 �C and above, the UTS remained above 650MPa in speci- mens irradiated at 625 �C. The very high tensile 0 200 400 600 800 1000 1200 1400 400 450 500 550 600 Temperat Y ie ld s tr es s (M P a) Figure 12 Yield stress and total elongation values at the irrad Breeder Reactor-II-irradiated Inconel 706. Based on data from R 38, 266–267. Temperat Y ie ld s tr es s (M P a) 0 200 400 600 800 1000 1200 1400 400 450 500 550 600 Yield at Ti Elong. at Ti Y E Figure 13 Yield stress and total elongation values at the irrad Breeder Reactor-II-irradiated Inconel 718. Based on data from R 38, 266–267. strengths exhibited at the lower irradiation tempera- tures were attributed to the instability of the (ordered body-centered tetragonal) g00 phase below 525 �C and its consequent dissolution, leading to the reprecipita- tion of nickel and niobium as (ordered face-centered cubic) g0 on dislocation loops. At higher irradiation temperatures, both g0 and g00 were stable, but 650 700 750 800 ure (°C) 0 2 4 6 8 10 12 14 16 18 20 To ta l e lo ng at io n (% ) Yield at Ti Yield at Ti+ 110 °C Elong. at Ti Elong. at Ti+ 110 °C iation temperature (Ti) and at Ti þ 110 �C for Experimental owcliffe, A. F.; Horak, J. A. Am. Nucl. Soc. Trans. 1981, ure (°C) To ta l e lo ng at io n (% ) 650 700 750 800 0 2 4 6 8 10 12 14 16 18 20 ield at Ti+ 110 °C long. at Ti+ 110 °C iation temperature (Ti) and at Ti þ 110 �C for Experimental owcliffe, A. F.; Horak, J. A. Am. Nucl. Soc. Trans.1981, 142 Radiation Effects in Nickel-Based Alloys precipitate coarsening resulted in lower tensile strength. Elongations to failure for tests carried out at the irradiation temperature were between 1.5% and 3% up to 625 �C, compared to >8% in unirradi- ated material. Irradiation embrittlement was gener- ally more severe in tests at Tiþ 110 �C, particularly at 610–735 �C where the lowest recorded ductility was 0.2%. Fractures in irradiated Inconel 706 were predominantly intergranular, with failure believed to be facilitated by the decohesion of Z phase (hexago- nal Ni3(Ti,Nb)) platelets which were formed at grain boundaries during the initial heat treatment. Rowcliffe and Horak’s data for ST Inconel 718 showed similar trends to Inconel 706. Precipitation of the g0 and g00 phases occurred during the irradia- tion of Inconel 718, resulting in yield strengths in excess of 1000MPa at irradiation temperatures up to 560 �C and above 800MPa at 625 �C. The ductility of Inconel 718 was reduced from more than 30% in the unirradiated condition to 0.2% or less in specimens which were irradiated at 500–560 �C and tested at Tiþ 110 �C. In contrast to Inconel 706, failures in irradiated Inconel 718 were reported to be predomi- nantly transgranular. Crack propagation in Inconel 718 appeared to have been via a ‘channel’ fracture mechanism, that is, with deformation occurring by highly localized planar slip and consequent linkage of radiation-induced voids. Temperat Y ie ld s tr es s (M P a) 0 200 400 600 800 1000 400 450 500 550 600 Figure 14 Yield stress and total elongation values at the irrad Breeder Reactor-II-irradiated Nimonic PE16. Based on data from of Radiation on Materials: 10th Conference; Kramer, D., Brager, and Materials: Philadelphia, PA, 1981; pp 326–351, ASTM STP 7 of Radiation on Materials, copyright ASTM International, 100 Ba Bajaj et al.87 examined the tensile properties of Nimonic PE16 irradiated in EBR-II to neutron fluences up to a maximum of �7� 1026 nm�2 (E> 0.1MeV), at temperatures in the range of 450– 735 �C. The alloy was in a STA (1 h at 900 �C plus 8 h at 750 �C) condition, and appears to have been the same low-Si heat of PE16 that was subsequently used in the AA-1 swelling experiment described by Garner and Gelles.22 Tensile tests were carried out at 232 �C (to simulate refueling conditions), at the irradiation temperature Ti and at Tiþ 110 �C (to simulate reac- tor transients), at a strain rate of 4� 10�4 s�1, and with a small number of tests at 4� 10�3 s�1. Irra- diated specimens tested at 232 �C generally showed a substantial increase in yield stress and a small increase in UTS over the unirradiated values (although samples irradiated at the highest tempera- ture of 735 �C exhibited some softening), and retained good levels of ductility with total elonga- tion values above 10%. Yield stress and total elon- gation data for PE16 at higher test temperatures are shown in Figure 14 for specimens irradiated to a fast neutron fluence of 4.3� 1026 nm�2 (enabling direct comparison with the data for the similarly irradiated Inconel alloys shown in Figures 12 and 13). Speci- mens tested at the irradiation temperature again showed strengthening at temperatures in the range of 450–625 �C and softening at 735 �C, with good ure (°C) To ta l e lo ng at io n (% ) 650 700 750 800 0 2 4 6 8 10 12 14 16 18 20 Yield at Ti Yield at Ti+ 110 °C Elong. at Ti Elong. at Ti+ 110 °C iation temperature (Ti) and at Ti þ 110 �C for Experimental Bajaj, R.; Shogan, R. P.; DeFlitch, C.; et al. In Effects H. S., Perrrin, J. S., Eds.; American Society for Testing 25. Reprinted, with permission, from ASTM STP725-Effects rr Harbor Drive, West Conshohocken, PA 19428. Radiation Effects in Nickel-Based Alloys 143 ductility at 450 �C but total elongations reduced to �3% at 560–625 �C. Tests at Tiþ 110 �C showed further increases in tensile strength (consistentwith the greater hardening expected from irradiation at a lower temperature) and more severe embrittlement with ductility levels at 670–735 �C reduced to 0.3% at a fluence of 4.3� 1026 nm�2 and to zero (i.e., failure before yield) in higher dose samples (7.1� 1026 nm�2). Tests at Ti at the higher strain rate resulted in an improvement in ductility by a factor of two or three. Examination of fracture surfaces showed that failures were predominantly intergranular in irradiated sam- ples tested above �550 �C, transgranular at 232 �C, and mixed mode at 450–550 �C. Bajaj et al. considered that the irradiation embrittlement of PE16 evident at high temperatures could simply be explained bymatrix hardening with little or no change in the grain bound- ary fracture strength – evidenced by increases in yield strength but no significant changes in true (as opposed to engineering) UTS values – so that mechanisms relying on the weakening of grain boundaries could be discounted for the test conditions studied. Sklad et al.50 reported tensile data for two aged conditions of Nimonic PE16 which were irradiated in EBR-II to 1.2� 1026 nm�2 (E>0.1MeV) at 500 �C and tested at strain rates from �3� 10�5 to 3� 10�3 s�1. There was no significant difference in the postirradiation properties of the two differently aged conditions, although one aging treatment (2 h at 800 �C plus 16 h at 700 �C) resulted in an unirradiated yield stress�25%higher than the other condition (1 h at 900 �C plus 8 h at 750 �C). No effect of strain rate on tensile properties was evident in tests at the irradiation temperature, where total elongations remained above 10%. Tests at higher temperatures were made only at the lowest strain rate, with failure elongations being reduced to 1.6% at 600 �C and 0.5% at 700 �C. The lowductility failureswere associatedwith an increased tendency toward intergranular fracture, and additional tests, in which samples irradiated to 4� 1026 nm�2 at 500 �Cwere fractured in situ in anAuger spectrometer, revealed helium release from samples which fractured intergranularly as well as the segregation of Ni, P, and S to grain boundaries. Helium release was estimated at�0.03 He atoms per grain boundary atom. No grain boundary helium bubbles were observable by TEM, and it was therefore considered that helium either remained in solution as a partial monolayer or was present in unresolved bubbles less than 1–2 nm in diameter. The presence of grain boundary helium bubbles in Nimonic PE16 was reported by Fisher et al.88 in sections of AGR (advanced gas-cooled reactor) tie bars irradiated at 512 �C and above. AGR tie bars, which are approximately 10m long and are under load only during charging and discharging of the fuel element stringers, operate at temperatures from 325 to 650 �C from bottom to top, with peak doses of �3 dpa occurring at around the 4m position. Stress-rupture testing at 600 �C at an applied stress of 500MPa showed a trough in properties (i.e., a minimum in failure times) and intergranular failures in sections of some tie bars which were irradiated at temperatures in the range of 350–400 �C where grain boundary helium bubbles were not generally observed. Even so, grain boundary cavitation was observed in a fractured tie bar section which was irradiated at 360 �C, with the cavities appearing to be nucleated (possibly at submicroscopic helium bubbles) at the intersection of slip bands with the boundary. The trough in stress-rupture properties occurred in tie bar sections which exhibited both high yield strengths (attributable to high concentra- tions of dislocation loops and small voids) and high levels of grain boundary segregation. EDX (energy dispersive X-ray) analyses showed a significant enrichment of Ni and Si, and a depletion of Fe, Cr, and Mo, at the grain boundaries of sections irradiated at 335–585 �C. In addition, high levels of Si were detected in sections irradiated at 335–512 �C in the g0 phase that precipitated at the surface of voids, with the Si content increasing with decreasing irradiation temperature. Although the presence of Si-enriched g0 phase at grain boundaries could not be confirmed, it was suggested that its formation may have contribu- ted to the minimum in stress-rupture life, which was thought to result from the weakening of the bound- aries relative to the matrix. Grain boundary helium bubbleswere also observed by Boothby and Harries89 and Boothby28 in PE16 irradiated in DFR and EBR-II at 535 �C and above. Tensile testing of DFR-irradiated PE16, exposed to �20 dpa at 465–635 �C, and strained at a rate of 2.5� 10�6 s�1 at temperatures approximating those of irradiation, revealed severe embrittlement with mini- mum elongations of�0.2% at 550 �C; TEM examina- tion of strained specimens provided evidence of intergranular cavitation, and the ductility data were interpreted using a model for the diffusion-induced growth of cavities nucleated at grain boundary helium bubbles.89 The postirradiation tensile properties and micro- structure of developmental g0 (D21, D25, and D66) and g0/g00 (D68) strengthened alloys were discussed 144 Radiation Effects in Nickel-Based Alloys by Yang et al.4 The alloys were all irradiated in a ST condition; additionally, D25 was irradiated in an aged (24 h at 700 �C) condition (STA), and D66 in a 30% cold-worked plus aged (11 h at 800 �C plus 2 h at 700 �C) condition (CWA). Specimens were irradiated at 450–735 �C to a fast neutron fluence of 4� 1026 nm�2 (E> 0.1MeV) in EBR-II, and were tested atTi,Tiþ 110 �C and 232 �C. Severe irradiation embrittlement was evident in the ST alloys and STA D25, particularly in tests atTiþ 110 �C. Zero ductility was recorded in the lower-Ni alloy D21 (25Ni–8Cr) irradiated and tested at 550 and 600 �C. Severe ductil- ity losses were associated with intergranular failures, which were attributed to irradiation-induced solute segregation and consequent precipitation of brittle g0 layers at grain boundaries. However, reasonable levels of ductility, ranging from 2 to 6%, coupled with trans- granular failures, were obtained at all temperatures in irradiated CWA D66 (45Ni–12Cr). The preirradiation grain boundary structure of this material, comprising a ‘necklace’ of small recrystallized subgrains plus large g0 particles and discrete Laves particles, remained stable with no indication of irradiation-induced g0 layers. Yang et al. considered that the radiation-induced segregation of g0 forming solutes to grain boundaries was inhibited by the introduction of a high density of dislocation sinks by cold working. Vaidyanathan et al.90 and Huang and Fish91 examined the embrittlement of EBR-II-irradiated, precipitation-hardened alloys, using ring ductility tests. In this test, small sections of tubing are com- pressed and the ductility, defined as the strain at the initiation of cracking, is deduced from the change in the sample radius of curvature at maximum load. Both experiments included Inconel 706 and Nimonic PE16 in ST conditions, while Vaidyanathan et al. also examined the developmental alloys D25 and D68 in ST and STA conditions. Peak fluences in these experiments were around 6–7� 1026 nm�2 (E> 0.1MeV) and irradiation temperatures were in the range 460–616 �C. All the materials exhibited low ductility failures at high test temperatures, particu- larly in tests at about Tiþ 110 �C where ductilities were generally below 0.1%, though Vaidyanathan et al. found that postirradiation heat treatments (typi- cally of 4 h at 785 �C) produced a moderate recovery in ductility. Based largely on observations reported by Yang81 for irradiated ST PE16, Vaidyanathan et al. and Huang and Fish considered that the irradiation- induced embrittlement of precipitation-hardened alloys could generally be attributed to the formation of brittle g0 layers at grain boundaries. However, the arguments presented were far from conclusive – microstructural examinations of the developmental alloys which were reported by Vaidyanathan et al. showed only weak indications of g0 precipitation in D25 even within the grains, and evidence for g0 pre- cipitation at grain boundaries in D68 was not found in the low ductility tested samples but only in material irradiated to a higher fluence. Yang81 examined the microstructure of a low Si (0.01%) heat of ST PE16, which was irradiated in EBR-II to doses of about 30 and 50 dpa at temperatures from 425 to 650 �C. Grain boundary g0 layers were observed in ST PE16 samples which were irradiated at 510 �C or above but not at 425 �C, and helium bubbles were detected at bound- aries in samples irradiated at 600–650 �C. It was considered by Yang that the irradiation-induced embrittlement of ST PE16 was mainly attributable to the cleavage fracture of grain boundary g0 layers and that any effects of helium were of secondary impor- tance. However, although grain boundary precipitation of g0 was observed by Boothby28 in PE16 irradiated to relatively high doses in EBR-II, there was no evidence for the formation of intergranular g0 layers in the aged conditions of PE16 which exhibited low ductility fail- ures following irradiation in DFR to �20 dpa.89 Thus, although it remains possible that the formation of grain boundary g0 layers may aggravate the embrittlement, it was considered by Boothby28 that the irradiation embrittlement of PE16 is primarily due to helium. A breach in solution-annealed Inconel 706 fuel pin cladding, irradiated to 5% burn-up in EBR-II, was reported by Yang and Makenas.92 The rupture occurred from 12.7 to 18.4 cm from the bottom of the pin, corresponding to irradiation at 447–526 �C at a fluence of 6� 1026 nm�2 (E> 0.1MeV). Microstruc- tural examinations revealed a brittle intergranular fracture, with failure being attributed to a combina- tion of matrix hardening due to g0 precipitation and grain boundary weakening due to the formation of interconnected Ni3(Ti,Nb) Z phase particles. In con- trast to the work of Rowcliffe and Horak86 where grain boundary Z phase was precipitated during a preirradiation aging treatment, this phase formed during the irradiation period in the solution- annealed cladding. Precipitation of Z was considered to be irradiation enhanced because it was not formed in long-term thermal annealing experiments at 480– 540 �C. Grain boundary precipitation of Z phase was also observed at the hot (650 �C) end of the fuel pin cladding, with both g0 and g00 in the matrix. Cauvin et al.93 and Le Naour et al.94 also attributed irradiation embrittlement effects in Inconel 706 Radiation Effects in Nickel-Based Alloys 145 cladding to the combined effects of matrix hardening and the precipitation of Z at grain boundaries. Inconel 706 fuel pin cladding, fabricated from four heats with Nb contents varying from 1 to 3% and in two heat-treated conditions (solution annealed or aged), was irradiated in the Phenix fast reactor up to a maximum of 100 dpa. Tensile tests on cladd- ing sections were carried out at a strain rate of 3� 10�4 s�1. Tensile tests performed at ambient tem- perature showed high UTS (>1000MPa) along the full length of the pins with peak values of�1500MPa in sections irradiated near 500 �C; ductility values (uniform elongations only were given) remained low (1MeV) of 1.7� 1024 nm�2 and a thermal flu- ence of 5.9� 1024 nm�2. The helium content of the reactor-irradiated specimens was estimated to be �45 appm, produced mainly from the thermal neutron reaction with 10B. Tensile tests were carried out at the implantation/irradiation temperature at a strain rate of �5� 10�4 s�1. The results showed sim- ilar trends in helium-implanted and neutron- irradiated specimens, with the total elongation values tending to decrease with increasing tensile strength. Variations in tensile strength for each alloy were largely attributable to variations in the initial heat treatment and working schedules. However, there were some indications of softening and reduced ductility in the neutron-irradiated specimens com- pared to those injected with helium. Overall, the g0- hardened alloy 7817 exhibited relatively high tensile strength (typically >700MPa) but low ductility fol- lowing helium implantation or neutron irradiation (with total elongation values generally 146 Radiation Effects in Nickel-Based Alloys Some additional, previously unpublished, data from a helium injection experiment conducted by Boothby and Cattle are given in Table 2. In this experiment, helium was implanted at ambient tem- perature to levels of 2, 10, and 50 appm into Nimonic PE16 that had been given a two-stage aging treat- ment (ST 1050 �C, aged 4 h at 800 �C plus 16 h 750 �C). As before, helium-doped samples were either retained in the as-implanted condition or given an additional aging treatment to coarsen the dispersion of gas bubbles. Tensile tests were carried out at 650 �C at strain rates of 3� 10�5 and 3� 10�6 s�1. The results show a significant loss of ductility even at 2 appm helium. In tests carried out at the higher strain rate, the total elongation values decreased progressively with increasing helium con- tent and were further reduced by postimplant aging. The ductility of as-implanted samples was generally lower but less sensitive to helium concentration in tests at a strain rate of 3� 10�6 than at 3� 10�5 s�1. However, there was little effect of the strain rate on the ductility of the postimplant aged samples. Figure 15 illustrates grain boundary structures in a tensile-tested PE16 sample which had been aged subsequent to helium injection. Failure in this case appeared to occur by the growth and coalescence of cavities which were nucleated at grain boundary gas bubbles.97 The nucleation of unstable cavities at grain boundary helium bubbles requires the application of a critical tensile stress, which is an inverse function of the bubble radius, normal to the boundary. Cavity growth then occurs via the stress-induced absorption of vacancies. In the as-implanted condition, however, Table 2 Tensile properties of helium-implanted Nimonic PE Strain rate (s�1) Helium (appm) 0.2% PS (MPa) UTS 3� 10�5 0 487 575 2 495 603 10 432 538 50 505 569 2a 445 500 10a 443 499 50a 403 483 3� 10�6 0 434 491 2 473 491 10 430 479 50 404 458 2a 425 438 10a 404 431 50a 430 439 aPostimplant aged 72 h at 750 �C. Unpublished data from Boothby, R. M.; Cattle, G. C. Development of g FPSG/P(91)9, with permission from AEA Technology PLC. the helium dispersion was too fine to enable grain boundary cavities to nucleate during tensile testing. Reduced ductility in the as-implanted samples appeared to be associated with grain boundary wedge cracking, where, as discussed by van der Schaaf and Marshall98 in relation to helium embrit- tlement of type 316 steel, the role of helium may be simply to decrease the effective surface energy for fracture. Although it is evident from simulation experi- ments that helium alone can largely account for irra- diation embrittlement, it is more difficult to assess the significance of other radiation-induced effects such as matrix hardening and grain boundary segregation and/or precipitation. One experimentwhich examined the effect of the radiation-induced precipitation of the Ni3Si g0 phase on the ductility of a binary Ni-8 at.% Si alloy was described by Packan et al.99 In this experi- ment, thin foil tensile specimens were bombardedwith either protons or a-particles to damage levels of 0.1– 0.3 dpa at 750K; irradiationwitha-particles resulted in the introduction of high helium concentrations of about 750 appm per 0.1 dpa. Proton and a-particle irradiations both resulted in the formation of g0 layers about 20–30 nm thick at the grain boundaries, but the material remained relatively ductile, exhibiting trans- granular failures, in tensile tests carried out at a strain rate of�3� 10�4 s�1 at room temperature and, for the proton-irradiated case only, 720K. Unfortunately, no tests were carried out at higher temperatures and sam- ples which were irradiated with a-particles only at 750K were not tested except at room temperature. However, low ductility intergranular failure was 16 at 650 �C (MPa) Uniform elongation (%) Total elongation (%) 3.9 35.4 5.2 19.3 4.9 11.7 3.7 7.1 4.2 5.4 2.8 4.4 2.2 2.4 2.3 31.8 0.8 6.8 1.2 7.2 1.1 6.0 0.6 5.8 2.3 4.3 0.4 1.3 0-hardened 25Ni-based Alloys for Fast Reactor Core Applications; 200 nm100 nm Figure 15 Transmission electron micrographs illustrating (left) bubble dispersion on a grain boundary parallel to the tensile axis and (right) cavitation on a boundary approximately normal to the tensile axis in Nimonic PE16 (implanted with 10 appm helium at ambient temperature, and subsequently aged for 60 h at 750 �C prior to tensile testing at 650 �C). Reproduced from Boothby, R. M. J. Nucl. Mater. 1990, 171, 215–222. Radiation Effects in Nickel-Based Alloys 147 induced in a test carried out at 720K in a samplewhich was preimplanted with 1000 appm He at 970K, then irradiated to 0.3 dpa, introducing an additional 2300 appm He at 750K. Preimplantation of helium at 970K produced grain boundary bubbles which were 10–20 nm in diameter, compared to 1.5–2 nm in mate- rial that was only irradiated with a-particles at 750K. The results of this experiment therefore indicated that the radiation-induced precipitation of g0 at grain boundaries did not give rise to embrittlement unless helium was also implanted into the specimens. 4.04.6 Concluding Remarks The effects of irradiation on the microstructure and mechanical properties of nickel-based alloys are com- plex and, although the main factors affecting their behavior have been identified, a full understanding of radiation-induced effects remains elusive. This is par- ticularly true of the precipitation-hardened alloys, typified by Nimonic PE16 and Inconel 706, where the role of the hardening phases – which confer high ther- mal creep strength, but are redistributed during irradi- ation andmay possibly influence swelling behavior and contribute to intergranular embrittlement – is unclear. The radiation-induced effects considered in this chapter – void swelling, irradiation creep, the evolution of precipitate anddislocation structures, and irradiation embrittlement – are interrelated in several ways, but particularly through the effects of point defect fluxes and the consequent redistribution of solute atoms. The beneficial effect of nickel on the swelling resistance of austenitic alloys is well known, but a clear explanation for the minimum in swelling that is found in alloys containing about 40–45% Ni has not been forthcoming. There is general agreement that the major influence of alloy composition on swelling arises through its effects on the effective vacancy diffusivity and on segregation via the inverse Kirkendall effect. However, on what appears to be the mistaken assumption that the swelling resis- tance of nickel-based alloys derives from an extended void nucleation period, swelling models have largely focused on factors affecting the nucleation rather than the growth of voids. Data for neutron-irradiatedNimo- nic PE16, for example, indicate that its swelling resis- tance is due to a combination of a comparatively low saturation void concentration, which is reached at a relatively low displacement dose, and a low void growth rate. Theminimum critical void radius concept appears to offer the most plausible explanation for the minimum in swelling found at intermediate nickel contents, although experimental data comparing the behavior of PE16 and a nonprecipitation hardenable alloy with a similar matrix composition indicate that, in addition to the Ni content of the alloy, the presence of Si and/or the g0 forming solutes Al plus Ti may also be important. The dependence of the void growth rate on Ni may be related to the effects of radiation-induced segregation on the bias terms for the absorption of point defects at sinks, though again there is evidence that minor solutes, including Si, B, and Mo, as well as the g0-forming elements, have a beneficial effect on the overall swelling behavior of nickel-based alloys. The irradiation creep behavior of nickel-based alloys generally appears to be similar to that of austenitic steels, though the higher thermal creep 148 Radiation Effects in Nickel-Based Alloys strength of precipitation-hardened alloys is an advantage at higher operating temperatures. The main drawback of precipitation-hardened nickel- based alloys for reactor core applications is per- ceived to be a high susceptibility to irradiation embrittlement. Although it has been suggested that a combination of matrix hardening and of grain boundary weakening due to the formation of brittle intergranular layers of g0 (e.g., in the case of Nimonic PE16) or Z phase (in the case of Inconel 706) is responsible for the irradiation embrittlement of these alloys, there is strong evidence, at least for g0- hardened materials, that helium is the primary cause of low ductility failures. Experimental data have shown that the implantation or generation of rela- tively small amounts of helium can give rise to low tensile ductility, with intergranular failures initiated by either the growth and linkage of cavities or by wedge cracking depending on test conditions and helium distribution, under conditions where grain boundary g0 layers are not formed. However, irre- spective of the details of the embrittlement mecha- nism, it is evident that this aspect of radiation damage does not preclude the in-core application of nickel-based alloys, as has clearly been demon- strated by the successful use of PE16 for fuel element cladding irradiated to high burn-ups in PFR. The long-term integrity of PE16 cladding is attributable to the dimensional stability of the alloy, arising from a combination of good swelling resistance and high creep strength, and relatively low operating stresses which allay irradiation embrittlement concerns. 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Hyde National Nuclear Laboratory, Harwell, Oxfordshire, UK � 2012 Elsevier Ltd. All rights reserved. 4.05.1 Introduction 152 4.05.2 Pressure Vessel Steels 154 4.05.3 Effect of Neutron Irradiation on Bulk Properties 155 4.05.3.1 Summary 159 4.05.4 Development of Mechanistic Insight of Factors Controlling Current Plant Lifetimes 160 4.05.4.1 Introduction 160 4.05.4.2 Radiation Damage Mechanisms 160 4.05.4.3 Mechanistic Framework 161 4.05.4.4 Cluster Development Under Irradiation: Cu 162 4.05.4.4.1 Introduction 162 4.05.4.4.2 CEC characterization 162 4.05.4.4.3 Development with increasing fluence 163 4.05.4.5 Development of Matrix Defects 166 4.05.4.5.1 Introduction 166 4.05.4.5.2 Nature 166 4.05.4.5.3 Development with flux and fluence and irradiation temperature 167 4.05.4.5.4 Effect of alloying 168 4.05.4.5.5 MD and hardening 168 4.05.4.6 Effect of Radiation Damage on Hardening 169 4.05.4.7 Segregation to Grain Boundaries 169 4.05.4.8 Summary 170 4.05.5 Development of Mechanistically Based DDRs 170 4.05.5.1 Introduction 170 4.05.5.2 DDRs for CMn Steels 171 4.05.5.3 US Mechanistically Guided DDRs 172 4.05.5.4 Japanese Embrittlement Correlations 174 4.05.5.5 Summary 175 4.05.6 Current Issues in the Development of DDRs 176 References 177 Abbreviations 3DAP Three-dimensional atom probe AES Auger electron spectroscopy ASME American Society of Mechanical Engineers ASTM American Society for Testing and Materials BWR Boiling water reactor CDB Coherent Doppler broadening CEC Cu-enriched cluster CRP Cu-rich cluster Cumatrix Level of Cu in matrix DBTT Ductile-to-brittle transition temperature DDR Dose–damage relationships DIDO DIDO was a MTR in the UK dpa Displacements per atom DS Change in PALA ‘S’ parameter DHv Change in Vickers hardness EFTEM Energy filtered transmission electron microscopy EOL End-of-life EONY Eason, Odette, Nanstad, Yamamoto 151 152 Radiation Damage of Reactor Pressure Vessel Steels EPRI Electric Power Research Institute (EPRI) EWO Eason Wright and Odette FEGSTEM/ EDX Field emission gun scanning transmission electron microscopy/ energy dispersive X-ray FIM Fragilisation par Irradiation Moyenne FIS Fragilisation par Irradiation Superieure HERALD HERALD was a MTR in the UK HSSI Heavy Section Steel Irradiation IAEA International Atomic Energy Authority IS Interstitial solute IVAR Irradiation Variable Program database on irradiation-induced changes in reactor Pressure Vessel Steels JEAC Japan Electric Association Code KMC Kinetic Monte Carlo (Models) KTA German Nuclear Standards Commission LEAP Local electrode atom probe LWR Light water reactor MD Matrix damage MMA Manual metal-arc welds MS Molecular statics MTR Materials Test Reactor ORNL Oak Ridge National Laboratory OSIRIS OSIRIS was a MTR in France PA Positron annihilation PALA Positron annihilation lineshape analysis PAS-OEMS Positron annihilation spectroscopy- orbital angular momentum distribution PA-t Positron Annihilation Lifetime PISA Phosphorus in Steel Ageing 5th Framework Programme of the EEC PLUTO PLUTO was a MTR in the UK PWHT Postweld heat treatment PWR Pressurized water reactor f Flux or dose rate ft Fluence or dose RPV Reactor pressure vessel SANS Small-angle neutron scattering SAW Submerged-arc weldment SIA Self-interstitial atom SMA Submerged-arc welds SMD Stable matrix defect SOL Start-of-life SRM Standard Reference Monitor TAGSI Technical Advisory Group on Structural Integrity TEM Transmission electron microscope UCRR Union Carbide Research Reactor UMD Unstable matrix defect USE Upper shelf energy USNRC US Nuclear Regulatory Commission VVER From the Russian. Translates as Water-Water Energetic Reactor 4.05.1 Introduction The ferritic steel reactor pressure vessel (RPV) of a light water reactor (LWR) is unique in terms of nuclear plant safety. This is because the RPV is a pressure boundary component whose catastrophic failure by brittle fracture could lead to severe core damage and, potentially, to the widespread release of radioactivity. In addition to the thermal, mechanical, and chemical degradation processes common to all primary circuit components, the ferritic steel RPV, because of its close proximity to the reactor core, also undergoes changes in mechanical properties because of radiation damage from the flux of fast neutrons arising from nuclear reactions in the fuel. The integ- rity of the RPV has to be assured throughout the life of the plant. Irradiation results in hardening and embrittlement processes, the most important effect of which is the rise in the ductile–brittle transition temperature (DBTT) and the decrease in the frac- ture toughness of the RPV. This can be an important issue in the region of the beltline which experiences the highest neutron fluence. These changes in mechanical properties make the RPV more suscepti- ble to brittle fracture as the plant ages. The primary purpose of this chapter is to demonstrate our current understanding of such radiation damage effects in ferritic pressure vessel steels. Low alloy ferritic steel pressure vessels are employed in all western LWRs (both boiling water reactors (BWRs) and pressurized water reactors (PWRs)), and also in VVER (LWR) reactors in Russia and a number of European countries (see, e.g., Steele and Sterne1). In the past, ferritic pressure vessel steels were also employed in gas-cooled Magnox reactors in the United Kingdom2 (Magnox reactors with steel vessels are now undergoing decommissioning3). Radiation Damage of Reactor Pressure Vessel Steels 153 Demonstrating the safe operation of such a plant has led to extensive international research over the last 40–50 years on the aging effects in ferritic steels. The need for such scientific understanding has been raised at the highest level. For example, Sir Alan Cottrell, then UK Government Chief Scientist, in a memorandum concerning the integrity of LWR pressure vessels, dated 22 January 1974, to the UK Parliament Select Committee on Science and Indus- try stated ‘‘The possible gradual growth of small cracks in highly stressed regions, by ageing and corro- sion effects during service needs further scientific investigation.’’4 The discussion here establishes that in the field of radiation damage, it is the direct application of fun- damental research to operating reactors that is signif- icant. This chapter demonstrates that developments in the understanding of the damage mechanisms have enabled an improved description of the in-service properties of RPVs of operating reactors. Indepen- dent peer review has been central to the process and particularly important from a regulatory perspective. Frequently, it is not simply the research community directly involved that has to assess any improved description of the degradation processes; for example, the safety authorities within the utility, operating the reactor (or fleet), or the national regulator will become involved in the process. Most importantly it has not always been possible to predict the end of life (EOL) vessel properties from the data obtained from materials irradiated as part of vessel surveillance programs. The vessel sur- veillance programs for commercial nuclear reactors are intended to monitor the irradiation-induced changes in mechanical properties of life-limiting structural materials subjected to significant neutron fluence. Thus, they are designed to provide advance information concerning the state of degradation in the mechanical properties of key structural compo- nents. However, because of the inevitable differences in neutron dose rate between the vessel wall and such surveillance samples (typically a factor of 5 or more), the scarcity of such data at the start of plant opera- tion, and complex and unexpected embrittlement dependencies on steel composition, it has become necessary to develop dose–damage relationships (DDRs) on the basis of mechanistic understanding that predict the embrittlement dependence on mate- rial and irradiation variables. The vast majority of investigations on the aging effects in ferritic steels over the last 40–50 years have been concerned with the effects of neutron irradiation over quite a narrow set of irradiation conditions, for example, irradiation temperatures of �250–300 �C (although there was interest in irradiation tempera- tures as low as 160 �C) and irradiation doses typical of in-service exposures ( 154 Radiation Damage of Reactor Pressure Vessel Steels 4.05.2 Pressure Vessel Steels In reviewing radiation damage effects in ferritic steels, it is important to recognize that a range of ferritic steels have been employed in commercial reactors. Such steels were necessarily employed in thick sec- tions and the fabrication of the vessels involved weld- ing of preformed plates or forgings (for a description of typical vessels, see Steele and Sterne1). Frequently, in early designs, the welds were located opposite the center of the reactor core and received the highest neutron dose. Commercially available ferritic steels were employed in the construction of the first reactors (both for Magnox and LWR designs). For LWRs, which have to contain higher pressure than the pressure vessels of Magnox reactors, the desirability of using steels of high toughness, ade- quate strength, and weldability in thick sections, combined with good service experience, has nar- rowed the choice to a number of low alloy steels; for example, those containing manganese, nickel, and Table 1 Chemical composition of Magnox vessel materials Material C Mn Si Plate 0.09–0.17 1.04–1.32 0.10–0. MMA 0.086 0.91 0.42 SMA 0.088 1.49 0.52 Forgings 0.18 1.30 0.36 Ni and Cr each Table 3 Chemical composition of surveillance materials for Sizewell ‘B’ PWR Wt% C Si Mn P S Ni Cr Mo Al V Cu A508 Class III forging 0.16 0.27 1.42 0.005 0.004 0.74 0.13 0.51 0.022 Table 4 Parameter range of interest for different reactor types Reactor type Temperature range Dose range (n cm�2)> 1MeV a Dose range (mdpa) Dose rate (n cm�2) E>1MeVs�1 Magnox 160–>390 �C 1016–2� 1018 fast n cm�2 (Ni doses) 0.02–4 PWR 270 �C–296 �C 6–8�1019 60 �1�1011 BWRs Predominantly 270 �C 1MeV (see Heatherly et al.15 and Eason et al.29). The irradiation temperature of the irradiation rig was restrained to Radiation Damage of Reactor Pressure Vessel Steels 157 are�1� 1012–1� 1013 n cm�2, E> 1MeV s�1, that is, one or two orders of magnitude higher than the PWR surveillance dose rates given in Table 4. Overall, it was found in all RPV steels of interest that embrittlement increases with increasing fluence, but the rate of embrittlement may decrease (with increasing fluence). Furthermore, embrittlement does not saturate in the fluence range of interest to power reactor applications. These trends are illustrated in Figure 2 for Magnox CMn steel SMAweld transition shifts and for a MnMoNi plate HSST-02. Irradiations in the 1960s demonstrated that com- position was a major factor controlling the response of the ferritic low alloy steels employed in operating reactors. By the mid-1960s, it was thought likely that residual elements in steels could be responsible for much of the observed scatter in the irradiation embrittlement response.31,32 The work suggested that reducing the residual element content of A302-B steel would markedly improve the resistance 0 20 40 60 80 100 120 Measured ΔT30 ft-lbs (�C) M ea su re d Δ T 3 0 ft -l b s (� C ) Fluence (n cm−2, E > 1 MeV) 0 1 � 1019 2 � 1019 3 � 1019 4 � 1019 5 � 1019 (a) (b) 50 0 100 150 DTMatrix 0 Tr an si tio n sh ift (� C ) 5 10 15 DTTotal Peak Cu precipitation Ö(dpa � 105) DTMeasured ¯ DTConverted ¯ Figure 2 (a) Magnox submerged-arc weld transition shift data. Reprinted with permission from Buswell, J. T.; Jones, R. B. In Effects of Radiation on Materials, 16th International Symposium, ASTM STP 1175; Kumar, A. S., Gelles, D. S., Nanstad, R. K., Eds.; American Society for Testing and Materials: Philadelphia, PA, 1994; pp 424–443. Copyright ASTM International and (b) HSST-02 reference plate irradiated in surveillance schemes in a number of US LWRs.22 to irradiation embrittlement at typical service tem- peratures of 550 oF (288 oC). In a pioneering set of studies, Hawthorne and coworkers undertook studies that explored the effect of material composition on irradiation embrittlement in a systematic manner. Studies were undertaken on materials irradiated at a controlled temperature of 288 �C in the Union Carbide Research Reactor (UCRR) and in the light water-cooled and moder- ated test reactor, UBR, at the Buffalo Materials Research Center.25 A series of small (150 kg) labora- tory melts were produced to the nominal plate steel specifications using pure elements. These were then split, generally into three blocks, two of which were remelted and selected residual element additions added, while the third was kept to provide a low residual element reference. Each steel was also com- pared to material obtained from normal commercial production. Potapovs and Hawthorne33 demonstrated that additions of Cu, and Cu and Ni, to a laboratory melt containing low level of residuals greatly increased the observed embrittlement (see Table 5). This must be regarded as a landmark paper in the understanding of the factors that control radiation damage in RPV steels. Note that it was �15 years before the underlying mechanisms were elucidated (see Section 4.05.4). The effect of different Cu and Ni levels in steels irradiated as part of US surveil- lance programs is illustrated in Figure 3. The effect of increased levels of Cu in steels of the same Ni level and the effect of increased levels of Ni at constant Cu level is clear (data taken from Eason et al.33). Table 5 Comparison of shift in 30 ft-lbs (41 J) transition temperature (DT30 ft-lb) due to irradiation at 288 �C for experimental and commercial weld deposits and the A543 reference plate studied by Potapovs and Hawthorne33 Material Composition (wt%) Fluence (1019 n cm�2) [E>1MeV] DT41 J (DT30 ft-lb )� C Cu Ni Expt. weld 1934 0.06 0.77 3.5 53 Expt. weld 1938 0.07 1.62 3.5 111 Expt. weld 1938 0.07 1.62 3.5 200 Expt. weld 1948 0.03 1.56 3.5 110 Commercial Filler (0.24Cu) 0.24 1.58 3.5 415 (a) (b) 0 1�1019 2�1019 3�1019 4�1019 5�10190 20 40 60 80 100 120 140 US weld 0.25 wt% Cu 0.54 wt% Ni US forgings 0.06 wt% Cu 0.7 wt% Ni C ha rp y sh ift Δ T 4 1J (� C ) Fluence n cm−2 (E > 1 MeV) 1�1019 2�10 19 3�1019 4�1019 5�1019 6�1019 7�10190 C ha rp y sh ift Δ T 4 1J (� C ) Fluence n cm−2 (E > 1 MeV) 0 50 100 150 200 250 US weld 0.23 wt% Cu 1.2 wt% Ni US weld 0.2 wt% Cu 0.06 wt% Ni US plate 0.2 wt% Cu 0.18 wt% Ni Figure 3 Charpy shift (DT41 J (�C)) for (a) a US weld and a US forging containing 0.25 and 0.06wt% Cu, respectively, and (b) US welds and a US plate containing �0.2wt% Cu and varying levels of Ni. A 0 100 (%P) 0.003 (%P) 0.003 0.015 0.015 *Irradiated 0.025 0.025 0.025 + 0.024% Sn 0.025 + 0.021% Sn 200 0 50 100 150 (D�C)(D�F) 300 C v4 1J (3 0 ft -l b ) t ra ns iti on t em p er at ur e in cr ea se Melt 67* Melt 68* I MeV Figure 4 Charpy 41 J transition temperature increases observed for plates from Melt 67 and Melt 68 showing that the phosphorus influence on radiation sensitivity depends on the copper content. Reprinted with permission from Hawthorne, J. R. In Irradiation Effects on Structural Alloys for Nuclear Applications, ASTM STP 484; American Society for Testing and Materials: Philadelphia, 1970; pp 96–126. Copyright ASTM International. 158 Radiation Damage of Reactor Pressure Vessel Steels The hardening of the low Cu forging illustrated in Figure 3(a) follows a square root dependence of embrittlement on fluence. Hawthorne and coworkers also studied the influ- ence of other elements.25 The isolation of the effect of phosphorus, using plates from split laboratory melts, is illustrated in Figure 4. In brief, the data revealed that the radiation sensitivity is strongly dependent on level of the phosphorus, but the magnitude of the effect is highly dependent on the amount of copper present. The contribution from increasing P is most pronounced when copper is low. A second observation from Figure 4 is that tin additions (0.023% vs. Weld SH Weld SG 0.0 0.010 100 100 S hi ft (� C ) 0.0100.0 Dose (dpa) Weld SF Welds SD and SL Dido OsirisHerald VT4 (low flux) Herald Figure 5 Comparison between results obtained at different dose rates and in different irradiations of quenched and tempered low alloy welds (DIDO, HERALD, and OSIRIS are all MTRs in which samples were irradiated). Reproduced from Williams, T. J.; Ellis, D.; Swan, D. I.; et al. In Proceedings of the 2nd International Symposium on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, Sept 1985; ANS: Monterey, CA, 1986. Radiation Damage of Reactor Pressure Vessel Steels 159 0.13%, respectively, by weight. Irradiation tempera- tures were reported to have been controlled to �1 �C34 and the fast neutron doses were restricted to a maximum of 2.5� 1017 n cm�2 (as measured by Ni monitors). The results between 100 and 350 �C exhibited a simple linear dependence of yield strength increase as a function of irradiation temper- ature. Jones and Williams35 carried out an analysis of the Barton data34 and another dataset on the irradia- tion temperature dependence of similar steels from Grounes,36 and pointed out that the combined data form a homogeneous data set with relatively little variability. The least squares regression is FTðDsÞ ¼ 1:869� 4:57� 10�3T ½1� This parameter has been important for the correla- tion of data from similar steels irradiated at different irradiation temperatures.2 Understanding the effect of flux, or dose rate, on radiation damage of ferritic steels has proved particularly important in the formulation of mecha- nistically derived DDRs. There was particularly strong interest in the United Kingdom in under- standing the effects of flux on bulk properties as Magnox RPVs operated within a range of fluxes.2 A number of experimental investigations have examined the dose rate dependence of hardening.37 In the absence of precipitation effects, no influ- ence of dose rate on irradiation hardening has been detected. Data obtained from over five orders of magnitude change in dose rate for C–Mn plate steels at an irradiation temperature of 200 �C, typical of Magnox applications,37 demonstrated no dependence on dose rate. However, there is agreement that there is a strong effect of flux on the embrittlement of Cu-containing steels.38 In these steels, it was found that at doses before the ‘saturation’ of embrittlement (see Section 4.05.4) the rate of embrittlement with fluence increased with decreasing dose rate.Williams et al. studied the effect of dose rate on the embrittlement in low Ni welds at preplateau doses19 (1� 1019 n cm�2, E> 1MeV is approximately 1.5� 10�2 dpa). They reported the irradiation-induced shift in the 41 J transition temper- ature of a number of Mn–Mo–Ni SMA welds after irradiation in MTRs at dose rates between 6� 10�10 and 2� 10�8 dpa s�1 and doses of less than�30mdpa. It was observed that for the welds SD and SL, in which theCu levels are low ( 160 Radiation Damage of Reactor Pressure Vessel Steels damage causing nonhardening embrittlement was also found in CMn steels. There are proven mechanical test techniques for determining the change in bulk properties and large irradiation programs have been performed. These have made use of both surveillance programs and MTRs. The level of the measured embrittlement depends on the fluence, flux, and irradiation temper- ature. The most important discovery was the sensi- tivity to steel composition, in particular a strong dependence of embrittlement on the levels of Cu, Ni, and P. 4.05.4 Development of Mechanistic Insight of Factors Controlling Current Plant Lifetimes 4.05.4.1 Introduction The previous section included a description of how the effect of radiation damage on the bulk properties of ferritic steels was established from the 1950s and how critical insight into the important role of Cu arose in the late 1960s and early 1970s. Equivalent advances in mechanistic understanding did not occur for another 10 or 15 years. The improved mechanistic understanding in the 1980s had its origin in the identification of the con- trolling variables that emerged from the experiments discussed in Section 4.05.3. This stimulated consider- able interest on the possible role of elements such as Cu in the embrittlement process. The other, possibly more important, reason was that in the 1970s there were only a few microstructural techniques available for characterizing irradiation damage. Transmission electron microscopy (TEM), the dominant technique, could not resolve the irradiation-induced damage in steels that resulted in a significant change in mechani- cal properties.14 However, in the mid-1980s there was an explosion of information resulting from the appli- cation of a range of different and improved micro- structural techniques. These techniques have now been applied to PWR, BWR, and Magnox steels irra- diated in surveillance locations of power reactors and to representative materials irradiated in materials test- ing reactors. A more recent advance that is highly relevant to developing detailed mechanistic insight is the advent of ‘multiscale’ modeling. Here, the power of modern computing tools is such that microstruc- tural development (and the resultant change in mechanical properties) can be modeled across the various length and timescales involved in RPV embrittlement. Such models are subject to intense R&D and the current capability can be seen in Wirth et al.,42 Soneda et al.,43 Becquart,44 and Domain et al.45 However, these models have not, as yet, made a direct impact on the development of DDRs and so the current state of multiscale model- ing is not reviewed here. 4.05.4.2 Radiation Damage Mechanisms Microstructural development of ferritic steels during service is driven by the interaction of neutrons and metal atoms. Collisions between the incident neutron and constituent atoms result in momentum transfer to the lattice atoms. If the transfer is above �40 eV, atoms can be permanently displaced from their lattice sites resulting in vacancy–interstitial pairs. Indeed, lattice atoms with tens of kiloelectron volts may be created and a branching, tree-like distribution of displaced atoms formed, termed a displacement cas- cade. Vacancy and interstitial clustering may occur within the cascade. Vacancies and interstitials escap- ing from the cascade give rise to concentrations of vacancy and interstitial point defects throughout the material. The fate of the point defects formed in the irradiation depends most sensitively on irradia- tion dose rate and temperature, and also material factors such as composition. This chapter focuses on the microstructural development of RPV steels. These steels experience a relatively low dose rate (and thus lifetime dose) at a relatively low temperature. The microstructure developed under these conditions is very different to that developed at high doses and high tempera- tures in the operating regime typical of fusion or fast reactors. The critical features are that at the temper- ature range of most operating pressure vessels both vacancy and interstitial point defects are mobile. Freely migrating defects (vacancies or interstitials) and mobile interstitial clusters escaping from the initial damage event may interact with point defect sinks, such as preexisting dislocations, recombine with each other, either directly or at solute traps, or cluster to form vacancy or interstitial clusters. The end result of the interactions described above is a microstructure comprising of a high density of small clusters in the matrix. These clusters may be point defect solute complexes, and, depending on steel composition, solute clusters formed from radiation-enhanced precipitation. At the dose, dose rate range, and temperature of interest to operating ng th o r tu re Total (high Cu) Radiation Damage of Reactor Pressure Vessel Steels 161 reactors, the clusters formed are usually thermally stable. Lastly, segregation of solutes to grain bound- aries or other sinks may occur. (a) In cr ea se in y ie ld s tr e C ha rp y te m p er a (ft)1/2 CEC for low Cu CEC for high Cu MD Total (low Cu) (b) Y ie ld s tr es s (s y) o r fr ac tu re s tr es s (s F) Temperature Embrittlement HardeningΔTT2 ΔTT3 ΔTT1 sF sy sy+ sF- Figure 6 (a) Schematic showing the dose dependence of matrix damage and copper clustering, and (b) variation of yield stress, sy, and fracture stress, sF, with temperature. 4.05.4.3 Mechanistic Framework The effect of radiation damage on ferritic RPV steels46–50 has for some time been considered in terms of � The formation of matrix damage (MD), that is, defect clusters and dislocation loops. It is well established that in low copper steels the shift in impact or yield strength properties depends on √dose. � The irradiation enhanced formation of copper- enriched clusters (CECs). (CEC are also referred to as CRPs (Cu-rich precipitates) as they were originally assumed to be strictly Cu-rich rather than simply Cu-enriched.) It has been demon- strated that, in many low-to-medium Ni steels and alloys, the yield strength change due to copper precipitation rises to a plateau value that is then unchanged by subsequent irradiation. � The irradiation induced/enhanced grain boundary segregation of embrittling elements such as P. The first two mechanisms contribute to embrittle- ment by increasing the steels’ hardness as illustrated in Figure 6(a). The third mechanism induces embrit- tlement without hardening. The latter mechanism is not necessarily found in all RPV steels under operating conditions. Indeed, for MnMoNi steels irradiated in surveillance schemes in Western LWRs, the observed embrittlement is associated with the first two mechanisms; that is, the total shift in the ductile to brittle transition, as measured at the Charpy 41 J level, is DT41J ¼ DT41JðCRPÞ þ DT41JðMDÞ ½2� where CRP¼Cu-rich precipitate and MD¼matrix damage, or equivalently the increases in yield strength, Dsy, is given by Dsy ¼ DsyðCRPÞ þ DsyðMDÞ ½3� The first two mechanisms serve to harden the mate- rial and increase the yield strength sy, while the third mechanism causes a drop in the fracture strength, sF. The effects of these changes on the fracture behavior are illustrated in Figure 6(b), where the temperature dependence of the yield stress, sy, and the fracture stress, sF, are plotted. It can be seen that the effect of irradiation in causing hardening or a change in sF is to cause a change in the transition temperature. In the figure, DTT1 is caused by an increase in sy, while segregation of P to grain boundaries can lower the fracture stress and result in a shift DTT2. If both mechanisms are operative, then a combined shift of DTT3 occurs. It is important to note that Ni and Mn are known to strongly influence hardening in steels containing low levels of Cu and also CEC hardening in Cu-containing steels. In Cu-containing steels, satura- tion of the cluster hardening has been demonstrated in steels containing up to �1wt% Ni. At steel Ni levels above �1.5 wt% (and with Mn 1.2–1.7 wt%), cluster hardening has not been observed to saturate. The precise Cu, Ni, and Mn levels at which the plateau is suppressed have not been fully char- acterized, and are the subject of current research. Similarly, the exact influence of Ni and Mn on the embrittlement of low Cu steels has not been fully established and is again a subject of ongoing research. (The term standard MnMoNi steels is used to refer to steels with typical Mn levels ( 162 Radiation Damage of Reactor Pressure Vessel Steels with some workers preferring a limit of 1.0 wt%, with others promoting a higher limit.) Indeed, as will be described in Section 4.05.6, there is concern about whether at high doses (typical of those achieved in plant with extended lives) there may be deviations from the simple framework established above. 4.05.4.4 Cluster Development Under Irradiation: Cu 4.05.4.4.1 Introduction The purpose of this subsection is to establish the level of understanding regarding the irradiation-induced formation of CECs, and how this understanding sup- ports the mechanistic framework outlined above. Insight has been developed by characterizing popula- tions of irradiation-induced CECs in Cu-containing steels and, in particular, the dependence of CEC structure, composition, size, and number density on material and irradiation parameters (e.g., steel com- position, fluence, flux, etc.). Measurements have also been made on the level of Cu in the matrix (Cumatrix) which is not associated with any CEC or precipitate. This is an important measure as, clearly, Cumatrix will decrease during irradiation as CECs are formed, and at SOL it may be lower than at the bulk level if Cu is precipitated during the final heat treatment. Several techniques have been successfully em- ployed to characterize the CECs formed during neu- tron irradiation of Cu-containing pressure vessel Table 6 Techniques for the characterization of irradiation-in Property Existing techniques Comme Shape and size of solute clusters 3DAP, SANS, TEM/FEGSTEM Analysi provid avera provid �60� Composition of solute clusters 3DAP, SANS, FEGSTEM/EDX 3DAP y obser in a s comp meas advan comp Number density (Nd) and volume fraction of solute clusters 3DAP, SANS, FEGSTEM/ EFTEM FEGST visible from a Level of matrix Cu 3DAP, FEGSTEM/EDX 3DAP p is not devel cluste steels (Table 6). TEM was first used to observe CECs, but the cluster sizes were close to the resolu- tion limit for most TEMs, so atom probe (3DAP) and small angle neutron scattering (SANS) became the most commonly used methods to characterize CECs. However, these are not the only techniques; for example, an important development has been the advent of experiments employing a positron annihi- lation technique, coincidence Doppler broadening (CDB) spectra of positrons annihilating in aged or irradiated Cu-containing alloys or steels. CDB pro- vides a means of identifying the elements around the annihilation site.52 It has become the standard practice to characterize the same as-irradiated specimens with a number of techniques.53,54 Part of the logic behind this is that all the techniques have limitations, either in terms of volume analyzed or complexity of interpret- ing experimental data, and that combining techniques provides better information. These microstructural techniques have been discussed in greater detail in English and Hyde.55,56 4.05.4.4.2 CEC characterization Evidence from various independent studies using techniques such as FEGSTEM or 3DAP has con- firmed the small solute clusters to be multi-alloyed with minor constituents of the steel. An example of the composition of a cluster formed in an irradiated lowNi A533B weld metal is shown in Figure 7. The results of the 3DAP study clearly show that the CECs are alloyed duced solute clusters nt s of the irradiation-induced scattering data from SANS es a feature size distribution (for clusters 1–10nm diameter) ged over a volume approximately 50mm3. 3DAP (LEAP) es a direct estimate of the size of clusters formed in a region 60�100nm3 ields a direct measure of the composition of each cluster ved. SANS measurements of irradiation-induced scattering trong magnetic field provide indirect evidence of cluster osition. Techniques are starting to be developed for uring the level of Fe in Cu-enriched precipitates using ced TEM techniques. At present, uncertainties on osition of small (�2 nm diameter) are large EM and 3DAP can directly measure the number density of clusters in the analyzed volume. In practice, Nd estimates ll techniques are subject to considerable uncertainty rovides a direct measure of the level of Cu in the matrix that in clusters. In the FEGSTEM techniques have been oped for measuring the level of matrix Cu from areas where rs are not observed51 Silicon Manganese Nickel Copper Figure 7 Example of single irradiation-induced cluster approximately 3 nm in diameter. Radiation Damage of Reactor Pressure Vessel Steels 163 with Mn, Ni, and Si; there is also evidence of an association with P near the cluster–matrix interface (not shown). Similar cluster compositions have been found by other workers using other atom probes and FEGSTEM (which is not sensitive to Fe in clusters), and there is good agreement between the techniques. The atom probe data indicate that the clusters are dilute in that the majority element is Fe.57–59 In contrast, analysis of SANS data is often undertaken assuming that irradiation-induced clusters are non- magnetic (implicitly Fe-free). The assumption of low levels of Fe in the Cu clusters was supported by thermodynamic calculations that predicted low levels of Fe in such precipitates.47 Furthermore, PAS-OEMS on as-irradiated (290 �C) Fe–0.9Cu or Fe–0.9Cu–1Mn indicate that the clusters in this con- dition are not magnetic and therefore very unlikely to contain Fe,60 although it should be noted that clusters in model alloys may be different from those found in RPV steels. Carter demonstrated that the scattering observed in SANS experiments is not inconsistent with the presence of magnetic clusters.57 More recently, Morley et al.61 attempted to characterize the extent of trajectory aberrations in the atom probe that might give rise to an incorrect estimate of the true Fe content in small clusters in thermally aged (330–405 �C) RPV steels. He concluded that the clusters do contain Fe, but the levels are lower than those measured. Furthermore, he found that the con- centration of Fe in the precipitate phase is a function of aging temperature with less Fe at higher aging temperatures. Consensus has not yet been reached on the precise Fe content in irradiation-induced clus- ters in ferritic steels and is the subject of ongoing research. 4.05.4.4.3 Development with increasing fluence The progressive precipitation of the Cu present in solution at the SOL as a result of irradiation has been recognized since the 1980s. There are a large number of studies that show that a high density of small CECs are formed under irradiation, and that the number density, size, and volume fraction are strongly depen- dent on the irradiation fluence, flux, and the material composition. For MnMoNi steels and Fe–Cu alloys, the number density and volume fraction increase rapidly with increasing fluence, and then there is an appearance of saturation, that is, a pattern of behavior that mir- rors the shape of the curve in Figure 6(a). This is illustrated in Figure 8 from the work of Odette et al. (see data presented in Eason et al.29 where the volume fraction, fp, and radius have been derived from SANS data). It can be seen that the radius increases with fluence in this example. Auger et al.62 had found a similar pattern of behavior in SANS and AP data from ten steels and two Fe–Cu alloys with 0.2 wt% Cu. Saturation occurred at a similar fluence to that of Odette et al., that is, �1� 1019 n cm�2 (for irradiation temperatures close to 290 �C). The rate at which the volume fraction increases with fluence is also strongly dependent on the irradi- ation flux and the composition of other elements such as Ni. Figure 9 shows the SANS measurements of Williams and Phythian63 on MnMoNi SAWs. It shows the effect of dose rate, Cu, and Ni on the development of CEC volume fractions with dose. Decreasing the Cu decreases both the volume fraction and CEC size. In the high Cu welds (seen most clearly at high dose rate), increasing the Ni clearly increases the total volume fraction of CECs formed at all doses. At the same time, the mean CEC diameter is somewhat decreased in the higher Ni weld (thus the volume fraction increase is associated with a large increase in cluster number density). It is also evident that, while the precipitated volume fraction appears to be satur- ating at the highest dose in the low Ni welds, there is no sign that saturation is close in the high Ni welds. A number of authors have found similar results.29,64,65 High Cu, High dose rate High Cu, Medium dose rate Low Cu, High dose rateHigh Cu, Low dose rate Dose (mdpa) Dose (mdpa) 0 (a) (b) Vo lu m e fr ac tio n (� 10 –3 ) Vo lu m e fr ac tio n (� 10 –3 ) 0 1 2 3 4 5 6 4 2 0 6 10 20 30 40 0 10 20 30 40 Figure 9 Effect of bulk composition and dose rate on CEC size and volume fraction in MnMoNi SAWs. In (a) data are from low Ni welds, high Cu �0.08Ni, 0.55Cu; low Ni, low Cu �0.08Ni, 0.15Cu; and in (b) data are from high Ni welds, high Cu �1.65Ni, 0.55Cu; high Ni, low Cu �1.5Ni, 0.05Cu (all wt%). In both figures could be seen high flux �5.5�10�9, medium flux �6.3� 10�10, and low flux �9�10�11 (all dpa s�1). Reprinted with permission from Williams, T. J.; Phythian, W. J. In Effects of Radiation on Materials, 17th International Symposium, ASTM STP 1270; Gelles, D. S., Nanstad, R. K., Kumar, A. S., Little, E. A., Eds.; American Society for Testing and Materials: Philadelphia, 1996; p 191. Copyright ASTM International. 0.01 0.1 ft (1023 n m−2) ft (1023 n m−2) 0.0 0.2 0 1 < R > (n m ) 2 0.4 0.6 f p (% ) 1 10 0.01 0.1 1 10 Figure 8 SANS data on volume fraction, fp, and, radius, rp, for 0.4wt% Cu, 1.25wt% Ni split melt model steel alloys (LD) irradiated at three flux levels between 0.6 and 10�1015nm�2 s�1 in IVAR at 290 �C, plotted as a function of fluence, ft. Reproduced from Eason, E. D.; Odette, G. R.; Nanstad, R. K.; Yamamoto, T.; EricksonKirk, M. T.A Physically Based Correlation of Irradiation-Induced Transition Temperature Shifts For RPV Steels; Oak Ridge Report ORNL/TM-2006/530, 2007. 164 Radiation Damage of Reactor Pressure Vessel Steels It is to be noted that Soneda65 found that when com- paring Cu clusters in low-fluence-irradiated steels formed at a flux of 109 and 1010 n cm�2 s�1 E> 1MeV, there was significant effect of dose rate on cluster size rather than number density. The effect of Ni was established early on,66 but it has only just been established that Mn also has an important effect. Odette et al. demonstrated from SANS studies that at constant Cu and Ni, increasing Mn decreased the size of the clusters but increased their number density as illustrated for 0.4 wt% Cu, 0.8 wt% Ni, and 3.4� 1019 E> 1MeV in Figure 10. In higher Ni steels, Burke et al.67 have also demon- strated that removing Mn from a steel significantly lowers the resultant embrittlement and the level of observable solute-related damage. It was thought for many years (e.g., Jones and Bolton2) that CEC-related hardening would reach a maximum level once all the Cu had precipitated, and remain at this level as the CEC size remained constant – probably as a result of a balance between cluster nucleation and growth, and cluster destruction in cascades, that is, overaging did not occur. Jones and Bolton2 reported measurements of Cu cluster diame- ter using SANS on unirradiated and irradiated C–Mn SMA welds. It was shown that under surveillance conditions, and at temperatures below about 300 �C, Cu clusters grew to about 2 nm in diameter. Even after subsequent accelerated irradiations (of the sur- veillance samples irradiated to the lowest doses) to doses between �200� 10�5 and 1200� 10�5 dpa, the mean precipitate diameter was still �2 nm. 3 Dose (�1019n cm−2) 2 1 0 5 4 3 2 1 0N um b er d en si ty (� 10 17 cm −3 ) R ad iu s of g yr at io n (n m ) 0 2 4 6 Doel-1 Doel-2 Figure 11 Radius of gyration and number density of CRPs for the RPV surveillance test specimens of Doel-1 (△) and Doel-2 (�). Reproduced from Toyama, T.; Nagai, Y.; Tang, Z.; et al. Acta Mater. 2007, 55, 6852–6860. C E C v o lu m e fr ac tio n (ft)1/2 Increasing Cu and Ni Decreasing flux Ni May decrease at high fluence Figure 12 Schematic of the effect of selected material and irradiation variables on CEC volume fraction. 2.0 1.8 1.6 1.4 1.2 1.0 6.0 5.0 4.0 3.0 2.0 1.0 0.0 0.5 1.0 0.4 wt% Cu 0.8 wt% Ni 290 �C 3.4 � 1023n m–2 Mn (%) < r> (n m ) N P (1 02 3 m −3 ) 1.5 2.0 0.0 0.5 1.0 Mn (%) 1.5 2.0 Figure 10 SANS data on cluster radius, rp, number density, Np, and volume fraction, fp for 0.4wt% Cu split melt model steels irradiated at high IVAR flux at 290 �C. Effects on Mn variations in alloys with �0.8wt% Ni. Reproduced from Eason, E. D.; Odette, G. R.; Nanstad, R. K.; Yamamoto, T.; EricksonKirk, M. T. A Physically Based Correlation of Irradiation-Induced Transition Temperature Shifts For RPV Steels; Oak Ridge Report ORNL/TM-2006/530, 2007. Radiation Damage of Reactor Pressure Vessel Steels 165 More recent observations at the onset of the pla- teau in CEC formation in a number of commercial steels have shown that only around half of the avail- able Cu had precipitated (Auger et al.)62 This suggests that particle coarsening and overaging could occur in irradiated RPV steels as well as in thermally aged steels, once the precipitation level was high. Particle coarsening has been reported in MnMoNi surveillance material from the Doel-1 and Doel-2 reactors,68,69 as shown in Figure 11. This particle coarsening should result in overaging (i.e., softening) after the hardness reaches a maximum value. Figure 12 illustrates schematically the influence of selected material and irradiation variables on the CEC volume fraction (nominally for a Cu-containing MnMoNi steel irradiated at �290 �C). The Cu level in the matrix at the SOL is an important parameter as it is the matrix Cu that is available for precipitation of CECs. This has been determined by either thermodynamic modeling or by direct measurement. For example, modeling calcula- tions were performed by Buswell and Jones70 to determine matrix Cu levels in Magnox SMA welds with bulk Cu contents between 0.13 and 0.31 wt%. They found that the precipitation of Cu during the final weld stress-relief heat treatment (at 590–600 �C), and also during the subsequent extremely slow cooling (5 �Ch�1) of the RPV before reactor operation, reduced the maximum Cu available to precipitate during irradiation to no more than 0.15–0.20wt%.70 Precipitation during the final weld stress relief also occurs in US steels. Indeed, a consensus has emerged that there is an upper limit to SOL-dissolved Cu, which is dependent on heat treatment. McElroy and 166 Radiation Damage of Reactor Pressure Vessel Steels Lowe have shown that even differences of 20–30 �C in the heat treatment temperature can markedly affect the dissolved Cu content,71 as can the slow cooling of real structures from the stress-relief temperature. 4.05.4.5 Development of Matrix Defects 4.05.4.5.1 Introduction Historically, the hardening observed in low Cu steels has been considered to arise from matrix defects. MD arises from the clustering of irradiation-induced point defects to form either vacancy or interstitial clusters, and/or solute–defect complexes (see, e.g., Odette and Lucas53). In addition, various solutes may diffuse to these clusters giving rise to complex defect–solute configurations. A critical factor in attempting to develop insight is that there is no one technique that allows a direct characterization of matrix defects. Insight has been obtained from indi- rect studies using techniques such as positron annihi- lation, where the nature of matrix defects can be inferred only after data analysis or modeling. Within these constraints, there are two aspects which should be discussed. First, the nature of MD, in particular whether it is vacancy or interstitial clus- ters that give rise to the observed hardening, and, second, the evolution of MD clusters at different fluxes and increasing fluence. The discussion also needs to include an assessment of whether MD is dependent on the presence of other solutes and whether MD can be treated independently of CEC formation. It should be noted that in generating insight into MD, studies of model alloys (e.g., Fe–Cu alloys) have been particularly important, including ion and electron irradiation studies. This is in contrast to the study of Cu precipitation where the majority of information has arisen from studies of Cu-containing neutron-irradiated steels. 4.05.4.5.2 Nature Experimental information on the nature of matrix defects necessarily splits into evidence for clusters formed by vacancies and interstitial point defects. Significant insight into potential vacancy clusters has come from positron studies of irradiated and model steels. Positrons are well established as a tool for probing vacancy-type defects in solids, where the defect concentration is typically 1018 cm�3. (Posi- trons are attracted to regions of the lattice which are more ‘open’ than average. The most obvious such regions are vacancies and vacancy clusters (the larger clusters being stronger positron traps). Less obvious open regions include the tensile parts of a dislocation strain field (even around an interstitial loop) or incoherent particle–matrix interfaces. The positron annihilation techniques have included both positron lifetime (PA-t) and lineshape (PALA) ana- lyses or, more recently, CDB. The latter is an impor- tant recent development. CDB (also known as positron annihilation spectroscopy-orbital angular momentum distribution, PAS-OEMS) measures the energy spectrum of the gamma rays produced at the positron annihilation sites, and thus the momenta of the electrons at those sites. The energy (or momentum) spectra characterize the elements sur- rounding the positron traps.72 Overall, the experimental data provide strong evi- dence for open-volume clusters that are sensitive to positrons, and for PA being a useful technique for studying MD (see, e.g., Buswell and Highton,73 Dai et al.,74 Carter et al.54). In model alloys, a number of authors report evidence for large vacancy clusters or microvoids after irradiation. Most interestingly, although the positron lifetime studies have indicated the presence of vacancy clusters in model alloys, such studies have not identified vacancy clusters in com- mercial steels (e.g., Dai et al.74). The study by Valo et al.75 is most convincing in this aspect, as the authors investigated both model alloys and RPV steels. Evidence for interstitial clusters has come primar- ily from studies of model alloys or steels irradiated to very high doses in excess of that of interest for most power reactor applications. Such studies have nor- mally employed TEM techniques that are sensitive to the strain field associated with small dislocation loops. The imaging of such features is difficult in ferritic materials because of the need to correct for the image distortion caused by the ferromagnetic behavior of the samples, and because of the contrast arising from surface oxide on the thinned specimens. Critically, the resolution for small dislocation loops is �2–3 nm in even well-prepared samples imaged in higher voltage microscopes. Krishnamoorthy and Ebrahimi76 and Hoelzer and Ebrahimi77 reported the formation of visible interstitial loops in Fe irra- diated to 4.63� 1023 nm�2; E> 1MeV at �280 �C; the loops increased in size and decreased in number density after annealing at 500 �C. There have been fewer studies on irradiated steels, but in MnNiMo steels little evidence for dislocation loops has been reported. Soneda and coworkers have undertaken weak-beam TEM obser- vations of RPV surveillance steels containing 0.06 and 0.12wt% Cu irradiated to 4� 1019 n cm�2.65 g g (b)(a) 20 nm Figure 13 Contrast analysis on dislocation loops: weak-beam images of the same area imaged (a) with diffraction vector g¼ (011) and (b) g¼ (200) close to the [011] pole, in the foil at 300nm depth. Reproduced from Fujii, K.; Fukuya, K. J. Nucl. Mater. 2005, 336, 323–330. Radiation Damage of Reactor Pressure Vessel Steels 167 Soneda reported the formation of interstitial dislo- cation loops, whose diameter and number density are 2–3 nm and of the order of 1021 m�3, respec- tively. At very high doses (and dose rates) a uniform density of loops has been observed. Fuji and Fukura78 undertook a weak-beam TEM study for MD in A533B RPV steel produced by 3MeV Ni2þ ion irradiation to a dose of 1 dpa at 290 �C. The MD was found to consist of small dislocation loops. The observed and analyzed dislocation loops have Burgers vectors b¼ a (Figure 13). The dis- location loops have a mean image size d¼ 2.5 nm and the number density is about 1� 1022m�3. Most of the loops are stable after thermal annealing at 400 �C for 30min. This indirect evidence suggests that their nature is interstitial. Kocik et al.79 examined the radiation damage microstructures in Cr–Mo–V surveillance base metal and weld containing (�0.06–0.07 wt% Cu and 0.012–0.014 wt% P) irradiated up to 6� 1024 nm�2 (E> 1MeV) for times up to 5 years at 265 �C. TEM examination of the irradiated materials revealed, in both the base metal and the weld metal, black dots, small (resolvable) dislocation loops, and small precipitated particles. Clouds of defects are formed along dislocations at higher neutron fluences, and it was only at the higher fluence that loops that may not be associated with dislocations could be seen. Interactions were observed between defects and (as-grown) dislocations that result in a rebuild of dislocation substructure. Miller et al.80 examined the radiation damage microstructures in similar Cr–Mo–V surveillance base metal and weld. They reported manganese-, silicon-, copper-, phosphorus-, and carbon-decorated dislocations and other features in the matrix of the neutron-irradiated base and weld materials. 4.05.4.5.3 Development with flux and fluence and irradiation temperature The most important inference from the mechanical test data is that hardening and embrittlement are proportional to the square root of fluence in low copper steels. Early theoretical and experimental work by Makin and coworkers81,82 demonstrated that a square root dependency on dose was consistent with the hardening arising from the cutting, by glide dislocations, of irradiation-produced obstacles, and that in the early stages of irradiation the number density of clusters is proportional to the irradiation exposure. Thus, in irradiated low Cu RPV steels, there is continuous production of hardening centers during irradiation. Further, the linear dependence of hardening on irradiation temperature from 150 to �300 �C in CMn steels and low Ni A533B weld- ments implies that thermal stability of MD clusters is important.83 There are relatively few studies that generate insight into the effect of flux and fluence on MD itself. Unsurprisingly, studies of model alloys tend to emphasize the increase in number density (and size) of the vacancy-rich clusters with increasing dose. Kampmann et al.84 found void-like features 1–2 nm diameter in Cu-free ternary Fe–Ni–P/Mn alloys irradiated 2–25� 1018 n cm�2. The authors considered that the microvoid numbers increase with dose up to �5� 1018 n cm�2, and then either remain constant or decrease. Analyzing positron annihilation data from annealing studies of neutron- irradiated A533B plate, A508–3 forging, and welds, Carter et al.54 considered that increasing the dose from 1� 1018 to 20� 1018 n cm�2 at 290 �C increased the volume fraction of vacancy clusters, probably via increasing both the number density and average size of the clusters. Increasing the flux from 6� 1011 to 5� 1012 n cm�2s�1 increased either the number den- sity or the mean radius, probably the radius. Postirradiation annealing has been shown to be a powerful means of investigating the nature of the MD further. A major development has been characterizing the matrix defect term as being due to two compo- nents; first, stable matrix defects (SMD) and second, at high fluxes, unstable matrix defects (UMD) (see, e.g., Mader et al.85). UMD are matrix defects that, although thermally unstable at the irradiation temperature, are frozen into the microstructure during the cooldown after irradiation. Such studies have also established that MD and hardening of low Cu steels will be dose rate dependent at high dose rates (>1–5� 1012 n cm�2 s�1, E> 1MeV).85 168 Radiation Damage of Reactor Pressure Vessel Steels Soneda65 modeled the effects of dose, dose rate, and irradiation temperature on the defect accumula- tion in bcc-Fe using the kinetic Monte Carlo (KMC) method.65,86 Jones and Williams83 proposed a model that describes the irradiation temperature depen- dence of the embrittlement of low Cu materials, DT¼ a� FT� (’t)1/2, where DT, a, and ’t are the transition temperature shift (TTS), constant coefficient, and dose, respectively, and FT¼ 1.869– 4.57� 10�3T (�C). This model was studied using a KMC simulation. The number densities of both vacancies and self-interstitial atom (SIA) clusters exhibited a linear temperature dependence with a slope equivalent to that of FT, and Soneda considered that the origin of the form of the FT term can be understood from the temperature dependence of point defect cluster formation. 4.05.4.5.4 Effect of alloying There have also been a number of studies of the effect of alloying on loop formation. These studies have not examined all the alloying elements of interest, but Mn and P have been shown to have an important influence on the cluster distributions observed by TEM in Fe binary or ternary alloys. For example, Ebraimi and coworkers76,77 examined the effects of adding Ni (and P). They found that a higher density of smaller loops was observed in a Fe–Mn alloy (as compared to pure iron irradiated under similar con- ditions), whereas P added to an Fe–Ni alloy caused an increase in loop size. Phosphorus dissolves substitutionally in iron and the solid solubility is 0.5 at.% (0.27 wt%) at 400 �C.87 Jones and Buswell,88 in reviewing the available micro- structural evidence, concluded that the hardening observed in low Cu steels could be attributable to precipitation hardening by M3P particles produced by the irradiation-induced segregation of phosphorus to defect sinks and the depletion of phosphorus in solid solution, as detected by TEM and AP methods. Nagai et al.89 have reported results from a CDB study of Fe–0.3 wt% Cu, Fe–0.15 wt% Cu, and Fe–0.05wt%Cu alloys irradiated at 8.3� 1018 n cm�2, E> 1MeVat�300 �C (the irradiation timewas 144 h). As a result of CDB and positron lifetime measure- ments on irradiated and annealed samples, the authors reported the formation of microvoids (�10 vacancies), dislocation loops, and Cu-mono-vacancy-Cu com- plexes. They considered that the microvoids were decorated with Cu in all the alloys studied, and that in all cases the microvoids were almost completely coated with Cu. After electron irradiation,90 vacancy clusters and single vacancies surrounded by Cu (v-Cun, where n 6) were observed in electron- irradiated Fe–Cu, and vacancy clusters were observed Fe–Ni and Fe–P, but no vacancy clustering in Fe–C, Fe–Si, or Fe–Mn was observed. A recent development of some importance is the observation (primarily using the LEAP) of MnNiSi clusters in irradiated low Cu steels. For example, Miller et al.91 characterized the irradiation-induced microstructure of low copper (0.05wt%) high nickel (1.26 and 1.78 wt%Ni) VVER-1000 forging and weld materials that were neutron irradiated to a total flu- ence of 1.38� 1023 nm�2 (E> 1MeV). Atom probe tomography revealed ultrafine Ni–Mn–Si-enriched clusters but no CECs. The number density of clusters in the VVER-1000 weld was estimated to be �1.5� 1023m�3, while the number density of clusters in the forging was estimated to be slightly lower at 1� 1023m�3. These ultrafine clusters may, or may not, be associated with vacancies. The observa- tions of such clusters may be interpreted as evidence of a mechanism not encompassed by the framework set out in this section. This is further discussed in the next section. There is strong evidence that interstitial solutes (ISs) such as C and N are attracted to the point defects produced by irradiation. ISs may well add to preexisting SIA clusters, and may even inhibit their growth. Conversely, they appear to encourage the formation of multiple-vacancy complexes. Little and Harries92 further demonstrated that the amount of free nitrogen, indicated by the height of the Snoek internal friction peaks, decreased with increasing irradiation fluence, such that it was zero with fluences of about 2� 1018 n cm�2. This was attributed to trapping of free nitrogen or precipitation of nitrides at point defects or defect clusters. 4.05.4.5.5 MD and hardening Various scientists have attempted to determine the nature of the defects which result in hardening. Soneda65 quoted evidence from Ortner93 showing that DHv (the change in Vickers hardness) and DS (related to the volume fraction of open-volume defects) increase after irradiation of a low Cu steel EP2, indicating that vacancy-type defects are formed by irradiation. During the postirradiation annealing, DS starts to recover at a lower temperature than DHv. This clearly indicates that the change in DS is unre- lated to the change in DHv, and thus, vacancy-type defects are not solely responsible for the observed irradiation-induced hardening. 0 50 Dsy,pred (MPa) Ds y, m ea s (M P a) 100 150 200 250 300 350 0 50 100 150 200 250 300 350 WV LC LD Figure 14 Measured versus predicted Dsy from CRPs based on SANS measurements of fp and rp used in a modified Russell–Brown precipitate hardening and computer simulation derived superposition model (WV is a high-Ni high-Cu weld, while LC and LD are two medium strength �0.4wt% Cu split melt alloys with varying Ni levels). Reproduced from Eason E. D.; Odette, G. R.; Nanstad, R. K.; Yamamoto, T.; EricksonKirk, M. T. A Physically Based Correlation of Irradiation-Induced Transition Temperature Shifts for RPV Steels; Oak Ridge Report ORNL/TM-2006/530, 2007. Radiation Damage of Reactor Pressure Vessel Steels 169 4.05.4.6 Effect of Radiation Damage on Hardening The small defects formed in irradiated steels and model alloys can act as barriers to dislocation move- ment and therefore result in an increase in yield strength and hardness. Particularly important is the hardening from the copper-enriched precipitates/ clusters formed during irradiation in the high copper steels which can be modeled using the Russell–Brown model.94 The Russell–Brown model of hardening due to copper precipitates is a modulus interaction the- ory, based on the reduction in energy of the segment of dislocation, which passes through a relatively soft copper particle in the iron matrix. As the energy of the dislocation is proportional to the modulus of the host material, an attractive force will act on the dislo- cation because the modulus of copper is less than that of iron. Russell and Brown estimated the attractive force as a function of copper volume fraction, and demonstrated that this could adequately describe hardening in Fe–Cu alloys. A key element in applying the Russell–Brown model is the estimation of the modulus. Three approaches have been employed, using the modulus for fcc Cu, or computing values for bcc Cu,95 or fitting to experimental data. The last one is the most common approach. The matrix hard- ening may be estimated from the response of low Cu steels (Cu< 0.1 wt%). The individual hardening contributions from CECs and the MD must be combined with one another, as well as with the hardening from the pre- existing microstructure. The limiting rules for such superposition are a linear sum (LS) law and a square root of the sum of the squares (RSS) law.96 Computer models can be employed to determine the exact superposition law to be employed.97,98 Figure 14 shows a scatter plot, where the measured Dsy is compared to the predicted values.99 It can be seen that excellent agreement can be achieved. Bacon and Osetsky100 carried out molecular static (MS) and molecular dynamics (MD) simulations of the passage of a dislocation through a bcc Cu precip- itate. The MS simulations led to a dependence of hardening on precipitate size which differed from that predicted by the Russell–Brown model. How- ever, Odette (see Section 2 of Eason et al.29) found that the Russell–Brown model gave slightly better agreement with the experimental data. It should be added that further insight into the parameters controlling the hardening is obtained from CECs by combining microstructural data with mechanical property data (particularly hardening or yield stress increase) where the MD has been sub- tracted from the total measured increase. 4.05.4.7 Segregation to Grain Boundaries It was pointed out in Section 4.05.4.1 that the segre- gation of certain impurities to grain boundaries could cause nonhardening embrittlement. This phenome- non has received less attention than the hardening from the production of small clusters. Several reliable techniques (AES, FEGSTEM, and atom probe) exist with which grain boundary segregation may be not only observed, but also quantified,41 and there have been a number of critical studies that have both measured and modeled the segregation of impu- rity elements under irradiation.39,101–105 (Extensive experimental programs on long-term aging have per- mitted the accumulation of segregation data on a variety of model alloys and steels. It has been possible to interpret these data in terms of the simple McLean theory of equilibrium segregation (McLean, D. Grain Boundaries in Metals ; Clarendon Press: Oxford, 1957). The success of the McLean model in describing 170 Radiation Damage of Reactor Pressure Vessel Steels segregation in these alloys and steels indicates that segregation is generally thermodynamically con- trolled, and defect gradients have no effect.) The segregation of P and C to grain boundaries in irradiated materials has received greatest atten- tion.41,100 Overall P segregation increases with irradi- ation dose in all of the model alloys and steel types examined. The rate of P segregation under irradia- tion appears quite variable, both in different classes of steel and within a given class. It is possible that P segregation under irradiation is slower in welds than in the CGHAZ microstructure, because of the presence of additional traps for P in the welds. Other causes of variability are less consistently observed. The behavior of C is less consistent. In the model alloys and the CMn steels, grain boundary C gener- ally decreases with fluence, but in the MnMoNi steels C segregation may either increase or decrease. Desegregation of C appears more likely to be related to carbide precipitation in these materials with rela- tively high free C than merely to trapping of C at matrix defects. Quantifying the data has been attempted in sev- eral cases.99,100 The majority of models indicate that P is dragged to grain boundaries during radiation by the flux of irradiation-induced defects to sinks. Consis- tency between themodels anddata neednot necessarily confirm the validity of the model, as all have adjustable parameters, and no data set is large enough or coherent enough to test the models with much stringency. Importantly, a conclusion from the European Commission 5th Framework PISA programme was ‘‘On the basis of the observations made here and else- where, it appears unlikely that nonhardening embrit- tlement will influence RPV condition during normal operation for homogeneous MnMoNi steels (i.e., A508 Class 3, A533B, 22NiMoCr37) of Radiation Damage of Reactor Pressure Vessel Steels 171 need to take advantage of the greatly improved understanding of embrittlement mechanisms in DDRs that enable interpolation or extrapolation with improved confidence to a parameter space poorly covered by a given (national) surveillance database. A common feature of such DDRs is that they follow the same mechanistic framework (described in Section 4.05.4), but give different weights to the parametric dependencies of radiation damage that have been described in the previous sections. Properly describing the effect of flux on the embrittlement of both low Cu and Cu-containing steels has been subject to extensive debate (see Section 4.05.6). A further common fea- ture is that the DDRs have been refined as new surveillance data have become available, frequently with changes in the form of the equations in order to accommodate more sophisticated mechanistic under- standing. It is important to note that they have been developed to describe embrittlement under relatively low dose rate conditions that apply to specific steel types, that is, CMn or MnMoNi steels. Two classes of steels have been described by such DDRs, namely, CMn RPV steels employed in Mag- nox reactors and MnMoNi steels used in Western LWRs (primarily in the United States and Japan). The mechanistically based DDRs for CMn steels were developed in the 1980s while it was not before �1998 that the first such DDR was published describing embrittlement in US RPV steels. 4.05.5.2 DDRs for CMn Steels DDRs used to predict the embrittlement of the C–Mn steels used in the UK Magnox RPV had been mechanistically based from the 1980s.2 These predict radiation-induced changes in yield stress (hardening) or embrittlement (Charpy impact energy transition temperature or fracture toughness transi- tion temperature) as a function of radiation dose and temperature. The approach adopted has been set out in Jones and Bolton2 and Wootton et al.3,114 The advantage of this approach was that the derived rela- tionships could be used with confidence when limited extrapolation was required into regions of neutron dose, dose rate, or irradiation temperatures that were not specifically included in the surveillance database. It is important to note that it was the advances in understanding that enabled the adoption of a mecha- nistic approach (rather than adopting an empirical approach which had been followed in all other embrit- tlement correlations of this time). More specifically, the seminal work of Fisher and coworkers in the early 1980s50 assumed that changes in yield stress arose from the combined effects of irradiation damage clusters and copper precipitates. Subsequently,2,3,114 a two-term relationship was finally adopted2 to model both hardening (Ds) and embrittlement (DT40 J) and had the following form: DT40J or Ds 9>= >; ¼ Dcopper þ Dmatrix ½4� This relationship follows the model of Fisher and coworkers,50 where Dcopper represents the contribution of nanoscale copper precipitation to the property change and Dmatrix the contribution from matrix hard- ening arising from the production of point defect clus- ters by neutron irradiation. A further simplificationwas made in developing a DDR that could be applied to operational Magnox reactors. Namely, under the con- ditions of irradiation dose and temperature of interest there was no overaging; that is, the contribution to hardening or embrittlement from Cu cluster formation would reach a peak and then remain constant. Fur- ther, the hardening from Cu clusters could be repre- sented by a constant at all doses of interest, clearly a conservative assumption at doses before which the hardening from Cu clusters had reached a peak. On this basis, mechanistically based DDRs of the form DT40J or Ds 9>= >; ¼ B þ AFT ffiffiffiffi D p ½5� were adopted. In this equation, B represented the material-specific copper precipitate contribution to the property change, with the MD contribution being given by AFT√D. In this term, A is a material specific constant, D is the dpa dose, and FT is the irradiation temperature dependence factor.2,35 The fact that B is a constant independent of the measured bulk Cu level is consistent with the effect of the low final stress-relief temperature on reducing the variation in the Cumatrix between different materials (see Buswell and Jones70). DDRs were derived for the different RPV materi- als over the years. They were revised as and when new Charpy impact energy or tensile test data became available or following revisions to the neu- tron doses accrued by the surveillance specimens.114 For example, it was found that SMA welds are much more susceptible to the occurrence of intergranular fracture effects, withmanualwelds, plates, and forgings 172 Radiation Damage of Reactor Pressure Vessel Steels showing minimal effects. DDRs had to be developed that accommodated a nonhardening embrittlement mechanism. In addition, it was established that thermal neutrons could make a significant contribution to the irradiation damage in side-core locations, and that they were not conservatively covered by the DDRs.115,116 This conclusion was reached from an analysis of sur- veillance data from samples irradiated in locations in reactors with different levels of thermal fluxes and also from a well-controlled irradiation in a heavy water moderated reactor in Halden. It was established that to allow for extra displacements from low-energy recoils (�500 eV), a thermal neutron effectiveness fac- tor (k) needed to be introduced to modify the dose term in each material DDR. This meant that the gen- eral form of the two-term DDRs for both embrittle- ment and hardening (eqn [5]) became DT or Ds 9>= >; ¼ B þ AFT ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi Df þ kDt p ½6� In this equation, the definitions of B, A, and FT remained unchanged, but the single dose term, D, was replaced by (Dfþ kDt), where Df and Dt are the doses of fast dpa (redefined to be from neutrons of energy E> 1 keV) and thermal dpa (from neutrons of energy (a) -50 -50 0 0 50 50 100 100 150 150 200 200 250 250 300 300 M ea su re d Δ T 4 1J (� C ) Predicted ΔT41J (�C) (b) 0 0 20 40 60 80 100 5 � 1018 1 � 1019 2 � 10192.5 � 1019 3 � 10193.5 � 10191.5 � 1019 HSST02 SRM (0.17 wt% Cu) Prediction EONY C ha rp y sh ift Δ T 4 1J (� C ) Fluence n cm−2 (E > 1 MeV) Figure 15 (a) Predicted values for DT41 J for all PWR data, and (b) comparison of data for the DT41J shift for the reference plate HSST02 with the predictions of the EONY model. Radiation Damage of Reactor Pressure Vessel Steels 173 Eason et al.123 This is the most explicitly mechanistic DDR for MnMoNi steels produced to date, referred to as ‘EONY’ for convenience after the authors. The DDR, which is much more complex than the mild steel DDRs discussed above, is DT ¼ MF þ CRP ðin�FÞ ½7� where MF ¼ Að1�0:001718TiÞ 1þ6:13PMn2:47 � � ffiffiffiffiffiffiffiffiffiffiffiðfteÞp ½8� and A¼ 1.140� 10�7 for forgings, 1.561� 10�7 for plates, and 1.417� 10�7 for welds Ti¼ irradiation temperature (�F); P¼ bulk P (wt%); Mn¼ bulk Mn (wt%) ðftÞe¼ ft for f 4:39�1010 ncm�2 s�1 ft 4:39�1010 f � �0:259 forf< 4:39�1010 ncm�2s�1 8>< >: 9>= >; ¼ effective ðflux-correctedÞ fluence and CRP ¼ B 1þ3:77Ni1:191� �f ðCue;PÞ gðCue;Ni;fteÞ ½9� where B¼ 102.3 for forgings; 135.2 for plates in vessels manufactured by Combustion Engineering (CE); 102.5 for non-CE plates; 155.0 for welds; 128.2 for plates of the standard reference materials (SRMs) Cue ¼ 0 for Cu < 0:072wt% min½Cuactual;Cumax� for Cu > 0:072wt% � � ¼ effective Cu level ½10� in which Cuactual¼ bulk Cu level (wt%), Cumax¼ 0.243 for typical (Ni> 0.5) Linde 80 welds, and 0.301 for all other materials. (Equations [7] and [8] are in �F, reflect- ing units in the original reference. It is to be noted that �F are employed in USNRC regulatory guides, rather than SI units.) f ðCue;PÞ¼ 0 for Cu 0:072 ½Cue�0:072�0:668 for Cu> 0:072 and P 0:008 ½Cu� 0:072þ1:359ðP�0:008Þ�0:668 for Cu> 0:072 and P> 0:008 8>>>>>>< >>>>>>: 9>>>>>>= >>>>>>; ½11� gðCue;Ni;fteÞ ¼ 1 2 þ1 2 tanh log10ðftÞeþ1:139Cue�0:448Ni�18:120 0:629 ½12� Overall, this DDR or embrittlement correlation provided a good description of the database (see Figure 15). As with the UK DDRs, this DDR contains two separate terms, referring to CRP (or CEC) precipita- tion and to MD. In the US expression, however, both terms develop with fluence and have a more complex dependence on flux and composition. A threshold for the effect of P and Cu are in keeping with earlier DDRs, and the different limits on Cumax reflect the differing PWHT used by US fabricators. The square root dependence of the embrittlement from the MD term matches the expectation from mechanical property data on low Cu steels. The composition (a) 0 20 40 60 80 100 1016 1017 1018 1019 Flux > 4.4 � 1010 n cm−2s−1 (E > 1 MeV) Flux = 1�109 n cm−2s−1 (E > 1 MeV) Flux = 1�1010 n cm−2s−1 (E > 1 MeV) P re d ic te d Δ T 4 1J (� C ) Fluence (n cm−2) (E > 1 MeV) 1016 1017 1018 1019 1020 (b) Decreasing flux Increasing Cu Decreasing Ni 20 40 60 80 100 120 140 P re d ic te d s hi ft fr om C R P t er m Δ T 4 1J (� C ) Fluence (n cm−2) (E > 1 MeV) Cu = 0.25 wt%, Ni = 1.0 wt% Flux > 4.4�1010n cm−2s−1 Flux = 1�109n cm−2s−1 Cu = 0.25 wt%, Ni = 0.6 wt% Cu = 0.15 wt%, Ni = 1.0 wt% Cu = 0.15 wt%, Ni = 0.6 wt% Flux > 4.4�1010 n cm−2s−1 Increasing Cu,Ni Figure 16 (a) Schematic of the effect of flux and fluence on the magnitude of the matrix feature term, and (b) schematic of the CRP term showing the effect of key variables (low flux is 109 and all others are 1011 n cm�2 s�1). 174 Radiation Damage of Reactor Pressure Vessel Steels dependence of both the matrix and CRP term is broadly consistent with the understanding outlined in the previous section. The concept of fte is partic- ularly important as it both provides a means of allow- ing for flux effects and gives a threshold below which flux effects might be expected.63 These trends are further illustrated in Figure 16. Overall, for a Cu-containing steel (say 0.2–0.3wt% Cu), the MD becomes a significant fraction of the damage only at doses beyond the plateau in the shift from CRPs. This is consistent with the hardening from MD inferred from microstructural data. Carter et al. examined the effect of irradiation on microstructure on a high copper Linde 80 flux weld BW2 (0.25 wt% Cu, 0.62 wt%Ni, 0.017wt% P),125 and concluded that out of a total hardness of DHvtot 40� 6 the hardness from MD was DHvMatrix 5–10 VPN. In Section 4.05.2, it was described how irradiation also caused a drop in the Charpy USE. It is to be noted that Eason et al.22 used the US surveillance power reactor database to investigate the dependence of the USE drop (DUSE) on a number of variables. They demonstrated that there was a strong correla- tion between the DUSE and Charpy TTS at 30 ft-lbs. Eason et al. derived a detailed set of equations that allowed the DUSE to be determined from the TTS for a number of product forms. 4.05.5.4 Japanese Embrittlement Correlations The first embrittlement correlation for the TTS of the Japanese RPV materials, JEAC 4201, was issued in 1991. Additional surveillance data have been com- piled since 1991 and in 2002 the Japanese electric power utilities started a project with CRIEPI to develop a new mechanistically based embrittlement correlation.126 Soneda and coworkers have adopted a two- step approach to developing a new correlation method.65,126,127 In the first step, the microstructural effects due to radiation damage are modeled, and the mechanical property changes engendered by such change are detailed. The microstructural changes, namely, the formation of solute atom clusters and MD features, due to irradiation are modeled using the following equations: @CSC @t ¼ x3 CmatCu þ e1 � � DCu þ e2 � � CMD þ x8 CavailCu DCu 1þ x7C0Ni � �� �2 ½13� @CMD @t ¼ x4F 2t x5 þ x6C0Ni � �2f� @CSC @t ½14� @CmatCu @t ¼ �vSC@CSC @t � v0SCCSC ½15� vSC ¼ x2 CavailCu DCu � �2 tr ½16� v0SC ¼ x1CavailCu DCu ½17� CavailCu 0 CmatCu CsolCu CmatCu � CsolCuCmatCu > CsolCu ( ½18� DCu ¼ DthermalCu þ DirradCu ¼ DthermalCu þ of� ½19� where Csc and CMD are the number densities of solute atom clusters and MD features, CmatCu and C 0 Ni are the 0 0.0E + 00 2.0E + 19 Fluence (n cm−2) 4.0E + 19 6.0E + 19 8.0E + 19 1.0E + 20 20 DT (º C ) 40 60 80 100 120 0.08Cu (CEC) 0.15Cu (CEC) 0.15Cu (MD) 0.24Cu (MD) 0.24Cu (CEC) 0.08Cu (MD) Figure 17 Partitioning of the total embrittlement of the materials with different copper content into copper-related contribution and matrix damage contribution in the CRIEPI correlation. Reproduced from Hiranumu, N.; Soneda, N.; Dohi, K.; Ishino, S.; Dohi, N.; Ohata, H. Mechanistic modeling of transition temperature shift of Japanese RPV materials. In Presented at the 30th MPA-Seminar in Conjunction with the 9th German-Japanese Seminar, Stuttgart, Germany, 2004. Radiation Damage of Reactor Pressure Vessel Steels 175 bulk chemical contents of Cu and Ni, DCu is the Cu diffusivity, f is the dose rate, t is the irradiation time, and tr is the relaxation time, respectively. Equations [13 and 14] represent the time evolu- tion of solute atom clusters and the MD clusters, respectively (see Hiranumu et al.126 for a full descrip- tion of the equations). In eqn [13], it is to be noted that solute atom clustering occurs with MD features as the nuclei. This process can occur without Cu atoms but is accelerated by their presence. In eqn [14], the formation of MD features is affected by the irradiation temperature and also the bulk Ni content. Equation [15] models the depletion of the matrix Cu content because of the formation and growth of Cu-enriched solute atom clusters. Note that the depletion of the matrix Cu reduces the formation rate of Cu-enriched solute atom clusters. Equation [19] gives an expression for the diffusivity of Cu atoms, which combines terms from both irradiation- induced vacancies and thermal vacancies. Mechanical property changes are correlated with the microstructural changes using the following equations: DTSC ¼ x16 ffiffiffiffiffi Vf p ¼ x16 ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi x15f CmatCu ;CSC � � g C0Ni � �þ hðftÞq ffiffiffiffiffiffiffiffiCSCp ½20� f CmatCu ;CSC � � ¼ x11C 0 Cu � CmatCu CSC þ x12 ½21� g C0Ni � � ¼ 1þ x13 C0Ni� �x14 � �2 ½22� hðftÞ ¼ x9ð1þ x10DSCÞft DSC � DCu ½23� DTMD ¼ x17 ffiffiffiffiffiffiffiffiffi CMD p ½24� DT ¼ ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi ðDTSCÞ2 þ ðDTMDÞ2 q ½25� where DTSC and DTMD are the contributions of solute atom clusters and MD features, which are calculated using eqns [20 and 24] as functions of CSC and CMD, respectively. In calculating the contri- bution of solute atom clusters, an empirical model, in which the TTS is proportional to the square root of the volume fraction of solute atom clusters, is used. The average volume per cluster, which is necessary for calculating the volume fraction, is modeled using eqns [21–23], which take into account the effect of chemical composition and the growth of the clusters during irradiation. The Greek characters in the above equations are coefficients, and were optimized using the surveillance database of Japanese commercial reactors.126,127 The partitioning of the total embrit- tlement between that due to copper clusters and MD features is shown in Figure 17. It can be seen that the MD has a weak dependence on the Cu level of the steel. Figure 18 shows a comparison between the calcu- lated and measured TTS. The standard deviation of the prediction error is smaller than that of the other correlation equations used in Japan and in the United States, as shown in Figure 10. When a plant-specific adjustment is applied to the initial transition temper- ature, the standard deviation of the prediction error becomes much smaller and is as low as 6 �C. A practical output of this approach is the develop- ment of a new embrittlement correlation method for Japanese RPV steels, and this method has been adopted in the JEAC 4201-2007. Thus, this study is a good example of how the understanding of a funda- mental mechanism can be applied in a real-world engineering application. 4.05.5.5 Summary It is clear from the discussion above that there has been successful development of mechanistically based DDRs for both CMn and MnMoNi steels. 0 0 20 -20 40 60 80 100 120 140 Method Std- dev. Mean error JEAC4201 RG1.99r2 EWO E900-02 CRIEPI CRIEPI adj 11.9 -1.3 -1.9 2.8 2.3 0.7 0.15.4 9.4 11.7 10.4 15.4 -20 20 Measured value (�C) P re d ic tio n (º C ) 40 60 80 100 120 140 -2s +2s w/o adjustment w adjustment 1:1 Figure 18 The comparison of predicted and measured transition temperature shifts. Plant-specific adjustment is performed by offsetting the initial values. Reproduced from Soneda, N. In Materials Issues for Generation IV Systems; Springer: The Netherlands, 2008; pp 254–262; NATO Science for Peace and Security: Physics and Biophysics, ISBN 18746500. 176 Radiation Damage of Reactor Pressure Vessel Steels Different DDRs have been developed in different countries to describe the hardening and embrittle- ment of the various RPV steels. The inevitably approximate nature of the DDR expressions, the lim- ited variation of different parameters in each surveil- lance database, and the limited amount of surveillance datamean that the effects of many parameters must be implicit. Different irradiation and compositional var- iable ranges in different surveillance schemes may contribute significantly to the forms of the DDRs and the strength of different dependences. The lim- itations in the form of the DDRs and the R&D into outstanding issues are the subject of the next section. 4.05.6 Current Issues in the Development of DDRs The DDRs for MnMoNi steels presented in the last section provide convincing examples of the applica- tion of fundamental insight to the prediction of changes in mechanical properties of operating RPVs due to radiation damage. Mechanistic understanding is continually developing as research continues and more data are obtained. Advances may lead to modifications in the form, or the values of, the fitting parameters. The major topics are the following: � The effect of flux � The role of Ni, Mn, and Si � The possibility of new mechanisms at fluences beyond the range for which there are data in the current surveillance databases There are two aspects of the effect of flux: first, the prediction of embrittlement at low fluxes and second, improvements in the general description of the effect of flux on embrittlement. It was described in the previous section that recent BWR data from the SSP capsules have greatly expanded the available BWR data, leading to an improved shift model. Carter et al.128 pointed out that, although this pro- vides a better description of BWR plate data, the model still tends to underpredict the embrittlement of BWR welds for measured DT41 J greater than 60 �C. This suggests that there may be further improvements necessary in the description of embrit- tlement in the low flux range. Indeed, there may be general improvements in the description of flux. Odette considers that there is a systematic flux effect in the range of 0.8–8� 1011 n cm�2 s�1 E> 1MeV in the IVAR database which is not predicted by the EONY model.30 Further analysis of the IVAR data- base may lead to improvements in the description of the flux dependence of embrittlement at both low (surveillance) fluxes and high (MTR) fluxes. The DDRs for MnMoNi steels discussed in the previous section really apply to only steels with Ni< 1.3 wt%. High Ni welds have been used in a limited number of civil PWRs, notably VVER 1000 reactors. High Ni welds were selected because vessel designers wished to take benefit from the greater hardenability and superior SOL properties (com- pared to lower Ni steels). At present the response of Cu-containing high Ni steels to irradiation doses of Radiation Damage of Reactor Pressure Vessel Steels 177 doses, Mn, Ni, and Si could form a new phase in RPV steel.47,130 This late-blooming phase (LBP) would produce an additional increment of hardening at high fluences, that is, late-onset embrittlement or anomalous hardening at high doses. If this is the case, then the DDRs described in the previous sec- tion may become nonconservative. Recently, there have been an increasing number of observations of MnNiSi clusters in irradiated low Cu steels (see, e.g., Auger et al.131 and Soneda et al.127), and there is intensive research aimed at establishing whether NiMnSi clusters represent segregation to small microstructural features (thereby lowering interfacial or strain energies) or represent precipitates of a distinct Ni–Si–Mn-enriched phase that is thermo- dynamically favored at RPV operating temperatures and RPV steel compositions. 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Leonard Oak Ridge National Laboratory, Oak Ridge, TN, USA Published by Elsevier Ltd. 4.06.1 Introduction 181 4.06.2 Niobium and Nb-Base Alloys 182 4.06.2.1 Introduction and History of Nb and Nb Alloys 182 4.06.2.2 Radiation-Induced Swelling of Nb and Nb-Base Alloys 183 4.06.2.3 Mechanical Properties of Irradiated Nb and Nb Alloys 185 4.06.3 Tantalum and Ta-Base Alloys 188 4.06.3.1 Introduction and History of Ta and Ta Alloys 188 4.06.3.2 Irradiation-Induced Swelling of Ta and Ta-Base Alloys 189 4.06.3.3 Mechanical Properties of Irradiated Ta and Ta-Base Alloys 190 4.06.4 Molybdenum and Mo-Base Alloys 194 4.06.4.1 Introduction and History of Mo and Mo Alloys 194 4.06.4.2 Irradiation-Induced Swelling and Physical Property Changes in Mo and Mo-Base Alloys 194 4.06.4.3 Mechanical Properties of Irradiated Mo and Mo Alloys 197 4.06.5 Tungsten and W-Base Alloys 206 4.06.5.1 Introduction and Irradiated Properties Database for W and W Alloys 206 4.06.5.2 Irradiation-Induced Swelling and Physical Property Changes in W and W Alloys 206 4.06.5.3 Irradiated Mechanical Properties of W and W Alloys 207 4.06.6 Outlook 209 References 211 Abbreviations bcc Body-centered cubic C-103 Nb–10Hf–1Ti alloy Cb-752 Nb–10W–2.5Zr alloy DBTT Ductile–brittle transition temperature FS-85 Nb–10W–28Ta–1Zr alloy HP High purity JIMO Jupiter icy moons orbiter LCAC Low-carbon arc cast NERVA Nuclear experiment for rocket vehicle applications ODS Oxide dispersion strengthened RIS Radiation-induced segregation SNAP Systems nuclear auxiliary power T-111 Ta–8W–2Hf alloy TZM Mo–0.5Ti–0.1Zr alloy UTS Ultimate tensile strengths UWMAK-III University of Wisconsin Madison fusion reactor concept Symbols D Thermoelectric power n Neutron particle T Temperature Tirr Irradiation temperature Tm Melting temperature a Alpha particle DV/V Volume fraction swelling w Fluence 4.06.1 Introduction Refractory metals and alloys offer attractive and promising high-temperature properties, including high-temperature strength, good thermal conductiv- ity, and compatibility with most liquid metal cool- ants, many of which are suitable for applications in nuclear environments. Though many of the refrac- tory alloys have been known for more than 60 years, there are significant gaps in the materials property database for both unirradiated and irradiated 181 182 Radiation Effects in Refractory Metals and Alloys conditions. In addition, significant issues related to low-temperature irradiated mechanical property degradation at even low neutron fluences restrict the use of refractory metals. Protection from oxidizing environments also restricts their use, unless suitable protection or a liquid metal coolants is used. Much of the early research on refractory metal alloys was centered on applications in aerospace as well as cladding and structural materials for fission reactors, with particular emphasis on space reactor applications. Reviews concerning the history of these programs and the development of many of the alloys whose irradiated properties are discussed in this chapter can be found elsewhere.1–5 Due to cancellations and reintroduction of new mission cri- teria for these space reactor programs, the materials database shows similar waves in the gains of intellec- tual knowledge regarding refractory alloy and irra- diated property behavior. Unfortunately, as seen in the subsequent sections of this chapter, much of the irradiated property database for refractory metals con- sists of scoping examinations that show little overlap in either material type, metallurgical conditions (i.e., grain size, impurity concentrations, thermomechanical treatments), radiation conditions (i.e., spectra, dose and temperature), or postirradiation test conditions or methods. The irradiation behavior of body-centered cubic (bcc) materials is known. Irradiation-induced swelling because of void formation in the material lattice is typical for temperatures between 0.3 and 0.6Tm, where Tm is the melting temperature. Maximum swelling in refractory metals is Radiation Effects in Refractory Metals and Alloys 183 governed the periodic scientific examinations of refractory alloys. The historical examination of Nb and its alloys is typical of this, with early studies of the irradiation properties exploring the potential uses of these alloys in fusion energy and fission type space reactor. While these alloys have favorable properties, such as elevated temperature capability and compati- bility with liquid alkali metals for energy applications, and attractive physical properties such as thermal conductivity, much of the work on Nb and Nb-base alloys has examined the nonirradiated properties. It is worth putting into perspective the relatively small commercial market for niobium-base alloys. Approximately, 75% of all niobium metal is used as minor alloying additions in steel, and only 1–2% is produced in the form of niobium-base alloys. The total market for niobium-base alloys in the mid-1990s was 100 ppm (DBTT of high purity Nb and Nb alloys is near 73K14,15). The effects of oxygen and carbon were less severe, but influential at levels of 100 ppm and greater. The effect of nitrogen on embrittlement also appears to be as severe as that of oxygen, though some uncertainty exists as to whether solubility limits have been exceeded in the data.1 The effect of interstitial impurities on the irradiated properties of Nb and Nb-base alloys is significant and has been examined, though the overall database for irradiated properties is limited. The interplay between the radiation-created defects and the interstitial impurity elements was investi- gated by Igata et al.16 for pure Nb (70wppm oxygen and 30wppm nitrogen) irradiated to 3.4� 1020 n cm�2 (E> 1MeV) at temperatures below 413K and post- irradiation annealed up to 973K. Increases in yield strength over the as-irradiated values following annealing were measured at 473 and 673K, attrib- uted to the interplay of the defect clusters trapping oxygen and nitrogen atoms, respectively. Above 773K, no difference between the annealed and as-irradiated yield stress was observed. Hautojärvi et al.17 and Naidu et al.18 examined the interaction between vacancies and interstitial 14 12 10 8 6 4 2 0 0.00 0.04 DV /V (% ) 0.08 0.12 0.16 0.20 Oxygen concentration (at.%) 30 MeV 58Ni+ irradiation Tirr=1225 K, 50 dpa 1240 K 1225 K 1200 K 1240 Tirr (K) 1220 1200 Nb + 2.4 at.% Ta Nb + 2.4 at.% Ni Nb + 2.3 at.% Fe Nb + 2.3 at.% Ti Nb + 2.4 at.% Zr Nb + 2.4 at.% Hf Nb + 2.4 at.% Mo Nb + 2.4 at.% V Figure 1 The dependence of void volume fraction (DV/V ) in 3MeV 58Ni+ ion-irradiated Nb on the concentration of oxygen and dilute solute additions. Reproduced from Loomis, B. A.; Gerber, S. B. J. Nucl. Mater. 1983, 17, 224–233. 184 Radiation Effects in Refractory Metals and Alloys impurities in irradiated Nb through positron annihi- lation studies. In high-purity Nb, vacancy clustering within the collision cascades is observed, starting as low as 160K, with vacancy migration peaking around 250K, but in materials with higher hydrogen content, the vacancy migration stage shifts to temperatures close to 400K.17 The irradiation exposure at inter- mediate temperatures (0.3–0.6Tm) can lead to void swelling, irradiation creep, and helium embrittlement through processes involved in (n,a) reactions or impurity gas atoms. Naidu et al.18 examined the effect of He and its interaction with vacancies in pure Nb, leading to the development of bubbles through a-irradiated specimens. At temperatures between 623 and 1023K, bubble growth occurs through the addition of He atoms and vacancies, followed by migration and coalescence at higher temperatures, eventually leading to the annealing out of the He bubbles and vacancy complexes above 1173K.18 The irradiation-induced swelling of pure Nb gen- erally appears at temperatures between 673 and 1323K with peak swelling near 873K (0.32Tm), though these limits are not clearly defined and are based on the very limited data available, compiled by Wiffen19 and Pionke and Davis.1 A maximum swelling of 4.8% fol- lowing irradiation to 2.5� 1022 n cm�2 at 858K was reported.1 However, the magnitude of swelling shows considerable scatter in the literature, possibly reflect- ing the influence of impurity concentrations and dif- ferences in irradiation conditions and microstructural interpretation of the materials.19 Fischer20 reported that void concentration increased four to seven times for a fourfold increase in flux for the same total fluence. This produced a reduction in void size with flux and therefore a reduction in the total swelling. Loomis and Gerber21–23 examined the influence of oxygen and substitutional binary alloy additions on the swelling of 3MeV 58Ni+ ion-irradiated Nb up to�50 dpa. Void formation and characteristics in size and morphology were found to be dependent on temperature, oxygen concentration, and the type of substitutional alloy addition. The average void diam- eter was found to increase with temperature as well as oxygen up to 0.02 at.%. Higher oxygen concentra- tions resulted in a decrease in void diameter to 0.1 at.% O, above which void diameters showed no significant changes. The number density of voids was found to decrease with temperature, but increase with oxygen concentration to �0.06 at.%, above which the number density showed no significant change. As the volume fraction of swelling (DV/V) is propor- tional to both the void number and the cube of the void diameter, the volume fraction is observed to increase with temperature and oxygen concentration to �0.04 at.%, followed by a decrease and plateau of the volume fraction above 0.1 at.%. The dependence of DV/V on temperature and oxygen concentration is illustrated in Figure 1. Microstructural examina- tion revealed an ordering of the voids into a lattice- type structure in the material irradiated at 1050K to �40 dpa and oxygen concentration �0.039 at.% oxygen. The higher temperature of the maximum swelling as compared to the neutron irradiation data is believed to be associated with the higher displace- ment damage rate of the ion-bombarded material,19 though the higher impurity levels may also provide an influence. The effect of dilute (�2.4 at.%) substitutional alloy addition on the swelling of 0.06 at.% oxygen- doped Nb was also examined for 3MeV 58Ni+ ion irradiation at 1225 K. The DV/V was determined to increase through the addition of Ta, but decreased with increasing effectiveness by the addition of Ti, Zr, V, and Hf. The addition of the reactive alloying elements to Nb suppresses void formation through the gettering of interstitial impurities that act as void nucleation sites. The DV/V was determined to be unaffected by the addition of Ni or Fe. The depen- dence of DV/Von temperature, oxygen, and substitu- tional addition is also shown in Figure 1. Radiation Effects in Refractory Metals and Alloys 185 Swelling in Nb–1Zr has been examined, though only scattered data are available in the examination of temperature and flux dependence. The available swelling data on Nb–1Zr, compiled by Powell et al.24 and Watanabe et al.25 presented in Figure 2, show the lack of data on the temperature range in which peak swelling appears. The swelling data shown in the figure were measured through electron microscopy, with the exception of the data by Powell et al.24 and Wiffen.26 Alloy impurity chemistry, in addition to interpretation and measurement error, may account for the scatter associated with the lower tempera- tures. The work of Watanabe et al.25 and Garner et al.27 indicates that irradiation-induced swelling is dependent on the thermomechanical history of the material. In that material, cold-working followed by solution anneal and aging exhibited swelling, while material not given the preirradiated cold-working showed some densification. The changes in density of the material are dependent on the phase-related transformations involving precipitation. Swelling in Nb–1Zr appears to be centered over a more narrow temperature range than in Nb, with a peak near 1073K that is higher than that of the pure metal. While the addition of Zr toNb appears to delay nucleation of voids to higher temperatures, the voids that form are of larger size than those appearing in pure Nb under comparable conditions. For example, 2.5 2.0 1.5 1.0 0.5 0 20 40 dpa 60 80 100 400 600 800 1000 1200 DV /V (% ) Irradiation temperature (K) 1400 Powell et al.24 Jang and Moteff145 Sprague et al.146 Wiffen28 Garner et al.27 Wiffen26 Michel and Smith147 Watanabe et al.25 Figure 2 Swelling as a function of irradiation temperature and dose for neutron-irradiated Nb–1Zr from available literature compiled by Powell et al.24 and Watanabe et al.25 following irradiation to 2.5� 1022 n cm�2 (E> 0.1 MeV) at 1063K, the diameter, concentration, and volume fraction of voids in Nb–1Zr was 57.5 nm, 1.8� 1020m�3, and 2.2%, respectively,1 whereas under similar conditions, the same void parameters in pure Nb were 18.6 nm, 2.8� 1021m�3, and 1.04%. While void formation and swelling in Nb and Nb–1Zr occurs, the total swelling is generally 10 dpa.3 The addition of Ti to Nb was found to increase void resistance and has been found to suppress void formation in V at concentrations as low as 3%.29 The combination of reactive alloy ele- ments and Nb in the C-103 alloy may suggest a greater void formation resistance than in pure Nb and Nb–1Zr. 4.06.2.3 Mechanical Properties of Irradiated Nb and Nb Alloys Little coverage of the changes in mechanical proper- ties following irradiation has been given to Nb and Nb alloys, with the majority of the data for tempera- tures below 800K. Some preliminary experimental work on the irradiated mechanical properties of Nb alloys Cb-752 (Nb–10W–2.5Zr)30 and FS-85 (Nb– 10W–28Ta–1Zr)31 is available. However, these alloys are not commercially produced and have shown indications of thermal aging instabilities, leading to grain boundary embrittlement.12,32,33 The irradiated mechanical properties of these alloys show similar radiation hardening as in the pure metal, but with mechanical properties more sensitive to thermal aging conditions. The bulk of the irradiated mechan- ical properties data is for the Nb–1Zr alloy as well as the pure metal, and is covered in this review. The irradiated mechanical properties of Nb and Nb–1Zr are strongly governed by irradiation temper- ature, determining whether the mechanical properties are controlled by dislocation loops or a combination of loops and voids in the microstructure. As cavity for- mation can be delayed or suppressed by higher irradi- ation temperatures in Nb–1Zr, mechanical property comparisons between the alloy and the base metal will reflect their irradiated microstructure. For Nb and Nb–1Zr irradiated to 3� 1022 n cm�2 at �728K, the pure metal contains both dislocation loops and voids, while the alloy exhibits no void formation.19 A compar- ison of the tensile properties of Nb and Nb–1Zr irra- diated under similar conditions is shown in Figure 3. The irradiated strength of both materials shows an increase in tensile strength above the unirradiated Ultimate Niobium Nb–1Zr Ultimate Ultimate Total Total Uniform Uniform Yield Yield and ultimate Yield Yield Control 3.0 � 1022n cm–2 At 460 �C 120 100 80 60 40 20 0 60 40 20 0 0 200 400 600 0 Test temperature (�C) 200 400 600 Total Total Uniform 800 600 400 200 0 S tr es s (1 00 0 p si ) S tr es s (M P a) E lo ng at io n (% ) Control 3.7 � 1022n cm–2 At 450 �C Figure 3 Comparison of tensile properties between Nb and Nb–1Zr tested under similar irradiation conditions. Reproduced from Wiffen, F. W. In Refractory Alloy Technology for Space Nuclear Power Applications, CONF-8308130; Cooper, R. H., Jr, Hoffman, E. E., Eds.; Oak Ridge National Laboratory: Oak Ridge, TN, 1984; pp 252–277. 4 200 0 0 400 600 S tr es s (M P a) 800 8 Elongation (%) 12 16 3.0 � 1026n cm–2 7.5 � 1024n cm–2 Tirr= 328 K Niobium Ttest= 298 K Tirr= 733 K Figure 4 Comparison of tensile curves between Nb irradiated at 328 and 733K. Yield instability is seen at 328K due to channeling of deformation dislocations through the irradiated dislocation loop structures. The higher irradiation temperature resulted in the development of small voids providing a barrier to dislocation movement. Reproduced from Wiffen, F. W. In Refractory Alloy Technology for Space Nuclear Power Applications, CONF-8308130; Cooper, R. H., Jr, Hoffman, E. E., Eds.; Oak Ridge National Laboratory: Oak Ridge, TN, 1984; pp 252–277. 186 Radiation Effects in Refractory Metals and Alloys condition, with Nb–1Zr showing a greater sensitivity to irradiation. As the mechanical properties of Nb–1Zr are dominated by the dislocation loop structures, yield instability is observed in the material, leading to the early onset of necking. This results in Radiation Effects in Refractory Metals and Alloys 187 relative strength increase over the unirradiated condi- tion. The higher irradiation temperature produced voids in themicrostructure, providing additional obsta- cles to deformation and higher uniform elongations and modest work hardening. Little is known with regard to the aging properties of Nb–1Zr or the combined thermal and radiation effects. The addition of 1 wt% Zr to Nb creates a dispersion-strengthened alloy, in which the Zr com- bines with interstitial impurities creating fine preci- pitates throughout the material. The development of these fine precipitates on aging at 1098K can increase the tensile strength between 50 and 100MPa over the annealed condition and provide an effective strength- ening greater than that observed through modest irradiation31 (Table 1). Irradiation of Nb–1Zr to 0.9 dpa at 1098K showed a modest increase in yield and ultimate tensile strength to 135 and 192MPa, respectively, over the annealed condition. This increase in tensile strength either through aging or irradiation results in a cor- responding decrease in uniform elongation from 15% to 3.5% and total elongation from 25% to 15%. Aging at temperatures above 1098K produces little effective hardening as the precipitates coarsen in the microstructure.33 Irradiation to 0.9 dpa at 1248 and 1398K of Nb–1Zr showed only a modest increase in the yield strength over the aged and annealed specimens, though ultimate tensile strength and elon- gation were unchanged or less. Irradiation to 1.88 dpa at 1223 K resulted in weaker tensile properties Table 1 Tensile property comparison illustrating the effec of Nb–1Zr Test/aged/irradiated temperature (K) Yield strength (MPa) Ultimate (MPa) As-annealed condition 298 185.00 281.00 1073 102.33 191.00 1223 89.00 215.00 1373 83.00 157.00 1100h aged 1073 185.67 248.67 1223 111.67 156.67 1373 82.50 130.50 Irradiated: 2.04�1021 n cm�2 (E>0.1MeV), 0.93dpa 1073 134.50 192.50 1223 149.50 182.00 1373 102.00 127.50 Irradiated: 4.13�1021 n cm�2 (E>0.1MeV), 1.88dpa 1223 104.50 166.00 Source: Busby, J. T.; Leonard, K. J.; Zinkle, S. J. Effects of neutron irr Oak Ridge National Laboratory: Oak Ridge, TN, Dec 2005. compared to the 0.9 dpa sample, believed to be due to further precipitate coarsening. The time under irradiation conditions for the 1.88 dpa sample was near 1100 h and produced similar tensile properties as that of the aged-only material. As discussed in the preceding paragraphs, the irradiated properties of Nb and Nb–1Zr are governed by their microstructure and are influenced by temper- ature, displacement damage rate, and neutron spec- trum. The tensile properties of neutron-irradiated Nb–1Zr for damage levels between 0.1 and 5 dpa (Horak et al.34 and Wiffen35) summarized by Zinkle and Wiffen3 are shown in Figure 5. At temperatures below 800K, a large increase in the tensile strength from irradiation is observed with the corresponding low uniform elongations. At higher temperatures, uniform elongation increases because of the presence of voids in the microstructure. However, the data plotted in Figure 5 show uniform elongations remain- ing low up to 1100K, while radiation hardening is relatively moderate, suggesting that impurities are the source of the reduced elongation values. No irradiated fracture toughness data exist for Nb or Nb–1Zr, though comparisons can be made from the larger irradiated vanadium alloy database, in which fracture toughness embrittlement becomes a concern when tensile strength exceeds 600–700MPa and there- fore at temperatures below 400K for Nb–1Zr.36 How- ever, if a conservative value is assigned to the critical stress to induce cleavage fracture of �400MPa (40% lower than that observed in vanadium alloys), ts of aging and irradiation on the mechanical properties tensile strength Uniform elongation (%) Total elongation (%) 18.60 34.60 16.03 23.90 15.05 23.70 10.50 44.80 8.07 13.27 8.07 13.27 6.80 28.65 7.10 18.55 3.55 17.80 9.55 33.65 8.10 20.90 adiation on refractory metal alloys, ORNL/LTR/NR-PROM1/05-38; 200 400 600 800 1000 Temperature (K) 1200 1400 1600 00 100U lti m at e te ns ile s tr en gt h (M P a) 200 300 400 500 600 700 5 10 15 20 25 30 35 U ni fo rm e lo ng at io n (% ) Irrad. UTS Unirrad. UTS Irrad. eU Figure 5 Unirradiated and irradiated (0.1–5dpa) ultimate tensile strengths (UTS) and uniform elongation (eU) of Nb–1Zr. Irradiated data represented by solid symbols and unirradiated by open symbols. Figure reprinted with permission from Zinkle, S. J.; Wiffen, F. W. Radiation effects in refractory alloys. In STAIF 2004, AIP Conference Proceedings; El-Genk, M. S., Ed.; 2004; Vol. 699, pp 733–740. Copyright 2004, American Institute of Physics. 188 Radiation Effects in Refractory Metals and Alloys fracture toughness becomes a concern at tempera- tures below 800K for Nb–1Zr.3 While irradiated ten- sile strength above 800K is close to the unirradiated values, uniform elongation values remain low until irradiation temperatures >1000K. Therefore, a con- servative approach towards engineering design needs to be taken with this alloy. The mechanical properties of irradiated refractory alloys can be influenced by the formation of He developed through the (n,a) reactions, leading to the grain boundary formation of bubbles and the eventual embrittlement of the material. Some scoping investiga- tions on the effect of He on the irradiated mechanical properties of Nb–1Zr have been performed. Wiffen37 investigated the high-temperature mechanical proper- ties of 50MeV a-irradiated Nb–1Zr. In tensile tests conducted at 1273 and 1473K, no significant effect of He on the strength or ductilityofNb–1Zrwas observed for samples containing 2–20 appm He. Later analysis of the creep ductility reductions was found to be dependent on the observed precipitate phase develop- ment through the pick-up of oxygen during implan- tation.38 He-implanted Nb–1Zr through 100MeV a-irradiations at 323 and 873K by Sauges and Auer39 found no significant effect on ductility up to 80 appm He. Wiffen19 observed that uniform elongations stayed around 1% between test temperatures of 723 and 1073K on 130 appm 10B doped Nb–1Zr irradiated in a fast reactor between 723 and 1223K up to 6� 1022 n cm�2. Thesewere slightly higher than those observed in undoped material; this is believed to be due to the formation of He bubbles in the grains of the material acting similar to voids in generating obstacles to dislocation channeling. In general, no detrimental effects on mechanical properties were reported for accelerator-injected He between 1273 and 1473K for He concentrations Radiation Effects in Refractory Metals and Alloys 189 However, the use of Ta–10W in space reactor appli- cations where liquid alkali coolants are considered was unacceptable because of the lack of oxide getter- ing elements such as Hf that form stable dispersion- strengthened structures. The T-111 (Ta–8%W–2% Hf) alloy, with its demonstrated compatibility with liquid alkali metals and improved strength over pure Ta while retaining ductility and weldability, has been a lead candidate alloy in space reactor systems since the 1960s.45 Though a considerable effort has been made on the Ta–10W and T-111 alloys, the irradiation properties database is very small. Irra- diated mechanical property behavior follows typical bcc alloys in which radiation hardening effects including limit ductility appear and are expected at temperatures �0.3Tm (987 K).3 4.06.3.2 Irradiation-Induced Swelling of Ta and Ta-Base Alloys Swelling data for Ta and its alloys are limited to a few studies.19 Void formation in pure Ta was experimen- tally observed through TEM examination of material irradiated to 2.5� 1022 n cm�2 (E> 0.1MeV) at tem- peratures between 673 and 1273K.46 An empirical estimation of the bulk swelling taken from microstruc- tural void size density data of that study is shown in Figure 6. Void concentrations in the material were highest at the peak swelling temperature anddecreased with higher irradiation temperature with an associated increase in cavity size. Orderingof the voids at the peak 3.0 2.5 2.0 1.5 1.0 0.5 0.0 Neutron fluence 2.5 � 1022n cm–2 (E > 0.1 MeV) 200 400 600 800 Irradiation temperature (K) 1000 1200 1400 1600 ΔV /V (% ) Bates and Pitner47 Wiffen46 Figure 6 Swelling data for pure Ta measured through microstructural void density measurements by Wiffen46 and from immersiondensitymeasurementsbyBates andPitner.47 swelling condition was reported to occur along the {110} planes in the bcc structure. A subject of consid- erable theoretical debate, the mechanisms of void ordering that have appeared in bcc and fcc metals have been examined,48–50 since the first reported occurrence in irradiated Mo.51 Disordered void struc- tures in the microstructure of the higher temperature irradiated Ta appear as the size of the voids increase, though some rafting, or grouping, was reported.46 The swelling data of Wiffen46 derived from micro- structural analysis correlate well with the immersion density data of Bates and Pitner47 (Figure 6), from which an empirical equation for percent swelling as a function of temperature, T (K), and fluence, F (in units of 1022 n cm�2, E> 0.1MeV), was developed, which is as follows: DV V ¼ ðFÞ0:4f1:69 exp½�ð0:018T � 16:347Þ2=a�g where a ¼ 14:87þ 44:57 exp½0:09ðT � 1338:71Þ� 1þ exp½0:09ðT � 1338:71Þ� ½1� The broader width of the swelling peak as a function of irradiation temperature for the calculation repre- sented by eqn [1] compared to the microstructural data of Wiffen46 is believed to be associated with errors in the accurate irradiation temperature of these early measurements. Experimental evidence of decreased swelling at higher fluences was reported by Murgatroyd et al.52 and attributed to the transmu- tation of Ta to W, resulting in a shift in the lattice constant. Similar effects have been more closely examined in Mo and TZM alloys, and attributed to impurity segregation at void surfaces leading to shrink- age of the voids.53 Swelling measurements in Ta–10W and T-111 alloys are limited specifically to work by Wiffen, from which a later summary was given.19 For irradia- tions at 723 and 873K to a fluence of 1.9� 1022 n cm�2 (E> 0.1MeV), no swelling in T-111 was observed, though a possible densification of up to 0.36% may have occurred as evidenced in length measurements. In companion irradiations to that of pure Ta already discussed, involving irradiations to 4.4� 1022 n cm�2 (E> 0.1MeV) at temperatures between 698 and 1323K,46 samples of Ta–10W were included with postirradiation examination involving TEM analysis. The microstructure of the irradiated Ta–10W con- tained fewer voids than the companion Ta samples, with a lower swelling assumed in the Ta–10W alloy but with values not accurately quantifiable.19 190 Radiation Effects in Refractory Metals and Alloys 4.06.3.3 Mechanical Properties of Irradiated Ta and Ta-Base Alloys The overall mechanical property data for irradiated Ta and Ta-base alloys are very limited, with most studies involving irradiation at temperatures 240 200 160 120 80 40 0 60 40 20 0 Control 2.5 dpa at 688 K 2.5 dpa at 913 K Control 2.5 dpa at 913 K 1.97 dpa at 663 K Ultimate Ultimate Ultimate Ultimate Ultimate Yield Yield YieldYield Yield Total Total Total TotalTotal Total Uniform Uniform Uniform Uniform Yield and ultimate Tantalum T-111 0 200 400 6000 200 500 700 900 500 700 900 1600 1200 800 400 0 400 600 Test temperature (�C) E lo ng at io n (% ) S tr es s (p si ) S tr es s (M P a) Test temperature (K) Figure 8 Comparison of tensile properties of neutron-irradiated Ta and T-111. Uniform elongations of 25.2 dpa 7.5 0 0 200 400 600 800 1000 10 20 30 Elongation (%) 40 50 4.4 0.7 2.0 5 10 Unirr. Unirr. Elongation (%) 15 200 0 200 400 600 800 E ng in ee rin g st re ss (M P a) E ng in ee rin g st re ss (M P a) E ng in ee rin g st re ss (M P a) 1000 1200 1400 1600 0.14 dpa 0.04 0.004 0.0004 0.00004 Unirr. 0 10 0.7 4.6 2.0 7.5 dpa Unirr. 20 30 40 600 400 200 0 800 1000Ta Neutron irradiated Tirr 333–373 K Ta–1W Neutron irradiated Tirr 333–373 K Ta–10W Proton and neutron irradiated 0.14 dpa 0.04 0.004 0.0004 0.00004 0 0 200 400 600 800 10 20 Elongation (%) 30 40 Proton and neutron irradiated Tirr 323–433 K (a) (b) (c) Tirr 323–433 K Figure 9 Room temperature tensile curves for irradiated (a) pure Ta, (b) Ta–1W, and (c) Ta–10W. Reproduced from Byun, T. S.; Maloy, S. A. J. Nucl. Mater. 2008, 377, 72–79. 192 Radiation Effects in Refractory Metals and Alloys The room temperature unirradiated tensile strength of Ta–10W is nearly double the value of the Ta–1W and triple that of pure Ta in the material investigated by Byun and Maloy,56 and also shows an increased sensitivity in radiation hardening over the pure metal (Figure 9(c)). This sensitivity is also clearly apparent at higher irradiation temperatures near 673K, as shown in the comparison of tensile curves that were compiled by Ullmaier and Carsughi58 of earlier work (Figure 11). Near room temperature irradiation of Ta–10W to the mixed proton and spallation neutron exposure by Byun andMaloy56 to doses between 2 and 25.2 dpa showed prompt necking following yielding. Total elongation values of 0.1MeV), 2.5 dpa, at 688 and 913K. The increase in radiation hardening is substantially greater than that observed in pure Ta irradiated under similar conditions. Yield and Plastic instability region Plastic instability region Fracture region Elastic regionElastic region Uniform plasticity (523 K)Uniform plasticity (523 K) Uniform plasticity (Trm)Uniform plasticity (Trm) 0.0 0 500 1000 Tr ue s tr es s (M P a) Tr ue s tr es s (M P a) 1500 2000 0.0001 0.001 0.01 0.1 Dose (dpa)(b) 1 10 Fracture region 0.0 0 200 400 600 800 1000 0.0001 0.001 0.01 0.1 Dose (dpa)(a) 1 Figure 10 Deformation mode map of (a) pure Ta and (b) Ta–1W for room temperature irradiations, illustrating fracture, plastic instability, uniform plasticity, and elastic regions as a function of stress and displacement dose. The increases in the uniform plasticity region for temperatures of 523K are superimposed. Reproduced from Byun, T. S.; Maloy, S. A. J. Nucl. Mater. 2008, 377, 72–79. 0 0 40 80 120 160 200 240 280 320 360 0.39 dpa 0.13 dpa S tr es s (M P a) Ta-10W Tirr~673 K 400 2 4 6 Strain (%) 8 Control Control Ta Figure 11 Comparison of the radiation hardening of Ta and Ta–10W irradiated at �673K to displacement doses of 194 Radiation Effects in Refractory Metals and Alloys 4.06.4 Molybdenum and Mo-Base Alloys 4.06.4.1 Introduction and History of Mo and Mo Alloys Molybdenum and its alloys are the perennial candi- dates for refractory metal alloy use in irradiation environments, due in part to their high melting temperature (2896 K), good thermal properties, high-temperature strength, and lower induced radioactivity (as compared to tantalum). The density of molybdenum (10.28 g cm�3) is also significantly lower than that of Ta andW, though greater than Nb. But like other refractory metal alloys, Mo can pres- ent difficulties in fabrication, low-temperature duc- tility, and low-temperature embrittlement from radiation damage. The TZM (Mo–0.5%Ti–0.1% Zr) and Mo–Re alloys were examined as part of the SP-100 and JIMO/Prometheus space reactor pro- grams, respectively, and offer additional benefits of improved high-temperature strength over the pure metal.5,19 Molybdenum and its alloys have also been examined for plasma facing and diverter components in fusion reactor designs due to the relatively low sputter yield, high thermal conductivity, and thermal compatibility with other structural materials.5–27,29–63 In addition, because of these benefits, Mo has also been examined for use as a grazing incident metal mirror in fusion diagnostic port designs.64,65 As in all other refractory metals, the mechanical properties are influenced by impurity concentra- tions, particularly through grain boundary weaken- ing. However, improvements in Mo ductility are achievable through grain refinement, impurity con- trol, and the addition of Re or reactive elements such as Ti and Zr. An upper limit to the acceptable level of C was also found to improve grain boundary strength. Low-carbon arc-cast molybdenum (LCAC-Mo) is one such example, in which oxygen impurities are reduced to tens of ppm, nitrogen to Radiation Effects in Refractory Metals and Alloys 195 to be inaccurate in determining the upper bound limit for maximum swelling.19 The swelling data col- lected from numerous sources,53,87–89 including those contained in the review work of Brimhall et al.90 for irradiated Mo as a function of dose and irradia- tion temperature, are provided in Figure 12. Void swelling was found to be 0.1MeVat temperatures between 673 and 1173K by Evans.53 Void swelling studied by Stubbins et al.88 in 3.1MeV 51V+ ion-irradiated Mo between 1173 and 1393K up to 50 dpa remained below 4%, while irradiations between 1523 and 1780Kwere near 10%. Void ordering has been observed in both neutron- irradiated28,89,91 and ion-irradiated Mo88 at tempera- tures between 700 and 1373K. Garner and Stubbins89 examined the irradiation and material conditions that contribute to void ordering. Irradiation tempera- tures near 700 K delineate the lower boundary tem- perature for void lattice formation at irradiations above 20 dpa. At lower doses, void lattice formation was not observed. The void superlattice constant, mea- sured as the distance between void centers along the direction in the material, is found to increase with temperature from �2.4 nm at 700K to 4.5 nm 4 3 2 DV /V (% ) 1 0 400 800 1200 160 Irradiation temperature (K) Garner and Stubbins89 (neutron Stubbins et al.88 (ion) Evans53 (neutron) Lee et al.87 (neutron) Brimhall et al.90 (neutron) Brimhall et al.90 (ion) Figure 12 Irradiation-induced swelling (DV/V ) as a function of for pure Mo. The irradiation source is marked in the key. Reprod 1976, 62, 115–117; Evans, J. H. J. Nucl. Mater. 1980, 88, 31–41; 64–77; Garner, F. A.; Stubbins, J. F. J. Nucl. Mater. 1994, 212–2 H. E. J. Nucl. Mater. 1973, 48, 339–350. at 1176K.91 Swelling is expected to reach a maximum of 3–4% on the development of the void lattice struc- ture, based on an attainment of an equilibrium ratio of void diameter to void superlattice parameter.92 At temperatures >1423K, void lattice formation is no longer observed, leading to the high values of swelling observed in the material ion irradiated to high doses.88 The onset of void growth in neutron-irradiated material appears to be accelerated in cold-worked materials compared to annealed materials, reaching a maximum in swelling at doses near 40 dpa for tem- peratures below 873K and 20 dpa at higher tempera- tures.89 At higher doses, swelling decreases through void shrinkage, with swelling values approaching those of annealed materials. Void shrinkage has also been reported by Bentley et al.93 and Evans53 to occur because of changes in the void sink bias89 presum- ably due to the segregation of transmuted species at the void surfaces, making them more attractive for interstitials. Irradiation-induced swelling in TZM has been reported53,94–97 and generally shows similar temper- ature dependence as the pure metal. The fluence and temperature dependence of swelling of TZM was examined by Powell et al.95 and Gelles et al.,94 with 0 0.01 0.1 1 dp a 10 100 ) irradiation temperature and displacement damage (dpa) uced from Lee, F.; Matolich, J.; Moteff, J. J. Nucl. Mater. Stubbis, J. F.; Moteff, J.; Taylor, A. J. Nucl. Mater. 1981, 101, 15, 1298–1302; Brimhall, J. L.; Simonen, E. P.; Kissinger, 196 Radiation Effects in Refractory Metals and Alloys results from the latter shown in Figure 13. Peak swelling in TZM following irradiation to 1.78� 1023 n cm�2 and 873K remained below 4%, though the data are limited to irradiation temperatures below 923K. Only limited data are available on direct com- parisons between TZM and pure Mo, with Bentley and Wiffen96 reporting 1% swelling in Mo–0.5%Ti and TZM alloys and 0.6% swelling in pure Mo under the same irradiation conditions. Similarly, 4% swelling was observed in TZM and 3% in pure Mo following irradiation to 5.4� 1022 n cm�2 at 923 K.97 In examining Mo and TZM of different preirra- diated material conditions, Evans53 observed equal or greater swelling in TZM compared to Mo follow- ing irradiation to 3.5� 1022 n cm�2 (E> 0.1MeV) at 823 and 873K. However, in the materials irradiated at 723 K for the same fluence, the TZM alloy showed lower swelling, except in the carburized condition. The Ti and Zr atoms not tied up as carbides are assumed to have played a role in reducing void size in the material at the lower temperature. There is little information on the swelling behav- ior of Mo–Re alloys. Measured swelling of 0.44% in Mo–50Re irradiated to 5.3� 1022 n cm�2 (E> 0.1 MeV) at temperatures which rose during irradiation from 1128 to 1329K was reported.26 For irradiated Mo–Re alloys, radiation-induced segregation (RIS) and transmutation can lead to precipitation of 1 1000 900T irr (K) 800 700 600 0.0 0.5 1.0 1.5 2.0 DV /V (% ) 2.5 3.0 3.5 4.0 Gelles et al.94 (11-20 dpa) Gelles et al.94 (36-53 dpa) Gelles et al.94 (47-65 dpa) Evans53 Figure 13 The swelling dependence on temperature and fluen Peterson, D. T.; Bates, J. F. J. Nucl. Mater. 1981, 103–104, 114 equilibrium or nonequilibrium phases, which can be detrimental to mechanical properties. This is examined in the next section. Electrical resistivity changes to 5.4 dpa irradiated Mo at 733K were examined by Zakharova et al.98 using single crystal samples. Increases in resistivity of 10–14% and 92–110% were measured at postirra- diation test temperatures of 298 and 77K, respec- tively. The largest resistivity changes were measured in the [100] direction. A residual 10% increase in resistivity was measured following annealing above 0.6Tm associated with the accumulation of trans- muted radionuclides. The changes in electrical resistivity of LCAC-Mo over a 353–1373K irradiation temperature range up to 3.3 dpa were examined by Li et al.99 and Cockeram et al.,100 with the latter examining the recovery of resistivity following isochronal anneals. The room temperature resistivity for 353 K irradiated LCAC- Mo rapidly increases between 0.01 and 0.1 dpa saturating near 0.2 dpa for an �42% increase over the unirradiated value.99 Increases in room tempera- ture resistivity of 10–12% were reported following 0.5–1.2 dpa irradiation at 543 K, and 3.3–5.3% after 1.4–2.4 dpa at 878 K. At irradiation temperatures �1208K, little ( Radiation Effects in Refractory Metals and Alloys 197 interstitials formed during irradiation to diffuse to sinks where annihilation occurs, reducing the electri- cal scattering effects that these defects have at lower irradiation temperatures. The small increases mea- sured at the higher irradiation temperatures were pri- marily due to transmutation products. As is shown in the next section, the changes in electrical resistivity with increasing irradiation temperatures also correlate with changes inmeasured hardness, though at a greater level of sensitivity. This is controlled by microstruc- tural changes, as the small dislocation loops and voids of high distribution density appearing at the lower irradiation temperatures coarsen into larger and fewer defects that have less interaction with deforma- tion dislocations. 4.06.4.3 Mechanical Properties of Irradiated Mo and Mo Alloys The mechanical property performance of pure Mo is strongly controlled by the grain size, oxygen, nitrogen, and carbon concentrations as well as alloy additions. This is true for unirradiated as well as irradiated properties.71,76,83,84 The sensitivity to embrittlement at irradiation temperatures �873K can be mitigated through a reduction in oxygen and nitrogen while keeping the carbon-to-oxygen ratio high to reduce the segregation of oxygen and nitrogen to the grain boundaries. A reduction in the grain size can further increase the number of sinks and reduce the mean distance that irradiation-induced defects must travel at temperatures at which mobility is limited. Irradiated mechanical properties of wrought LCAC-Mo in both the recrystallized and stress- relieved conditions have been examined over several decades.81,82,84–86,99–106 In general, LCAC-Mo un- dergoes significant increases in tensile strength through radiation hardening �873K, which produces reductions in ductility and high DBTT values. A summary of tensile properties as a function of irra- diation temperature and dose is shown in Figure 14. Irradiated stress-relieved LCAC-Mo shows less radia- tion embrittlement than as-crystallized materials at temperatures 773K are observed for irradiated LCAC-Mo over a range of fluences for irradiation temperatures �873K.26,82,85,103,105,108 A summary of DBTT values for LCAC-Mo is presented in Figure 16, along with data from high-purity grain-refined LCAC-Mo, TZM, and ODS-Mo, which is discussed next. In general, recovery in the DBTT of LCAC-Mo is not observed until irradiation temperatures above 873– 973K, depending on material conditions. A reduced sensitivity to low-temperature irradiation embrittle- ment of LCAC-Mo is observed in materials with reduced levels of impurities. A high-purity form of LCAC-Mo (HP-LCAC-Mo) was developed through the use of 1873K hydrogen atmosphere annealing of LCAC-Mo plates prior to further arc casting, extru- sion, and rolling into sheet stock.82 Levels of oxygen and nitrogen were HP-LCAC (0.11-1.29 dpa, 573 K) LCAC (0.6 dpa, 573 K) LCAC (13.1 dpa, 833-1057 K) LCAC (0.9-1.4 dpa, 878 K) LCAC (0.9 -1.4 dpa, 878 K) EB (0.23 dpa, 873 K) LCAC (3.3 dpa, 1373 K) LCAC (3.3 dpa, 1373 K) LCAC (12.3 dpa, 567 K) LCAC (13.1 dpa, 833-1057 K) LCAC (unirr.) LCAC (unirr.) HP-LCAC (unirr.) HP-LCAC (unirr.) LCAC (12.3 dpa, 567 K) 12001000800600 Test temperature (K) 4002000 0 2000 1500 1000 500 25 20 15 10 5 Y ie ld s tr es s (M P a) To ta l e lo ng at io n (% ) EB (0.23 dpa, 873 K) HP-LCAC (0.11-1.29 dpa, 573 K) Figure 14 Yield stress and total elongation as a function of test temperature for LCAC-Mo, HP-LCAC-Mo and electron beam (EB) melted Mo. Irradiation dose and temperature (dpa, K) is marked. Unirradiated samples (unirr.) are also presented. Reproduced from Cockeram, B. V.; Smith, R. W.; Leonard, K. J.; Byun, T. S.; Snead, L. L. J. Nucl. Mater. 2008, 382, 1–23; Cockeram, B. V.; Smith, R. W.; Snead, L. L. J. Nucl. Mater. 2005, 346, 165; Abe, K.; Takeuchi, T.; Kikuchi, M.; Morozumi, S. J. Nucl. Mater. 1981, 99, 25–37. 198 Radiation Effects in Refractory Metals and Alloys The majority of the irradiated mechanical prop- erty database for TZM is limited to displacement damage 873 K 573 K 1173 K C ha ng e in h ar d ne ss (M P a) 1 h an ne al t em p er at ur e (K ) 400 300 200 100 0 -100 0.00 2.00 4.00 6.00 8.00 Displacement dose (dpa) 10.00 12.00 14.00 75-100% 50-75% 25-50% 0-25% No change Recovery of hardness >1573 K 0.6dpa 543 K LCAC 12.3dpa 573 K ODS 3.9dpa 877 K ODS 3.9dpa 882 K TZM 13.1dpa 833 K ODS 13.1dpa 833 K TZM 12.3dpa 573 K TZM 1.4 dpa 877 K LCAC 12.3dpa 567 K LCAC 13.1dpa 833 K LCAC Dose Tirr 500 700 900 1100 1300 1500 1700 LCAC-Mo (Tirr= 573 K) ODS-Mo (Tirr= 573 K) TZM (Tirr= 573 K) LCAC-Mo (Tirr= 873 K) ODS-Mo (Tirr= 873 K) TZM (Tirr= 873 K) LCAC-Mo (Tirr= 1173 K) ODS-Mo (Tirr= 1173 K) TZM (Tirr= 1173 K) (a) (b) Figure 15 (a) Change in hardness as a function of neutron dose for LCAC-Mo, TZM, and ODS-Mo irradiated between 573 and 1173K up to 13.1dpa. (b) Recovery of hardness as a function of isochronal annealing temperature for material irradiated �573K. Adapted from Cockeram, B. V.; Smith, R. W.; Byun, T. S.; Snead, L. L. J. Nucl. Mater. 2009, 393, 12–21. Radiation Effects in Refractory Metals and Alloys 199 0 -100 100 200 300 400 500 600 700 800 0 100 200 300 400 500 Irradiation temperature (°C) D B TT (° C ) 600 700 800 900 10001100 12001300 LCAC-Mo Wiffen26 (13.3-23.4 dpa) Chakin and Kazakov108 (2.1-3 dpa) Hasegawa et al.106 (10.5-50 dpa) Abe et al.104,105 (0.6-1.4 dpa) Webster et al.148 (7.4-10 dpa) Smith and Michel149 (5.3 dpa) Scibetta et al.109 (0.15-0.19 dpa) Cockeram et al.85 (12.3-13 dpa) Cockeram et al.82 (1.29 dpa) Cockeram et al.85 (12.3-13 dpa) Wiffen26 (16 dpa) Cockeram et al.85 (12.3-13.1 dpa) Cockeram et al.82 (1.29 dpa) HP-LCAC-Mo TZM ODS-Mo Figure 16 Summary of DBTT values as a function of neutron irradiation temperature for LCAC-Mo and TZM. Adapted from Cockeram, B. V.; Smith, R. W.; Leonard, K. J.; Byun, T. S.; Snead, L. L. J. Nucl. Mater. 2008, 382, 1–23; Cockeram, B. V.; Smith, R. W.; Snead, L. L. J. Nucl. Mater. 2005, 346, 165. 200 Radiation Effects in Refractory Metals and Alloys displacement damage to the TZM alloy produces a large increase in the yield strength of the material with a corresponding drop in total elongation, with the level of change increasing with dose and/or lower irradiation temperatures. Some recovery of properties begins to be observed at test conditions above 973K for materials irradiated at lower temperatures. A compila- tion of earlier and more recent tensile data for irra- diated TZM is provided in Figure 17. Irradiation �1.2 dpa and temperatures>800K showed little effec- tive strengthening above the unirradiated values, for tensile tests below the irradiation temperature, 85,109,110 though higher displacement doses resulted in a signifi- cant increase.81 Total elongation in material irradiated 1400 1200 1000 800 Unirradiated range 600 400 200 200 (a) (b) 400 600 800 1000 Test temperature (K) Y ie ld s tr en gt h (M P a) 1200 1400 1974 Kasakov et al.: unirradiated 2005 Cockeram et al. (3.9 dpa, T irr= 833 K) 2005 Cockeram et al./Byun et al. (unirradiated) 2005 Cockeram et al. (12.3 dpa, T irr= 567 K) 2008 Byun et al. (3.9 to 13.1 dpa, T irr= 567 - 833 K) 2008 Byun et al. (3.9 to 13.1 dpa, T irr= 1057 - 1209 K) 1974 Kasakov et al. (1.6 dpa, T irr= 823 K) 1974 Kasakov et al. (1.2 dpa, T irr= 1223 K) 20 18 16 14 12 10 8 6 4 2 0 200 300 400 500 600 700 800 Test temperature (K) To ta l e lo ng at io n (% ) 900 1000 1100 1200 Unirradiated range 1974 Kasakov et al.: unirradiated 2005 Cockeram et al. (3.9 dpa, T irr= 833 K) 2005 Cockeram et al./Byun et al. (unirradiated) 2005 Cockeram et al. (12.3 dpa, T irr= 567 K) 1976 Steichen (2.12 dpa, T irr= 644 K) 1976 Steichen (4.79 dpa, T irr= 661 K) 1974 Kasakov et al. (1.6 dpa, T irr= 823 K) 1974 Kasakov et al. (1.2 dpa, T irr= 1223 K) Figure 17 Neutron-irradiated tensile data for TZM (a) yield stress versus test temperature and (b) total elongation versus test temperature. Reproduced from Byun, T. S.; Li, M.; Cockeram, B. V.; Snead, L. L. J. Nucl. Mater. 2008, 376, 240–246; Cockeram, B. V.; Smith, R. W.; Snead, L. L. J. Nucl. Mater. 2005, 346, 165; Kasakov, V. A.; Kolesnikov, A. N.; Krassnoselov, V. A.; et al. Effect of neutron irradiation on properties of potential structural material for thermonuclear reactors, USSR-US Exchange on CTR Materials, Nov 1974; Steichen, J. M. J. Nucl. Mater. 1976, 60, 13–19. Radiation Effects in Refractory Metals and Alloys 201 the DBTT remained unchanged with irradiation, despite a �50% increase in Vickers microhardness. More surprisingly, the Mo–1%TiC sample increased in toughness following 0.8 dpa irradiation attributed to grain boundary strengthening by (Ti, Mo)C radiation-enhanced precipitates. Developed to improve the low-temperature duc- tility and weld characteristics of unalloyed Mo, the Mo–Re alloys have gained considerable attention over the past decade for use in nuclear applications. Single-phase solid solution a-Mo phase field extends up to �42wt% Re, above which the s-MoRe2 phase precipitates. At higher Re concentrations, the w-MoRe3 phase is present. However, the exact phase boundaries are not well delineated at temperatures below 1773K113,114 mainly because of the slow kinet- ics in phase development.115 The Mo–Re alloys show a hardening response to irradiation stronger than that of the pure metal and TZM following irradiation.116–118 The hardening response of Mo–Re alloys ranging in composition from 2 to 13 as well as 41wt% Re following 200 0 0.1 0.2 To ta l a b so rb ed e ne rg y (J m m -3 ) 0.3 0.4 0.5 0.6 0.7 0.8 0.9 250 300 350 Test temperature (K) 400 450 500 TZM Mo-1%TiC Mo-0.5%TiC Mo-0.1%TiC Figure 18 Total absorbed energy versus test temperature for TZM and 0.1–1.0wt% TiC additions to Mo irradiated to 0.08dpa at 573–773K. Reproduced from Kitsunai, Y.; Kurishita, H.; Narui, M.; Kayano, H.; Hiraoka, Y. J. Nucl. Mater. 1996, 239, 253–260. 202 Radiation Effects in Refractory Metals and Alloys irradiation up to 20 dpa at temperatures between 681 and 1072Kwas examined by Nemoto et al.116 A linear increase in hardness with Re concentration was ob- served for the unirradiated controls as well as samples irradiated at temperatures �874K. For samples irra- diated at 1072K, little variation was observed with increasing Re content, though hardness values remained nearly double that of the unirradiated mate- rial. The dependence of hardness on irradiation tem- perature and fluence for Mo–5Re and Mo–41Re in comparisonwith LCAC-Mo is presented in Figure 19. The high degree of radiation hardening at tem- peratures 200 0 200 400 600V ic ke rs h ar d ne ss 800 1000 1200 1400 1600 1800 400 600 LCAC-Mo Unirrad. Mo-5Re Mo-41Re Irradiation temperature (K) 800 1000 1200 Mo (Nemoto et al.116 18-21 dpa) Mo-5Re (Nemoto et al.116 18-21 dpa) Mo-41Re (Nemoto et al.116 18-21 dpa) Mo-5Re (Hasegawa et al.117 6.8-34 dpa) Mo-41Re (Hasegawa et al.117 6.8-34 dpa) Mo-10Re (Nemoto et al.116 18–21 dpa) Mo-5Re (Hasegawa et al.117 6.8-34 dpa) Figure 19 Vickers hardness as a function of neutron irradiation temperature and dose for LCAC-Mo, Mo–5Re, and Mo–41Re alloys. Displacement damage levels are provided in the key. Reproduced from Nemoto, Y.; Hasegawa, A.; Satou, M.; Abe, K.; Hiraoka, Y. J. Nucl. Mater. 2004, 324, 62–70; Hasegawa, A.; Ueda, K.; Satou, M.; Abe, K. J. Nucl. Mater. 1998, 258–263, 902–906. 0 0 500 1000 S tr es s (M P a) 1500 1.46 dpa 1.46 dpa 2000 5 10 15 20 Annealed 1100 h aged 1100 h aged 25 30 35 500 1000 0.72 dpa 0.72 dpa 1.46 dpa 1500 1.46 dpa Annealed 0.72 dpa 0.72 dpa Mo-41Re Mo-47.5Re 0 5 10 15 20 Strain (%)Strain (%) 25 30 35 40 Figure 20 Comparison of stress–strain curves for neutron irradiated, 1100h and as-annealed Mo–41Re and Mo–47.5 Re samples at 1073K (Tirr ¼ Ttest). Adapted from Busby, J. T.; Leonard, K. J.; Zinkle, S. J. Effects of neutron irradiation on refractory metal alloys, ORNL/LTR/NR-PROM1/05-38; Oak Ridge National Laboratory: Oak Ridge, TN, Dec 2005; Busby, J. T.; Leonard, K. J.; Zinkle, S. J. J. Nucl. Mater. 2007, 366, 388–406. Radiation Effects in Refractory Metals and Alloys 203 600 0 5 10 15 500 U lti m at e te ns ile s tr en gt h (M P a) To ta l e lo ng at io n (% ) 1000 1500 2000 2500 Irrad. UTS Unirrad. UTS Irrad. UTS (brittle fracture) 700 800 Brittle fracture 900 Temperature (K) 1000 1100 1200 1300 1400 Mo-41 (Busby) Mo-Re (Fabrietsiev) Mo-5Re (Hasegawa) Open symbols: unirrad. Closed symbols: irrad. Mo-47.5Re (Busby) Figure 21 Tensile data comparisons of Mo–Re alloys detailing the upper limits for irradiation embrittlement. Adapted fromBusby, J. T.; Leonard, K. J.; Zinkle, S. J. Effects of neutron irradiation on refractory metal alloys, ORNL/LTR/NR-PROM1/ 05-38; Oak Ridge National Laboratory: Oak Ridge, TN, Dec 2005; Busby, J. T.; Leonard, K. J.; Zinkle, S. J. J. Nucl. Mater. 2007, 366, 388–406. 204 Radiation Effects in Refractory Metals and Alloys hardening to levels over twice the as-annealed condi- tion was observed for the alloys irradiated at 1223 and 1373K; however, total elongation was between 4 and 12%. Analysis of the fractured surfaces of these samples revealed intergranular failure, with the sever- ity increasingwith irradiation temperature. A compari- son of mechanical property data of Mo–Re samples from the sources discussed is shown in Figure 21. The degree of RIS influencing the properties of Mo–Re alloys varies with temperature, dose rate, and total fluence. At temperatures 0.5Tm, a reduced driving force for segregation occurs because of the high thermal defect concentrations. At intermediate temperatures (�850–1430K for Mo–Re), the radiation generated point defects diffuse to defect sinks such as grain boundaries or dislocations. Any preferential coupling of vacancies or interstitial defects fluxes with solute atoms, including transmuted species, will create enrichment at the defect sinks. This is observed in the nucleation of Re-rich phases in the microstruc- tures of neutron-irradiated samples116,121 and the degradation in mechanical properties and transition to intergranular fracture in higher Re concentration alloys.62,120 Further information on RIS can be found in Chapter 1.18, Radiation-Induced Segregation. Through modeling and experimental work, Erck and Rehn123 showed that the degree of segregation per dpa reaches a maximum for Mo–30 at.% Re (�45wt% Re) near 1223K and that for Mo–7 at.% Re (�13 wt% Re) near 1473K. While the Mo–5Re alloys irradiated up to 20 dpa show some limited ductility,109,118,119 the maximum irradiation tempera- tures were 100 0 5 10 15 To ta l e lo ng at io n (% ) Y ie ld s tr es s (M P a) 20 200 400 600 800 1000 1200 1400 1600 1800 300 500 Unirradiated 700 Test temperature (K) 900 1100 1300 1500 Tirr= 567 K, 12.3 dpa Tirr= 882 K, 3.9 dpa Tirr= 1143 K, 1.2 dpa Tirr= 1143 K, 3.4 dpa Unirradiated: longitudinal, stress-relieved Figure 22 Yield stress and total elongation as a function of test temperature for lanthanum oxide ODS-Mo neutron irradiated up to 13dpa at temperatures between 567 and 1209K. Adapted from Cockeram, B. V.; Smith, R. W.; Snead, L. L. J. Nucl. Mater. 2005, 346, 165–184. Radiation Effects in Refractory Metals and Alloys 205 The Mo–(41 and 47.5)Re alloys irradiated at 1073– 1373K120 showed indications of RIS even at relatively low damage levels, part of which may have been a contribution of a thermal aging component, which in the unirradiated as-aged Mo–41Re and Mo–47.5Re showed increases in Re at the grain boundaries, lead- ing to precipitation of s- and w-phases at the grain boundaries in the 47.5Re containing alloy.115 Utilizing Mo–Re alloys with a more moderate Re content may improve the irradiation performance of these alloys, especially when considering higher doses and/or lon- ger irradiation times at temperatures at which thermal precipitation effects may further compound RIS influ- ences on mechanical properties. Additional information on the fracture toughness data for Mo–Re alloys is also needed. Preliminary data by Scibetta and coworkers109 on precracked com- pact tension specimens of Mo–5Re showed reduc- tions in fracture toughness from unirradiated values of 17–23MPa√m at room temperature and 623K, to �11MPa√m for 0.35 dpa irradiated at 313K, and 15MPa√m for 0.29 dpa irradiated at 643K. These low irradiated values, at which no ductile crack growth was observed in the specimens, are a concern. Recent work examining the irradiated properties of wrought, commercially available, ODS-Mo contain- ing lanthanum oxide particles has shown promising results.81,82,124 The fabrication methods produce a microstructure consisting of elongated grains with appreciable texturing and alignment of the oxide par- ticles. The high degree of working associated with fabrication produces a 206 Radiation Effects in Refractory Metals and Alloys ODS-Mo are available. Unirradiated fracture tough- ness values are between 23 and 38MPa√m, depending on the grain orientation tested.83 1.8 4.06.5 Tungsten and W-Base Alloys 4.06.5.1 Introduction and Irradiated Properties Database for W and W Alloys Despite the recurring interest in the use of tungsten as a structural material for very-high-temperature applications and for use as plasma facing components in fusion devices, the database on irradiated properties is very limited and based primarily on fast neutron irradiation experiments. Similar to all bcc materials, tungsten is susceptible to low-temperature embrit- tlement at T� 0.3Tm (Tm¼ 3695K) for fluences >1� 1024 nm�2, which makes this material even more limited in toughness and ductility. Improvements in the unirradiated mechanical properties of tungsten are observed with the addi- tion of Re, which is found to increase ductility at elevated temperatures125 as well as fracture tough- ness126 through the reduction in the DBTT. How- ever, as discussed in this section, the gains in performance through added Re content do not nec- essarily hold for irradiated materials. 600 0 0.2 0.4 0.6 DV /V (% ) 0.8 1.2 1.4 1.6 1 800 1000 1200 Irradiation temperature (K) 1400 1600 1800 2000 Tungsten: Matolich et al.131 5.5 ´ 1022n cm–2 (E > 0.1 MeV) 4 to 6 ´ 1022n cm–2 (E > 0.1 MeV) 5.5 ´ 1022n cm–2 (E > 0.1 MeV) Tungsten: Wiffen19 W-25Re: Matolich et al.131 Figure 23 Irradiation-induced swelling measured through immersion density methods of W and W–25Re by Matolich et al.131 and Wiffen.19 4.06.5.2 Irradiation-Induced Swelling and Physical Property Changes in W and W Alloys Early investigations into the behavior of irradiated tungsten59,125,127–129 examined the defect formation and recovery defects. Much of this initial work was through the examination of electrical resistivity fol- lowing irradiation. Increases in electrical resistivity of pure annealed tungsten of up to 24% following irradiation to 2� 1022 n cm�2 in a fast reactor and �14% in a mixed spectrum reactor to 1021 n cm�2 have been reported.59 Keys et al.127,128 examined the recovery of neutron- irradiated tungsten through isochronal resistivity studies following irradiation at 343 K to dose levels of 1.5� 1021 n cm�2 (E> 1MeV). The beginning of saturation in resistivity observed in their studies appears just below 1020 n cm�2 and correlates with the work by Lacefield et al.130 on the appearance of defect clusters by 2.4� 1019 n cm�2 identified through TEM examination. The work by Keys et al.127,128 identified distinct stages of recovery such as self-interstitial migration occurring near 0.15Tm, with a less pronounced recovery at 0.22Tm through divacancy and impurity migration, which was followed by vacancy migration above 0.31Tm. The residual resistivity, not recovered following anneals above 0.4Tm after irradiation to fluences>3.3� 1019 n cm�2, was due to the development of Re in the tungsten from transmutation reactions with thermal neutrons. Very little data exist on irradiation-induced swell- ing inWand its alloys. Data on pureWare restricted to two reported series of experiments concerning the temperature dependence of swelling. The irradiation- induced swelling measured by Matolich et al.131 and Wiffen19 using immersion density methods is shown inFigure 23. It should be noted that there is an order of magnitude difference in fluences between the two stud- ies. No other systematic examination of the swelling dependence on fluence and temperature is available. Though swelling data for W–Re alloys are also lim- ited, work by Matolich et al.131 for W–25Re irradiated to 5.5� 1022 n cm�2 revealed no significant amount of swelling. The data are also shown in Figure 23. Microstructural examination of W–Re alloys with concentrations of 5%, 11%, and 25%Re showed no cavity formation for fluences between 4.3 Radiation Effects in Refractory Metals and Alloys 207 and 6.1� 1021 n cm�2 (E> 0.1MeV) at temperatures between 873 and 1773K,132 while void cavities have been experimentally observed in irradiated pure Wover similar fluences.133,134 The effects of increas- ing Re or Os content inWwere experimentally shown to decrease the density and radius of dislocation loops and voids in 0.15 dpa proton and neutron- irradiated material by He et al.135 This reduction in size and number density is the result of the restricted mobility of the radiation-induced defects by the lat- tice dilations from the Re and Os solute. The transmutation ofW to Re and Re to Os during irradiation can have an effect on microstructural, physical, and mechanical properties of the material. The transmutation of the material, which results in the shifting of solute concentrations to higher levels, may result in precipitation in alloys that are nomi- nally in a single-phase region. One example of micro- structural and physical property changes because of irradiation is the decalibration of type-C (W–3%Re/ W–25%Re or W–5%Re/W–26%Re) thermocou- ples, such as that used in fuel element centerline temperature measurements. Reviews of early experi- mental work on W/Re thermocouples and on the dependence of decalibration on the neutron fluence have previously been discussed.136,137 While displa- cive neutron damage may result in material changes such as vacancy clusters or dislocation loops, the maximum theoretical changes expected in emf out- put of the thermocouple is �1 mV �C�1,138 whereas the changes associated with transmutation effects can result in more significant decreases. A �300 �C drift in temperature following 6000 h irradiation under 2.7� 1022 n cm�2 thermal and 8� 1021 n cm�2 fast fluence was reported for a W–3%Re/W–25%Re couple.139 These changes can also be significant in fast reac- tor irradiations. Experimental work by Williams et al.132 showed that for a W–5%Re/W–25%Re ther- mocouple irradiated to 6.1� 1021 n cm�2 fast fluence at 1173K, the precipitation-induced changes in the Seebeck coefficient were�6.6 and�0.02mV �C�1 for the 5 and 25%Re alloys, respectively. Calculated final compositions following 6.1� 1021 n cm�2 (14MeV) irradiation of W–5(wt%)Re produce W–5.130Re– 0.021Os–0.150Ta alloy, while a W–26Re alloy trans- mutes to W–25.955Re–0.107Os–0.117Ta.140 Postirradiation examination of the microstructure of the irradiated 5, 11, and 25% Re alloys in Williams et al.132 revealed w-phase precipitation at irradiation temperatures above 1373K, though unidentifiable precipitation was apparent in the alloys at 1173 K. The development of the w-phase over the equilibrium s-phase in irradiated samples, but not in the unirradi- ated annealed samples, is the result of irradiation- induced solute segregation to defect sinks. The development of the w-phase was also reported in microstructural studies of W–26Re irradiated up to 11 dpa at temperatures between 646 and 1073K.141 It should be pointed out that the change or tem- perature shift under irradiation is proportional to the degree of localized transmutation and local tempera- ture gradients and therefore dependent on the pro- files of the temperature and irradiation fields to which the thermocouple is exposed. Therefore, experimental work typically involves the irradiation of the entire cable, while in reactor applications, significant variations in temperature and fluence may result. The changes in thermoelectric power (D) as a function of irradiation fluence can be mod- eled by the following:140 D¼ 0 for 0� f� 0:25� 1021 n cm�2 D¼ 100½1� e0:067ð0:25�fÞ� for 0:25� f� 1� 1021 n cm�2 D¼ 100½1� e0:104ð0:52�fÞ� for f> 1� 1021 n cm�2 ½2� Though significant radiation-induced decalibration may occur in fission reactors, this effect may not be readily observed in fusion reactors where the thermal flux is much lower. In addition, typical end-of-life estimates of total neutron fluence of 208 Radiation Effects in Refractory Metals and Alloys properties,59 likely through the development and coarsening of interstitial impurities into precipitate formations that reduce grain boundary sensitivities. For the irradiated tensile properties of tungsten, two works are typically referenced that make up the bulk of the data available. In the work of Steichen,111 the properties of wrought tungsten, stress-relieved at 1273K, irradiated at 658K to fluences between 0.4� 1022 and 0.9� 1022 n cm�2 (E> 0.1MeV), were examined, the results of which are shown in Figure 24. The irradiated yield strength of the material increased to approximately twice that of the unirradiated values, 200 To ta l e lo ng at io n (% ) 0 5 10 15 20 1.0 ´ 102 Gorynin Brittle failure 0.9 ´ 1022n Steichen11 Unirradiated Steichen111 0 200 400 600Y ie ld s tr es s (M P a) 800 1000 1200 1400 1600 400 600 Test tem Figure 24 Temperature-dependent tensile properties of irradi Gorynin, I. V.; Ignatov, V. A.; Rybin, V. V.; et al. J. Nucl. Mater. 1 1976, 60, 13–19. while significant reduction in ductility was observed. Only in tensile tests well above the irradiation temper- ature did ductility values approach that of the unirra- diated material. In the work by Gorynin et al.,59 pure W consoli- dated through powder metallurgy was irradiated at temperatures of up to 1073 K in both a mixed and fast reactor up to 2� 1022 n cm�2 (E> 0.1MeV). Samples irradiated and tested at temperatures near 573 K showed brittle failures at low stress levels, while some ductility and appreciable hardening were observed for samples tested and irradiated at 2n cm–2 Tirr= 603–723 K et al.59 1.5 ´ 1022n cm–2 Tirr= 773 K Gorynin et al.59 0.5 ´ 1022n cm–2 Tirr= 644 K Steichen111 cm–2 Tirr= 661 K 1 800 perature (K) 1000 1200 1400 ated and unirradiated tungsten. Reproduced from 992, 191–194, 421–425; Steichen, J. M. J. Nucl. Mater. Radiation Effects in Refractory Metals and Alloys 209 1073 K. Limited recovery of strength and ductility was observed in postirradiated material annealed at 1473 K for 1 h. Embrittlement following irradiation due to radia- tion hardening and loss of grain boundary strength due to impurities resulted in increased DBTT for the aforementioned work. The DBTT is dependent on the test conditions in addition to material conditions prior to irradiation and should be used with caution. The DBTT in samples examined by Steichen111 increases from �333K in the unirradiated condition to 503K, following 1–2 dpa irradiation at �653K, while DBTT values increased from 673Kunirradiated to 873K after 1 dpa at 373K for sintered W.59 Increased DBTT values with irradiation were also reported by Krautwasser et al.142 in powder metallurgy to form W, W–10Re, and Densimet 18 (W–3.4Ni–1.6Fe) bend-test bars irradiated between 525 and 575K up to 5.6� 1021 n cm�2 (E> 0.1MeV) (see Figure 25). While the addition of Re toWresults in improved nonirradiated mechanical properties,142 the increased DBTT in the irradiated W–10Re is more severe than in pure W. In the case of the former, the possible development of the w-phase may be responsible for the higher DBTT values and the general increased sensitivity to radiation hardening. The w-phase observed inW–26Re irradiated from 2 to 9.5 dpa at temperatures between 373 and 800 �C141 is reported as precipitating as plate-like particles on the {110} planes of the W matrix, therefore, restricting slip in the material. 5.1 9.2 Dose (´1020n cm–2), Tirr= 523-573 K 20.9 44-560 0 200 400 600D B TT (K ) 800 1000 1200 1400 W-pure W-10Re Figure 25 Ductile-to-brittle transition temperature (DBTT) as a function of neutron fluence (E > 0.1MeV) of W and its alloys. Reproduced from Krautwasser, P.; Heinz, D.; Kny, E. High Temp. High Press. 1990, 22, 25–32. It should be noted that due to the limited mechan- ical property data available for W and W–Re alloys, particularly the lack of irradiated data at ele- vated temperatures, accurate determination of the DBTT cannot be made. Nonetheless, increases in DBTT between 200 and 500 K for�1 dpa of damage reported for the various grades of pure tungsten create limitations on its use, particularly at low irradiation temperatures.59,109,142 Based on irradia- tion data for Mo alloys, the minimum irradiation temperature which avoids severe radiation embrit- tlement is >0.3Tm or �1300 K for tungsten at neutron fluences >0.03 dpa or 1� 1021 n cm�2 (E> 0.1MeV),3 which correlates with the irradia- tion defect recovery data on tungsten compiled by Keys et al.127 Recent work in the development of ultra-fine grained tungsten incorporating TiC additions has shown promising results in reducing the sensitivity to radiation-induced degradation of properties.143,144 The grain size refinement, in the range of 50–200 nm, depending on TiC additions and process, theoreti- cally reduces the effective size of weak grain bound- aries that can act as crack initiators. In addition, significant reductions are observed in the density of void formation in the materials relative to pure W at irradiations conducted at 873 K and 2� 1020 n cm�2, though interstitial loop densities are unchanged. While unirradiated room temperature tensile prop- erties still show brittle fracture behavior, the fracture stress is up to four times higher in the W–TiC sam- ples than in pure W in addition to showing 100K lower DBTT in impact testing. In microhardness measurements following irradiation, the W–TiC samples exhibited no radiation hardening compared with pure W. The change in Vickers hardness follow- ing irradiation for the W–TiC material of Kurishita et al.143 compared to neutron- and proton-irradiated Wand W–Re alloys135 irradiated to similar tempera- tures and doses is shown in Figure 26. The reduced sensitivity of the W–TiC alloy to radiation hardening offers the potential for further development of these alloys for nuclear applications. 4.06.6 Outlook The use of refractory metal alloys in radiation envir- onments can offer high-temperature capabilities not matched in other alloy categories. Refractory metal alloys also offer exceptional compatibility with liquid metal coolants. As described in some detail in this 0.15 dpa Proton 773 K W-pure Dose Type Tirr Material 0 50 100 150 200 250 0.15 dpa Neutron 873 K W-pure 0.15 dpa Proton 873 K W-3Re 0.15 dpa Neutron 873 K W-3Re 0.15 dpa Neutron 873 K W-5Re 0.08 dpa Neutron 873 K W-pure 0.08 dpa Neutron 873 K W-0.5TiC He et al.135 Kurishita et al.143 C ha ng e in V ic ke rs H ar d ne ss w ith ir ra d ia tio n H V irr ad – H V un irr ad (k g m m –2 ) Figure 26 Comparison of the increase in Vickers Hardness for tungsten and tungsten alloys for similar dose and irradiation temperatures. Reproduced from He, J. C.; Hasegawa, A.; Abe, K. J. Nucl. Mater. 2008, 377, 348–351; Kurishita, H.; Kobayashi, S.; Nakai, K.; et al. J. Nucl. Mater. 2008, 377, 34–40. 210 Radiation Effects in Refractory Metals and Alloys chapter through mechanical property comparisons, these materials are sensitive to impurity contami- nation during metallurgical processing as well as in-service exposures that can lead to grain bound- ary embrittlement issues. The inherent irradiation response of bcc-structured materials also limits refractory metal use at temperatures >0.3Tm, with significant degradation in material properties with displacive irradiation doses as low as 0.03 dpa.3 Improvements in the irradiated mechanical prop- erties of refractory metal alloys have been observed in recent experimental work, even at low irradia- tion temperatures. This is in part through improved control over impurity levels and also through ther- momechanical processing techniques that result in microstructures with reduced sensitivity to radi- ation embrittlement. This was discussed with refer- ence to LCAC molybdenum,100 where samples irradiated in the stress-relieved condition showed improvement over material in the recrystallized condition up to the recrystallization temperature. Further development of HP-LCAC molybdenum has resulted in higher aspect ratio grain morpholo- gies that led to plain strain conditions in the grain lamellae during deformation.82 In addition, reduced grain sizes or higher aspect ratios decrease dis- tances to defect sinks, further reducing irradia- tion sensitivity. While Mo has traditionally been used to study the behavior of W, the microstruc- tural changes and purity control that have been employed for irradiation studies of Mo have not been incorporated into W. The control over precipitate formation in the pre- irradiated condition appears to result in changes to some physical material properties, specifically, swelling and densification in Nb–1Zr,25,27 that may lead to variations in mechanical properties. An understanding of the effect of preirradiation thermo- mechanical processing or in-service microstructural changes that occur during irradiation may lead to improved properties or the ability to avoid dangerous embrittlement issues that can occur through precipi- tate development. This may be of particular interest in Nb and Ta-base alloys that incorporate Zr or Hf additions that react with impurity elements and pro- duce precipitates. Alloying Mo and W with Re results in improved mechanical properties of unirradiated alloys, in- creased radiation hardening, and radiation-induced embrittlement.62,120 However, much of this work is Radiation Effects in Refractory Metals and Alloys 211 on recrystallized, high Re concentration material, the purity of which may not be ideal. The effect that RIS has on the degradation of properties of Mo–Re alloys is a matter of concern. Further work is needed on higher purity, lower Re (5–20 wt% Re) concentration material with reduced grain size, or that with a tai- lored aspect ratio similar to that of LCAC-Mo. Initial results show improvements to the irradiated properties of Mo and W through the incorporation of either rare earth oxide124 or TiC additions.112,143 These additions aid in restricting grain growth, pro- vide sinks for radiation-induced defects, and act as obstacles to or deflection points for crack propagation. 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Mater. 1977, 66, 125–142. 4.07 Radiation Effects in SiC and SiC–SiC L. L. Snead and Y. Katoh Oak Ridge National Laboratory, Oak Ridge, TN, USA T. Nozawa Japan Atomic Energy Agency, Rokkasho, Aomori, Japan � 2012 Elsevier Ltd. All rights reserved. 4.07.1 Introduction 215 4.07.2 Irradiation-Induced Swelling and Microstructure of Pure SiC 216 4.07.3 Irradiation-Induced Thermal Conductivity Degradation of Monolithic SiC 221 4.07.4 Effect of Irradiation on the Mechanical Properties of Monolithic SiC 224 4.07.4.1 Elastic Modulus of Monolithic SiC 224 4.07.4.2 Hardness of Monolithic SiC 226 4.07.4.3 Fracture Toughness of Monolithic SiC 226 4.07.4.4 Strength and Statistical Variation in Strength for Monolithic SiC 227 4.07.5 Irradiation Creep of SiC 233 4.07.6 Silicon Carbide Composites Under Irradiation 234 References 239 Abbreviations ATR Advanced test reactor BSD Black spot dot BSR Bend stress relaxation CVD Chemical vapor deposition CVI Chemical vapor infiltration dpa Displacement per atom DuET Dual-beam facility for energy science and technology ETR Engineering test reactor FB Fluidized bed fcc Face-centered cubic HFBR High-flux beam reactor HFIR High Flux Isotope Reactor HFIR-METS High-flux isotope reactor – mapping elevated temperature swelling HNLS Hi Nicalon Type S HP Hot-pressing JMTR Japan materials testing reactor Kd Thermal conductivity by defect scattering Kgb Thermal conductivity by grain boundary scattering Kirr Irradiated thermal conductivity Knonirr Nonirradiated thermal conductivity Krd Thermal conductivity by radiation Ku Thermal conductivity by Umklapp scattering PLS Proportional limit stress PS Pressureless sintering PW Plain weave PyC Pyrolytic carbon SEM Scanning electron microscopy SiC/SiC composite Silicon carbide fiber reinforced silicon carbide matrix composite SW Satin weave TEM Transmission electron microscopy Tirr Irradiation temperature TRISO TRIstructural ISOtropic TySA Tyranno SA UTS Ultimate tensile stress 4.07.1 Introduction Silicon carbide (SiC) has been studied and utilized in nuclear systems for decades. Its primary use was, and still is, as the micro pressure vessel for high- temperature gas-cooled reactor fuels. For these so-called TRI-ISOtropic (TRISO) fuels, the SiC is deposited via a gas-phase decomposition process over two layers of pyrolytic graphite surrounding the fuel kernel. In addition to being strong enough to with- stand the pressure buildup from the fission product gas liberated, this SiC layer must also withstand chemical attack from metallic fission products such as palladium and the mechanical loads derived from irradiation-induced dimensional changes occurring in the pyrolytic graphite. More recent nuclear appli- cations of SiC include its use as structural composites 215 216 Radiation Effects in SiC and SiC–SiC (i.e., SiC/SiC) for high-temperature gas-cooled reac- tors and for fusion power systems. The possibility of using composite and monolithic SiC thermal insula- tors for both fusion and fission systems is also being investigated. Moreover, both monolithic and compos- ite forms of SiC are being investigated for use in advanced sodium fast, advanced liquid salt-cooled, and advanced light water reactors. In this chapter, the effects of neutron irradiation on relatively pure, radiation resistant forms of SiC are discussed. This chapter has been limited to the effects of irradiation on the microstructure, and the mechanical and thermal properties of SiC, although it is recognized that environment aspects such as oxida- tion and corrosion will also be factors in eventual nuclear application. These areas are not discussed here. 4.07.2 Irradiation-Induced Swelling and Microstructure of Pure SiC The neutron-induced swelling of SiC has been well studied for low and intermediate temperatures (�293– 1273K). Originally, this material was investigated in 200 400 600 800 1000 1200 1400 1600 0.1 1 Fluen Ir ra d ia tio n te m p er at ur e (� C ) 1. Price (1973)4 2. Yano (1998)17 3. Senor (2003)18 4. Iseki (1990)19 5. Katoh, neutron (2006)15 6. Katoh, ion (2006)15 7. Snead (2007)16 1 1 1 1 1 1 1 3 3 4 5 55 7 7 7 7 7 66 66 6 6 66 6 6 6 Black spot defects(BSD) and/or unidentified small loops Frank loops Figure 1 Updated summary of the microstructural developme Reproduced from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S support of nuclear fuel coating1–9 and more recently, for various nuclear applications such as structural SiC composites.10 Before proceeding, it is important to distinguish neutron-induced effects on high-purity materials, such as single crystal and most forms of chemical vapor deposited (CVD) SiC, from those on lower purity forms such as sintered with additives, reaction-bonded, or polymer-derived SiC. It is well understood that the presence of significant second phases and/or poorly crystallized phases in thesemate- rials leads to unstable behavior under neutron irradia- tion,11–14 as compared to stoichiometric materials, which exhibit remarkable radiation tolerance. Discus- sion and data for this section refer only to high purity, stoichiometric, near-theoretical density SiC, unless otherwise specified. Rohm and Haas (currently Dow Chemicals) CVD SiC is an example of such material. The irradiation-induced microstructural evolu- tion of CVD SiC is roughly understood and has been reviewed recently by Katoh et al.15 An updated version of the microstructural evolution map is shown in Figure 1. However, the contribution of the defects themselves to the swelling in SiC is less understood. Below several hundred Kelvin, the observable 10 100 1000 ce (dpa) BSD and/or unidentified small loops Frank loops Unfaulted loops and/or network Voids 1 1 2 2 2 2 7 6 6 6 6 6 Larger loops dislocation network voids nt in cubic SiC during neutron and self-ion irradiation. .; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377. Tirr= 300 °C, 6 dpa Tirr= 800 °C, 7.7 dpa Dot number density » 2.2e + 24 m-3 Mean dot diameter = 1 nm Loop number density » 3.3e + 23 m-3 Mean loop diameter = 3.0 nm g = 200 g = 200 40 nm 50 nm Figure 2 Microstructure for CVD neutron irradiated at 573 and 1073K. Radiation Effects in SiC and SiC–SiC 217 microstructure of neutron-irradiated SiC is described as containing ‘black spots, which are most likely tiny clusters of self-interstitial atoms in various indeterminate configurations. For irradiation tem- peratures less than about 423 K, accumulation of strain due to the irradiation-produced defects can exceed a critical level above which the crystal becomes amorphous. This has been shown in the case of both self-ion irradiation and fast neutron irra- diation.20–22 As shown by Katoh et al.,23 the swelling at 323K under self-ion irradiation increases logarithmi- cally with dose until amorphization occurs. The swelling of neutron- and ion-amorphized SiC has been reported to be 10.8% for 343K irradiation.22 However, there is evidence that the density of amor- phous SiC will depend on the conditions of irradiation (dose, temperature, etc.)24 For temperatures above the critical amorphiza- tion temperature (423 K), the swelling increases logarithmically with the dose until it approaches saturation, with a steady decrease in the saturation swelling level with increasing irradiation tempera- ture. The dose exponents of swelling during the logarithmical period are in many cases close to two- thirds, as predicted by a kinetic model assuming planar geometry for interstitial clusters.25 This tem- perature regime is generally referred to as the point- defect swelling regime and can be roughly set between 423 and 1273K. As an example of how these ‘black spot’ defects mature in the point-defect swelling regime, Figure 2 shows neutron-irradiated microstructures at 573 and 1073 K for doses consistent with a saturation in density. While these microstructural features are generically classified as ‘black spots,’ the defects formed at 1073K are clearly coarser compared to those formed under 573K irradiation. The approach to saturation swelling is shown for High Flux Isotope Reactor (HFIR) neutron irra- diated Rohm and Haas CVD SiC in Figure 3. In this figure, the swelling is depicted in both logarith- mic (Figure 3(a)) and linear (Figure 3(b)) plots. In addition to the approach to saturation, this figure highlights two other characteristics of neutron- induced swelling of SiC. First, the swelling of SiC is highly temperature dependent. For the data given in Figure 3, the 1 dpa and saturation values of swelling at 473K are approximately five times that for 1073K irradiation. This reduced swelling with increasing irradiation temperature is primarily attrib- uted to enhanced recombination of cascade- produced Frenkel defects due to lower interstitial clustering density at higher temperatures. The sec- ond characteristic swelling behavior to note is that the swelling saturates at a relatively low dose. For damage levels of a few dpa (typically months in a fission power core), the swelling in the point-defect recombination range has found its saturation value. At higher temperatures such as 1173–1673K,4,18,26 Frank faulted loops of the interstitial type become the dominant defects observed by transmission elec- tron microscopy (TEM). It has also been reported that Frank faulted loops appear for lower tempera- ture neutron irradiation at extremely high doses.27 0 0.01 0.1 1 2 0.1 1 10 30 0 0.5 1 1.5 S w el lin g (% ) S w el lin g (% ) 2 2.5 3 CVD SiC CVD SiC Tirr= 500 °C Tirr= 400 °CTirr= 200 °C Tirr= 300 °C Tirr= 600 °C Tirr= 800 °C 200 °C 300 °C 400 °C 600 °C 800 °C 800 °C 600 °C 400 °C 200 °C 650 °C 1 2 3 4 Neutron dose (dpa)(a) (b) Neutron dose (dpa) 5 6 7 8 Figure 3 Swelling of SiC in the intermediate temperature point defect swelling regime. Reproduced from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377. 218 Radiation Effects in SiC and SiC–SiC Under silicon ion irradiation at 1673K, the develop- ment of Frank loops into dislocation networks through unfaulting reactions at high doses is reported.26 The volume associated with dislocation loops in irradiated SiC has been estimated to be on the order of 0.1%.28 At temperatures where vacancies are sufficiently mobile, vacancy clusters can be formed. Three- dimensional (3D) cavities (or voids) are the only vacancy clusters known to commonly develop to large sizes in irradiated SiC. The lowest temperature at which void formation was previously reported under neutron irradiation is 1523 K.4 Senor reported the lack of void production after neutron irradiation to 0.9 dpa at 1373 K, although voids were observed after subsequent annealing at 1773K for 1 h.18 Under silicon ion irradiation, voids start to form at 1273K at very low density and become major contributors to swelling at irradiation conditions of 1673K at >10 dpa.29 Positron annihilation and electron para- magnetic resonance studies have shown that the silicon vacancy in cubic SiC becomes mobile at 1073–1173 K.30,31 Therefore, it would not be sur- prising for void swelling to take place at as low as�1273 K at high doses, particularly for low damage rate irradiations. As previously mentioned, data on swelling of SiC in the high-temperature ‘void swelling’ regime has been somewhat limited. Recently, however, work has been carried out in the �1173–1773K range for Rohm and Haas CVD SiC irradiated in HFIR. Of particular significance to that experiment is the confidence in irradiation temperature owing to the melt-wire passive thermometry.32 Recent TEM imaging by Kondo28 clearly shows the evolution of complex defects. As an example, Figure 4 indicates sparse void formation on stacking faults for material irradiated at 1403 K. Significant growth of voids commences at 1723 K. The well-faceted voids appeared to be tetrahedrally bounded by planes, which likely provide the lowest surface energy in cubic SiC. In many cases, voids appeared to be aligned on stacking faults at all temperatures. However, intra- granular voids unattached to stacking faults were also commonly observed at 1723K. The evolution of dis- location microstructures at 1403–1723K is shown in Figure 5. In this temperature range, dislocation loops are identified to be Frank faulted loops of inter- stitial type. Evolution of the dislocation loops into dislocation networks was confirmed for irradiation at 1723K. 1130 �C, 1.8 dpa 1450 �C, 1.8 dpa 1450 �C, 5.0 dpa 1450 �C, 8.5 dpa 1130 �C, 8.5 dpa 1280 �C, 5.0 dpa (a) (b) 20 nm (d) (e) (f) (c) Figure 4 Evolution of voids in high-temperature irradiated CVD SiC. (a) g g g (b) (c) (d) (e) (f) 30 nm 1130 �C, 1.8 dpa 1450 �C, 1.8 dpa 1450 �C, 5.0 dpa 1450 �C, 8.5 dpa 1130 �C, 8.5 dpa 1280 �C, 5.0 dpa Figure 5 Evolution in dislocation networks for high-temperature irradiated CVD SiC. Radiation Effects in SiC and SiC–SiC 219 Figure 6 plots both historical data, recently pub- lished, and unpublished data on the swelling behavior of SiC over a wider range of temperature.16,33 This plot is limited to literature data on high-purity CVD SiC. A divergence from point-defect ‘saturated’ swelling to unsaturated swelling is observed in the 1273–1473 K range, although additional data in this temperature range as a function of fluence would be required to precisely define such behavior. Above 1373K, there exists a clear unsaturated swelling behavior for CVD SiC. The three divergent curves drawn in Figure 6 represent data taken at nomi- nally �1.75, 5.0, and 8.5� 1025 nm�2 (E> 0.1MeV) (assumed 1.75, 5.0, and 8.5 dpa). In the 1373–1473K temperature range, volumetric swelling is apparently at a minimum, although it increases from �0.2% to �0.4% to �0.7% for �1.75, 5.0, and 8.5 dpa, respectively. Clearly, the swelling in this temperature range has not saturated by 10 dpa. Above this minimum in swelling, the data indicates a continual swelling increase to the highest irradiation temperature of �1773–1873K. At �1773K, measured swelling Snead 2006 Price 1973 Price 1969 Blackstone 1971 Price 1973,#2 Senor 2003 Price 1973Snead 2006 Snead, unpub. Amorphization regime Saturable regime point defect swelling Nonsaturable regime void swelling 0.1 0.2 0.3 0.5 0.7 S w el lin g (% ) 1 2 3 5 7 10 20 0 200 400 600 800 Irradiation temperature (°C) 1000 1200 1400 1600 8.5 dpa 5 dpa 1.75 dpa Figure 6 Irradiation-induced swelling of SiC to high irradiation temperatures. Reproduced from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377. 220 Radiation Effects in SiC and SiC–SiC was �0.4, 1.0, and 2.0% for �1.75, 5.0, and 8.5 dpa, respectively. It was also noted in the study by Snead et al.33 that at �1773K, surface reaction between SiC and the graphite holder had taken place. However, a loss of silicon from the surface cannot be ruled out. Figure 6 includes historical data for swelling above 1273K.3,4,18,22,34,35 Specifically, Senor et al.18 report swelling for the same type of CVD SiC irra- diated in this study when irradiated in a water- moderated fission reactor (the ATR) as well. Their maximum dose, irradiation temperature, and swelling data were �1 dpa, �1373� 30K, and 0.36� 0.02%, respectively. The irradiation temperature quoted in Senor et al.’s work was a best estimate, although the authors also provide an absolute bound of 1073–1473 K for their experiment. The maximum swelling in their work (0.36� 0.02% at �1 dpa) is somewhat higher than the �0.25% swelling at 2 dpa, �1373 K, of the trend data in Figure 6. This is seen from the rightmost figure of Figure 6. Also seen in the figure is the high-temperature swelling of Price.3,4,34 The Price data, which are in the dose range of about 4–8 dpa, are in fair agreement with the measured swelling of the Snead data16,33 of Figure 6. The highest swelling material (�1523K, �6 and 10 dpa) shows the largest discrepancy, although if the temperature error bands quoted by the various authors are taken into account, the data are conceivable more in alignment. It is also noted that the Price material may have had some excess silicon leading to higher swelling as compared to stoichiometric material. As mentioned earlier, the microstructural evolu- tion of irradiated SiC is roughly understood, at least for temperatures up to �1373K. The swelling near the critical amorphization temperature (�423K) is classically described as the differential strain between the single interstitial, or tiny interstitial clusters, immobile vacancies, and antisite defects. As the tem- perature increases above the critical amorphization temperature, the number of defects surviving the postcascade thermally activated recombination is reduced and the mobility of both silicon and carbon interstitials becomes significant. For temperatures exceeding �1273K, microstructural studies have noted the presence of both Frank loops and tiny voids, indicating limited mobility of vacancies. Radiation Effects in SiC and SiC–SiC 221 The apparent increase in swelling with dose in the 1373–1873K range seen in Figure 6 and the observed production of voids are interesting considering that the maximum irradiation temperature (�1773 K) in Figure 6 is �0.65 of the melting (dissociation) temperature (Tm) for SiC. Here, we have assumed Olesinski and Abbaschian’s36 value of 2818K where stoichiometric SiC transforms into Cþ liquid phase. This value of 0.65Tm is high when viewed in compar- ison to fcc metal systems where void swelling typi- cally begins at �0.35Tm, goes through a maximum value, and decreases to nil swelling by �0.55Tm. (It is noted that the melting and dissociation temperatures of SiC are somewhat variable in the literature. How- ever, even considering this variability, the previous statement is accurate). If, as the swelling data seems to indicate, the voids in SiC are continuing to grow in SiC irradiated to 1773 K, the energies for diffusion of either the Si or C vacancy or both must be quite high, as are the binding energies for clustered vacancies. This has been shown through theoretical work in the literature.37–40 However, it is to be noted that the defect energetics obtained from this body of work, and in particular those of the Si and C vacancies within SiC, vary widely. Perhaps, the work of Bockstedte et al.,39 which follows an ab initio approach, is the most accurate, yielding a ground state migra- tion energy of 3.5 and 3.4 eV for Si and C vacancies, respectively. It was also noted by Bockstedte et al.41 that the assumed charge state of the vacancy affects the calculated migration energy. Specifically, the car- bon vacancy in the þ1 and þ2 charge state increases from 3.5 to 4.1 and 5.2 eV, respectively, and that of silicon in the –1 and –2 charge state decreases from 3.4 to 3.2 and 2.4 eV, respectively. Several papers discuss the vacancy and vacancy cluster mobil- ity measured experimentally. The silicon monova- cancy has been shown to be mobile below 1273K. Using electron spin resonance, Itoh et al.30 found the irradiation-produced T1 center in 3C–SiC disappear- ing above 1023K.TheT1 center was later confirmed to be an Si vacancy.31 Using electron spin resonance, the carbon vacancy in 6H–SiC is shown to anneal above 1673K.42 Using isochronal annealing and positron lifetime analysis, Lam et al.40 have shown a carbon– silicon vacancy complex to dissociate above �1773K for the same 6H single crystal materials studied here. From a practical nuclear application point of view, the swelling data for CVD SiC can be broken down into the amorphization regime (1273K. From the data of Figure 6, it is still unclear where the actual transition into the unsaturated swelling begins. Furthermore, while there is an increase in swelling in the 1273–1773 K range, as the dose is increased from �1.75, 5.0, and 8.5� 1025 nm�2 (E> 0.1MeV), swelling is close to linear with neutron doses, and it is unclear how swelling will increase as a function of dose above 10 dpa. For example, swelling by voids estimated from the TEM examina- tion accounts for only relatively small fractions of the total swelling even in the void swelling regime. Analo- gous to the typical swelling behavior in metals, void growth may cause steady-state swelling after a certain transition dose regime. However, dose dependence of the swelling due to the nonvoid contribution remains to be understood. Extrapolation of swelling outside of the dose range of Figure 6 is therefore problematic. 4.07.3 Irradiation-Induced Thermal Conductivity Degradation of Monolithic SiC According to Lee et al.,43 the effect of neutron irradi- ation on the specific heat of SiC was negligibly small. The specific heat of SiC is therefore assumed to be unchanged by neutron irradiation, although this has not been verified at high dose. A single study5 also indicated that stored energy (Wigner energy) occurs in SiC irradiated in the point defect regime. The relative amount of stored energy appears to be less than that of graphite.44 Because of a low density of valence band electrons, thermal conductivity of most ceramic materials, SiC in particular, is based on phonon transport. For a ceramic at the relatively high temperature associated with nuclear applications, the conduction heat can be generally described by the strength of the individual contributors to phonon scattering: grain boundary scattering (1/Kgb), phonon–phonon interaction (or Umklapp scattering 1/Ku), and defect scattering (1/Kd). Because scattering of each of these types occurs at differing phonon frequencies and can be considered separable, the summed thermal resis- tance for a material can be simply described as the summation of the individual components; that is, 1/K¼ 1/Kgbþ 1/Kuþ 1/Kd. As seen in Figure 7, the unirradiated thermal conductivity of SiC is highly dependent on the nature of the material (grain size, impurities, etc.) and the temperature. While materials can be optimized for low intrinsic defect, impurity, Note 0 100 200 300 400 500 0 500 1000 1500 2000 Temperature (K) Th er m al c on d uc tiv ity (W m –1 K –1 ) Legend Reference Material Note N/R Single Crystal CVD Grain size ~5µm CVD Morton CVD CVD Morton CVD CVD Grain size ~10µm CVD Grain size ~3µm CVD Grain size >10µm CVD Grain size 0.001 5 10 30 50 R oo m t em p er at ur e th er m al c on d uc tiv ity (W m −1 K –1 ) 100 300 Nonirradiated thermal conductivity 381 ± 26 W m−1K−1 (except Rohde, 108 W m−1K−1) CVD SiC Tirr= 750-850 °C Tirr= 600 °C Tirr= 80 °C Tirr= 800 °C Tirr= 200 °C Tirr= 480-550 °C Tirr= 400-480 °C Tirr= 375 °C Tirr= 200 °C 0.01 0.1 Fast neutron dose (´ 1025n m-2 E > 0.1 MeV) 1 10 Figure 8 Degradation in room-temperature thermal conductivity for CVD SiC. (Rohde data designated as �.) Reproduced from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377. Radiation Effects in SiC and SiC–SiC 223 conductivity of any high-purity CVD SiC, indepen- dent of the starting thermal conductivity. The accu- mulation in thermal defect resistance generated from the data of Figure 8 is shown in Figure 9. Another result of the previously reported analysis on irradiated CVD SiC16,52 is that the thermal defect resistance appears to be directly proportional to the irradiation-induced swelling, although the data-set for making the previous assertion was somewhat lim- ited. A compilation plot including the previous data- set as well as the new data of Figure 9 is shown in Figure 10. It is clear from this plot that a linear relationship exists between swelling and thermal defect resistance. Moreover, there does not appear to be any effect of irradiation temperature on this result. The fact that the thermal defect resistance is proportional to the irradiation-induced swelling allows a rough estimate of thermal conductivity. As measurement of thermal conductivity for the SiC TRISO shell is not practical, while measurement of density is routine, this finding allows an indirect determination of thermal conductivity by measure- ment of the density change in the TRISO SiC shell by means of a density gradient column or some other technique. The thermal conductivity degradation discussed up to this point has been for irradiation temperature associated with the point defect regime. For irra- diation above this temperature (the nonsaturable void swelling regime), the thermal properties are not expected to saturate (at least at low dpa). The primary reason for this is that the formation of voids and other complex defects in the high-temperature regime (which contributes to the unsaturated swelling as seen in Figure 6) contributes to phonon scattering, and these defects will not saturate. Moreover, it has been shown that the linear relationship that existed between swelling and thermal defect resistance (as seen in Figure 10) does not exist in this elevated temperature irradiation regime.16,52 This underlines the fact that the phonon scattering and swelling are not controlled by the same defects in the lower temperature ‘saturable,’ and elevated temperature ‘nonsaturable’ irradiation regimes. A compilation plot of room-temperature thermal conductivity as a function of irradiation temperature for the saturable and nonsaturable temperature regimes is given in Figure 11. By comparison to the unirradiated room- temperature conductivity value of �280Wm�1 K�1, 0.001 0.001 0.01 0.1 0.01 0.1 Fast neutron dose (´ 1025n m-2E > 0.1 MeV) Th er m al d ef ec t re si st an ce (m K W –1 ) 1 10 T irr = 80 °C T irr = 200 °C T irr = 375 °C T irr = 400-480 °C T irr = 480-550 °C T irr = 600 °C T irr = 750-850 °C Rhode, T irr = 80 °C T irr = 80 °C T irr = 200 °C T irr = 375 °C T irr = 400-480 °C T irr = 480-550 °C T irr = 600 °C T irr = 750-800 °C Figure 9 Thermal defect resistance for stoichiometric CVD SiC as a function of neutron dose. (Rohde data designated as �.) Reproduced from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377. 224 Radiation Effects in SiC and SiC–SiC it is clear that the thermal conductivity degradation in the highest temperature regimes is less dramatic, even though the swelling is rapidly increasing (see Figure 6). This is opposite to the behavior in the lower temperature, saturable regime, where high swelling corresponds to extreme reduction in thermal conductivity. Unfortunately, the data on thermal con- ductivity reduction in the nonsaturable regime is limited, and given the lack of knowledge of the spe- cific defects governing the phonon scattering, it is not possible to accurately predict behavior outside of the data-set of Figure 6. Data presented thus far has been limited to mea- surement of thermal conductivity at room tempera- ture. As described in Figure 7, there is a dramatic dependence of thermal conductivity on measure- ment temperature. The temperature dependence of irradiated materials can be found by applying the temperature dependence of unirradiated SiC (the Umklapp thermal resistance term) to the as-neutron-degraded room-temperature values. This approximation (dashed lines) is compared to actual data (solid lines) in Figure 12 and shows fair corre- spondence. However, it is clear that such a treatment systematically underestimated the thermal conduc- tivity degradation. This implies temperature depen- dence on the defect scattering that is not presently understood. 4.07.4 Effect of Irradiation on the Mechanical Properties of Monolithic SiC 4.07.4.1 Elastic Modulus of Monolithic SiC Figure 13 summarizes the irradiation temperature dependence of the elastic modulus change. Irradia- tion generally reduces modulus to a greater extent for lower temperature irradiation. The modulus reduction becomes negligible when irradiation temperature reaches or exceeds �1273K. There seems to be an indistinct stage between 1073 and 1273K. As expected, the elastic modulus trends with ‘point defect swelling’ of SiC. Point defect swelling is an isotropic volume X X X Tirr= 200 °C Tirr= 300 °C Tirr= 375 °C Tirr= 400-480 °C Tirr= 480-550 °C Tirr= 750-850 °C Tirr= 80 °C Tirr= 600 °C 100 10 0 0.5 1 1.5 Swelling (%) 2 2.5 0 0.02 0.04 0.06 Th er m al d ef ec t re si st an ce (m K W -1 ) R oo m t em p er at ur e th er m al c on d uc tiv ity (W m -1 K –1 ) 0.08 Open symbols Right Closed symbols Left 0.1 0.12 0.14 Figure 10 The room-temperature thermal conductivity and thermal defect resistance as a function of irradiation-induced density change. Reproduced from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377. Radiation Effects in SiC and SiC–SiC 225 expansion that is believed to occur by lattice relaxation due to accumulated isolated point defects and small point defect clusters during irradiation at tempera- tures where vacancies are not readily mobile. In SiC irradiated in the point defect swelling regime, a fairly good agreement between dimensional expansion and lattice spacing has been confirmed by X-ray diffrac- tometry studies. In contrast, the data in the nonsatur- able swelling regime is somewhat limited, although the data suggest that there is little reduction in elastic modulus in spite of the swelling being relatively large. However, in this regime, the defects responsible for swelling are voids and other relatively large defects, which would have less of an effect on elastic modulus as compared to point defects. An estimation of the influence of lattice relaxa- tion on elastic modulus was attempted using the Tersoff potential.54 The result predicted a linear lattice swelling of 1% causing approximately 10% reduction in elastic modulus (Figure 14). The predicted elastic modulus change could be varied by more than 10% depending on a selection of interatomic potential, with the Tersoff potential giving a relatively high sensitivity of modulus to the interatomic distance among various empirical interatomic potential functions. Therefore, the measured elastic modulus changes observed in this experiment are generally greater than the theoreti- cal prediction. It is noted that the methods applied for generating the data of Figure 14 are various and of differing quality. Typically, elastic modulus as measured by nanoindentation, which sometimes is the only alternative for miniature specimens, tends to give widely scattered and less reliable data than themechan- ical or sonic modulus methods. Nonetheless, it is clear that the lattice expansion is a major cause of the irradiation-induced elastic modulus reduction in SiC. 0 0 20 40 60 R oo m t em p er at ur e th er m al co nd uc tiv ity (W m -1 K -1 ) 80 100 120 Amorphization regime Saturable regime Unsaturable regime 200 400 600 800 1000 Irradiation temperature (°C) 1200 1400 ~8.5 dpa ~5 dpa ~1.75 dpa Nonirradiated conductivity ~280 W m-1K-1 Rohm and Haas CVD SiC 1600 Figure 11 Room-temperature thermal conductivity in the saturable and nonsaturable regime. Reproduced from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377. 226 Radiation Effects in SiC and SiC–SiC 4.07.4.2 Hardness of Monolithic SiC The irradiation effect on nanoindentation hardness of Rohm and Haas CVD SiC in a fluence range of 0.1–18.7 dpa is summarized in Figure 15. It is interest- ing to note that the nanoindentation hardness exhibits relatively small scatter for the individual experiments, and the trend in data as a function of temperature is uniform. This observation is in contrast to both the flexural strength and the indentation fracture tough- ness data, which indicate a broad peak at an intermedi- ate temperature and a relatively large scatter. It is worth noting that nanoindentation hardness of brittle ceramics is, in general, determined primarily by the dynamic crack extension resistance in the near surface bulk material, and therefore should be more relevant to fracture toughness than to plastic deformation resis- tance. However, surface effects of the original sample affect the nanoindentation hardness less, as the samples are generally polished prior to testing. 4.07.4.3 Fracture Toughness of Monolithic SiC The effect of irradiation on the fracture toughness of Rohm andHaas CVD SiC is summarized inFigure 16. This compilation plots data using the Chevron notched beam technique, although the bulk of the data sets report Vicker’s or nanoindentation gener- ated data.55–57 The general trend is that the irradiation-induced toughening seems to be signifi- cant at 573–1273K for the indentation fracture tough- ness data, in spite of the decrease in elastic modulus, which confirms the increase in fracture energy caused by irradiation. The scatter of the indentation fracture toughness data among different experiments is likely caused by both the condition of the sample surface and the lack of standardized experimental procedures. Typ- ically, indentation should be applied on the polished surfaces, but conditions of polishing are not always provided in literature.Moreover, the crack length mea- surements are done using optical microscopy, conven- tional scanning electron microscopy (SEM), or field emission SEM, all of which may give very different crack visibility. In addition, a few different models have been used for derivation of the fracture toughness. In conclusion, indentation fracture toughness techni- ques can be used only for qualitative comparisonwithin a consistent set of experiments. It is noted that the experiment employing the Chevron notched beam technique also indicates the irradiation-induced tough- ening, although scatters of toughness values were even greater. These results lead to the conclusion that, in the intermediate irradiation temperature range, the increase of the fracture toughness of SiC exists. + +++ ++ ++ 0 10 (a) 20 30 40 50 60 80 100 200 300 Rohm and Haas CVD SiC irradiated in HFIR at 800 ˚C Nonirradiated 0.05 dpa 0.5 dpa 1.94 dpa 4.3 dpa Calculated based on Umklapp scattering Fit to data 400 500 200 400 Irradiation and measurement temperature ( ˚C) Th er m al c on d uc tiv ity (W m -1 K –1 ) 600 800 1000 0 (b) 0 20 40 60 80 100 120 100 200 300 Rohm and Haas CVD SiC irradiated in HFIR Fit to data Calculation based on Umklapp scattering 7.5 dpa Tirr= 1500 �C 1.75 dpa Tirr= 1500 �C Irradiation and measurement temperature ( �C) Th er m al c on d uc tiv ity (W m -1 K –1 ) 400 500 600 700 800 + + + + + + + + + + + + 0 (c) 0 5 10 15 Th er m al c on d uc tiv ity (W m -1 K –1 ) 20 25 30 35 40 100 200 300 400 Irradiation and measurement temperature ( ˚C) Rohm and Haas CVD SiC irradiated in HFIR 500 Fit to data Calculation based on Umklapp scattering 600 700 800 7.5 dpa 1.75 dpa 8.5 dpa 5 dpa 1.75 dpa Tirr= 1060 �C Tirr= 1050 �C T irr = 1020 �C 5 dpa Figure 12 Effect of temperature on the conductivity of irradiated SiC. (a) Tirr¼1073K, (b) Tirr¼1773K, and (c) Tirr¼1293–1333K. Reproduced from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377. Radiation Effects in SiC and SiC–SiC 227 4.07.4.4 Strength and Statistical Variation in Strength for Monolithic SiC There have been several studies on the effect of neu- tron irradiation on the strength of various types of SiC forms including reaction-bonded, sintered, pressure- less sintered, and CVD SiC materials.1,11,13,14,58,60–65 The strength of SiC depends significantly on stoichi- ometry under neutron irradiation. Both the sintered SiC and the reaction-bonded SiC forms exhibit significant deterioration in strength by neutron irra- diation (Figure 17).13 The presence of impurities such as sintering additives for sintered SiC and excess Osborne (1999)55,HFIR, 2 dpa Nanoindentation 4pt. bend 200 0.60 0.70 0.80 0.90 1.00 1.10 1.20 400 600 800 1000 Irradiation temperature (K) R el at iv e Yo un g’ s m od ul us 1200 1400 1600 1800 2000 Sonic resonance Nogami (2002)56,HFIR + HFBR, 0.15–7.7 dpa Park (2003)57,DuET 5.1 MeV Si, 3 dpa Katoh (2005)58,HFIR – 14J, 6.0–7.7 dpa Snead (2007)16, HFIR – METS, 0.7–8.6 dpa Price(1977)59,ETR, 2.8–12.2 dpa Snead (2007)16, HFIR, 0.7–4.2 dpa Figure 13 Irradiation temperature dependence of irradiated elastic modulus of CVD SiC, at ambient temperature, normalized to unirradiated values. The error bars are showing standard deviations for all the neutron data points and ranges of data scatter for the ion data points. Reproduced from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377. 1.20 1.10 1.00 0.90 0.80 0.70 0.0 0.5 1.0 1.5 2.0 Volumetric swelling (%) 2.5 3.0 3.5 Osborne (1999)55,HFIR, 2 dpa Snead (2007)16, HFIR - METS, 1.7–8.6 dpa Price (1977)59, ETR, 2.8–12.2 dpa Nogami (2002)56,HFIR + HFBR, 0.15–7.7 dpa Park (2003)57,DuET 5.1MeV Si, 3 dpa Katoh (2005)58,HFIR - 14J, 6.0–7.7 dpa Snead (2007)16, HFIR, 0.7–4.2 dpa Model (Tersoff potential) Nanoindentation 4pt. bend Sonic resonance R el at iv e Y ou ng ’s m od ul us Figure 14 Irradiation-induced change of elastic modulus versus swelling of CVD SiC. An estimation of the influence of lattice relaxation on elastic modulus is calculated using Tersoff potential. Reproduced from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377. 228 Radiation Effects in SiC and SiC–SiC Irr ad ia te d na no -i nd en ta tio n ha rd ne ss no rm al iz ed to u ni rr ad ia te d va lu es Katoh (2005)58, HFIR-14J, 6.0–7.7 dpa Nogami (2002)56, HFIR+HFBR, 0.2–7.7 dpa Osborne (1999)55, HFIR, 0.1–5.0 dpa Park (2003)57, DuET 5.1 MeV Si, 3 dpa 1.6 1.4 1.2 1 0.8 0.6 0.4 0.2 0 200 400 600 800 1000 Irradiation temperature (K) 1200 1400 1600 Figure 15 Nanoindentation hardness of CVD SiC at ambient temperature as a function of irradiation temperature. Reproduced from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377. Nogami (2002)56, HFIR+HFBR, 0.2–7.7 dpa Osborne (1999)55, HFIR, 0.1–5.0 dpa Park (2003)57, DuET 5.1 MeVSi, 3 dpa Snead (2007)16, HFIR, 0.7–4.2 dpa 0 200 400 600 800 1000 1200 1400 1600 1800 0.2 0.4 0.6 0.8 1 1.2 1.4 1.6 1.8 2 Irradiation temperature (K) Ir ra d ia te d fr ac tu re t ou gh ne ss no rm al iz ed t o un irr ad ia te d v al ue s Figure 16 Indentation fracture toughness of CVD SiC at ambient temperature as a function of irradiation temperature. Reproduced from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377. Radiation Effects in SiC and SiC–SiC 229 Si for reaction-bonded SiC, which typically segregate to grain boundaries during sintering, tends to have a significant influence on strength under neutron irradiation. For the case of sintered SiC with boron compounds as sintering additives, the reaction of 10B(n, a)7Li causes the accumulation of helium bub- bles at and near the grain boundary phases under neutron irradiation.60–63 In contrast, unmatched 1.5 Hot pressed and sintered SiC forms N or m al iz ed s tr en gt h 1.0 0.50 0.0 0.1 1 10 Dose (dpa) 100 Matheney (1979)64 Matthews (1974)11 Iseki (1990)19 RB(reaction bonded) Iseki (1990)19 PLS(pressureless sintered) Iseki (1990)19 HP(hot pressed) Price (1982)63 NC-430 Price (1982)63 Carborundum Dienst (1991)62 0.75% B Dienst (1991)62 1.9% B Corelli (1983)60 NC430 0 Figure 17 Fluence dependence of irradiated flexural strength of hot-pressed and sintered SiC normalized to unirradiated strength. Irradiation is variable but in the saturable swelling regime. Reproduced from Newsome, G. A.; Snead, L. L.; Hinoki, T.; Katoh, Y.; Peters, D. J. Nucl. Mater. 2007, 371, 76–89. 2.0 1.5 CVD silicon carbide 1.0 0.50 0.0 0.1 1 10 100 N or m al iz ed s tr en gt h Dose (dpa) Snead Price Dienst Newsome 300 �C Newsome 500 �C Newsome 800 �C Figure 18 Flexural strength of CVD SiC at ambient temperature as a function of irradiation dose. Reproduced from Newsome, G. A.; Snead, L. L.; Hinoki, T.; Katoh, Y.; Peters, D. J. Nucl. Mater. 2007, 371, 76–89. 230 Radiation Effects in SiC and SiC–SiC swelling between Si and SiC for reaction-bonded SiC causes disruption at the grain boundary, severely reducing the strength.11,19,60–64 Meanwhile, the high- purity materials such as CVD SiC exhibit superior irradiation resistance. The irradiation effect on flexural strength of Rohm and Haas CVD SiC in a fluence range of 0.15–30 dpa is summarized in Figure 18. In comparing Figure 18 with Figure 17, it is clear that CVD SiC retains stability in strength to a much higher dose than the sintered and reaction-bonded forms of SiC. It is to be noted that in Figure 18, the data of Dienst does indicate a significant as-irradiated degradation in strength around 15 dpa. However, such degradation is not seen for the �30 dpa irradiation of Snead. It is speculated that the degradation in the Dienst data may Radiation Effects in SiC and SiC–SiC 231 have been due to statistical limitations of the study and/or due to issues with sample handling postirradi- ation. This issue is discussed in the Dienst reference.65 A compilation of strength data as a function of irradia- tion temperature is given in Figure 18, indicating no apparent correlation for the dose and temperature ranges studied. However, as with the fracture tough- ness data, irradiation-induced strengthening seems to be significant at 573–1273K. The large scatter in flex- ural strength of brittle ceramics is inevitable, as the fracture strength is determined by the effective frac- ture toughness, morphology, and characteristics of the flaw that caused the fracture. Irradiation possibly modifies both the flaw characteristics and the fracture toughness through potential surface modification, relaxation of the machining-induced local stress, modifications of elastic properties, and fracture energy. A typical means of describing the failure of cera- mics is through the use of Weibull statistics, which is a departure from the analysis of data that is assumed to follow a normal Gaussian distribution. In the two- parameter Weibull formalism, sometimes referred to as a weakest-link treatment, the failure probability F is described as FðxÞ ¼ 1� e�ðxm=x0Þ where m is the Weibull modulus and xo is the distribution size parameter. A change in the Weibull statistics, indicating a higher scatter in as-irradiated flexural strength has been observed by previous authors, although the point could not be made convincingly because of limitations in the number of tests observed. In the earliest work known to the authors, Sheldon66 noted a 14% decrease in crushing strength of highly irradiated CVD SiC shells with an increase of the coefficient of variation from 8% to 14%. Price63 went on to a 4-point bend test using relatively thin (�0.6mm) strips of CVD SiC deposited onto a graphite substrate. In his work, the flexural strength following a �9.4� 1025 nm�2 (E> 0.1MeV) irradiation was unchanged within the statistical scatter, but the scatter itself increased from about 10 to 30% of the mean flexural strength as described assuming a normal distribution. Unfortu- nately, there were not sufficient samples in Price’s work to infer Weibull parameters. In more recent work by Dienst,65 the Weibull modulus was reported to decrease from about 10 for irradiation of�1� 1026 nm�2 (E> 0.1MeV). However, it is worth noting that the Dienst work used a very limited sample population (about 10 bars.) Statistically meaningful data sets on the effect of flexural strength of CVD SiC have been re- ported by Newsome and coworkers14 and Katoh and coworkers.58,67 Figure 19 shows a compilationWeibull plot of the flexural strength of unirradiated and irra- diated Rohm and Haas CVD SiC taken from the two separate irradiation experiments carried out by New- some and cowokers14 and Katoh and coworkers.58,67 The sample population was in excess of 30 for each case. In Figure 19(a), the data was arranged by irradia- tion temperature, including data for unirradiated samples and 1.5–4.6� 1026 nm�2 (E> 0.1MeV) dose range. It is likely that the Weibull modulus decreased by irradiation, appearing to be dependent on irradia- tion temperature. This is not easily visualized through inspection of Figure 19(a) unless one notes that there are significantly more low stress fractures populating the 573K population. The scale parameters of flexural strength of unirradiated materials and materials irra- diated at 573, 773, and 1073Kwere 450, 618, 578, and 592MPa, respectively. The Weibull modulus of the flexural strength of unirradiated materials and materi- als irradiated at 573, 773, and 1073Kwere 9.6, 6.2, 5.5, and 8.7, respectively, with significant uncertainty. The work of Katoh, on identical material irra- diated at the same temperature as in the Newsome work, is at a slightly higher irradiation dose than the data of Newsome. As seen in Figure 19(b), the effect on the Weibull modulus undergoes a trend similar to that of Newsome, although the modulus for the 773K and 1073 K irradiation of Katoh remained almost unchanged. Given the data discussed on the effect of irradiation on theWeibull modulus and scale parameter of CVD SiC bend bars, it is clear that the material is somewhat strengthened and that the Weibull modulus may undergo irradiation-induced change, with the greatest decrease occurring for the lowest temperature irradiation. The fracture strength and failure statistics of tubular SiC ‘TRISO surrogates’ have been deter- mined by the internal pressurization test and the results are plotted in Figure 20. Thin-walled tubular SiC specimens of 1.22mm outer diameter, 0.1 mm wall thickness, and 5.8mm length were produced by the fluidized-bed technique alongside TRISO fuels.68 The specimens were irradiated in the HFIR to 1.9 and 4.2� 1025 nm�2 (E> 0.1MeV) at 1293 and 1553K. In the internal pressurization test, tensile hoop stress was induced in the wall of the tubular specimens by compressively loading a polyurethane insert.68,69 In Figure 20, Weibull plots of the flexural strength and internal pressurization fracture strength -6 -5 -4 -3 -2 -1 0 5.0 5.5 6.0 6.5 7.0 5.0 5.5 6.0 6.5 7.0 1 2 3 200 300 400 500 600 800 1000 200 300 400 500 600 800 1000 ln (ln (1 /( 1– F i ))) ln (ln (1 /( 1– F i ))) ln(si) ln(si) si (MPa) si (MPa) (a) (b) Nonirrad. m = 9.6 300 ºC, 2.0 dpa m = 6.2 800 ºC, 2.0 dpa m = 8.7 500 ºC, 2.0 dpa m = 5.5 -6 -5 -4 -3 -2 -1 0 1 2 3 Nonirrad. m = 9.9 300 ºC, 6.0 dpa m = 5.5 500 ºC, 6.0 dpa m = 10.8 800 ºC, 7.7 dpa m = 7.9 Figure 19 Weibull plots of flexural strength of unirradiated and irradiated CVD SiC in the dose range of (a) 1.5–4.6�1025 nm�2 (E > 0.1MeV) from Newsome14 and (b) 7.7� 1025 n cm�2 (E > 0.1MeV) from Katoh.58 232 Radiation Effects in SiC and SiC–SiC of unirradiated and irradiated samples are presented. As with the Newsome and Katoh data, the sample population is large enough to be considered statistically meaningful. From this data, the mean fracture stress of tubular specimens is seen to increase to 337MPa (from 297MPa) and the Weibull modulus slightly decreased to 3.9 (from 6.9) after irradiation to 1.9� 1025 nm�2 (E> 0.1MeV) dpa at 1293K. The mean fracture stresses and Weibull moduli at 4.2� 1025 nm�2 (E> 0.1MeV) were similar to those at 1.9 dpa. The results for 4.2 dpa irradiation indicate that by increasing the irradiation temperature from 1293 to 1553K, no discernible change in fracture stress distribution occurred. The horizontal shift indicates a simple toughening or an increase in frac- ture toughness alone. While the data for these surro- gate TRISO samples, irradiated through internal compression, are somewhat limited, the findings indi- cate that the trend in strength and statistics of failure are consistent with those found for the bend bars. Therefore, the general findings of the bend bar irra- diation on strength and Weibull modulus appear -5 -4 -3 -2 -1 0 1 5.0 5.5 6.0 6.5 7.0 2 3 Nonirrad. m = 7.6 1020 �C, 1.9 dpa m = 4.4 1280 �C, 4.2 dpa m = 3.8 1020 �C, 4.2 dpa m = 5.4 si (MPa) ln (ln (1 /( 1– F i ))) ln(si) 200 300 400 500 600 800 1000 Figure 20 Weibull statistical fracture strength of CVD SiC measured by the internal pressurization test. Reproduced from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377. Radiation Effects in SiC and SiC–SiC 233 appropriate for application to TRISO fuel model- ing. Specifically, a slight increase in the mean strength is expected (although it may be less signifi- cant at higher temperatures), and the statistical spread of the fracture data as described by the Weibull modulus may broaden. Unfortunately, a precise description of how the Weibull modulus trends with irradiation dose and temperature is not yet possible, although within the dose range and temperature covered by the data in Figures 19 and 20, a modest reduction is possible. 4.07.5 Irradiation Creep of SiC Irradiation creep is defined as the difference in dimensional changes between a stressed and an unstressed sample irradiated under identical condi- tions. Irradiation creep is important for structural materials for nuclear services as it is a major contributor to the dimensional instability of irra- diated materials at temperatures where thermal creep is negligible. However, studies on irradiation creep of SiC(-based materials) have so far been very limited, although it is of high importance for the behavior of the SiC TRISO shell. Price published the result of the irradiation creep study on CVD SiC in 1977.59 In this work, elastically bent strip samples of CVD SiC were irra- diated in a fission reactor, and the steady-state creep compliance was estimated to be in the order of 10–38 (Pa dpam�2 (E> 0.18MeV))�1 at 1053– 1403K. Scholz and coworkers measured the transient creep deformation of SCS-6 CVD SiC-based fiber, which was torsionally loaded under penetrating pro- ton or deuteron beam irradiation.70–73 They reported several important observations including the linear stress and flux dependency of the tangential primary creep rate at 873 K, and the negative temperature dependence of primary creep strain at the same dose. Recently, Katoh et al. determined the bend stress relaxation (BSR) creep in Rohm and Haas CVD SiC and Hoya monocrystalline 3C–SiC during irradia- tion in HFIR and JMTR at 673–1353K.74 The results reported for CVD SiC are summarized in Table 1. In the BSR irradiation creep experiment by Katoh et al., the creep strain for CVD SiC exhibited a weak temperature dependence at Table 1 Irradiation creep data for CVD SiC from bend stress relaxation experiments Tirr ( �C) Fluence (dpa) Reactor Initial/final bend stress (MPa) Initial/final bend strain (�10�4) Creep strain (�10�4) BSR ratio m Average creep compliance �10�6 (MPadpa)�1 CVD SiC 400 0.6 JMTR 82/60 1.80/1.39 0.41 0.77 0.97 600 0.2 JMTR 81/57 1.80/1.31 0.49 0.73 3.5 600 0.6 JMTR 81/46 1.80/1.05 0.75 0.58 2.0 640 3.7 HFIR 87/36 1.95/0.83 1.12 0.42 0.50 700 0.7 HFIR 102/72 2.27/1.64 0.63 0.72 1.1 750 0.6 JMTR 80/55 1.80/1.27 0.53 0.71 1.3 1030 0.7 HFIR 85/61 1.94/1.42 0.52 0.73 0.97 1080 4.2 HFIR 101/8 2.29/0.19 2.10 0.08 0.91 3C–SiC 640 3.7 HFIR 87/30 1.94/0.68 1.26 0.35 0.59 700 0.7 HFIR 102/90 2.27/2.06 0.21 0.87 0.34 1030 0.7 HFIR 86/57 1.94/1.31 0.63 0.67 1.2 1080 4.2 HFIR 101/1 2.29/0.02 2.27 0.01 1.1 234 Radiation Effects in SiC and SiC–SiC early domination of the transient irradiation creep. The transient creep is speculatively caused by the rapid development of defect clusters and the structural relaxation of as-grown defects during early stages of irradiation damage accumulation. At �1353K, irradi- ation creep mechanisms, which are common to metals, are likely operating. In metals, steady-state irradiation creep rates are generally proportional to the applied stress and neutron (or other projectiles) flux, f,75,76 and there- fore, irradiation creep compliance, B, has been con- veniently introduced75: eic ¼ sðBfþ DSÞ where S is void swelling and D is a coefficient of swelling–creep coupling. Ignoring the swelling– creep coupling term (valid in the saturable swelling regime), preliminary estimations of the steady-state irradiation creep compliance of CVD SiC were given to be 2.7� 2.6� 10–7 and 1.5� 0.8� 10–6 (MPa dpa)�1 at �873–�1223K and �1353 K, respectively. If linear-averaged, creep compliances of 1–2� 10–6 (MPa dpa)�1 were obtained for doses of 0.6–0.7 dpa at all temperatures. Monocrystalline 3C–SiC samples exhibited a significantly smaller transient creep strain by 0.7 dpa and a greater subsequent deformation when loaded along direction. To better define the irradiation creep behavior of SiC and the underlying physical mechanisms, it will be essential to further examine the stress depen- dence, dose dependence, effect of crystallographic orientation, microstructures of the crept samples, and the coupling between irradiation creep and swelling. 4.07.6 Silicon Carbide Composites Under Irradiation SiC composites are a family of materials of varied constituents and architectures. Up to the point of writing this chapter, nuclear-grade SiC composites (those specifically developed for application in fast neutron environments and exhibiting neutron irradi- ation damage resistance) are more precisely defined as continuous fiber-reinforced ceramic composites. The history of development for these materials has been reviewed in a number of publications.29,77–79 The primary constituents of these nuclear-grade com- posites are the continuous SiC fiber, a fiber/matrix interphase material that can be SiC or pyrolytic graph- ite or a combination of the two, and a matrix of SiC infiltrated into the woven fiber preform. The most common matrix material is derived from chemical vapor infiltration (CVI), and is essentially identical in structure, properties, and irradiation response to the CVD SiC discussed in previous sections. While there has been little direct study on the effects of irradiation on the material properties of the SiC interphase, it can be assumed that it would also behave in a similar manner to the SiC matrix. However, the effect of neutron irradiation on pyrolytic graphite interphase (if used) will be substantially different from that on both matrix and fiber. While the effect of irradiation on the underlying properties of graphite interphase has not been well studied, it can be assumed that the interphase will behave in a similar manner to nuclear graphite (discussed inChapter 4.05, Radiation Dam- age of Reactor Pressure Vessel Steels). (a) (b) Y Y 300 mm Z X PyC Hi-Nicalon type-S fiber SiC matrix SiC/PyC multilayer 1 mm (c) Figure 21 Example of braided nuclear-grade SiC/SiC composite. Fiber: Hi-Nicalon™ Type-S; Interphase: Multilayer SiC with pyrolytic carbon; Matrix: CVI SiC deposited through an isothermal process. Reproduced from Nozawa, T.; Lara-Curzio, E.; Katoh, Y.; Shinavski, R. J. Tensile properties of advanced SiC/SiC composites for nuclear control rod applications. Wiley: 2007; pp 223–234. Radiation Effects in SiC and SiC–SiC 235 An example of an SiC/SiC composite that has been developed for high-temperature gas-cooled reactor control rod applications is shown in Figure 21. The basic textile weaving of the composite is evident on inspection of Figure 21(a). In this case, a �55� weave is depicted. For the polished section of Figure 21(b), large voids, which are an unavoidable characteristic of chemical vapor infiltrated materials and also the primary reason why it is difficult to produce gas- impermeable SiC/SiC composite, are clearly observed. In Figure 21(c), the complicated structure of the interphase is seen. In this case, alternating layers of SiC and pyrolytic graphite have been applied. The pyrolytic graphite layer between the SiC layers is quite thin (tens of nanometer), with a relatively thick graphite layer in contact with the fiber itself. From the earliest study of SiC/SiC composites under irradiation, it was clear that the fiber was the key to performance. As with the impure forms of SiC monolithic ceramics (cf. Figure 17), the impure and oxygen-rich early grades of SiC fiber (trade name Nicalon™) were quite unstable under neutron irra- diation.12,80,81 Researchers were able to directly link an irradiation-induced shrinkage of the SiC-based fibers with debonding of the fiber–matrix interface that severely compromised the ability to load the high-strength fibers.80 Composite mechanical prop- erties such as strength suffered appreciably. With continued evolution of the fiber systems to increasingly pure, stoichiometric materials, the irradiation stability improved significantly. Presently, there are two commercial fiber systems used in nuclear-grade composites, both of which have rela- tively low impurity contents and are approaching a 1:1 stoichiometry. Specifically, the �11 micron Hi-Nicalon™ Type-S fiber has the nominal chemis- try of SiC1.05, 0.2%-O, while the �7.5 and �10 mm Tyranno™ SA-3 fibers have the nominal chemistry of SiC1.07, 0.5% Al. Study has revealed that these ‘near stoichiometric’ fibers exhibit irradiation- induced swelling similar to that of CVD,82 thus avoiding the debonding phenomenon mentioned in the previous paragraph. For this reason, composites fabricated from these materials are superior under irradiation to their predecessors. Consistent with the discussion of properties of irradiated monolithic SiC, the following discussion will be limited to the more pure, near stoichiometric fiber materials. 0 0 2 4 6 8 10 1 2 3 4 Neutron dose (dpa) W ei b ul l m ea n st re ng th (G P a) Offset for nonirradiated 800 950 470 280 280 470 770 950 300 500 Hi-Nicalon™ type-S, Katoh (2010)67 Tyranno™-SA3, Katoh (2010)67 Hi-Nicalon™ type-S, Nozawa (2004)83 Figure 22 Effect of neutron irradiation on fiber strength. Data labels indicate the nominal irradiation temperature (in �C) for Hi-Nicalon™ Type-S (upright) and Tyranno™ SA-3 (oblique) fibers. Reproduced from Katoh, Y.; Snead, L. L.; Nozawa, T.; Kondo, S.; Busby, J. T. J. Nucl. Mater. 2010, 403, 48–61. 236 Radiation Effects in SiC and SiC–SiC The effect of neutron irradiation on the Weibull mean strength of individual ‘near stoichiometric’ fibers is given in Figure 22.83,84 Within inherent sta- tistical scatter, no change in strength is observed for either the Hi-Nicalon™ Type-S or the Tyranno™ SA-3 bare fibers. The numbers inset to the figure indicate the irradiation temperature of the SiC fibers, with no apparent function of irradiation temperature on strength observed. From the same study, the effect of irradiation on composite properties is also observed. Figure 2367 gives the proportional limit stress for which the load departs from elastic behavior and the ultimate tensile strength. As with the fiber data, and the data for monolithic CVD SiC (Figure 18), the composite strength does not exhibit any statistically meaningful change. Supporting studies14,82,83,85–87 on the strength in tension or bending of neutron-irradiated stoichiometric fiber composites support the fact that at least up to �40 dpa, composite strength is not significantly affected by irradiation. A recent study88 on the fracture toughness of irradiated and unirradiated Hi-Nicalon™ Type-S composites also reports no appreciable change. However, a minor dif- ference in the fracture surface (length of fiber pull out) and a trend in the fiber–matrix interphase properties are reported,89 suggesting that mechanical property evolution may occur at higher doses. In the unirradiated state, the thermal conductivity of SiC composites is dependent on variables includ- ing the fibers and matrix constituents, processing, and the level of porosity. For the nuclear composite considered here, there is considerable thermal conduc- tivity anisotropy and temperature dependence typical of all ceramics. This is demonstrated in Figure 24, which gives the measured and calculated thermal conductivity for the two nuclear-grade SiC compo- sites.90 Presented are Hi-Nicalon™ Type-S fiber and Tyranno™ SA fiber composites, each matrix infiltrated through CVI.58 Architectures included balanced (1:1:1 for x:y:z) and unbalanced (1:1:4) 3D forms and 2D laminates (SW: satin weave, PW: plain weave.) In each case, a pyrolytic graphite interphase was applied. The conductivity for all materials is presented in the through thickness direc- tion (perpendicular to the plate and the fabric for the 2D composite.) This typically represents the low- conductivity direction. As evident from Figure 24 and the supporting analysis by Katoh,90 the fiber makes a significant con- tribution to the thermal conductivity of these highly stoichiometric fiber composites, and this conductivity is a fairly strong function of temperature. However, the absolute conductivity is only a fraction of that for the highest thermal conductivity CVD SiC (cf. Figure 7.) 400 Nonirradiated 460 480 690 760 780 1000 PLS UTS610 420 400 380 350 220 530 480 570 610 300 200 100 0 0 1 2 3 Neutron dose (dpa) Te ns ile s tr es s (M P a) 4 5 6 Hi-Nicalon-S CVI UTS Hi-Nicalon-S CVI PLS Tyranno-SA3 CVI UTS Tyranno-SA3 CVI PLS Figure 23 Effect of neutron dose on tensile proportional limit and ultimate tensile stresses for composites. Data labels indicate the nominal irradiation temperature in �C for Hi-Nicalon™ Type-S (upright) and Tyranno™ SA-3 (oblique) composites. Reproduced from Katoh, Y.; Snead, L. L.; Nozawa, T.; Kondo, S.; Busby, J. T. J. Nucl. Mater. 2010, 403, 48–61. 80 70 60 50 40 30 20 10 0 0 200 400 600 Temperature (�C) 800 1000 1200 3D 1:1:4 TySA/PyC, through-thickness, model 3D 1:1:4 TySA/PyC through-thickness, experiment 3D 1:1:1 TySA/PyC through-thickness, experiment 3D 1:1:1 TySA/PyC, through-thickness, model 2D-PW TySA/PyC, through-thickness, model 2D-PW TySA/PyC through-thickness, experiment 5HSW HNLS/PyC, through-thickness, model 5HSW HNLS/PyC through-thickness, experiment Th er m al c on d uc tiv ity (W m -1 K –1 ) Figure 24 Thermal conductivity of representative nuclear-grade SiC/SiC composite in unirradiated condition. Reproduced from Katoh, Y.; Nozawa, T.; Snead, L. L.; Hinoki, T.; Kohyama, A. Fus. Eng. Des. 2006, 81, 937–944. Radiation Effects in SiC and SiC–SiC 237 As with the CVD SiC discussed in section 4.07.3, silicon carbide composite also undergoes significant degradation in thermal conductivity because of neutron irradiation. The data is somewhat limited; however, Figure 25 gives the ambient through- thickness thermal conductivity for a plain weave Hi-Nicalon™ Type-S, multilayer SiC interphase, and CVI SiC matrix composite. It is noted that, in 10 Knonirr= 10.1 ± 2.2 8 6 4 2 0 0.001 0.01 0.1 Dose (dpa) Nicalon Type-S fiber composite 1 10 Th er m al c on d uc tiv ity a t am b ie nt (W m -1 K –1 ) Tirr~ 200 �C Tirr~ 800 �C ~ 600 �C ~ 400 �C Figure 25 Effect of neutron irradiation on the through-thickness thermal conductivity of Hi-Nicalon™ Type-S, CVI matrix composite. 0.1 0.01 Tirr~ 200 �C composite ~ 800 �C composite Tirr~ 200 �C CVD SiC Tirr~ 800 �C CVD SiC 0.001 Th er m al d ef ec t re si st an ce (W m -1 K -1 )- 1 0.001 0.01 0.1 Dose (dpa) 1 10 1 Figure 26 Comparison of the thermal defect resistance for neutron irradiated CVD SiC and Hi-Nicalon™ Type-S, CVI matrix composite. 238 Radiation Effects in SiC and SiC–SiC Radiation Effects in SiC and SiC–SiC 239 comparison to the conductivity shown in Figure 24 (second from lowest curve), the ambient through- thickness thermal conductivity for the material of Figure 25 is relatively low (10.2� 2.2Wm�1 K�1). This is mostly ascribed to the large porosity for that composite. Nevertheless, the figure clearly shows a significant, irradiation temperature-dependent reduc- tion in thermal conductivity as a function of irradiation dose. The fact that this is temperature dependent suggests that the degradation is due to the produc- tion of stable point defects and clusters, as discussed in Section 4.07.3, although this may not be the sole factor determining the degradation. Figure 26 provides the accumulated thermal defect resistance at the lowest and highest irradiation temperature for the composite materials of Figure 25, compared with high-conductivity CVD SiC. It is interesting to note that the thermal defect resistance for the composite, while accumulating in the same manner as that of the CVD SiC, is about an order of magnitude greater than that of CVD SiC at a given dose (at least prior to saturation.) This greater accumulation of thermal defect resistance has been recently observed by Katoh.67 The reason for this is unclear, although it is plausible that, in addition to defect production, propa- gation of internal interfaces (e.g., cracks) in the com- posite is occurring under irradiation. It is also possible that the defects population responsible for phonon scattering for the composite material is stabilized at a higher level than that of the highly pure CVD SiC.90 References 1. CEGA NP-MHTGR Material Models of Pyrocarbon and Pyrolytic Silicon Carbide; CEGA-002820, Rev 1; July 1993. 2. Blackstone, R.; Voice, E. H. J. Nucl. Mater. 1971, 39, 319–322. 3. Price, R. J. J. Nucl. Mater. 1969, 33, 17–22. 4. Price, R. J. J. Nucl. Mater. 1973, 48, 47–57. 5. Primak, W.; Fuchs, L. H.; Day, P. P. Phys. Rev. 1956, 103(5), 1184–1192. 6. Balarin, M. Phys. Stat. Sol. 1965, 11, K67–K71. 7. Pravdyuk, N. F.; Nikolaenko, V. A.; Kapuchin, V. I.; Kusnetsov, V. N. 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All rights reserved. 4.08.1 Introduction 242 4.08.2 Nanoscale Oxide Particle Control 242 4.08.2.1 Dissociation and Precipitation 242 4.08.2.2 Structure and Coherency 243 4.08.3 Martensitic 9Cr-ODS Steels 244 4.08.3.1 Chemical Composition and Microstructure 244 4.08.3.2 Residual Ferrite Formation and Strength Characterization 245 4.08.3.2.1 Mechanically alloyed powder characterization 245 4.08.3.2.2 Pinning of a–g interface by oxide particles 245 4.08.3.2.3 Strength characterization 247 4.08.3.3 Cladding Manufacturing 248 4.08.3.3.1 Continuous cooling transformation diagram 248 4.08.3.3.2 Manufacturing process 249 4.08.3.4 Creep and tensile properties 250 4.08.4 Ferritic 12Cr-ODS Steels 250 4.08.4.1 Strength Anisotropy 250 4.08.4.2 Recrystallization Tests 252 4.08.4.3 Cold-Rolling Cladding Manufacturing 252 4.08.4.4 Internal Creep Rupture Property 253 4.08.5 Al-Added 16Cr-ODS Steels 255 4.08.5.1 Application and Technical Issues 255 4.08.5.2 Thermal Aging Embrittlement Due to High Cr Content 256 4.08.5.3 Mechanical Properties 256 4.08.5.4 Cladding Manufacturing 257 4.08.6 Existing ODS Steel Cladding 258 4.08.7 Corrosion and Oxidation 259 4.08.7.1 Sodium Compatibility 259 4.08.7.2 LBE Compatibility 259 4.08.7.3 SCPW Compatibility 262 4.08.7.4 Oxidation 262 4.08.8 Irradiation 263 4.08.8.1 Simulated Irradiation 263 4.08.8.2 Neutron Irradiation of Materials 264 4.08.8.3 Fuel Pin Irradiation 268 4.08.8.3.1 9Cr- and 12Cr-ODS steel cladding in BOR-60 268 4.08.8.3.2 12Cr-ODS steel cladding in EBR-II 269 4.08.8.3.3 DT2203Y05 in Phénix 269 4.08.9 Summary 269 References 270 Abbreviations CTT Continuous cooling transformation CEN-SCK Centre d’Etude de l’énergie Nucleaire – Studiecentrum voor Kernenergie CVN Charpy V-notch EFTEM Energy-filtered transmission electron microscopy EPMA Electron probe microanalysis 241 242 Oxide Dispersion Strengthened Steels FFT Fast Fourier transformation HIP Hot isostatic pressing HRTEM High-resolution transmission electron microscopy INCO International Nickel Company JAEA Japan Atomic Energy Agency LBE Lead–bismuth eutectic LFR Lead fast reactor LMP Larson–Miller parameter MA Mechanical alloying MOX Mixed oxide ODS Oxide dispersion strengthened PMW Pulse magnetic Welding PRW Pressurized resistance welding SCPW Super critical pressurized water SEM Secondary electron microscopy SFR Sodium fast reactor TIG Tungsten inert gas welding UTS Ultimate tensile strength 4.08.1 Introduction Recent progress in oxide dispersion strengthened (ODS) steels produced by mechanical alloying (MA) techniques allows them to be used as fuel cladding in sodium-cooled fast reactors (SFR). The thermally stable oxide particles dispersed in the ferritic matrix improve the radiation resistance and creep resistance at high temperature. As a result, ODS steels have a strong potential for high burnup (long-life) and high- temperature applications typical for SFR fuels. The attractiveness of ODS steels is due not only to the nanosize oxide particles composed of Y–Ti–O atoms but also to their controlled micron-size grain mor- phology. We review existing knowledge on the crys- talline structure and lattice coherency of these nanosize particles with their surrounding matrix, since these factors dominate the dispersion and strength-determining mechanism through dislocation interaction. The development of manufacturing pro- cesses is a principal issue for hardened ODS steels to realize long, thin-walled ODS steel cladding on pro- duction scales. Therewas the long-standing problem in low hoop strength due to the extremely elongated fine grains parallel to the rolling direction. To soften hard- ened cold-rolled products and modify their grain morphology, martensitic 9Cr-ODS steels and ferritic 12Cr-ODS steels have been developed. Current prog- ress in the development of these ODS steel claddings, including their relevant mechanical properties, for example, tensile and creep rupture strengths in the hoop directions, and irradiation performance, is reviewed. The development of Al-added high Cr- ODS steel cladding is also addressed, with a focus on superior resistance to oxidation and corrosion in a lead– bismuth eutectic (LBE), and supercritical pressurized water (SCPW) in the international Generation IV advanced nuclear power system. Nanocluster ODS steels,1 for example, 14YWT, etc., for fusion blanket structure materials, are not addressed in this chapter. 4.08.2 Nanoscale Oxide Particle Control 4.08.2.1 Dissociation and Precipitation The fine distribution of Y2O3 particles, which is essential to improving the high temperature strength of ODS steels, is attained by the dissociation of oxide particles during MA processing.2 The thermodynam- ically stale Y2O3 particles are forcedly decomposed into the ferritic steel matrix during the MA process. Subsequent annealing induces oxide particles to pre- cipitate finely at elevated temperature of around 1000 �C. The co-addition of Ti during MA proces- sing promotes the decomposition of Y2O3 and then the precipitation of Y–Ti complex oxide particles through an annealing heat treatment.3,4 A field emis- sion ion micro-probe (FIM) analysis confirmed that this type of complex oxide is constituted of several nanometer-sized Y–Ti–O compounds.5–7 The precipitation process of the decomposed Y2O3 was investigated by means of a small angle neutron scattering (SANS) experiment.8 The neutron-scattering cross-section (dS/dO) versus scattering vector (q2) plots for the milled U14YWT(Fe–14Cr–0.4Ti–3W– 0.25Y2O3) are shown in Figure 1(a). They indicate that the hot isostatic pressing (HIP) of U14YWT at 850 �C leads to the precipitation of a high number density of nanoclusters, as designated by Odette. Figure 1(b) shows the effects of HIP (filled symbols) and powder annealing (open symbols) at tempera- tures of 700, 850, 1000, and 1150 �C. The increase in magnitude and decrease in slope of the dS/dO versus q2 curves indicate that the radius of nanoclus- ters decreases and their number density increases with decreasing temperature at HIP and powder annealing. Annealing at 700 �C produces the highest scattering and lowest sloping, which indicates that the smallest-sized nanoclusters precipitate with the high- est number density at lower temperatures. In terms of an X-ray diffraction experiment using Super Photon 10 1 850 �C HIP (No Y2O3) 850 �C HIP (U14YWT) As-MA (U14YWT) As-MA (No Y2O3) 0.1 0.01 20 (a) (b) 4 q2 (nm−2) d S/ d W (c m sr ad )− 1 10 1150 �C 1150 1000 1000 �C 850 �C 850 �C Controls (No Y2O3) 700 �C Decreasing Increasing Nd 1 HIPed materials Powder anneals 0.1 0.01 d Σ/ d Ω (c m sr ad )− 1 6 8 10 20 4 q2 (nm−2) 6 8 10 Figure 1 Results of a SANS experiments for as-mechanically alloyed powders and after HIP and annealing in U14YWT (Fe–14Cr–0.4Ti–3W–0.25Y2O3). (a) As-MA, 850 �CHIP and annealing and (b) HIP (filled) and annealing (annealing) at 1150, 1000, 850, and 700 �C . Reproduced from Alinger, M. J.; Odette, G. R.; Hoelzer, D. T. J. Nucl. Mater. 2004, 382, 329–333. 3.06 Å (a) (b) (0 1 − 1) (2 2 − 2) (2 2 − 2 − ) (1 1 − 0) 3.06 Å (0 1 1 − ) (1 0 1 − ) ( − 1 1 0) ( − 2 2 2) ( − 2 2 2 − ) ( − 1 0 1)70.5 � Figure 2 HRTEM micrograph of the Y2O3 particle with surrounding matrix (a) and FFT image of micrograph (b). The diffraction spots from Y2O3 particle of {2 2 2} type form the rectangle, whereas diffraction spots from the matrix of {1 1 0} type form the hexagon at the [1 1 0] zone axis and [1 1 1] of matrix. Reproduced from Kliniankou, M.; Lindau, R.; Moslang, A. J. Nucl. Mater. 2004, 329–333, 347–351. Oxide Dispersion Strengthened Steels 243 ring-8 eV (Spring-8) constructed in Japan, Kim et al. recently reported that nanoclusters could be in a noncrystalline state and can be transformed to nano- crystalline oxide particles at around 1000 �C.9 4.08.2.2 Structure and Coherency With regard to ODS steels without Ti, high resolu- tion (HR) TEM investigations were performed by Klimiankou to investigate the structure of Y2O3. 10 The crystallographic lattice of the metal matrix cor- responds to a-Fe with a bcc structure and a lattice constant of a0¼ 0.287 nm.11 The Y2O3 has a crystal- line bcc structure with a 1.06 nm lattice constant.11 Figure 2 shows an HRTEM image taken from an Y2O3 particle that is surrounded by the matrix (M) lattice. This image was taken from the grain, oriented with [1 1 1]M zone axis to the electron beam. A fast Fourier transformation (FFT) of the image shows the matrix lattice as a hexagonal pattern with diffraction spots of the {1 1 0} type and dM(110)¼ 0.203 nm dis- tance. In the FFT image, the Y2O3 (YO) lattice is rectangular, with diffraction spots of the {2 2 2} type and a corresponding atomic planes distance of dYO(2 2 2)¼ 0.306 nm. The angle of 70.5� between diffraction spots of the {2 2 2}YO type marked in Figure 2(a) confirms that the Y2O3 particle is oriented with the [1 1 0] zone axis, and consequently [1 1 0]YO// [1 1 1]M. The orientation correlation of both lattices is (1 1� 1�)YO//(1 1� 0)M. Therefore, the following Kurdjumov–Sachs orientation relationship12 is satisfied: ð11�1�ÞYO==ð11�0ÞM; ½110�YO==½111�M: ½1� The interfacial coherency between the Y2O3 particle and the ferritic matrix can be estimated by Klimiankou as follows: 3dMð1 1 0Þ � 2dYOð2 2 2Þ 2dYOð2 2 2Þ � 0:5%: ½2� This result suggests that a coherency could be satisfied for the Y2O3 particles surrounded by a ferritic matrix. Concerning the Y–Ti–O complex oxide particles formed in Ti-added ODS steels, Figure 3(a) shows an energy-filtered (EF) TEM micrograph from a Y–Ti–O particle in which two atomic planes are simul- taneously visible.10 Figure 3(b) shows the (0 0 4) and (2 2� 2) atomic planes of Y2Ti2O7 cubic (a0¼ 1.01 nm) phases with the [1 1 0] zone axis. In fact, the measured 0 0 4 5 nm (a) (b) 2 − 2 2 2 − 2 2 − Figure 3 EFTEM images of Y2Ti2O7 particles. Reproduced from Kliniankou, M.; Lindau, R.; Moslang, A. J. Nucl. Mater. 2004, 329–333, 347–351. Residual ferrite Tempered martensite 20 μm Figure 4 Microstructure of 9Cr-ODS steel showing residual ferrite and tempered martensite. 244 Oxide Dispersion Strengthened Steels data are equal to the following data calculated from the Y2Ti2O7 structure: d2 2 2¼ 0.29 nm and d0 0 4¼ 0.25, and an angle between the (0 0 4) and (2 2� 2) atomic planes of 54.7� . The analysis of EFTEM results defini- tively shows that Y–Ti–O particles have a Y2Ti2O7 composition. These findings suggest that nano-oxide particles precipitate from the ferritic matrix, maintaining crys- talline coherency or partial-coherency with a ferritic matrix. In general, the nucleation and growth of pre- cipitates proceeds, as both interfacial and strain ener- gies become minimal. In the case of ODS steels, interfacial coherency could be maintained between thermodynamically stable nanoparticle precipitates and the ferritic matrix in order to decrease the free energy in the system from the extremely high energy state induced by MA. Elucidation of the details of the nanoscale precipitation is important not only as basic materials science research but also as the develop- ment of high-strength engineering materials. 4.08.3 Martensitic 9Cr-ODS Steels 4.08.3.1 Chemical Composition and Microstructure 9Cr-ODS steels are being developed by the JAEA ( Japan Atomic Energy Agency) for application to SFR fuel cladding. Their standard chemical compo- sition is 9Cr–0.13C–0.2Ti–2W–0.35Y2O3 (wt%). The chromium concentration was determined to be 9wt% in terms of ductility, fracture toughness, and corrosion resistance based on a series of irradiation data of ferrite steels. The addition of titanium pro- duces the nanoscale dispersion of oxide particles, which leads to a markedly improved high-temperature strength. If titanium is added to excess, however, it creates too much strength, which negatively impacts cold rolling and cold workability. To achieve a bal- ance between strength and workability, a value of 0.2 wt% was selected. Tungsten of 2 wt% is also added in order to improve high-temperature strength by means of solid solution hardening. The microstructure of 9Cr-ODS steels13–18 can be easily controlled by a reversible a–g transformation with a remarkably high driving force of a few hundred MJm�3, as compared with a driving force of irrevers- ible recrystallization with a fewMJm�3 for 12Cr-ODS steels. By inducing reversible a–g transformations, 9Cr-ODS steel cladding for fast-reactor fuel elements is currently being manufactured at the JAEA. The microstructure of 9Cr-ODS steel cladding is basically tempered martensite. However, it has been recognized that 9Cr-ODS steel cladding manufactured in an engineering process possesses a dual-phase struc- ture that comprises both tempered martensite and ferrite phases. An example of their microstructure is shown in Figure 4. The ferrite phase appears white, and the elongated phase is indicated by arrows. Their size is about 30–60 mm in length �3–10 mm in width. The formation of a ferrite phase in 9Cr-ODS steel is somewhat unusual, because only the full martensite phase can be expected in 9Cr-ferritic steel without yttria under normalizing and air-cooling conditions. Moreover, the high-temperature strength of manu- factured 9Cr-ODS steel is significantly improved by the presence of the ferrite phase.19–21 This is obvious from the creep rupture data shown in Figure 5.20 200 50 60 70 80 90 100 10 100 1000 10 000 Time to rupture (h) U ni -a xi al s tr es s (M P a) Solid: with residual ferrite Open: without residual ferrite Mechanically milled without Y2O3 700 �C Figure 5 Uni-axial creep rupture strength of 9Cr-ODS steels at 700 �C after the normalizing-and-tempering (1050 �C � 1 h, Ar-gas cooling (AC) = > 780 �C � 1 h, AC) with and without residual ferrite. Reproduced from Ohtsuka, S.; Ukai, S.; Fujiwara, M.; Kaito, T.; Narita, T. Mater. Trans. 2005, 46, 487. 0 1500 1400 1300 1200 1100 1000 900 800 700 600 500 400 300 0.1 C content (wt%) Te m p er at ur e (� C ) 0.2 δ γ + δ γ + δ +TiC γ + TiC L γ + TiC + M23C6 α + M23C6+ TiC α + M23C6+ laves + TiC α +TiC α +TiC +Laves A3 A1 0.3 Figure 6 Computed phase diagram with respect to carbon content for 9Cr–xC–0.2Ti–2W system without Y2O3. Oxide Dispersion Strengthened Steels 245 Therefore, the control of ferrite phase formation is a key to the realization of high-temperature strength in 9Cr-ODS steel cladding. 4.08.3.2 Residual Ferrite Formation and Strength Characterization 4.08.3.2.1 Mechanically alloyed powder characterization The computed phase diagram of the Fe–0.13C–2W– 0.2Ti system without Y2O3 is shown in Figure 6 with respect to carbon content. For a carbon content of 0.13wt%, a single austenite g-phase containing TiC carbide exists at a normalizing temperature of 1050 �C. The equilibrium g/g + d-phase boundary at this temperature corresponds to a carbon content of 0.08wt%, beyond which d-ferrite is not stable. The specimens without and with 0.1wt% Y2O3 exhibit the full martensite structure, whereas the specimens with 0.35 and 0.7wt% Y2O3 exhibit a dual phase compris- ing both martensite and ferrite phases. Digital image analyses show that the area fraction of the ferrite phase is �0.2 for specimens with 0.35 and 0.7wt% Y2O3. High-temperature X-ray diffraction measurement at 950 �C showed a considerable difference; the specimen without Y2O3 shows diffraction peaks that correspond only to the austenite g-phase, whereas specimens with 0.35 and 0.7 wt% Y2O3 show diffraction peaks corresponding not only to an austenite g-phase but to a ferrite phase as well. The austenite g-phase transforms to the martensite phase, but the ferrite phase remains unchanged by quenching. Considering that the ferrite phase is formed only in the specimens containing 0.35 and 0.7wt% Y2O3, and that four types of ODS steels have an identical chemical composition except for Y2O3 content, the Y2O3 particles could suppress the a–g reverse transformation. Figure 722 shows the results of dilatometric mea- surement when 9Cr–0.13C–2W–0.2Ti is heated without and with 0.35 wt% Y2O3. In the case of the specimen without Y2O3, the linear thermal expansion begins to decrease from anAC1 point of 850 �C to anAC3 point of 880 �C, due to the reverse transformation of a–g-phase, which corresponds reasonably well with the computed phase diagram. The addition of 0.35wt% Y2O3 induces an increase up to an AC3 point of 935 �C. By comparing both curves, it was found that the specimen with 0.35wt% Y2O3 exhibits a smaller degree of reduction in linear thermal expansion during the reverse transformation of the a–g-phase; this obser- vation indicates that the entire a-phase could not be transformed to a g-phase. This untransformed ferrite phase was designated as a residual ferrite. 4.08.3.2.2 Pinning of a–g interface by oxide particles Alinger’s results indicate that the mechanically alloyed powder annealed at 700 �C shows the smallest radius and highest density in Y–Ti complex oxide particles,8 as shown in Figure 1. Considering that Li ne ar t he rm al e xp an si on (Δ L/ L, % ) Temperature (ºC) 1 700 750 800 850 900 950 1000 1050 1100 1 0.6 0.7 0.8 0.9 1.1 1.2 1.3 1.4 0.6 0.7 0.8 0.9 1.1 1.2 1.3 1.4 700 750 800 850 900 950 1000 1050 1100 0.35 mass % Y2O3 Without Y2O3 AC1 AC3 AC1 AC3 Figure 7 Results of linear thermal expansion measurement between 700 and 1100 �C at temperature rising of 0.33 �Cs�1 for 0 mass % and 0.35 mass % Y2O3 in 9Cr–0.13C–2W–0.2Ti specimens. Reproduced from Yamamoto, M.; Ukai, S.; Hayashi, S.; Kaito, T.; Ohtsuka, S. Mater. Sci. Eng. A 2010, 527, 4418–4423. D riv in g fo rc e (M J m –3 ) Temperature (°C) –2 0 2 4 6 8 10 12 800 900 1000 1100 1200 1300 ΔG(0.2 mass % C) ΔG(0.13 mass % C) F(0.1 mass % Y2O3) F(0.35 mass % Y2O3) F(0.7 mass % Y2O3) Figure 8 Comparison of the driving force (DG) for a to g reverse transformation derived by using Thermo-Calc code and pinning force (F ) due to oxide particles for 0.1 mass %, 0.35 mass %, and 0.7 mass % Y2O3 in Fe–0.13C–2W–0.2Ti specimens. Driving force (DG) for 0.13 mass % C and 0.2 mass % C is shown. Reproduced from Yamamoto, M.; Ukai, S.; Hayashi, S.; Kaito, T.; Ohtsuka, S. Mater. Sci. Eng. A 2010, 527, 4418–4423. 246 Oxide Dispersion Strengthened Steels Y2O3 particles are decomposed duringMA, subsequent annealing results in the formation and precipitation of Y–Ti complex oxide particles at elevated temperatures of 700 �C or higher. Since the reverse transformation of a–g-phase takes place at a temperature over 850 �C, which is higher than the precipitation temperature of Y–Ti complex oxide particles, it is possible that the retention of the residual a-ferrite can be attributed to the presence of Y–Ti complex oxide particles in 9Cr-ODS steels. These particles could block the motion of the a–g interface, thereby partly suppres- sing the reverse transformation from a- to g-phase. This section presents a quantitative evaluation of this process. The chemical driving force (DG) for the reverse transformation from a- to g-phase in the Fe–0.13C– 2W–0.2Ti system without Y2O3, can be evaluated in terms of Gibbs energy versus carbon content curves at each temperature. These curves were derived using the Thermo-Calc code and the TCFE6 database. The result of the calculation is presented in Figure 8.22,23 The peak value of the driving force for the reverse transformation from a- to g-phase reaches 4MJm�3 at 1000 �C in the case of 0.13wt% C. The pinning force (F ) against the motion of the a–g interface can be expressed as the following equation, which was derived from the modified Zener equation of Mishizawa et al.24 F ¼ 3sf 2=3 p 8r ; ½3� where, s( Jm�2) is the interfacial energy between a- and g-phases, and its value was selected to be 0.56 Jm�2.25 The character r represents the radius of the oxide particles (m) in the a-phase; its value was determined as 1.5 nm by using TEM observation. The character fp represents the volume fraction of dispersed oxide particles (�), and was derived on the basis of the experimental evidence that oxide particles consist of Y2Ti2O7. By substituting these values into the afore- mentioned equation, the value of pinning force F was determined for 0.1, 0.35, and 0.7wt% Y2O3, which are also shown in Figure 8.22,23 The value of F increases with the amount of Y2O3 added according to the relation of f 2=3 . The velocity of the a–g interface motion (v) is proportional to the difference between F and DG, as shown in the following equation: v ¼ MðDG � FÞ: ½4� M is the mobility of the interface. DG and F are competitive, and DG > F indicates a positive velocity for the interface motion, that is, the reverse trans- formation from a- to g-phase. On the other hand, DG< F indicates that the a–g interface can be 2.0 3.0 4.0 5.0 6.0 H ar d ne ss (G P a) Average covering residual ferrite and tempered martensite Tempered martensite Residual ferrite NT 550 ºC 1 h 750 ºC 1 h 800 ºC 7 h 800 ºC 58 h FC Figure 10 Hardness change at room temperature as a function of tempering conditions for the residual ferrite and tempered martensite. NT: normalizing and tempering; FC: furnace cooling. Ukai, S.; Ohtsuka, S.; Kaito, T.; Sakasegawa, H.; Chikata, N.; Hayashi, S.; Ohnuki, S. Mater. Sci. Eng. A 2009, 510–511, 115–120. AC1 α γ γ α AC3 Carbide Oxide particle ・ ・ ・ ・ ・ ・ ・ ・ ・ ・ ・ ・ ・ ・ ・ ・ ・ ・ ・ ・ ・ ・ ・ ・ Figure 9 Formation process of residual ferrite in 9Cr-ODS steel (Fe–0.13C–2W–0.2Ti–0.35Y2O3). Reproduced from Yamamoto, M.; Ukai, S.; Hayashi, S.; Kaito, T.; Ohtsuka, S. Mater. Sci. Eng. A 2010, 527, 4418–4423. Oxide Dispersion Strengthened Steels 247 pinned by oxide particles so that the a-phase is, thus, retained. The results of the calculation shown in Figure 822 reveal that in the case of Y2O3 con- tents of 0.35 and 0.7 wt%, the pinning force is larger than the driving force for 0.13 wt% C. These results are reasonably consistent with our observation of the retainment of residual ferrite during a–g reverse transformation. On the basis of the aforementioned discussion, the formation process of the residual ferrite in Fe–0.13C– 2W–0.2Ti–0.35Y2O3 is schematically illustrated in Figure 9. At the AC1 point, the carbide begins to decompose, and a–g inverse transformation takes place in the area of higher carbon content around the decomposed carbide, where the driving force of the reverse transformation exceeds the pinning force because the carbon content may be >0.2 wt% (see Figure 8). The g-phase could be enlarged by these processes. Approaching the AC3 point, the matrix carbon content achieves equilibrium at 0.13 wt%, where the pinning force (0.35Y2O3) exceeds the driving force (0.13C), and the velocity of the a–g interface motion is markedly reduced due to dragging by the oxide particles. Thus, the a-ferrite could be retained even beyond the AC3 point. 4.08.3.2.3 Strength characterization Nanoindentation measurements were conducted in order to evaluate the mechanical properties of the residual ferrite itself. The trace of a Berkovich tip can be placed within the interiors of the residual ferrite regions, while conventional micro-Vickers diamond tips using 100-mN loads cover 7� 7 mm. Figure 10 shows the hardness change in the individual phases measured by this nanoindentation technique as a parameter of the tempering conditions.26 The decrease in hardness is significantly restricted in the residual ferrite as compared to that of the martensite phase in terms of increasing the tempering conditions. The overall hardness measured by the micro-Vickers tester is also shown by the broken line which covers both the residual ferrite and martensite, therefore, representing the average hardness of both phases. Hardness Hv is correlated with yield stress sy using the relationship provided by Tabor.27 For tempering conditions at 800 �C for 58 h, which is equivalent to tempering at 700 �C for 10 000 h based on the LMP (Larson–Miller parameter), hardness can be converted to yield stress at room temperature for the individual phases: 1360MPa for the residual ferrite and 930MPa for the tempered martensite. The yield strength of the residual ferrite is 1.5 times higher than that of mar- tensite at tempering at 700 �C for 10 000 h. A full ferrite ODS steel and full martensite ODS steel were manufactured, and the oxide particle dis- tribution in both ODS steels was measured by TEM. The results are shown in Figure 11.28 It is obvious that a few nanometer-sized oxide particles are finely distributed in the full ferrite ODS steel, whereas their size is coarsened in the bi-modal distribution in the martensite ODS steel. Considering that the residual ferrite phase belongs to full ferrite ODS steel, resid- ual ferrite contains fine (nanosized) oxide particles which are responsible for higher strength in residual ferrite containing ODS steels. In regard to the bi- modal distribution of oxide particles in martensite 248 Oxide Dispersion Strengthened Steels ODS steels, the a–g-phase transformation could induce the coarsening of oxide particles by disturbing the interface coherency between these particles and the g-phase matrix. 4.08.3.3 Cladding Manufacturing 4.08.3.3.1 Continuous cooling transformation diagram The preparation of a CCT (continuous cooling trans- formation) diagram is essential to the microstructure 100 200 300 400 500 600 700 800 900 1000 103102101 Time from 800 Te m p er at ur e, T (� C ) γ α+γ γ + m m α + γ + m α+ m No residua 9Cr-ODS ferritic steel Containing residual-α 18 000 (K/h) Austenite (γ) => Martensite 3000 ( Figure 12 CCT diagram of 9Cr-ODS steel. Reproduced fromO Mater. 2006, 351, 241. 20 nm 20 nm(a) (b) Figure 11 TEM photograph of the oxide particles: (a) finely distributed oxide particles in full ferrite ODS steel and (b) bi-modal distribution of oxide particles with larger size in the full martensiteODS steel. Yamamoto,M.; Ukai, S.; Hayashi, S.; Kaito, T.; Ohtsuka, S. J. Nucl. Mater. 2011, 417, 237–240. control of 9Cr-ODS steels. Figure 12 exhibits a CCT diagram that was experimentally constructed for 9Cr-ODS steel.21 The minimum cooling rate for the matrix phase in order to fully transform to martensite is extremely higher in 9Cr-ODS steel (solid circular symbol) than in mechanically milled EM10 (open diamond symbol) that does not contain added Y2O3. 29 Residual ferrite plays an important role in the process of continuous cooling transformation. The minimum cooling rate is known to increase with a decrease in the size of prior austenite (g) grains. This smaller size of prior g grains provides more nucleation sites (grain boundaries) for a g–a-phase transformation, so that a higher cooling rate is required to enable steel with small prior g grains to fully transform to a. The presence of residual ferrite restricts the growth of g grains; the prior grain size of residual ferrite-containing steel is roughly 5 mm, thus increasing the minimum cooling rate to produce a full martensite matrix. In steel that does not contain residual ferrite and the mechanically milled EM10, the size of the prior g grains is roughly 10 mm and 35 mm, respectively. The results shown in Figure 12 can be explained by the relationship between the size of prior g grains and the minimum cooling rate.21 As for the normal- izing heat treatment used in commercial furnaces, the cooling rate would be roughly 3000 �Ch�1, so that 300 500 700 900 1100 1300 105104 �C, t (s) Te m p er at ur e, T (K ) α l-α EM10 K/h) 30 (K/h) Austenite (γ) => Ferrite (α) htsuka, S.; Ukai, S.; Fujiwara, M.; Kaito, T.; Narita, T. J. Nucl. 9Cr–0.13C–2W–0.2Ti–0.35Y2O3 Cladding tube At intermediate heat treatment At final heat treatment Cold rolling (pilger mill) Hot extrusion (1423 K) MA powder Low carbon steel Elemental powders Yttria powder Mechanical alloying (MA) Figure 13 Cladding tube manufacturing process developed for 9Cr-ODS steel. Heat treatment H ar d ne ss (H v) Cold rolling (Rd = 50%) 1st 2nd 3rd 4th 1st 2nd 3rd 4th Mother tube 300 350 400 450 Figure 14 Hardness change in the process of cold rolling and intermediate and final heat treatments for cladding tube manufacturing of 9Cr-ODS steels. Oxide Dispersion Strengthened Steels 249 the matrix phase of 9Cr-ODS steel cladding consists of residual ferrite, martensite, and a small amount of transformed ferrite from the g-phase. 4.08.3.3.2 Manufacturing process 9Cr-ODS steels are promising materials to enable fast reactor fuel cladding to realize a high burnup of 200GWd t�1 at 700 �C, since they have superior radiation resistance and high temperature strength. Figure 13 shows a series of manufacturing processes of fuel cladding that is 8.5mm in diameter by 0.5mm in thickness by 2m in length. The element powders and yttria powder are mechanically alloyed for 48 h in an argon gas atmosphere using an attrition type ball mill with a capacity of 10 kg batch. The mechanically alloyed powders are sealed in hollow-shaped cans and degassed at 400 �C in a 0.1 Pa vacuum for 2 h. The hollow shape of the bars is consolidated by hot- extrusion at an elevated temperature of 1150 �C to the dimensions of 32mm in outer diameter, 5.5 mm in wall-thickness, and 4m in length. After machining to the precise dimensions, claddings are produced at their final dimension (8.5 mm in outer diameter, 0.5mm in thickness, and 2m in length) by four-pass rolling with about a 50% reduction ratio on each pass by using a pilger mill. Without heat treatment, it is too difficult to man- ufacture cladding for ODS steels by the cold-rolling process. Using the CCT diagram of 9Cr–0.13C–2W– 0.2Ti–0.35Y2O3, as shown in Figure 12, a cooling rate of about 150K h�1 was applied to the intermedi- ate heat treatment in order to induce the ferrite phase at room temperature without martensite transforma- tion. This phase has a lower degree of hardness. Hard- ened cladding due to the accumulation of cold deformation can be sufficiently softened by this inter- mediate heat treatment, and cold rolling can then be continuedwith the softened ferrite structure. Figure 14 represents the typical hardness change of 9Cr-ODS steel in the process of cladding manufacturing by repeated cold rolling and intermediate heat treatment. The elongated grain structure induced by the fourth cold rolling can ultimately be made into equi-axed grain structure by the final heat treatment, which 250 Oxide Dispersion Strengthened Steels consists of normalizing at 1050 �C for 1 h, followed by tempering at 800 �C for 1 h. 4.08.3.4 Creep and tensile properties The lifetime of a fast reactor fuel pin is most strongly determined by the internal creep rupture strength of the cladding induced by the internal pressure of the fission gas at a temperature of around 700 �C. For 9Cr-ODS steel cladding, internal creep rupture data at 650, 700, and 750 �C are shown in Figure 15.30 Additionally, the best fit lines for hoop stress versus rupture time at each temperature are shown by solid lines. These results confirmed that creep rupture strengths in the hoop and longitudinal directions of cladding are almost the same, due to their equi-axed grains. The corresponding creep rupture curves for HT931 and austenitic PNC31632 are also presented for comparison. PNC316 is a typical austenitic clad- ding developed by JAEA in the fast reactor program. Notably, superior performance in rupture time is shown in 9Cr-ODS steel cladding. The slope of PNC316 is steeper, and there is a cross-over before 1000 h at 750 �C and before 10 000 h at 700 �C. The stress condition of the fast reactor fuel pin gradually increases due to the accumulation of fission gases and reaches around 120MPa at its final service milestone of 75 000 h at 700 �C. In this stress range, it is obvious that 9Cr-ODS steel cladding is of advantage. Time to rupt Stress range for SFR fuel cladding 500 400 300 200 100 90 80 70 60 50 10 102 103 H oo p s tr es s (M P a) HT9 (1023 K) HT9 (973 K PNC316 (923 K) PNC316 (973 K) PNC316 (1023 K) Figure 15 Creep rupture curves of 9Cr-ODS steel claddings in at temperatures of 650, 700, and 750 �C, compared with those o Nanstad, R. K.; Samaras, M.; Ukai, S. Mater. Res. Soc. Bull. 200 The ultimate tensile strength (UTS) of 9Cr-ODS ferritic cladding in the hoop direction as measured in a temperature range from room temperature to 850 �C, is shown in Figure 16, together with the corresponding data for the ferritic–martensitic stain- less steel (PNC-FMS)19 that is conventionally used as fast reactor fuel cladding. The strength of 9Cr-ODS steel is superior to that of conventional PNC-FMS. The uniform elongation that takes place from room temperature to 900 �C is also shown in Figure 16. In the temperature range from 400 to 700 �C at which a fast reactor is commonly operated, the measured uniform elongation exhibits adequate ductility. This advantage of superior elongation in the produced clad- dings can probably be ascribed to the pinning of dis- locations by oxide particles, which retard recovery and sustain work-hardening. 4.08.4 Ferritic 12Cr-ODS Steels 4.08.4.1 Strength Anisotropy When JAEA started to develop ODS steels in 1985, the ferritic type of ODS steels was applied.3,33 These are similar to MA957,34 which is single ferrite phase and does not include the martensite. Based on the results of R&D conducted for several years, three kinds of claddings, 63DSA, 1DK, and 1DS, were manufactured in 1990. Their chemical compositions ure (h) 9Cr-ODS (923 K) 9Cr-ODS (973 K) 9Cr-ODS (1023 K) 104 105 ) HT9 (923 K) hoop direction by using internally pressurized specimens f HT9 and PNC316. Reproduced from Allen, T.; Burlet, H.; 9, 34(1), 20–27. Oxide Dispersion Strengthened Steels 251 are 13Cr–0.02C–3W–0.7Ti–0.46Y2O3 (63DSA), 13Cr– 0.05C–3W–0.5Ti–0.34Y2O3 (1DK), and 11Cr–0.09C– 3W–0.4Ti–0.66Y2O3 (1DS). Themanufacturing process is almost the same as the process shown in Figure 13, except for the rolling process and intermediate heat treatment, because cold-rolling processing can be hardly applied to these ODS steels. In the case of the 1DK cladding, six warm drawings at 800–850 �C, followed by four warm rolling passes at 500 �C with intermediate annealing at 1080 �C, were repeated to manufacture the thin-walled cladding in the dimension of 7.5mm outer PNC-FMS 200 400 600 800 1000 1200 500 0 1000 Te ns ile s tr en gt h (M P a) Temperature (K) 1500 Figure 16 Tensile strength and uniform elongation of 9Cr-OD Reproduced from Ukai, S.; Kaito, T.; Otsuka, S.; Narita, T.; Fujiw MA 957 (see figure 29) (DT220 1000 600 500 400 300 200 100 1 10 100 Time to ru S tr es s (M P a) Figure 17 Creep rupture strength of 1DK, 1DS, and 63DSA cl specimens at 650 �C. Reproduced from Ukai, S.; Harada, M.; O 1993, 204, 65–73. diameter, 0.4mm thickness, and 1m length. In the case of the 63DSA and IDS claddings, only six warm rolling passes at 650–700 �C with intermediate annealing at 1100 �C were conducted. The temperature of the final heat treatment of 1DK cladding was 1150 �C for 60 s, and 63DSA and 1DS claddings at 1100 for 3.6 ks. The uni- and bi-axial creep rupture strengths of the manufactured claddings at 650 �C are shown in Figure 17, where the uni-axial corresponds to the hot working direction and bi-axial belongs to the internal hoop direction.3 It was found that 200 400 600 800 1000 1200 Temperature (K) 5 0 10 15 PNC-FMS U ni fo rm e lo ng at io n (% ) S steel cladding in hoop direction by the ring specimens. ara, M.; Kobayashi, T. ISIJ Int. 2003, 43, 2038. Mol-ODS 3Y05, see figure 28) 1000 10 000 100 000 pture (h) Bi-axialUni-axial 1 DK 1 DS 63 DSA Uni-axial Bi-axial addings in hoop direction by using internally pressurized kada, H.; Inoue, M.; Nishid, T.; Fujiwara, M. J. Nucl. Mater. 252 Oxide Dispersion Strengthened Steels there is strong strength anisotropy, and the bi-axial creep rupture strength is considerably lower than that of the uni-axial direction. Microstructure obser- vations of these claddings exhibited the elongated grains like a bamboo structure in parallel to the working direction. The strength degradation in the bi-axial/internal hoop direction, which is essential for the fuel elements, should be mainly attributed to the grain boundary sliding and crack propagation due to stress concentration. 4.08.4.2 Recrystallization Tests Based on the aforementioned finding in ODS steels, the recrystallization processing was extensively stud- ied to change the substantially elongated grain struc- ture to the equi-axed grain structure. The Y2O3 content should be Table 1 Chemical composition of F1–F4 specimens with different titanium and yttria contents (mass %) in 12Cr-ODS steels Specimen no. Chemical composition (mass %) C Cr W Ni Ti Y2O3 O N Ar F1 0.065 11.8 1.92 0.03 0.13 0.08 0.08 0.010 0.005 F2 0.054 11.8 1.94 0.03 0.13 0.13 0.05 0.010 0.004 F3 0.065 11.8 1.93 0.03 0.22 0.22 0.09 0.012 0.005 F4 0.056 11.7 1.92 0.03 0.31 0.24 0.04 0.010 0.004 Source: Reproduced from Ukai, S.; Okuda, T.; Fujiwara, M.; Kobayashi, T.; Mizuta, S.; Nakashima, H. J. Nucl. Sci. Technol. 2002, 39(8), 872–879. Lo ng itu d in al F1 F2 F3 F4 Tr an sv er se 20 mm 20 mm Figure 19 Optical microstructure of the F1, F2, F3, and F4 specimens in the final claddings. Reproduced from Ukai, S.; Okuda, T.; Fujiwara, M.; Kobayashi, T.; Mizuta, S.; Nakashima, H. J. Nucl. Sci. Technol. 2002, 39(8), 872–879. Oxide Dispersion Strengthened Steels 253 repeated twice with a reduction ratio of about 50% per rolling. The intermediate heat-treatment to soften the cold-rolled cladding was performed at 1100 �C for 30min, and the final heat-treatment was performed at 1150 �C for 0.5 h. Figure 19 shows the optical microstructures of the manufactured claddings in the longitudinal and transverse directions.36 All of the specimens seem to be recrystallized. However, the extent of recrys- tallization depends on the yttria and titanium con- tents. In the transverse cross-section, the grain size becomes slightly finer with increasing yttria and titanium contents from F1 to F4. In the F4 speci- men, the elongated grains can be still seen and the aspect ratio in the longitudinal (L) and transverse (T) directions is large, whereas the aspect ratio of specimen F1 appears to be nearly unity. These findings show that F4 specimen with higher yttria and titanium contents did not attain the perfectly recrystallized grain structure by the annealing of 1150 �C for 0.5 h. 4.08.4.4 Internal Creep Rupture Property The internal creep rupture properties of the manu- factured 12Cr-ODS steels at 700 �C are shown in Figure 20.36 Increasing the yttria and titanium con- tents improves the internal creep rupture strength (F4 > F3 > F2 > F1). The uni-axial creep rupture strength for F4 is also plotted; there is the strength anisotropy between the uni-axial and internal hoop directions. This strength anisotropy can be associated with the slightly elongated grain structure shown in Figure 19. The stress–strain rate relationship was investi- gated for ODS ferritic claddings to evaluate the creep deformation mode. The results of the analyses are given in the log–log plot in Figure 21.36 In general, the creep strain rate in the steady-state con- dition is expressed using applied stress s as: _" ¼ Asn ½5� where n is the stress exponent and A is the temperature-dependent coefficient.37 In the case of F1 F2 F3 F4 F4 (uni-axial) PNC-FMS 50 10−5 10−6 10−7 10−8 10−9 10−10 60 70 80 90 100 Stress (MPa) S tr ai n ra te (s −1 ) 200 300 Figure 21 Stress–strain rate relationship for internal creep of specimens F1–F4 and PNC-FMS, and for uni-axial creep of specimen F4 at 700 �C. Reproduced from Ukai, S.; Okuda, T.; Fujiwara, M.; Kobayashi, T.; Mizuta, S.; Nakashima, H. J. Nucl. Sci. Technol. 2002, 39(8), 872–879. 10 100 100 1000 H oo p s tr es s (M P a) 1000 10 000 Time to rupture (h) Uni-axial Internal, bi-axial direction F1 F2 F3 F4 F4 (uni-axial, this work) Figure 20 The creep rupture strength in hoop direction for pressurized F1 to F4 specimens at 700 �C. Reproduced from Ukai, S.; Okuda, T.; Fujiwara, M.; Kobayashi, T.; Mizuta, S.; Nakashima, H. J. Nucl. Sci. Technol. 2002, 39(8), 872–879. 254 Oxide Dispersion Strengthened Steels the uni-axial creep mode, a significantly high stress sensitivity of n¼ 43.7 appears. This stress exponent value is typical for an ODS alloy.37 The applied stress that initiates the strain is clearly located around 250MPa; this stress corresponds to the so-called threshold stress for deformation. On the other hand, the stress exponent, n, is 10.4 for the internal creep mode of F4, and a higher strain rate is found even below a stress of 200MPa. A transverse section of this specimen shows finely equi-axed grains of 5–10mm (Figure 19). Apart frompinning the gliding dislocations due to oxide particle-dislocation interaction, the defor- mation mechanism associated with grain morphology may be the dominant factor that induced accelerated strain in the hoop stress mode of the tubular specimen. In order to characterize the high temperature strength of manufactured 12Cr-ODS steel cladding, its strength mechanism was evaluated from the view- point of the interaction between Y2O3 particles and dislocations. This interaction could be formu- lated by the void-hardening mechanism proposed by Srolovitz,38 in which oxide particles were re- placed by voids. The oxide particle-hardening stress sp can be evaluated by the following equation based on Scattergood and Bacon’s equation,39 which takes into account the interaction between the branches of the bowed-out dislocation around a Y2O3 particle: sp=G ¼ AMb=ð2plÞ½lnðD=r0Þ þ B�; ½6� for screw dislocation; A ¼ ð1þ v sin2’Þcos ’=ð1� vÞ; B ¼ 0:6 for edge dislocation; A ¼ 1� v sin2’=ð1� vÞ� � cos ’; B ¼ 0:7 where G is the Shear modulus, v is Poisson’s ratio, M is the Taylor factor, b is the magnitude of Burgers vector, and r0 is the inner cut-off radius of the dislo- cation core. The value of ’ is the critical angle at which the dislocation detaches from the particles. This value was estimated to be ’¼ 46� for screw dislocations and ’¼ 19� for edge dislocations. Fur- ther, l is the average face-to-face distance between particles on a slip plane and is given as a function of the average particle radius rs and the average center- to-center distance ls between the particles by l ¼ 1:25ls � 2rs; ½7� where the averages are calculated by considering the size distribution of the particles. The factor 1.25 is the conversion coefficient from regular square distri- bution to random distribution.40 The characters ls and rs represent the results of the measurement of oxide particles by means of TEM. D is the harmonic mean of 2rs and l. The values of l were calculated, and the oxide particle-hardening stress was estimated by Oxide Dispersion Strengthened Steels 255 substituting l, M ¼ 3.0,41 n ¼ 0.334, b ¼ 2.48 � 10�10 m, and G ¼ 50 600MPa, at 700 �C. Figure 22 shows the results of analyses in rela- tion to the face-to-face distance between particles.36 The oxide particle-hardening stress levels estimated by using the aforementioned equations at 700 �C are represented by vertical bars, with the upper and 0 400 350 300 250 200 150 100 50 0 50 100 150 200 250 300 Face-to-face distance between particles, l (nm) Stress at strain rate of 10−9 S −1 in the internal hoop direction F4 F3 F2 F1 F4 (uni-axially longitudinal) Oxide particle-hardening stress (s p) from particle distribution by TEM H oo p s tr es s (M P a) Figure 22 Comparison of oxide particle-hardening stress estimated from dispersion parameters of F1, F2, F3, and F4 specimens, uni-axially longitudinal creep strength of F4 specimen, and internal creep strength in hoop direction at a strain rate of 10�9s�1 for F1, F3, and F4 specimens, as functions of face-to-face distance between particles. Each stress was obtained at 973K. Note that internal creep strength is located below the oxide particle-hardening stress due to the grain boundary sliding in the hoop stress mode. Reproduced from Ukai, S.; Okuda, T.; Fujiwara, M.; Kobayashi, T.; Mizuta, S.; Nakashima, H. J. Nucl. Sci. Technol. 2002, 39(8), 872–879. 0 200 400 600 800 1000 1200 1400 0 Te m p er at ur e (� C ) Generations II–III 50 Displacem GFR LFR VHTR SCWR Figure 23 Temperature and dose regimes for Generation IV ad reactor; SCWR: supercritical water-cooled reactor; GFR: gas fas SFR: sodium fast reactor. Reproduced from Guerin, Y.; Was, G. lower bars derived from an estimate of edge and screw dislocations, and with the uncertainty of r0 ranging from b to (3 � b). The measured stress in the uni-axial mode of the F4 specimen is shown by an open circle. These results imply that the higher oxide particle-hardening stress for specimen F4 is due to its shortened face-to-face particle distance l of 70 nm. The lower band represents the stress corresponding to a strain rate of 10�9 s�1 in the internal hoop direc- tional mode. For the F1 specimen, as a stress level corresponding to a strain rate of 10�9 s�1 approaches the oxide particle-hardening stress, the strong anisot- ropy tends to disappear. However, for the F3 and F4 specimens with a shortened distance between par- ticles, stress levels for a strain rate of 10�9 s�1 in the hoop direction are degraded from the oxide particle- hardening stress. The strong anisotropy still remains in the F4 specimen. The accelerated deformation in the internal hoop direction could be the result of grain boundary sliding, since finely equi-axed grains with a small size of 5–10mm are formed, and the grain bound- aries occupy a large fractional area in the transverse cross-section of the F4 specimen (see Figure 19). Based on these results, it seems to be difficult to control internal creep rupture strength by recrystallization processing in 12Cr-ODS steel cladding. 4.08.5 Al-Added 16Cr-ODS Steels 4.08.5.1 Application and Technical Issues Generation IV advanced nuclear power systems are proposed; the temperature and dose regimes for their operation are shown in Figure 23.42 Among them, the supercritical water-cooled reactor (SCWR) and the lead fast reactor (LFR) require a higher neutron dose 200100 150 ent per atom MSR SFR vanced nuclear power plants. VHTR: very high temperature t reactor; LFR: lead fast reactor; MSR: molten salt reactor; S.; Zinkle, S. J. Mater. Res. Soc. Bull. 2009, 34(1), 10–14. 12 0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 14 16 Cr content (mass%) 100 h As-received Aged at 500 �C A b so rb ed e ne rg y (J ) 18 20 22 1000 h 4300 h 10 000 h Figure 24 Aging embrittlement of high Cr-ODS steels with respect to Cr content. Absorbed fracture energy was measured at room temperature with the use of miniaturized Charpy V-notch (CVN) specimen which measures 1.5mm square with 20mm length. Reproduced from Kimura, A.; 256 Oxide Dispersion Strengthened Steels at an operating temperature of 600 �C. It is known that 9Cr-ODS steels have superior compatibility with sodium, but their corrosion resistance is not adequate for SCPW and LBE at a temperature >600 �C. Thus, the most critical issue for the application of 9Cr-ODS steels to SCWRand LFR is to improve their resistance to corrosion. It has been reported that the addition of chromium (>13wt%) and aluminum (4wt%) to ODS steels quite effectively suppresses corrosion in an SCPWand LBE environment. In general, however, an increase in the Cr content often results in increased susceptibility to ther- mal aging embrittlement. Furthermore, the addition of Al significantly reduces steel strength at high tempera- tures. Recent progress in R&D of high Cr–Al-added ODS ferritic steels is summarized in the proceedings of the International Conference of Advanced Power Plants (ICAPP) 2009. The oxidation and corrosion performance of Al-added 16Cr-ODS steels in SCPW and LBE environments is described in Section 4.08.7. Kasada, R.; Iwata, N.; et al. In Proceedings of ICAPP ’09, Tokyo, Japan, May 10–14, 2009; Paper 9220. 4.08.5.2 Thermal Aging Embrittlement Due to High Cr Content High Cr concentration often increases susceptibility to aging embrittlement through the formation of Cr-rich secondary phases. The trade-off between cor- rosion resistance and aging embrittlement caused by increasing Cr content is one of the critical issues facing the developers of high-CrODS steels. The aging effects of ODS steels with different Cr content were investigated by measuring their impact fracture energy at RTafter aging at 500 �C up to 10 kh. The results are shown in Figure 24.43 The fracture energy decreases with increasing Cr content before aging. Aging, then, causes a reduction in the fracture energy. ODS steels with a Cr content >18wt% show a significant reduc- tion in fracture energy after aging for 100 h. In contrast, 16Cr–4Al ODS steel showed a small reduction in frac- ture energy even after aging for 10 kh. Microstructure observation by TEM revealed that fine secondary phases were formed in high density after aging for 1000 h at 500 �C. These secondary phases are consid- ered to be Cr-rich phases. In order to reduce suscepti- bility to aging embrittlement, the Cr content could be 0 1400 1200 1000 800 600 400 200 0 1 2 Al concentration (mass %) U TS av e. (M P a) 3 Fe–16Cr–0W–0.1–0.35Y2O3 4 700 �C 450 �C 1000 800 600 400 200 0 Cr concentration (mass %) U TS av e. (M P a) Fe–4Al–0W–0.1–0.35Y2O3 12 13 14 15 16 17 18 19 700 �C 450 �C Figure 25 The UTS with respect of Al and Cr content in Al added high Cr-ODS steels at 450 �C and 700 �C. Reproduced from Furukawa, T.; Ohtsuka, S.; Inoue, M.; et al. In Proceedings of ICAPP ’09, Tokyo, Japan, May 10–14, 2009; Paper 9221. Time to rupture (h) 700 �C 101 103 102 10 102 103 104 105 S tr es s (M P a) 0Al (16Cr–0.1Ti–0.35Y2O3) Standard (15Cr–4Al–2W–0.1Ti–0.35Y2O3) 0.63Zr (15Cr–4Al–2W–0.63Zr–0.1Ti–0.35Y2O3) 0.62Hf (15Cr–4Al–2W–0.62Hf–0.1Ti–0.35Y2O3) Figure 26 Creep rupture strength of various Al added high Cr-ODS steels in hoop direction by using pressurized specimens at 700 �C. Reproduced from Furukawa, T.; Ohtsuka, S.; Inoue, M.; et al. In Proceedings of ICAPP ’09, Tokyo, Japan, May 10–14, 2009; Paper 9221. Oxide Dispersion Strengthened Steels 257 fracture elongation and reduction of area are slightly higher than those of Al-free steel. Under microstruc- tural observation byTEM, oxide particles consisting of Y3A5O12, YAlO3, and Al2O3, were observed in typical Al-added ODS steel, whereas this Y–Al complex oxide can be changed to Y2Hf2O7 in Hf-added ODS steel. The improved creep rupture strength in Hf-added ODS steel could be attributed to the nanosize disper- sion of the Y2Hf2O7 complex oxide. 4.08.5.4 Cladding Manufacturing 16Cr–4Al-ODS steels exhibit a full ferrite structure withouta–g-phase transformation, and themanufacture of their cladding is effected by means of cold-rolling and recrystallization annealing, which is the same as 12Cr-ODS steel described in Section 4.08.4. From mother tubes of 16Cr–2W–0.1Ti–4Al– 0.35Y2O3 with an 18mm outer diameter and 3mm thickness, cold-rolling was repeated four times using a pilger mill; the reduction rate for each rolling reached about 45%. Annealing is required to soften the deformed structure in order to make possible the next round of cold rolling. A recrystallized structure cannot be reproduced by the final heat treatment once recrystallization has taken place in the intermediate annealing.45 A heat treatment test was therefore con- ducted to determine an appropriate annealing tem- perature to induce softening by dislocation recovery but without recrystallization. The temperature for the final heat treatment was selected in order to realize recrystallized grains. Figure 27 illustrates an orientation image map (OIM) measured after heat treatment at 900 �C for 1 h, 1000 �C for 1 h, and 1150 �C for 1 h, from which an evolution of recrystallization can be clearly identi- fied in the cladding.46 The primary recrystallization beginswith an {111} orientation,which is desig- nated by the color blue at 900 �C annealing. When the temperature increases to 1000 �C, a {110} Goss orientation, designated by green, partially appears. The primary recrystallization orientation of {111} is almost replaced by a {110} Goss orientation when the temperature is elevated to 1150 �C, which corresponds to a secondary recrystalli- zation. Thus, the final heat treatment was conducted at 1150 �C for 1 h in order to create a perfectly recrys- tallized structure. The manufactured cladding has sec- ondary recrystallized grains with a {110} Goss orientation, which was formed by annealing from the cold rolled {112} orientation through primary recrystallization of the {111} orientation.47 258 Oxide Dispersion Strengthened Steels 4.08.6 Existing ODS Steel Cladding The basic chemical composition of the representative ODS steels is summarized inTable 2. ODS steels are divided into two groups which have either been 200 μm (a) 200 μm (b) 200 μm (c) Figure 27 OIM for 16Cr–4Al–ODS steels heat-treated at 900 �C for 1 h (a), 1000 �C for 1 h (b), and 1150 �C for 1 h (c). Reproduced from Ukai, S.; Ohnuki, S.; Hayashi, S.; et al. In Proceedings of ICAPP’09, Tokyo, Japan, May 10–14, 2009; Paper 9232. Table 2 Basic chemical composition of ODS steels (mass % Steels Cr Mo W Ti Al Dis Turbine, combustion Incoloy MA956 20 – – 0.5 4.5 0.5Y PM2000 19 – – 0.5 5.5 0.5Y Fast reactor fuel DT2203Y05 13 1.5 – 2.2 – 0.5Y 0. DT2906 13 1.5 – 2.9 – 1.8T Incoloy MA957 14 0.3 – 1 – 0.25 9Cr-ODS steel 9 – 2 0.2 – 0.35 12Cr-ODS steel 12 – 2 0.3 – 0.23 16Cr–4Al-ODS steel 15.5 – 2 0.1 4 0.35 SM: Special metals, former International Nickel Company; JAEA: Japa SCK�CEN: Centre d’Etude de l’énergie Nucleaire – Studiecentrum voo commercialized or are under development. The first group includes Incoloy MA956 and PM2000. The former is produced by what was formerly the Inter- national Nickel Company (INCO) and is now the Special Metals (SM) Company. The latter is a prod- uct of the Plansee Company of Austria. MA956 and PM2000 are 20% Cr-ODS steels containing 5% Al, which exhibit superior resistance to oxidation and corrosion in hot gases at temperatures >1000 �C. Tubes, sheets, and bars made from these steels are commercially used in various stationary and high- temperature components in turbines, combustion chambers, diesel engines, and burners. The second group is devoted to the application of fuel cladding for nuclear fast reactors, anticipating its superior resistance to radiation resistance, and its excellent creep strength and dimensional stability at an elevated temperature of 700 �C. As shown in Table 2, DT2906 contains Ti2O3 dispersoids, and DT2203Y05 is strengthened by Ti2O3 and Y2O3. Both steels have been developed by SCK�CEN (Centre d’Etude de l’énergie Nucleaire – Studiecentrum voor Kernenergie) Mol (Belgium).48–50 The elementary metallic powders and Y2O3 or TiO2 powder are mechanically alloyed by means of a pilot scale ball mill with a capacity of 9.2 kg per batch. Mechanically alloyed powders are hot-compacted into billets, which are subsequently hot-extruded into the hol- lows of 20/17mm. A plug drawing is applied to manufacture the cladding tube from the hollows. Intermediate annealing is carried out at 1050 �C by using induction heating after a certain number of drawing passes. The entire cold drawing is composed of 15–20 passed and three intermediate annealing ) persoid Fe Others Development 2O3 Bal SM/US 2O3 Bal Plansee/Austria 2O3, 9Ti2O3 Bal SCK�CEN Mol/Belgium i2O3 Bal SCK�CEN Mol/Belgium Y2O3 Bal SM/US Y2O3 Bal 0.13C, martensite + residual ferrite JAEA/Japan Y2O3 Bal JAEA/Japan Y2O3 Bal 0.6Hf or 0.6Zr KU/Japan n Atomic Energy Agency; KU: Kyoto University. r Kernenergie. Oxide Dispersion Strengthened Steels 259 steps. The final annealing is performed at 1050 �C and 800 �C to precipitate an w-phase (70%Fe, 15%Cr, 7% Ti, and 6%Mo). More than 1000 cladding tubes were manufactured. For defect control, this cladding is non- destructively tested using eddy currents and ultrason- ics which employ specified artificial reference defects which define the rejection level for naturally defective cladding. For example, the creep rupture strength of DT2203Y05 cladding in the hoop direction is shown in Figure 28.51 For the fabrication of fuel pins with DT2203Y05 cladding, a special resistance welding machine was designed at SCK�CEN, because ODS steels can hardly be welded by conventional fusion welding methods such as tungsten inert gas (TIG) or electron beam welding, since they result in an oxide particle-free zone. Fuel and blanket pellets were filled into the cladding, and resistance welding with an end- plug was performed in a glove box at Belgonucleaire. The two fuel assemblies were fabricated for Phenix irradiation. Incoloy MA957 was developed by the International Nickel Company (INCO) for application to fast reac- tor fuel cladding. It is strengthened by a very fine, uniformly distributed yttria dispersoid. Its fabrication involves a MA process and subsequent extrusion, which ultimately results in a highly elongated grain structure. An extruded bar with a diameter of 25.4mm was gun-drilled in order to generate a tube hollow with a 4.75mm thick wall. Extensive cladding fabrica- tion tests were conducted on the tube hollow using a rolling and plug draws in the United States, France, and Japan. It can be said that MA957 is too hard to perform satisfactorily on a small scale without faults. The structure of the fabricated MA957 cladding is highly anisotropic with equi-axed grains in the 10 10 100 1000 100 600 �C 650 �C 700 �C 750 �C Time to rupture (h) H oo p s tr es s (N m m –2 ) 1000 10 000 Figure 28 Creep rupture strength of DT2203Y05 cladding in a hoop direction. Reproduced from Huet, J. J.; Coheur, L.; De Bremaecker, A.; et al. Nucl. Technol. 1985, 70, 215–219. transverse direction, but with highly elongated grains with a bamboo-like structure in the longitudinal or working direction. Therefore, it turned out that the creep rupture strength of MA957 cladding is signifi- cantly degraded in the hoop direction, which is essen- tial for fuel pins. Some of the stress rupture data are shown in Figure 29.52 The pulsed magnetic welding (PMW) method was developed in the United States for MA957 for the manufacture of fuel elements. 4.08.7 Corrosion and Oxidation 4.08.7.1 Sodium Compatibility It is essential to evaluate the environmental effects of sodium on the mechanical strength properties of ODS steels to ensure their structural integrity throughout their design life-time in SFR. ODS-steels basically display superior compatibility with sodium. For 9Cr- ODS steel (M93) and 12Cr-ODS steel (F95), which are potential cladding materials for SFR, their UTS at 700 �C after exposure to sodium in a stagnant state is shown in Figure 30.53 Both show almost constant strength after exposure to sodium, and it was confirmed that there is no degradation up to 10 000 h. For conven- tional ferritic steel without Y2O3, a clear strength reduction occurs above 600 �C due to decarburization phenomena in sodium. ODS steel does not show such a clear strength reduction because the fine Y2O3 oxide particles remain stable in steel, therebymaintaining the strength of the steel. Figure 31 shows the results of creep-rupture tests with internally pressurized specimens in a stagnant sodium environment.54 The creep-rupture strength of 9Cr-ODS steel (M11) in sodium is equal to its strength in air, and no impact from a sodium environ- ment was observed. However, under a flowing sodium condition of 4.5m s�1, the element nickel penetrates the surface of ODS steel cladding, where an increase in nickel concentration and decrease in chromium con- centration were observed at 700 �C. These results suggest that the effects of a sodium environment can be ignored under stagnant conditions; however, as fuel cladding is utilized in an environment with a high flow rate of sodium, the effects of the microstruc- ture change associated with nickel diffusion into the cladding surface need to be considered.53 4.08.7.2 LBE Compatibility Molten LBE has a high solubility of nickel, iron, and chromium,which are themost important alloyelements 10 10 100 100 650 �C 704 �C 760 �C tr (h) s ( M P a) 1000 1000 10 000 39011045.6 10 000 650 �C 700 �C 760 �C STC TR PNC STC ORT → → → → → → → Figure 29 Creep rupture strength of IncoloyMA957 cladding. Reproduced fromHamilton, M. L.; Gelles, D. S. PNNL-13168, Feb 2000. 0 20 40 60 80 U TS As-received Sodium exposure time (h) 0 200 400 600 800 1000 1200 M93: 700 �C F95: 700 �C M93: 650 �C Sodium flow < 0.001 m s–1 F95: 650 �C Figure 30 UTS of 9Cr-ODS steel (M93) and 12Cr-ODS steel (F95) in hoop direction after sodium exposure. Reproduced from Yoshida, E.; Kato, S. J. Nucl. Mater. 2004, 329–333, 1393–1397. 260 Oxide Dispersion Strengthened Steels in austenitic stainless steels. Thus, nickel super alloys and austenitic stainless steels cannot be used as the structural materials for LBE-cooled systems, especially at temperatures >500 �C. Ferritic steels have been con- sidered more appropriate for LBE application. Exposure of 9Cr-ODS steels to an LBE environ- ment at 530 �C was carried out in the DELTA Loop of the Los Alamos National Laboratory. The molten alloy flow velocity in the loop is 1.2 m s�1, and oxygen sensors were used to measure and maintain an oxygen concentration of about 1 � 10�6wt%. Samples were exposed for 200, 400, and 600 h, in order to study the early stages of oxide formation and growth. A cross- sectional view and the distribution of elements are shown in Figure 32.54 In a short time, the 9Cr-ODS steel formed a protective duplex oxide layer consisting 100 1000 10 100 1000 10 000 100 000 Time to rupture (h) H oo p s tr es s (M P a) Material: M11 650 �C 700 �C 700 �C : In air/Ar 700 �C : In sodium Sodium flow rate; < 0.001 m s–1 650 �C 650 �C Figure 31 Creep-rupture strength of 9Cr-ODS steel (M11) in hoop direction under sodium exposure at 650 �C and 700 �C. Reproduced from Yoshida, E.; Kato, S. J. Nucl. Mater. 2004, 329–333, 1393–1397. OKa, 41 CrKa, 26 FeKa, 64 PbMb, 124 Diff. zone Bulk 10 μmX 2000 24 36 BEC15 kV Oxide LBE Figure 32 Backscatter cross-section secondary electron microscopy (SEM) image and Energy dispersive X-ray spectrometry (EDS) map of 600 h 9Cr-ODS steel, showing much thinner Cr-rich oxide but a thicker diffusion zone. Reproduced from Machut, M.; Sridharan, K.; Li, N.; Ukai, S.; Allen, T. J. Nucl. Mater. 2007, 371, 134–144. Oxide Dispersion Strengthened Steels 261 of an outer magnetite (Fe3O4) layer and an inner Fe–Cr spinel ((Fe,Cr)3O4) layer, which is sometimes accom- panied by an O-enriched and Fe-depleted diffusion zone at the oxide–bulk interface. Over time, the outer magnetite layer is removed and the underlying spinel layer serves to mitigate more catastrophic corrosion degradation such as dissolution and liquid metal attack along the grain boundaries. Very thin oxides are not particularly protective in regard to loss of metal, as manifested by the thick diffusion zones associated with them. Furukawa pointed out that at tempera- tures above 600 �C, the thickness of the oxide layer diminishes with increasing temperature. This behavior can be ascribed to a change in the stable form of iron oxide frommagnetite towustite at 570 �C. Beyond this temperature, dissolution attack was observed on some portions of 9Cr-ODS steel, and the oxide layer’s adhesion to the material began to weaken.55 It has been reported that the addition of aluminum to steel effectively prevents LBE corrosion. Figure 33 shows the appearance of ODS steel specimens after a corrosion test in LBE for 1 � 104 h at 650 �C.43 The 18wt% Cr-ODS steel without the addition of Al dissolved markedly into LBE, while those ODS specimens containing 4wt% Al almost completely maintained their shape even in Al-added 14Cr- and 16Cr-ODS steels, indicating a very high resistance to LBE corrosion. It is noteworthy that this corrosion resistance was independent of Cr concentration from 13 to 19 wt% in Al-added ODS steels. From the distribution of elements across the cladding surface, we deduce that LBE corrosion can be prevented by the formation of an Al enriched film.56 It was demonstrated that Al-added 16Cr-ODS steel (16Cr– 2W–4Al–0.1Ti–0.35Y2O3) has superior corrosion resis- tance at 650 �C for 5000 h. 14Cr–4Al 16Cr–4Al 19Cr–4Al 18Cr Figure 33 The appearance of Al added high Cr-ODS steel specimens after corrosion test in LBE for 1 � 104 h at 923K (DO: 1 � 10�6wt%). Reproduced from Kimura, A.; Kasada, R.; Iwata, N.; et al. In Proceedings of ICAPP ’09, Tokyo, Japan, May 10–14, 2009; Paper 9220. W ei gh t ga in (m g d m –2 ) 14 16 18 Cr content (wt%) 6 4 2 0 With 4 wt% Al 1800 h 600 h 100 h 0 2 4 Al content (wt%) 21 14 7 0 With 16 wt% Cr 1800 h 600 h 100 h (a) (b) W ei gh t ga in (m g d m –2 ) Figure 34 Weight gain of Al added high Cr-ODS steels with Cr content (a) and Al content (b) after exposure to SCPW at 500 �C with 8ppm of dissolved oxygen under a pressure of 25MPa (10dm = 1m). Reproduced from Lee, J. H.; Kimura, A.; Kasada, R.; et al. In Proceedings of ICAPP ’09, Tokyo, Japan, May 10–14, 2009; Paper 9223. 262 Oxide Dispersion Strengthened Steels 4.08.7.3 SCPW Compatibility Figure 3457 shows the effects of Cr and Al content on the weight gain of ODS ferritic steels after exposure to SCPW at 500 �C with 8 ppm of dissolved oxygen. Increasing the Cr content from 14 to 17wt % does not affect corrosion resistance if ODS ferritic steels contain 4wt % Al. For 16wt% Cr, the addition of A1 increases corrosion resistance in 16Cr-ODS steels. As shown in Figure 35,43 tested at SCPW (510 �C, 25MPa) for 600 h, the addition of 4wt% Al did not significantly influence corrosion resistance in 19Cr- ODS steel, though a rather dense chromia film was observed on the specimen surface. The 16wt% Cr is not large enough to form homogeneous and stable chromia on the entire surface of the specimen, whereas a very thin alumina film covers the entire surface of the specimen with the Al addition of 2wt%. Thus, the addition of Al effectively improves corrosion resis- tance in 16Cr-ODS steel. As shown in a comparison with 9Cr-ODS steel in Figure 35, its weight gain is much larger than 16Cr-ODS steel, indicating that 9Cr-ODS steel is not adequate for application to SCWR. The suppression of SCPW corrosion by the addition of Al to 16Cr-ODS steel is due to the for- mation of a very thin alumina film on the surface. 4.08.7.4 Oxidation Oxidation tests for 9Cr-ODS and 12Cr-ODS steels were performed using pickled specimens in a controlled atmosphere of dry air. Weight measure- ment to evaluate the degree of oxidation was per- formed at intervals of 50, 100, 400, 1000, and 2000 h, at temperatures of 650, 750, and 850 �C. The results of the measured weight gain due to oxidation at 750 �C are shown in Figure 36.58 For 9Cr-ODS and 12Cr-ODS steels, the weight gain due to oxida- tion was quite small and comparable to that of PNC316 containing 17 wt% Cr. Their weight gain is limited to below 0.1mgmm�2. On the other hand, a quite large oxidation of 0.8mgmm�2 was observed in PNC-FMS. The measured results on SUS430, which show a greater weight gain than that of ODS steels, show that advanced oxidation resistance is attained with ODS steels, even when compared to higher 17wt% Cr containing stainless steel. The element distribution obtained by Electron probemicroanalysis (EPMA) showed a scale consisting Testing time (h) 1500 25002000500 10000 0.0 0.2 0.4 0.6 0.8 1.0 W ei gh t ga in b y ox id at io n (m g m m –2 ) 9Cr-ODS 12Cr-ODS (fine grain) SUS430 PNC-FMS PNC316 Figure 36 Weight gain of 9Cr-ODS and 12Cr-ODS steels by the oxidation at 750 �C. Reproduced from Kaito, T.; Narita, T.; Ukai, S.; Matsuda, Y. J. Nucl. Mater. 2004, 329–333, 1388–1392. 12Cr–ODS (fine grain) Cr supply through grain boundary diffusion 12Cr–ODS (large grain) 923 K � 50 h Y2O3 effects PNC–FMS 0.00W ei gh t ga in b y ox id at io n (m g m m –2 ) 0.02 0.06 0.08 0.10 0.04 Figure 37 Weight gain of 12Cr-ODS steel and PNC-FMS oxidized at 650 �C for 50 h. Reproduced from Kaito, T.; Narita, T.; Ukai, S.; Matsuda, Y. J. Nucl. Mater. 2004, 329–333, 1388–1392. Al content (mass %) 0 0 0.6 1.2 10.0 20.0 30.0 19Cr-ODS 16Cr-ODS 9Cr-ODSSUS430 (16Cr) SCW 773 K.25 MPa. 600 h 2 4 W ei gh t ga in (g m –2 ) Figure 35 The dependence of the weight gain on the Cr and Al contents in 16Cr- and 19Cr-ODS steels. SUS430 is a ferritic steel containing 16 mass % Cr and 4 mass % Al. Reproduced from Kimura, A.; Kasada, R.; Iwata, N.; et al. In Proceedings of ICAPP ’09, Tokyo, Japan, May 10–14, 2009; Paper 9220. Oxide Dispersion Strengthened Steels 263 of Fe-rich oxide in the outer layers andCr-rich oxide in the inner layers. At the interface between ODS steel and the oxide scale, there was a thin layer (a few micrometers) of further Cr-enriched oxide. Raman spectroscopy measurement indicated that the outer Fe-rich and inner Cr-rich layers correspond to a-Fe2O3 and spinel type (Fe, Cr)3O4, respectively. It was also confirmed that a-Cr2O3 is formed at the matrix–scale interface. In oxidation tests, Fe, which is a major constituent in steel, tends to be easily oxidized at an early stage, but further oxidation can be suppressed by the formation of a protective a-Cr2O3 layer. This a-Cr2O3 formation is generally controlled by the rate at which Cr is supplied to the reaction front. It is known that a high Cr content in steel, as well as an increasing diffusion flux through the grain boundary, that is, finer grains, accelerates both the Cr supply and the formation of a-Cr2O3. A short-term oxidation test, whose results are shown in Figure 37, was conducted to investigate the mechanism of suppressing oxidation in ODS steels.58 The decrease in oxidation in fine grain 12Cr-ODS ferritic steel can be attributed to the enhanced rate at which Cr was supplied throughout the accelerated grain boundary diffusion. In both cases of fine/large grains in 12Cr-ODS steels, Raman spectroscopy detected protective a-Cr2O3 at the interface between the matrix and scale. Comparing 12Cr-ODS large grain and PNC-FMS, the Cr content is similar, and the grain size is rather smaller in PNC-FMS. Never- theless, protective a-Cr2O3 cannot be detected by Raman spectroscopy, and oxidation is enhanced in PNC-FMS, implying that the suppression of oxidation in 12Cr-ODS with large grains could be due to the effects of the Y2O3 oxide particles themselves. Chen et al. showed some TEM images of Y-rich oxides on grain boundaries that may be part of the explanation.59 4.08.8 Irradiation 4.08.8.1 Simulated Irradiation Testing that involves the simulated irradiation of 9Cr-ODS steel was conducted by Allen et al. at the 264 Oxide Dispersion Strengthened Steels Environmental and Molecular Science Laboratory at Pacific Northwest National Laboratory, using 5MeV Ni ions at 500, 600, and 700 �C with a damage rate of 1.4 � 10�3 dpa s�1. The results regarding measured particle size distribution as a function of dose are plotted in Figure 38 for irradiation at 500, 600, and 700 �C.60 Due to TEM’s limited resolution of the images, particles smaller than 2 nm were not detected. At all temperatures, the size of the oxide particles decreases as the dose increases. At higher tempera- tures (600–700 �C), the average size appears to reach a value of �5 nm. At all three temperatures, the density increases as the radiation dose increases. The decrease in size takes place faster at 600 and 700 �C than at 500 �C, indicating that the reduction in size is not strictly a ballistic effect and that a diffusion-based mechanism is also involved in the dissolution. Allen extensively reviewed previous papers that presented different approaches to the irradiation of ODS ferritic–martensitic steels that employed various ion beams, electrons, and neutrons; the results are summarized in Table 3.61 A great many findings asserted that oxide particles are stable under radiation. However, as shown inTable 4, the dissolution of oxide particles at higher temperatures and doses has been reported in other studies. Dubuisson62 and Monnet63 reported that small oxides dissolved under radiation at higher temperatures and doses, but did not dissolve at a lower irradiation dose. Their data will be dis- cussed in detail in the following section. In material irradiated in the JOYO fast reactor at temperatures 450–561 �C to doses of 21 dpa, Yamashita found that small particles disappear and average particles increase slightly in size with increasing temperature or dose.64 Monnet supplemented neutron radiation studies with the electron irradiation of yttrium oxides and magnesium oxides in the EM10 alloy at tempera- tures between 300 and 550 �C, and to doses of 100 dpa. In these studies, the yttrium oxides were stable at 400 �C when irradiated with 1.0MeV electrons, but dissolved under 1.2MeV electron irradiation. Allen59 pointed out that the displacement energy for Y and O in yttrium oxide is 57 eV65,66 while that for iron is 40 eV. Assuming similar displacement energies in the Y–Ti–O oxide, the radiation-induced vacancy concentration should be larger in the metal matrix, providing a driving force for a net vacancy flux to the precipitate. This could drive the precipi- tate mass loss if vacancy absorption frees a precipitate atom. From a comparison between electron irradia- tion (Frenkel pairs) and ion irradiation (displacement cascades), Monnet63 also concluded that the ballistic ejection of atoms alone cannot be responsible for the loss of diameter in oxide particles. Free point defects and their diffusion-based mechanism are therefore of major importance and play a dominant role in the dissolution of oxide particles. 4.08.8.2 Neutron Irradiation of Materials The 5mm-wide ring-tensile specimens with a 1.5mm-wide gauge section were prepared from the cladding of 12Cr-ODS steels (F94, F95, and 1DS) and 9Cr-ODS steels (M93).67 This type of specimen makes it possible to test mechanical properties in the hoop direction of the cladding. These ring-tensile samples were irradiated in the experimental fast reac- tor JOYO using the material irradiation rig at tem- peratures between 400 and 534 �C to fast neutron fluences ranging from 5.0 � 1025 to 3.0 � 1026 nm�2 (E > 0.1MeV). The yield strength of the irradiated samples as a function of test temperature is shown in Figure 39, together with that of the unirradiated ones.67 After irradiation, the yield strength of irra- diated F94, F95, and M93 cladding, is modestly higher ( Particle size (nm) 20 30 40 10 0 Fr ac tio n (% ) 20 30 40 10 0 Fr ac tio n (% ) 20 30 40 10 0 Fr ac tio n (% ) 20 30 40 10 0 2 4 6 8 10 12 14 16 18 20 22 Fr ac tio n (% ) 0 T = 700 �C, dose = 5 dpa T = 700 �C, dose = 50 dpa T = 700 �C, dose = 150 dpa Unirradiated condition Particle size (nm) Particle size (nm) 20 15 10 5 0 Fr ac tio n (% ) 20 15 10 5 0 Fr ac tio n (% ) 20 15 10 5 0 Fr ac tio n (% ) 20 15 10 5 0 Fr ac tio n (% ) 2 4 6 8 10 12 14 16 18 20 220 2 4 6 8 10 12 14 16 18 20 220 20 30 40 10 0 Fr ac tio n (% ) 20 30 40 10 0 Fr ac tio n (% ) 20 30 40 10 0 Fr ac tio n (% ) T = 500 �C, dose = 50 dpa T = 500 �C, dose = 150 dpa T = 600 �C, dose = 50 dpa T = 600 �C, dose = 5 dpa Unirradiated conditionUnirradiated condition T = 500 �C, dose = 5 dpa Figure 38 Particle size (diameter) distribution for 9Cr-ODS steel irradiated at 500, 600, and 700 �C to doses of 0, 5, 50, and 150dpa. Reproduced from Allen, T. R.; Gan, J.; Cole, J. I.; et al. J. Nucl. Mater. 2008, 375, 26–37. Oxide Dispersion Strengthened Steels 265 Table 3 Historical survey of yttrium–titanium-oxides reported to be stable under radiation Author Material Irradiation particle (dpa) Temperature (�C) Dose (dpa) Dose rate (dpa s�1) Result Pareige et al.74 12YWT 150 keV Fe 300 0.7 1.9 � 10�4 Stable dispersions Asano et al.75 MA957 1MeV He (þ) 450 150 2 � 10�3 Stable oxides 4MeV Ni 650 Hide et al.76 MA957 42 keV He 475 200 1.0–1.4 � 10�2 Stable oxides at 25 �C (þ) 525 200 keV C� 575 625 Hide et al.76 MA957 220 keV He at 25 �C 475 150 3.0 � 10�3 Stable oxides (+)3MeV Niþ 525 Little77 DT2203YO5 52MeV Cr6 þ (þ) 475 50 3.0 � 10�4 Stable oxides 4MeV He Saito et al.61 13Cr–0.5TiO2 –0.2Y2O3 1MeV electron 400 12 2.2 � 10�3 Stable oxides 500 Kinoshita et al.78 13Cr-ODS (+) Nb, V, Zr 1MeV electron 350 15 2 � 10�3 Stable oxides 450 Akasaka et al.79 9Cr and 12Cr-ODS JOYO 330 7.0 Not reporteda Stable oxides 400 2.5 450 14.0 500 15.0 Mathon et al.80 MA957 Thermal neutrons 325 0.8, 2.0 1 � 1014 n cm�2 Stable dispersions (OSIRIS) 3.5, 5.5 (E > 1MeV) Monnet et al.63 DY EM10þY2O3EM10 þ MgO 1MeV Helium 400 0.05 Not reported No change in oxide particles Kimura et al.81 (13–19)Cr–4Al-ODS 300 20 1 � 10�4 No reported change in oxide size –500 aTypical fast reactor displacement rates in the driver fuel portion of the core are 1� 10�6 dpa s�1. Source: Reproduced from Allen, T. R.; Gan, J.; Cole, J. I.; et al. J. Nucl. Mater. 2008, 375, 26–37. 2 6 6 O x id e D is p e rs io n S tre n g th e n e d S te e ls Table 4 Historical survey of yttrium–titanium-oxides reported change size under radiation Author Material Irradiation Temperature (�C) Dose (dpa) Dose rate (dpa s�1) Result Yamashita et al.64 IDS (11Cr) JOYO 450–561 21 Not reporteda Small particles disappear. Average particles increase slightly with increasing temperature or dose. IDK (13Cr) Dubuisson et al.62 DT2203YO5 Phenix 400–580 81 Not reporteda Oxide particles are totally dissolved (small oxides) or reduced in size and were surrounded by a halo of smaller oxides (large oxides). Monnet et al.63 DT2203YO5 Phenix 400–580 81 Not reporteda Disappearance of small oxides and significant halo of smaller oxides at higher temperatures and doses. Monnet et al.63 DY EM10 + Y2O3 1MeV and 1.2MeV 300–550 100 3–6 � 10�3 Oxides stable at 400 �C under 1.0MeV electrons but dissolve under 1.2MeV. EM10 + MgO Electron aTypical fast reactor displacement rates in the driver fuel portion of the core are about 1� 10�6 dpa s�1. Source: Reproduced from Allen, T. R.; Gan, J.; Cole, J. I.; et al. J. Nucl. Mater. 2008, 375, 26–37. Fluence Test temperature (K) (Fluence; � 1026 n m–2 (E > 0.1 MeV) Y ie ld s tr en gt h (M P a) 600 (0.5) (1.35) (3.56) (0.45) (2.8) (3.0) (1.4, 2.5) 600 400 200 650 700 750 800 800 900 1000 1200 850 F94 F95 M93 F94 unirrad. M93 unirrad. 1DS unirrad. F95 unirrad. Y2O3(wt%) 0.24 0.24 0.35 0.40 F95 M93 1DS F94 Figure 39 Yield strength of 9Cr-ODS steel (M93) and 12Cr-ODS steels (F94, F95, 1DS) in hoop direction by ring specimens before and after irradiation. Reproduced from Yoshitake, T.; Abe, Y.; Akasaka, N.; Ohtsuka, S.; Ukai, S.; Kimura, A. J. Nucl. Mater. 2004, 329–333, 342–346. Fluence Test temperature (K) (Fluence; � 1026 n m–2(E > 0.1MeV) U ni fo rm e lo ng at io n (% ) 600 (0.5) (1.35) (3.56) (0.45) (2.8) (3.0) (1.4, 2.5) 6 4 2 650 700 750 800 8 900 10 0 850 F94 F95 M93 F94 unirrad. M93 unirrad. 1DS unirrad. 1DS F95 unirrad. Figure 40 Uniform elongation of 9Cr-ODS steel (M93) and 12Cr-ODS steels (F94, F95, 1DS) in hoop direction by ring specimens before and after irradiation. Reproduced from Yoshitake, T.; Abe, Y.; Akasaka, N.; Ohtsuka, S.; Ukai, S.; Kimura, A. J. Nucl. Mater. 2004, 329–333, 342–346. Oxide Dispersion Strengthened Steels 267 xenon and krypton tag gases was enclosed. The irradi- ation temperatures were 700, 725, and 750 �C, and the hoop stress ranged from 45 to 155MPa. Themaximum neutron dose reached 20 dpa. It was confirmed that in- pile creep rupture time is located within the out-of- pile data band, and there is no degradation in creep strength due to irradiation.68 MA957 and MA956 were irradiated in Fast Flux Test Facility (FFTF)-Materials Open Test Assembly (MOTA) at 420 �C up to 200 dpa.69 No voids were seen in this area, but precipitates did appear, which were expected to be a0. The results regarding the radiation damage resistance of ODS steels were highly encourag- ing. Evidence was apparent in both MA956 andMA957 0.2 mm Figure 41 Longitudinal cross-sectional structure in the vicinity of welded section by PRW (9Cr-ODS steel cladding and endplug). Reproduced from Ukai, S.; Kaito, T.; Seki, M.; Mayorshin, A. A.; Shishalov, O. V. J. Nucl. Sci. Technol. 2005, 42(1), 109–122. 268 Oxide Dispersion Strengthened Steels of a0 precipitation, and in regions where recrystalliza- tion occurred before irradiation in MA957, a few voids were slightly observed. Gelles69 pointed out that these could be overcome by employing suitable alloy design and that ODS steel microstructures, when properly manufactured to provide a uniform oxide dispersoid in a structure, appear to be completely resistant to radiation damage at doses as high as 200 dpa. Figure 42 Optical micrograph of 9Cr-ODS fuel pin after irradiation at 700 �C, 5 at.% burnup and 25dpa in BOR-60. Reproduced from Kaito, T.; Ukai, S.; Povstyanko, A. V.; Efimov, V. N. J. Nucl. Sci. Technol. 2009, 46(6), 529–533. 4.08.8.3 Fuel Pin Irradiation 4.08.8.3.1 9Cr- and 12Cr-ODS steel cladding in BOR-60 In order to weld 9Cr- and 12Cr-ODS steel claddings with end-plugs for the manufacture of fuel pins, the PRW method was developed in JAEA, which makes joining possible in the solid state condition.70 This method is based on the electrical resistance heating of the components, while maintaining a continuous force sufficient to forge-weld without melting. The appropriate conditions, for example, electric current, voltage, and contact force, were selected. For the PRW- welded specimens, tensile, internal burst, and creep rupture tests, were conducted and their integrity was confirmed. In addition, a nondestructive ultrasonic inspection method was developed to assure the integ- rity of the weld between the cladding and end-plug. Using this PRW method, upper end-plugs were welded for two types of 9Cr-ODS steel cladding (Mm13) and 12Cr-ODS steel cladding (F13) at JAEA. Figure 4171 shows a cross-section of the welded part between the 9Cr-ODS steel cladding and end-plug. The ODS steel cladding welded to the upper end- plug was shipped to the fuel production facility of the Institute of Atomic Reactor (RIAR) in Russia where the MOX and UO2 granulated fuels, as well as uranium metal getter particles, were vibro-packed into the ODS steel cladding, and the lower end-plug was welded by the TIG end-face method. The TIG-welded part at the lower end-plug ensured that its integrity would be maintained at a lower temperature of 400 �C. The inspection and quality control of the fabricated ODS fuel pins were done through X-ray analysis, gamma scanning, and leak testing, etc., which confirmed that the fuel pins satisfied BOR-60 requirements. The fuel pins were loaded into two dismountable experimental assemblies to satisfy the cladding middle wall tempera- ture within 700 �C and 650 �C, and irradiation was conducted in the BOR-60 up to 5 at.% burnup and 25 dpa as the collaborative work between JAEA in Japan and RIAR of Research.71 The results of the postirradiation examination are shown in Figure 42 in the optical micrographs of the upper part of the fuel column of 9Cr-ODS steel fuel; no obvious corrosion inside the cladding was observed.72 The maximum depth of corrosion of 25 mm is partially confirmed in the upper part of the fuel column. The inner corrosion of the ODS cladding can be reduced by using a lower O/M (a) (b) (c) 100 nm 1 μm 1 μm Figure 43 Precipitation occurring during the in-pile service (a) a0-phases at 400 �C, �0dpa, (b) w-phases at 523 �C to 78.8dpa, and (c) Laves phases at 580 �C to 30.5dpa. Reproduced from Dubuisson, P.; Schill, R.; Higon, M.-P.; Grislin, I.; Seran, J.-L. In Effects of Radiation on Materials: 18th International Symposium; Nanstad, R. K., Hamilton, M. L., Garner, F. A., Kumar, A. S., Eds.; American Society for Testing and Materials: Philadelphia, PA; p 882, ASTM STP 1325. Oxide Dispersion Strengthened Steels 269 ratio fuel, even in lower Cr content cladding such as 9Cr-ODS steel. 4.08.8.3.2 12Cr-ODS steel cladding in EBR-II JAEA manufactured 12Cr-ODS steel cladding (1DK and 1DS) and Argonne National Laboratory in the United States qualified a welding process that employs PRW. Fuel pins composed of 12Cr-ODS steel cladding andMOX fuel pellets were successfully fabricated and qualified, and irradiated up to 35 dpa at EBR-II.73 The ODS cladding with high smear density solid pellet MOX fuel did induce some diametral strain, demon- strating some in-core ductility. This program demon- strated the viability of ODS steel as a potential cladding material for long-life advanced FRs. 4.08.8.3.3 DT2203Y05 in Phénix Fuel pins with DT2203Y05 cladding were irradiated in an experimental capsule placed in a special subas- sembly in Phénix. The process by which they were manufactured was described in Section 4.08.6. The dose reached at midplane was 81 dpa and the temper- ature along the fuel pin ranged from 400 to 580 �C. It was observed by TEM that the uniform distribu- tion of fine oxides totally disappeared, and a few large oxides were also fragmented into smaller ones. The recoil resolution of particles is a process where the atoms that compose particles are ballistically ejected by an impinging neutron. Dubuisson63 pointed out that the atoms ejected from oxides by ballistic dissolu- tion depend on radiation-enhanced solute diffusivity and enhanced solubility under irradiation. A uniform distribution of tiny particles 270 Oxide Dispersion Strengthened Steels information concerning the irradiation embrittle- ment of DT2203Y05 cladding produced by CEN-SCK Mol due to the formation of a0-phase below 500 �C and w-phase above 500 �C, 9Cr-ODS and 12Cr-ODS steels containing low Cr and low Ti were developed. 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All rights reserved. 4.09.1 Welding Defects 273 4.09.1.1 Supersolidus Cracking 274 4.09.1.1.1 Solidification cracking 274 4.09.1.1.2 Liquation cracking 278 4.09.1.1.3 Hot tearing 278 4.09.1.2 Subsolidus Cracking 279 4.09.1.2.1 Precipitation-induced cracking 279 4.09.1.2.2 Segregation-induced cracking 282 4.09.1.3 Other Welding Defects 285 4.09.2 Stresses and Strains in Welds 285 4.09.2.1 Quantification of Residual Stresses and Strains 287 4.09.2.1.1 Elastic stress 287 4.09.2.1.2 Plastic strain 287 4.09.3 In-Service Performance 289 4.09.3.1 Environmentally Assisted Cracking 289 4.09.3.2 Microchemical Changes 291 4.09.3.3 Microstructural Changes 293 4.09.4 Weldability of Specific Alloy Systems 294 4.09.4.1 Low-Alloy Steels 294 4.09.4.2 Austenitic Stainless Steels 294 4.09.4.3 Nickel-Based Alloys 295 4.09.4.4 Zirconium Alloys 295 References 296 Abbreviations AMIS Average intragrain misorientation CTE Coefficient of thermal expansion EAC Environmentally assisted cracking GTAW Gas tungsten arc welding HAZ Heat-affected zone PIC Precipitation-induced cracking PMZ Partially melted zone SCC Stress corrosion cracking SIC Segregation-induced cracking SMAW Shielded metal arc welding SS Stainless steel Zr-4 Zircaloy-4 4.09.1 Welding Defects Freedom from cracks or other sharp discontinuities is of primary concern in weld quality and performance. Crack-like defects can degrade component lifetime by eliminating the initiation stage of phenomena such as fatigue or stress corrosion. Similarly, other flaws (e.g., lack of fusion defects, gas porosity, and inclusions) can act to magnify global stresses, pro- duce locally aggressive environments via their occluded geometry or composition, and initiate cracking. In order to mitigate these flaws, it is critical to differentiate between defect types. The first step in the prevention of cracking is understanding the temperature range over which the cracking occurs. The primary measure is to determine whether defects are ‘hot cracks’ or ‘cold cracks,’ that is, whether they form above or below the solidus temperature of the alloy. Secondly, the loca- tion of the crack in the weld (composite region, unmixed zone, partially melted zone, heat-affected zone) and in the microstructure (solidification boundary, crystallographic boundary, etc.) must be determined.1 Once these distinctions are made, 273 274 Welds for Nuclear Systems strategies to eliminate cracking can be developed via changes to the welding process, weld parameters, filler metal, joint design, fixturing, and/or postweld heat treatment. The supersolidus/subsolidus distinction, com- bined with the unifying concept of homologous tem- perature, illustrates the commonality of welding defects and degradation mechanisms across alloy sys- tems as shown in Figure 1. Supersolidus ‘hot crack- ing’ defects include solidification cracking, liquation cracking, and hot tearing. Subsolidus ‘cold cracking’ includes precipitation, transformation, and segregation- induced cracking (SIC) mechanisms. 4.09.1.1 Supersolidus Cracking Cracks that occur between the liquidus and the soli- dus are commonly termed ‘hot cracks.’ Hot cracks can be further differentiated depending on whether they occur in the composite region of the weld bead on cooling (solidification-type) or in the partially melted or heat-affected zone where a composition gradient or low melting phase acts to locally depress the solidus of the alloy (liquation-type). Additionally, a third type of ‘hot cracking,’ that is, hot tearing, can be Homologous temperature (T/Tmelt) Hydrogen embrittlement (low temp. crack propagation, hydriding) Liquid and solid metal embrittlement 0.1 0.2 0.3 0.4 0.5 0.6 Radiation hardening Impurity segregation v Ordering reactions/brittle Subsolidus Segregation-induced cracking: hydrogen Radiation-induced segregation Precipitation-induced crac ductility dip cracking, strain- cracking, subsolid Weld cracking Environmental deg Figure 1 Comparison of the typical temperature ranges for th environmental degradation common to nuclear power systems ( the homologous temperature of the alloy under consideration. distinguished from solidification and liquation cracks. Hot tears are primarily mechanical in nature, driven by the geometry and stresses in the weldment. A schematic of the locations and representative micrographs of the different types of supersolidus cracking are shown in Figure 2. 4.09.1.1.1 Solidification cracking Solidification cracks occur in the mushy zone of a weld bead on cooling, as the strains that develop exceed the ductility of the (solid þ liquid) mixture. The modern theory of solidification cracking was developed by Borland,2 who highlighted the impor- tance of the quantity and distribution of the liquid near the terminal phase of solidification, as well as the stresses that act on that liquid. The primary factors that affect hot cracking are summarized in texts by Kou, Messler, and others.3–6 These factors are listed below: 1. The solidification temperature range : The larger the solidification temperature range, the more exten- sive the solid þ liquid mushy zone, which is sus- ceptible to cracking. While large solidification temperature ranges may promote crack healing via Creep-rupture Supersolidus 0.7 0.8 0.9 ia diffusion second phases precipitate 1.0 king: also known as age cracking, reheat us cracking radation Liquation -type Solidification -type ‘Hot tearing’ (mechanical) e different types of weld cracking (top) and forms of bottom). All temperature ranges are approximate, based on Liquation WeldHAZ 500μm Notches Multipass Gas Tungsten Arc Weld Solidification 200μm Tearing 500μm Figure 2 Illustration of the different forms of hot cracking. Liquation cracks occur in the partially melted and/or heat-affected zone (HAZ) of the weld bead being deposited (top). Solidification-type cracks occur in the composite region of the weld during solidification (middle), and hot tears are dominated by mechanical forces and occur at macroscopic notches (bottom). Note that the cracks are all from nickel–chromium alloy welds but are not all from the same weldment. Welds for Nuclear Systems 275 backfilling, solute-rich ‘backfill’ may have degraded properties relative to the bulk weld metal. The approximate solidification temperature ranges of several alloys used in nuclear construction are shown in Figure 3. 2. The solidification path : Solidification crack suscep- tibility is markedly influenced by the type and distribution of solid phases, for example, initial d-ferrite formation from the liquid in austenitic stainless steel welds imparts hot crack resistance by breaking up the solidification structure and by scavenging tramp elements (e.g., sulfur and phosphorous). Conversely, eutectic-type reac- tions during terminal solidification (e.g., liquid ! g þ Laves in nickel-based alloys) are notably detrimental to solidification cracking resistance.7 Figure 4 illustrates the calculated solidification path and hot cracking resistance of two nickel– chromium filler metals. The more solute-rich filler metal that forms Ni2(Nb, Mo)-type Laves phase is 92 87 82 157 117 262 102 102 247 187 232 677 1000 1100 1200 1300 1400 1500 1600 1700 1800 1900 A5 33 B ER 10 0S -1 31 6L 30 8L A6 00 EN 82 A6 90 EN 52 EN 52 M A6 25 Zr -2 .5 Nb Zr -4 C al cu la te d s ol id ifi ca tio n T ra ng e (� C ) Note: Scheil calculations are approximate and overpredict the actual solidification temperature range Figure 3 Comparison of the calculated solidification temperature ranges of some common materials used in nuclear power systems (JMatPro, Version 4.1). For a given alloy class, hot cracking is promoted by larger solidification temperature range and low solidus temperature. Note that the compositions used are ‘typical’ values and significant variability exists within each alloy’s specification range. 276 Welds for Nuclear Systems much more prone to hot cracking than the Laves- free alloy. 3. The surface tension of the terminal liquid : Low surface tension liquids wet solidification boundaries and promote cracking by increasing the amount of interface incapable of supporting appreciable ten- sile strains. 4. The metallurgical structure of the weld: Large colum- nar dendritic grains are more susceptible to solid- ification cracking than finer equiaxed structures. Coarse solidification structures result in longer crack paths and less grain boundary area to dis- tribute elements that lower the solidus and/or embrittle the boundary. Columnar grains may exacerbate hot cracking by promoting wetting of the grain faces and can result in linear solidifica- tion boundaries near the centerline of the weld bead where tensile stresses are often the highest. 5. The mechanical forces that act on the weld: High tensile strains during the terminal stages of solidification promote cracking. Given the complexity of the factors that contribute to solidification cracking, it is difficult to predict its occurrence in production welds. However, weld- ability tests such as the transvarestraint test enable a quantitative ranking of alloys with respect to solidification cracking susceptibility and offer a standardized methodology to optimize welding para- meters.8–10 Results from transvarestraint tests on sev- eral corrosion-resistant alloys are shown in Figure 5, which compares the maximum crack distance (i.e., the extent of the mushy zone when the crack distance becomes insensitive to the applied strain) and hence the intrinsic susceptibility of the alloy to solidifica- tion cracking. A representative transvarestraint sam- ple is shown in Figure 6, which illustrates the locations of solidification- and liquation-type cracks. Note that solid-state cracks can also be produced in this test.10 In general, alloying additions that are rejected into the liquid (i.e., whose equilibrium segregation coeffi- cient, k, is 1050 1100 1150 1200 1250 1300 1350 1400 1450 0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1 1.1 Fraction solid Te m p er at ur e (º C ) γ MN MC γ MN MC σ Laves δ Ni–Cr–Fe filler metal Ni–Cr–Fe–Mo–Nb filler metal 20μm (a) (b) (c) Mo L 100 mils 5 mm 5 mm 5 mmGrey Nb L0 48 115 182 0 57 100 mils Figure 4 Illustration of the effect of solidification path on cracking resistance. (a) Two Ni–30Cr filler metals show markedly different hot crack susceptibilities. (b) Scheil modeling predicts that the more solute-rich alloy has a larger solidification temperature range and can form s, Laves and d in the terminal solid and (c) SEM investigation confirms Nb, Mo-rich Laves near solidification cracks. Welds for Nuclear Systems 277 coefficients EN52MSS>EN52i>EN82H > 68HP, which is a direct result of the decreasing alloying content of the strongmelting point depressants molybdenum and niobium.7,10–15 Applied strain (%) 0 1 2 3 4 5 6 7 8 9 10 11 12 13 14 M ax im um c ra ck d is ta nc e (m m ) 0.00 0.50 1.00 1.50 2.00 2.50 3.00 Different symbols of the same color denote different heats Alloy 625 EN52MSS EN52i EN82H SANICRO 68HP 347 SS 308L SS Transvarestraint test data GTAW weld buildups 180 amps, 12.7 volts, 5 ipm ~45.7 kJ in.–1 Figure 5 Comparison of transvarestraint test data for several low-strength corrosion-resistant alloys used in nuclear power systems. Note that for a given alloy class, increasing solute content generally increases the susceptibility to solidification cracking as shown by comparing the plateau portions of the curves. 278 Welds for Nuclear Systems 4.09.1.1.2 Liquation cracking In contrast to solidification cracks, liquation cracks occur in the partially melted zone and the heat- affected zone and can be either interdendritic or intergranular in nature. An example of an interden- dritic liquation crack in a nickel–chromium alloy is given in Figure 2 (top), while intergranular liquation cracks in a pressure vessel steel are shown in Figure 8. The liquation cracks in Figure 8 are caused by the presence of sulfur-rich inclusions that liquate in the partially melted and heat-affected zones of the weld. Another variation of liquation-type cracking can occur via the partial dissolution of second-phase particles, that is, the constitutional liquation mecha- nism proposed by Savage and shown experimentally by Pepe and Savage.1,16,17 In this type of cracking, the heat from welding partially solutionizes the second- phase particles in the heat-affected zone. The result- ing concentration gradient around the particle lowers the solidus (e.g., the effect of niobium on nickel-based alloys from NbC or Ni2Nb) locally. 4.09.1.1.3 Hot tearing ‘Hot cracking’ can also be primarily mechanical in nature; the restraint, constraint, and geometry of the weld act to pull apart the weld metal at temperatures near the solidus. This type of cracking may be trans- granular or interdendritic and is favored by mechan- ical notches and partial penetration weld joints. Note that the hot tear shown in Figure 2 (bottom) was the only crack present in that multipass weld, illustrating the dominant effect of the notch at the weld root. Liquid at the time of straining Solid at the time of straining HAZ/PMZ Solidification cracks Mushy zone at the time of straining 2.5 mm Welding direction 500 mm Liquation cracks Autogenous weld bead Figure 6 Illustration of solidification- versus liquation-type cracking in a transvarestraint sample of Alloy 625 tested at 10% strain. Solidification cracks (bottom right) form on-cooling in the mushy zone behind the solid/liquid interface. Liquation cracks (upper left) are in the partially melted zone (PMZ) and/or heat-affected zone adjacent to the autogenous weld bead. Welds for Nuclear Systems 279 4.09.1.2 Subsolidus Cracking 4.09.1.2.1 Precipitation-induced cracking Solid-state cracks in welds often occur near the time/ temperature regime of a phase transformation in which the local stress or strain produced from the phase transformation interacts with global stresses in the weldment and results in cracking. This basic phenomenon has several different names based on the alloy system it occurs in and includes ‘ductility dip’ cracking in low-strength nickel-based alloys and stainless steels,10,18 ‘strain-age’ cracking in precipita- tion hardenable nickel- and iron-based alloys,19–21 ‘reheat cracking’ in 2¼Cr–1Mo-type steels,22 and ‘subsolidus cracking’ in titanium alloys.23 Ductility dip cracking has been studied in detail by Young and Capobianco, who provide a good example of how this phenomenon occurs.18 The cracking derives its name from the corresponding MLTS-2 (27Cr)Alloy 625 EN82H 347 SS68HP 308L SS 5 mm Figure 7 Comparison of the grain structures and the solidification cracks produced in the transvarestraint test at constant heat input and 5% strain. Note the long cracks in the solute-rich nickel-based alloys that are relatively susceptible to solidification cracking. 280 Welds for Nuclear Systems loss of tensile ductility in the homologous tempera- ture range (�0.4–0.9Tm) that corresponds to the time/temperature regime of the precipitation of a partially or fully coherent second phase. In low- strength nickel–chromium alloys, the ductility dip occurs during on-cooling from a peak temperature high enough to solutionize existing carbides and cause intergranular precipitation of the detrimental phase (M23C6 carbides, in this case). The relationship between the precipitation kinet- ics of the detrimental phase and the macroscopic tensile ductility is shown in Figure 9, which com- pares a calculated TTT plot for M23C6 precipitation in a Ni–29Cr–9Fe–0.01C (wt%) alloy (i.e., an analog to EN52/Alloy 690), with experimental on-cooling tensile ductility data for the alloy.10 As shown, if very rapid cooling suppresses precipitation, there is no ductility loss (region 1). The ductility minimum occurs near the nose of the precipitation curve when the local strain contribution from intergranular carbide precipitation is maximized (region 2). Duc- tility recovery occurs as precipitation progresses because local misfit strains decrease as chromium depletion occurs and as misfit dislocations are gener- ated (region 3). Ductility is restored when precipita- tion is complete (region 4). In Figure 10, the stages of ductility dip crack formation are outlined, in which (often in reheated weld metal of a multipass weld or in the base metal heat-affected zone) (Cr,Fe)23C6 carbides preferen- tially nucleate during on-cooling on grain boundaries with partial, cube-on-cube coherency (Figure 10(a)). Due to misfit strains, tension develops between the carbides, producing intermittent microscopic cracking (Figure 10(b)). Upon the development of global stres- ses (e.g., from thermal strains on-cooling or applied during hot ductility testing), these cracks often link up and form the classic ‘ductilitydip’ crack (Figure 10(c)), that is, an intergranular crack that typically extends�1 grain in length. Compared to a solidification-type crack, the fracture surfaces of these solid-state cracks show less evidence of the underlying dendritic struc- ture and are littered with (Cr,Fe)23C6-type carbides. 24 Figure 10(d) illustrates how the misfit strain between the carbide andmatrix increaseswith increasing chro- mium concentration in the alloy. In part, this explains why 30 wt% alloys (A690 and EN52) are more sus- ceptible to this defect than their lower chromium counterparts (A600/E-182). The transient nature of ductility loss with time and temperature, which are important dependencies cannot be explained by other proposed mechanisms for this solid-state cracking.25–32 Specifically, in the Ni–Cr alloys of interest to nuclear systems, neither impurity segregation (at least at ‘typical’ levels of 3 4 2 S 1 3 21 HAZEB house 1c 1% nital etch10 mil Weld 9/21/104 10.0 keV 2.1 kX 10.0 mm 9/20/104 MnC 10.0 keV 2.1 kX 10.0 mm 9/20/104 10.0 keV 2.1 kX 10.0 mm 9/20/104 10.0 keV 2.1 kX 10.0 mm 4 F F F F Figure 8 Illustration of intergranular liquation-type cracks in a low-alloy steel. The cracks occurred in both the partially melted zone and heat-affected zone along prior austenite grain boundaries (top) and subsequent analysis of crack surfaces via Auger electron microscopy identified bands of MnS-type sulfide inclusions (bottom). Welds for Nuclear Systems 281 notion that solid-state cracking is caused by sliding, that is, the uniformity of slip with some strain accu- mulation (yellow and red areas) near the ductility dip cracks. If grain boundary sliding played a role in ductility dip cracking, there would be less strain contrast near the cracks, not more, as this technique is sensitive only to the diffraction pattern rotation produced by dislocations. The scenario of weld cracking occurring in the time/temperature regime of precipitation of a par- tially or fully coherent second phase is also well recognized as controlling strain-age or ‘reheat’-type cracking in g0/g00-strengthened alloys.10,18,20,21,33 While susceptibility to reheat cracking is often plot- ted as a function of the aluminum and titanium content of the alloy, a more fundamental correlation based on the transformation kinetics and precipitate/ matrix mismatch of the alloy is possible.18 As shown in Figure 12, alloys with high susceptibility to strain- age cracking display fast transformation kinetics and a large negative precipitate/matrix mismatch (i.e., tension develops between precipitates). Time (s) 0 5 10 15 20 25 30 35 Te m p er at ur e (� C ) 500 (a) (b) 600 700 800 900 1000 1100 Hold time at 870 �C prior to straining (s) R ed uc tio n in a re a (% ) 35 40 45 50 55 60 65 Finish of grain boundary M23C6 precipitation Start of grain boundary M23C6 precipitation Hot ductility of EN52 On cooling from 1330 �C Strained at 870 �C 1 1 2 2 3 3 4 4 M23C6 TTT kinetics Estimated from JMatPro 4.1 Ni–29Cr–9Fe–0.01C alloy 0 5 10 15 20 25 30 35 Figure 9 Comparison of on-cooling hot ductility tests on EN52 (b) with the estimated M23C6 time–temperature- transformation behavior (a). The local ductility minimum correlates with the onset of grain boundary M23C6 precipitation and ductility recovery with the completion of precipitation. 282 Welds for Nuclear Systems As transformation stresses are displaced to longer times and precipitation-induced stresses become compressive, weldability is improved. However, while weldability is increased by displacing precipitation- induced stresses to longer times, hot workability is often degraded for the same reasons. It is notable that some titanium alloys undergo a tensile ductility loss when tested near the b! a tra- nsformation.23,34,35 While the authors do not know of equivalent research on zirconium alloys, these alloys could also be susceptible to this form of precipitation- induced cracking (PIC). Mechanistically, this could be caused by the nucleation of a from the b-phase if the (110)bk(0001)a is significant, or from some other phase with partial coherency (e.g., HCP Laves on HCP-a). 4.09.1.2.2 Segregation-induced cracking The most technologically important manifestation of SIC is hydrogen-induced cracking (also known as ‘cold cracking’) caused by the numerous sources of hydrogen in both fabrication and service, the relative ease of hydrogen entry into metals, its high diffusiv- ity, and its ability to weaken metallic bonds or form brittle second phases.36–39 Hydrogen cracking in steels and hydride-type cracking of zirconium alloys –0.7% –0.6% –0.5% –0.4% –0.3% –0.2% –0.1% 0.0% 0.1% 3.550 3.555 3.560 3.565 3.570 3.575 3.580 Lattice parameter of alloy (Å) E st im at ed s tr ai n fr om M 23 C 6 p re ci p ita tio n Alloy 690 ~29 wt% Cr EN82H ~20 wt% Cr Alloy 600 ~15 wt% Cr Partially coherent grain boundary Cr23C6 Cracking between carbides ‘Ductility Dip’ Crack Tension between carbides Small cracks coalesce 0.1μm (Cr,Fe)23C6 μm μm Grain boundary 30 Cr alloy a≈3.58 Grain 1 Grain 2 Grain boundary (a) (b) Acc.V 20.0 kV 3.5 10000x SE 12.7 Spot Magn Det WD 2 mm (c) (d) 1μm 100 μm (Cr,Fe)23C6 Cr23C6 a≈10.66 Figure 10 Illustration of the mechanism of ductility dip cracking in Ni–Cr alloys. (a) Partially coherent (Cr,Fe)23C6 carbides form in reheated weld metal, which have misfit with the grain boundary. (b) Precipitation generates local grain boundary tensile stresses between carbides which (c) results in ductility dip cracking when sufficiently large global stresses are present during welding or are applied during tensile testing. (d) The increased misfit between the carbide and the matrix with increasing chromium concentration helps explain the susceptibility of alloy 690/EN52 and the resistance of A600/EN82 to DDC. Welds for Nuclear Systems 283 =10 mm; BC; Step = 0.1 mm; Grid274 x 207 =20 mm; BC; Step = 0.3333 mm; Grid350 x 372 (d) (b) (c) (a) DDC =10 mm; LocMis2; Step = 0.1 mm; Grid274 x 207 1 2 1 2 30 =20 mm; LocMis2; Step = 0.3333 mm; Grid350 x 372 Figure 11 Comparison of the local misorientation (left) and band contrast (right) images for EN52 strained to 5% ((a) and (b), respectively) and to 10% strain ((c) and (d)) during cooling. Note the generally uniform plasticity with some strain accumulation at the grain boundary. Calculated nose of the TTT curve (s) 0.1 1 10 100 1000 P re ci p ita te /m at rix la tt ic e m is m at ch (% ) – 0.8 – 0.6 – 0.4 – 0.2 0.0 0.2 0.4 0.6 0.8 IN 100 Astroloy Udimet 700 Rene 41 Waspaloy X-750 718 Increasing resistance to strain-age cracking Tension between precipitates Compression between precipitates Figure 12 Correlation of the subsolidus cracking susceptibility of selected superalloys with the misfit and kinetics of second-phase precipitation (g0 or g0 0). Adapted from Young, G. A., et al. Welding J. Res. Suppl. 2008, 31S–43S; Prager, M.; Shira, C. S. Welding Res. Council Bull. 1968, 128; calculations done with JMatPro, Version 4.1. 284 Welds for Nuclear Systems Welds for Nuclear Systems 285 have been treated recently in the literature and more extensive reviews are found elsewhere.40,41 However, it should be highlighted that low- strength austenitic alloys are resistant but not immune to hydrogen-induced cracking. Figure 13 shows a hydrogen-induced crack in a Ni–20Cr–3Mn– 2.5Nb–1Fe weld metal (EN82) that was produced by the combination of poor welding practice and the use of hydrogen-bearing shield gas. The 95%Ar–5%H2 shield gas helps minimize surface oxides and interpass grinding but results in �12wt ppm hydrogen dis- solved in the filler metal. ‘Refuse welding’ or remelt- ing beads in an attempt to improve the tie-in and contour increases the plastic strain in the joint and can trigger cold cracking.42,43 4.09.1.3 Other Welding Defects In addition to cracking, defects such as lack of fusion between weld beads or the weld bead and the side- wall, variable penetration, or second-phase inclusions can degrade weld quality. Lack of fusion defects is a notable concern when welding high-alloy nickel- based materials, which have notably ‘sluggish’ weld pools and are difficult to wet and tie into adjacent material (Figure 14(a)). Inclusion-type defects are another concern and can be grouped into at least two types: (1) those that result from alloying additions or slag and (2) those that form via reaction with the environment. An example of the first type is given in Figure 14(b), which shows an unmelted iron–niobium Laves phase that is an intentional alloying addition to the flux coating of a shielded metal arc electrode. While alloying in this manner is a cost-effective way to tailor the composition of the electrode, it can lead to brittle second phases that also affect the local composition. In this case, the Nb-rich Laves phase is a strong melting point suppressant which can lead to either solidification- or liquation-type cracking. The second type of inclusion is generally oxide or nitride-type particles that form via reaction with air. The corrosion-resistant alloys used in nuclear power systems (i.e., Fe-based stainless and Ni-based alloys) are especially prone to oxide-type defects as the nature of their corrosion resistance depends on the formation of stable, tenacious oxide films. An example of an aluminum–titanium-rich eutectic- type oxide that formed in a poorly shielded Alloy 690 fusion weld is shown in Figure 14(c). Another consideration of this oxide formation is that whatever metallurgical effect these alloying elements have is lost if they oxidize prior to solidification (e.g., the grain nucleating effect of Ti(C,N)-type particles). Recent research shows that control of oxygen is critical to the weld puddle flow and wetting in nickel- based filler metals.10 In practice, this often translates into careful wire drawing practice so as to minimize the extent of embedded oxides or wire drawing lubri- cants into the filler metal. Figure 15 shows variability in the bead contour and tie-in of two filler metals welded under identical conditions, which was later traced to wire cleanliness. Additionally, separate test- ing shows that�100 wt ppm levels of oxides can have macroscopic detriment on the regular flow and con- tour of Ni–30Cr-type filler metals.10 4.09.2 Stresses and Strains in Welds Fusion welding leads to residual stresses and strains (distortion) via thermally induced stresses, solidifica- tion shrinkage, and phase transformations. Thermal stresses arise from the large temperature gradients inherent to fusion welding and from differences in the coefficients of thermal expansion (CTE) between materials that make up the weldment (Table 1). Thermal contraction generates stress and distortion during on-cooling, with the maximum residual stress often being the flow stress at which the lowest tem- perature distortion occurs.5,45 Dissimilar metal welds are regions of special con- cern for nuclear power systems as the residual stres- ses are often higher than for similar metal welds.18 For example, in ‘safe end’-type welds, the CTE dif- ference between a pressure vessel low-alloy steel and a corrosion-resistant austenitic alloy leads to higher stresses in these welds and in fact, these locations are known to be at increased risk of stress corrosion cracking for this reason.46–48 Solidification shrinkage is a lesser effect, as the liquid cannot support appre- ciable stress but can affect the local bead contour, the deformation at the weld face, and lead to stress raisers (e.g., concave beads or cracks).5 Phase transformations can also have appreciable effects on the sign, magnitude, and distribution of residual stresses in welds. Becker et al. have shown that for accurate prediction of the residual stresses in pressure vessel-type steels, it is critical to account for the on-cooling phase transformations that occur from welding.49 Furthermore, phase transformations dur- ing postweld heat treatments and from service expo- sure must also be considered. For example, nickel alloys can be susceptible to the development of 524´D 25.0 kV 100 mm AMRAY #0000* 2,540´D 25.0 kV 10 mm AMRAY #0000* M -3, 8:1 P hos. =1500 mm; BC + GB; Step = 5 mm; Grid800 ´ 100 =1500 mm; E123; Step = 5 mm; Grid800 ´ 100 =1500 mm; BC + GB; Step = 5 mm; Grid800 ´ 100 =1500 mm; E123; Step = 5 mm; Grid800 ´ 100 =1500 mm; BC + GB; Step = 5 mm; Grid800 ´ 100 =1500 mm; E123; Step = 5 mm; Grid800 ´ 100 0.5 cm Figure 13 Example of a hydrogen crack produced in EN82H from the use of 95%Ar–5%H2 shielding gas and abusive welding practice (refuse welding). The top figures show the crack in cross-section and the corresponding electron backscatter diffraction strain maps. The bottom fractographs show the intergranular/interdendritic nature of the hydrogen crack. 286 Welds for Nuclear Systems 50μm (a) (b) (c) 10 mil JP1-100 Acc.V Spot Magn Det WD Exp 20.0 kV 5.0 500x BSE 6001 Acc.V Spot Magn Det WD Acc.V Spot Magn Det WD Exp Exp 20.0 kV 4.0 1200x BSE 4.8 6001 Acc.V Spot Magn Det WD Exp 20.0 kV 5.0 150x SE 5.2 6001 20 mm 20 mm 500μm 50μm 50μm 200μm 50μm 4.0 Figure 14 Illustration of some welding defects in commercially available Ni–30 wt% Cr alloys. (a) Lack of fusion defects in an SMAW (left) and a GTAW (right), (b) unmelted NbFe2 alloying addition in an SMAW (left) and a slag inclusion from an SMAW (right) and (c) surface (left) and internal (right) oxides from a laser weld. Welds for Nuclear Systems 287 short- and long-range order. The ordered structure typically has a smaller lattice parameter than the bulk alloy and can lead to increased residual stresses with service exposure.10,50–54 Another point to note is that the welds typically have considerable texture that can be a significant factor in both the macroscopic and microscopic (intergrain) stresses and strains in welds. 4.09.2.1 Quantification of Residual Stresses and Strains 4.09.2.1.1 Elastic stress Stresses in welds can be determined via several computational and experimental techniques. Compu- tational methods are generally based on finite element methods, while experimental techniques include X-ray and neutron diffraction, hole drilling, and sur- face deformation mapping (e.g., slitting). Details of the application of these techniques can be found in several research proceedings55–57 and recent books.58,59 4.09.2.1.2 Plastic strain The evolution of automated electron backscatter dif- fraction analysis has made the mapping and quantifi- cation of plastic strains in welds accessible via the scanning electron microscope.43,60–64 Strains can be visualized qualitatively via the intragrain misorienta- tion (Figures 11 and 19(c)) of the diffraction pattern or quantitatively (Figure 13) via the average intra- grain misorientation (i.e., the ‘AMIS’ parameter) of many grains and an appropriate calibration curve. Calibration curves from uniaxial tensile samples for several nickel-based alloys are given in Figure 16. For reference, the measured plastic strain in sev- eral different welds and a heat-affected zone are compared in Table 2. Appreciable plastic strains (2–4%) occur even in unconstrained bead-on-plate welds and a wide range of strains (�2% to almost 30%) may be found, depending on the precise weld geometry, constraint, and welding practice. An example of the effect of welding practice is Table 1 Comparison of some physical properties for elements of interest to nuclear power systems Metal or alloy Melting point (K) Thermal conductivity (W mK�1) Coefficient of thermal expansion, �293–373 K (10�6 K�1) Reference Fe 1809 80.4 11.8 Ni 1726 74.9 13.3 [44] Zr 2125 21.1 5.0 900.0X 50.0 mm Ca OC SEM 2 cm 10.0 keV 10.0 keV900.0X 50.0 mm F 11/28/101 F 11/28/101 10.0 keV 900.0X 50.0 mm F 11/28/101 10.0 keV 900.0X 50.0 mm F 11/28/101 Figure 15 An example of the effect of wire cleanliness on weld quality. The top picture shows two GTAWwelds made under identical conditions. The right hand weld displayed poor flow and tie-in and was later traced back to impurities (likely drawing lubricant) embedded in the weld wire. The bottom figure shows an SEM image of the wire surface and Auger maps of embedded calcium, carbon, and oxygen contamination. 288 Welds for Nuclear Systems given in Figure 17 for a 2 in. thick, Alloy 690 narrow groove weld made with EN82H filler metal via automatic gas tungsten arc welding (A-GTAW).42,43 If welded with no ‘repairs’ (i.e., autogenous remelting of beads to improve bead- to-bead tie-in), it shows �5.5% plastic strain near the weld root. This plastic strain increases if the beads above the weld are remelted as shown in the y = 4.9776x – 2.8443 R2= 1 (EN82H - annealed) y = 2.8332x2- 3.9127x + 6.3447 R2= 0.9844 (A690) y = 1.9331x2- 1.498x + 1.3342 R2= 0.9482 (HP nickel - annealed) y = 0.5754x2+ 3.3861x - 0.9728 R2= 0.9949 (A600) y = 6.7568x - 2.1824 R2 = 1 (304L and 316L) 0 5 10 15 20 25 30 35 40 0 1 2 3 4 5 6 AMIS parameter (degrees) Te ns ile s tr ai n (% ) A600 (as-received plate) A690 (as-received plate) HP nickel (annealed) EN82H (annealed) 304L and 316L (GE data) Linear (304L and 316L (GE data)) Figure 16 Tensile data used to calibrate the ‘AMIS’ parameter for several austenitic alloys. The ‘GE Data’ are from Angeliu, T. In Tenth International Conference on Environmental Degradation of Materials in Nuclear Power Systems; NACE: Lake Tahoe, NV, 2000. Table 2 Comparison of the experimentally measured ‘AMIS’ parameter and the calculated plastic strain for several nickel-alloy welds and a heat-affected zone Weld Range of AMIS measured (degrees) Approximate plastic strain in weld or HAZ (%) EN82H, unconstrained bead-on-plate, A-GTAW 0.9–1.3 2–4 A600/EN82H pipe weld, near root, M-GTAW 1.1–1.9 3–8 A690/EN82H narrow groove weld near root, A-GTAW Best practice 1.0–3.2 2–13 A690/EN82H narrow groove weld, A-GTAW Abusive weld practice 3.0–6.2 12–28 A600 HAZ (unconstrained, E-182 SMAW) 1.0–1.5 3–6 Welds for Nuclear Systems 289 graph with strains of �11.5%, 15.0%, and 16.5% with 1, 2, and 3 simulated ‘repairs’ above the weld. As expected, high levels of plastic strain lead to increased yield strength, decreased ductility, and increased susceptibility to stress corrosion cracking. 4.09.3 In-Service Performance 4.09.3.1 Environmentally Assisted Cracking Welds and their heat-affected zones have long been known to be the areas of concern for environmentally assisted cracking (EAC) because of their propensity for as-fabricated flaws, high residual stresses, ele- vated plastic strains, chemical heterogeneity, and microstructural differences relative to base metals (Chapter 5.04, Corrosion and Stress Corrosion Cracking of Ni-Base Alloys; Chapter 5.05, Corro- sion and Stress Corrosion Cracking of Austenitic Stainless Steels; Chapter 5.06, Corrosion and Environmentally-Assisted Cracking of Carbon and Low-Alloy Steels; Chapter 5.02, Water Chem- istry Control in LWRs; and Chapter 5.08, Irradia- tion Assisted Stress Corrosion Cracking). Common EAC concerns in the nuclear industry include corro- sion fatigue of low-alloy steels, hydride-induced 0 2 4 6 8 10 12 14 16 18 Best practice P la st ic s tr ai n ne ar w el d r oo t (% ) 2� thick, A690 Narrow groove weld, A-GTAW Weld filler metal is EN82H. ‘Repairs’ are autogenous remelting of welds beads. 1 repair 2 repairs 3 repairs Figure 17 Illustration of the effect of simulated repairs (i.e., autogenous remelting of beads) on the plastic strain developed in a 2 in. thick Alloy 690/EN82H narrow groove weld. 290 Welds for Nuclear Systems cracking of zirconium alloys, and stress corrosion cracking of corrosion-resistant structural alloys. Specifically, stress corrosion of austenitic stainless steel65–68 and nickel-alloy welds and their heat- affected zones has been a topic of considerable research.9,66,69–72 In austenitic stainless steels, sensiti- zation- and welding-induced plastic strains are the key factors in stress corrosion resistance. In nickel-alloy welds, the bulk chromium con- centration, the solidification segregation, and the as-fabricated plastic strain are critical factors for understanding their stress corrosion performance. The stress corrosion cracking growth rate of nickel- alloy filler metals in high-temperature, high-purity water is shown as a function of their bulk chromium concentration in Figure 18(a). Note the strong decrease in stress corrosion crack growth rate at bulk chromium levels near 22 wt%. This decrease is likely associated with a change in crack tip oxide from NiO-type to a more stable spinel (NiCr2O4) or corundum (Cr2O3) structure. 73–75 However, the bulk chromium concentration does not explain the extreme resistance of EN52 weld metal compared to other�30 wt% bulk chromium alloys (Figure 18(a)). Consideration of how solidification segregation affects the grain boundary chromium concentration is critical to understanding the stress corrosion resis- tance of the high-alloy weld metals.76 Specifically, niobium- and molybdenum-bearing alloys tend to deplete the solidification grain boundaries in chromium, while the Nb- and Mo-free EN52 grain boundaries are enriched in chromium as shown in the graph in Figure 18(b). In Alloy 600 heat-affected zones, the increased susceptibility to SCC in high-temperature deaerated water is due, in large part, to the lack of intergranular chromium carbides.77,78 Figure 19(a) shows a cross- section of a stress corrosion crack grown in an Alloy 600 heat-affected zone and the flat grain boundary topography (GBT) in the HAZ, which is an indica- tion of a low degree of intergranular chromium car- bide precipitation.78 Additionally, Figure 19 shows the different chromium concentration profiles in the HAZ and base metal (Figure 19(b)), the increased strain in the weld and HAZ relative to the base metal (Figure 19(c)), and the transmission electron micro- graphs of the grain boundaries in the HAZ (showing sparse (M7C3- and M23C6-type carbides) versus the large continuous Cr7C3 carbides in the unaffected base metal (Figure 19(d)). The diffraction patterns in Figure 19(d) identify the M23C6 (left) and M7C3 (right) carbides. Stress corrosion crack growth rate predictions for Alloy 600 heat-affected zones are shown in Figure 20, which illustrates the strong temperature dependence as well as the effects of the applied stress intensity factor and the electrochemical potential. Figure 20 is based on eqn [1], which describes the crack growth rate of A600-type alloys exposed to high-temperature, high-purity water,77 in which A0, n, m, b, x0, and c are the Bulk chromium (wt%) 10 15 20 25 30 35 10 15 20 25 30 35 S C C r at e (m ils d ay –1 ) 100 10–1 10–2 10–3 10–4 10–13 10–12 10–11 10–10 10–9 10–8 102 101 S C C r at e (m ils d ay –1 ) 100 10–1 10–2 10–3 10–4 102 101 S C C r at e (m s– 1 ) 10–13 10–12 10–11 10–10 10–9 10–8 S C C r at e (m s– 1 ) EN82H* EN82H* EN52 100X Grain boundary chromium (wt%) EN52 100X E-182* E-182* EN52i : SCC benefit with weldability of EN82H 360 �C, Initial KI= 50 MPa m ½ 20 SCC H2 per kg H2O Adjusted to 100% engagement 360 �C, Initial KI= 50 MPa m ½ 20 SCC H2 per kg H2O Adjusted to 100% engagement EN52i : SCC benefit with weldability of EN82H Figure 18 (a) The effect of chromium concentration on the stress corrosion crack growth rates of nickel–chromium weld metals. While a dramatic improvement in SCC resistance is seen near 22 wt% bulk chromium, high chromium (�30 wt%) alloys appear to show variable resistance. Understanding how solidification segregation affects the grain boundary chromium level (b) is key to understanding the different stress corrosion resistance of these alloys. E-182 and EN82H SCC rates are from model predictions, E-182 grain boundary chromium concentration estimated. Welds for Nuclear Systems 291 experimentally determined constants, KI is the stress intensity factor, sYS is the yield strength, DEcP is the electrochemical potential relative to the Ni/NiO phase transition, Q Effective is the apparent activation energy, R is the gas constant, and T, the temperature. The appropriate parameters for Alloy 600 HAZs are given in Table 3. _a ¼ A0 � K nI � smYS � 1þ b � exp �0:5 � ð�EcPNi=NiO � x0Þ c � �2" #)( � exp �Q Effective R � T � � ½1� 4.09.3.2 Microchemical Changes In addition to solidification segregation, the local com- position of welds can change in-service via thermal exposure and via radiation-induced transmutation and radiation-induced segregation (RIS). Thermally induced embrittlement is most notable in low-alloy steels that are high in tramp elements, in locations where welding produces local compositional enrich- ment, and in higher nickel grades (e.g., A508 Gr4N), which are intrinsically more susceptible because of the cosegregation of nickel and phosphorous.79–82 RIS is a phenomenon in which irradiation-created defects cause spatial redistribution of alloying ele- ments as they diffuse to, and get trapped in, sinks Distance from grain boundary (μm) –2 0 1 2 C hr om iu m c on ce nt ra tio n (w t% ) 5 6 7 8 9 10 11 12 13 14 15 16 17 A600 Base Metal A600 Heat Affected Zone Far from weld EN82H GTA weld Misorientation (°) A600 HAZ Low strain High strain Hardness indents HAZ Base Flat GBT End of precrack SCC 0.5 mm(a) (b) (c) (d) 01 1378 6 018000 01 1352 6 010500 EN82H Weld A600 heat affected zone –1 1 mm 1 mm Figure 19 Comparison of Alloy 600 heat-affected zone and base metal structure, chemistry, and strain: (a) cross-section of stress corrosion sample showing the location of the cracking, (b) grain boundary chromium profiles for base metal (blue) and the HAZ (red), (c) qualitative strain map for the base metal, HAZ/weld interface, and (d) typical grain boundary microstructure for the HAZ (sparse M23C6 and M7C3) and base metal (extensive M7C3). 292 Welds for Nuclear Systems (e.g., voids and grain boundaries) (Chapter 1.18, Radiation-Induced Segregation). RIS can occur when a point defect flux interacts preferentially with a certain elemental species in the alloy, causing that element to be enriched or depleted near the defect sinks. This preference can be driven kineti- cally (migration barriers) and/or thermodynamically (binding/ordering). RIS is segregation that occurs Table 3 Fitted parameters for the Alloy 600 heat-affected zone stress corrosion crack growth rate data presented in Figure 20 ln (A0 ) n m B x0 (mV) c (mV) Q (kJ mol �1) Best estimate 22.607 0.869 0a 3.604 �15.61 42.79 136.0 95% confidence �3.729 �0.349 – – �20.71 �19.25 �18.2 aInsufficient data were available to determine this dependence. 360 �C 305 �C 250 �C 1.0E – 12 1.0E – 11 1.0E – 10 1.0E – 9 1.0E – 8 10 20 30 40 50 60 70 80 –100–75 –50–25 025 5075 P re d ic te d c ra ck g ro w th r at e (m s– 1 ) ΔEcP (mV) A600 HAZ K I (M Pa m ½ ) Figure 20 Illustration of the predicted crack growth rates for Alloy 600 HAZ material as a function of potential, stress intensity factor, and temperature. Welds for Nuclear Systems 293 in addition to thermal segregation. Like thermal seg- regation, this local change in composition can result in detrimental changes to mechanical and corrosion properties.83 RIS has been a concern in the nuclear industry for over 30 years and is considered one of the many factors that lead to irradiation-assisted stress corrosion cracking.6–9 This phenomenon was first predicted by Anthony84 and has been observed in a number of dif- ferent alloys and steels used in nuclear reactors.85–87 In both the iron- and nickel-based face-centered cubic Fe–Ni–Cr alloys, experimental RIS trends at grain boundaries are generally chromium depletion, nickel enrichment, and possible compensation through iron enrichment or depletion.85,86 In body-centered cubic ferritic–martensitic steels, both chromium enrichment and depletion have been reported.88,89 RIS in welds has not been extensively researched but the possibilityof grain boundary depletion of chromium in corrosion-resistant alloys that would act in addition to the chromium depletion that occurs during solidification segregation (e.g., Ni–Cr–Nb and Ni–Cr–Mo alloys) is of significant concern. 4.09.3.3 Microstructural Changes In addition to the relatively short-timemicrostructural changes that can occur on-cooling or with postweld heat treatment (typically 294 Welds for Nuclear Systems formation of long-range order can lead to increased residual stresses and decreased resistance to EAC,most notably to hydrogen embrittlement.50,52,53,90 4.09.4 Weldability of Specific Alloy Systems 4.09.4.1 Low-Alloy Steels Low-alloy steels generally have good weldability with the main concern being liquation cracking near impurities (Figure 8), the propensity of some grades (esp. Cr–Mo steels and some pressure vessel grades) to reheat-type cracking (discussed earlier) and to hydrogen-induced cracking.1,91,92 The suscep- tibility to hydrogen-induced cracking is controlled by four major considerations: 1. The composition of the steel 2. The mobile hydrogen concentration 3. The stresses in the weldment 4. The thermal management of the weld In general, the more hardenable the steel (i.e., the more easily martensite is formed), the more suscep- tible it is to cracking. Since hardenability generally increases with carbon content and alloying additions, several parameters have been developed to gauge susceptibility to hydrogen-induced cracking, includ- ing the carbon equivalent (Ceq) in eqn [2] and the Ito- Bessyo ‘cold cracking’ parameter (Pcm) (eqn [3]). 93 The concentrations in eqns [2] and [3] are for weight percentage. Ceq ¼ C þ CrþMoþ V 5 þMn 6 þNiþ Cu 15 ½2� Pcm ¼ 5 B þ C þ V 10 þMo 15 þMnþ Crþ Cu 20 þ Si 30 þNi 60 ½3� The hydrogen concentration in the weldment can be minimized by proper cleanliness, preheat, and post- weld heat treatment.92,94 Additionally, microstruc- tural hydrogen traps can provide significant benefit by preventing hydrogen redistribution to regions of high stress.37 As the mobile hydrogen concentration is most detrimental, significant work has gone into standards to accurately assess the amount of diffusible hydrogen in steels.37,92,94 Minimizing residual stres- ses and avoiding geometric stress concentrators (e.g., notches) in the weldment also impart resistance to hydrogen-induced cracking. Thermal management of the weldment (e.g., pre- heating, interpass temperature, bead tempering, and postweld heat treatment) is also critical to the mitiga- tion of hydrogen-induced cracking. Preheating of components and interpass temperature control act to outgas hydrogen or hydrogen-bearing compounds (e.g., water) and lower the cooling rate. Additionally, careful control of heat input can produce hydrogen- resistant microstructures (i.e., bead tempering). Postweld heat treatment can act to lower residual stresses, produce beneficial hydrogen traps, and remove dissolved hydrogen from the weld.3,4,95,96 Several sources provide guidelines to mitigate hydro- gen cracking in specific grades of steel.95–97 4.09.4.2 Austenitic Stainless Steels The main concerns in austenitic stainless steel weld- ing are solidification, liquation, and PIC. Addition- ally, avoiding grain boundary chromium depletion via carbide precipitation (i.e., sensitization or ‘knife- line’ attack) is critical to maintaining in-service corrosion resistance (Chapter 2.09, Properties of Austenitic Steels for Nuclear Reactor Applica- tions). A key factor in solidification cracking resis- tance is proper control of the weld chemistry to form delta ferrite in the initial solid.98–100 Delta ferrite formation breaks up long, linear solidification bound- aries and acts to scavenge tramp elements from the liquid and prevent their concentration at interden- dritic boundaries during terminal solidification.101,102 Delta ferrite levels are typically controlled to �5–10% by volume to impart solidification cracking resistance but retain the mechanical properties of a face-centered cubic alloy. Specifically, the body- centered cubic ferrite is susceptible to cleavage at low temperatures and to spinodal decomposition of the ferrite into iron-rich (a) and chromium-rich (a0) phases at intermediate temperatures. Control of delta ferrite levels for different alloys is given in handbooks and can be predicted via the experimen- tally based Schaeffler or Delong diagrams, or com- putationally, by multicomponent phase diagrams, as shown in Figure 21.103–107 As a practical example, it is often difficult to weld over fully austenitic metals (e.g., nickel-based alloys) with austenitic filler metals (e.g., 308), as the increased nickel (from dilution) promotes the primary austenite solidification mode. As with most fusion welds, liquation cracking can be controlled by minimizing solidification segrega- tion (e.g., faster cooling rates and finer, more equiaxed structures) and by using lower heat input. fcc_A1 bcc_A2 Liquid+fcc_A1 Liquid+bcc_A2 14 15 16 17 18 19 20 21 22 23 24 25 26 Liquid Primary austenite forms Primary ferrite forms 70Fe-XCr-(30-X)Ni Chromium concentration (wt%) 1500 1475 1350 1375 1400 1425 1450 Te m p er at ur e (�C ) fcc_A1+bcc_A2 Figure 21 Illustration of the effect of chromium concentration on the solidification behavior of model ‘austenitic’ stainless steels. Formation of primary (d) ferrite provides hot cracking resistance. This pseudo-binary phase diagram was generated with Pandat Version 8.1, after Lippold, J. C.; Savage, W. F. Welding J. 1979, 58, 362s–374s. Welds for Nuclear Systems 295 As discussed earlier, stainless steels can be susceptible to PIC via partially coherent M23C6 carbide precipita- tion. One distinction relative to nickel-based alloys is that cracking may more likely occur on heating or with longer in-service exposures, as carbonmay be in solution as-welded, and precipitation kinetics are usually slower than for high-chromium nickel-based alloys.108 4.09.4.3 Nickel-Based Alloys Like austenitic stainless steels, single-phase nickel- based alloys are susceptible to solidification and liquation cracking (Chapter 2.08, Nickel Alloys: Properties and Characteristics). Common melting point suppressants that must be controlled to avoid cracking include tramp elements such as sulfur and intentional alloying additions such as boron, zirconium, niobium, and molybdenum.7,10–15,18,109,110 Addition- ally, both low-strength and precipitation hardenable grades can be susceptible to PICmechanisms (ductility dip and strain-age cracking) as discussed previously. 4.09.4.4 Zirconium Alloys Zirconium alloys are readily weldable but as with all reactive metals, need special precautions to prevent pickup of interstitial elements such as oxygen, carbon, and nitrogen that can degrade both the mechanical properties and the corrosion performance of the weld (Chapter 2.07, Zirconium Alloys: Properties and Characteristics).92,111,112 Vacuum or inert gases (argon, helium, or Ar–He mixtures) can be used to shield zirconium, but again, care needs to be taken to ensure sufficient vacuum level or gas purity to pre- vent contamination.92,111 Zirconium alloys can be susceptible to both super- solidus and subsolidus (i.e., hydride-type) cracking. Supersolidus cracking is typically solidification-type and many common alloying elements and/or poten- tial contaminants promote susceptibility. For exam- ple, iron, nickel, chromium, and copper all stabilize low-temperature eutectic reactions, and small con- centrations can greatly increase the solidification temperature range. An example of this is shown in Figure 22, which compares the maximum crack length in transvarestraint tests for a Zr–Cr alloy, Zircaloy-4 (Zr-4), and Zr–2.5Nb welded under iden- tical conditions. As shown, the Zr–Cr alloy exhibits the most susceptibility to solidification cracking. Acknowledgments The Authors are indebted to the welders, technicians, specialists, and engineers of the Welding & Materials Process Development Unit at Knolls Atomic Power Applied strain (%) 0 1 2 3 4 5 6 7 M ax c ra ck le ng th (m m ) 0.0 0.5 1.0 1.5 2.0 2.5 Zr-4 Zr-1.0Cr-0.5Sn-0.5Fe Zr-2.5Nb Transvarestraint tests Wrought plate GTAW, 8 kJ in.–1 Figure 22 Comparison of transvarestraint test data for three zirconium alloys. The strong effects of Cr and Fe on the solidification temperature range likely explain the increased susceptibility of the Zr–Cr alloy relative to Zr-4 and Zr–2.5Nb. 296 Welds for Nuclear Systems Laboratory, whose dedication and expertise made this work possible. They are also grateful to Dr. David S. Knorr of General Electric for his important con- tributions to the manuscript. References 1. Savage, W. F. 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Burchell Oak Ridge National Laboratory, Oak Ridge, TN, USA Published by Elsevier Ltd. 4.10.1 Introduction 300 4.10.2 Nuclear Graphite Manufacture 302 4.10.3 Graphite-Moderated Reactors 303 4.10.4 Displacement Damage and Induced Structural and Dimensional Changes in Graphite 305 4.10.5 Neutron-Induced Property Changes 310 4.10.5.1 Wigner Energy 310 4.10.5.2 Mechanical and Physical Properties 311 4.10.6 Irradiation Creep 315 4.10.6.1 The Relevance of Creep to Reactor Design and Operation 315 4.10.6.2 The Irradiation-Induced Creep Mechanism (In-Crystal) 316 4.10.6.3 Review of Prior Creep Models 317 4.10.6.3.1 Linear viscoelastic creep model 317 4.10.6.3.2 The UK creep model 317 4.10.6.3.3 The Kennedy model 317 4.10.6.3.4 The Kelly and Burchell model 318 4.10.6.3.5 The M2 model 319 4.10.6.4 Deficiencies in Current Creep Models at High Neutron Doses 321 4.10.7 Outlook 323 References 323 Abbreviations AGR Advanced gas-cooled reactor ASTM American Society for Testing and Materials CEN Centre European Nuclear CP-1 Chicago Pile No. 1 CTE Coefficient of thermal expansion DWNTs Double-walled carbon nanotubes Esu Elastic strain unit HFR High flux reactor HOPG Highly oriented pyrolytic graphite HRTEM High-resolution transmission electron microscope IV Interstitial–vacancy MHTGR Modular high-temperature gas-cooled reactor NGNP Next Generation Nuclear Plant ORNL Oak Ridge National Laboratory PGA Pile grade A PKA Primary knock-on atom SKA Secondary knock-on atom STM Scanning tunneling microscope TEM Transmission electron microscope USSR Union of Soviet Socialist Republics Symbols a Constant in viscoelastic creep model a Crystallographic a-direction (within the basal planes) b Burger’s vector b Constant in viscoelastic creep model B Empirical fitting parameter, analogous to the steady-state creep coefficient c Crystallographic c-direction (perpendicular to the basal planes) c Flaw size C Specific heat Cp Specific heat at constant pressure d«c/dg Initial secondary (steady-state) creep rate E Elastic modulus E0 Initial (preirradiated) Young’s modulus 299 300 Radiation Effects in Graphite Ed Displacement energy for a carbon atom from its equilibrium lattice position Ep Young’s modulus after initial increase due to dislocation pinning Eg Young’s modulus at dose g Fx Pore generation term F 0x Pore generation term for a crept specimen G Fracture toughness or strain energy release rate gx Rate of change of dimensions in the x-direction with respect to neutron dose g 0x Rate of change of dimensions in the x-direction for a crept specimen with respect to neutron dose (dimensional change component) h Planck’s constant k Boltzmann’s constant k Steady-state creep coefficient k0(g) Modified steady-state creep coefficient k0 Initial secondary creep coefficient k1 Primary creep dose constant k2 Recoverable creep dose constant La Mean graphite crystal dimensions in the a-direction Lc Mean graphite crystal dimensions in the c-direction R Gas constant (8.314 J mol�1K) S(g) Structure factor (E/Ep) T Temperature W Oxidation change factor (1/Xa) (dXa/dg) Rate of change of crystallite dimensions perpendicular to the hexagonal axis (1/Xc) (dXc/dg) Rate of change of crystallite dimensions parallel to the hexagonal axis XT Crystal shape parameter Z Atomic number a Coefficient of thermal expansion aa Crystal coefficient of thermal expansion in the a-direction ac Crystal coefficient of thermal expansion in the c-direction ax Coefficient of thermal expansion in the x-direction áx Coefficient of thermal expansion of a crept specimen in the x-direction a(f) Crystal coefficient of thermal expansion at angle f to the c-direction «̇ Strain rate «c Creep strain «́c Apparent creep strain «d Dimensional change strain «e Elastic strain «p Primary creep strain «s Secondary creep strain «t Thermal strain «Total Total strain f Fast neutron flux g Fast neutron fluence l Empirical fitting parameter m Empirical constant (�0.75) n Dislocation velocity uD Debye temperature r Density r0 True density st Tensile strength s Tensile stress v Frequency of vibrational oscillations V Mobile dislocation density j Empirical fitting parameter 4.10.1 Introduction There are many graphite-moderated, power- producing, fission reactors operating worldwide today.1 The majority are in the United Kingdom (gas-cooled) and the countries of the former Soviet Union (water- cooled). In a nuclear fission reactor, the energy is derived when the fuel (a heavy element such as 92U 235) fissions or ‘splits’ apart according to the following reaction: 92U 235 þ0 n1 !92 U236� ! F1 þ F2 þ n þg� energy ½I� An impinging neutron usually initiates the fission reac- tion, and the reaction yields an average of 2.5 neutrons per fission. The fission fragments (F1 and F2 in eqn [I]) and the neutron possess kinetic energy, which can be degraded to heat and harnessed to drive a turbine- generator to produce electricity. The role of graphite in the fission reactor (in addition to providing mechani- cal support to the fuel) is to facilitate the nuclear chain reaction by moderation of the high-energy fission neu- tron. The fission fragments (eqn [I]) lose their kinetic energy as thermal energy to the uranium fuel mass in which fission occurred by successive collisions with the fuel atoms. The fission neutrons (n in eqn [I]) give up Radiation Effects in Graphite 301 their energy within the moderator via the process of elastic collision. The g-energy given up in the fission reaction (eqn [I]) is absorbed in the bulk of the reactor outside the fuel, that is, moderator, pressure vessel, and shielding. The longer a fission neutron dwells in the vicinity of a fuel atom during the fission process, the greater is its probability of being captured and thereby causing that fuel nucleus to undergo fission. Hence, it is desirable to slow the energetic fission neu- tron (E� 2MeV), referred to as a fast neutron, to lower thermal energies (�0.025 eV at room temperature), which corresponds to a velocity of 2.2� 1015 cm s�1. The process of thermalization or slowing down of the fission fast neutron is called ‘moderation,’ and the material in a thermal reactor (i.e., a reactor in which fission is caused by neutrons with thermal energies) that is responsible for slowing down the fast fission neutrons is referred to as the moderator. Good nuclear moderators should possess the following attributes: � do not react with neutrons (because if they are captured in the moderator the fission reaction cannot be sustained); � should efficiently thermalize (slowdown) neutrons with few (elastic) collisions in the moderator; � should be inexpensive; � compatible with other materials in the reactor core; � meet the core structural requirements; and ideally � do not undergo any damaging chemical or physi- cal changes when bombarded with neutrons. In the fast neutron thermalization process, the maxi- mum energy lost per collision occurs when the target nucleus has unit mass, and tends to zero for heavy 0 5 10 15 20 25 0.2 0.4 0.6 0.8 1 1.2 1.4 1.6 1.8 Mass number (M) S lo w in g- d ow n po w er (c m –1 ) Figure 1 Moderator Figures of merit for several candidate mod H2O (M ¼ 18); heavy water, D2O (M ¼ 20). target elements. Low atomic number (Z) is thus a prime requirement of a good moderator. The density (number of atoms per unit volume) of the moderator and the likelihood of a scattering collision taking place must also be accounted for. Frequently used ‘Figures of merit’ for assessing moderators are the ‘slowing down power’ and the ‘moderating ratio.’ Figure 1 shows these Figures of merit for several candidate moderator materials. The slowing down power accounts for the mean energy loss per collision, the number of atoms per unit volume, and the scatter- ing cross-section of the moderator. The tendency for a material to capture neutrons (the neutron capture cross-section) must also be considered. Thus, the sec- ond figure of merit, the moderation ratio, is the ratio of the slowing down power to the neutron absorption (capture) cross-section. Ideally the slowing down power is large, the neutron capture cross-section is small, and hence the moderating ratio is also large. Practically, the choices of moderating materials are limited to the few elements with atomic number Raw petroleum or pitch coke Calcined coke Blended particles Coal tar binder pitch Calcined at 1300 �C Crushed, ground, and blended Mixed Cooled Extruded, molded, or 302 Radiation Effects in Graphite the cost of separating the heavy hydrogen isotope is large. Beryllium and beryllium oxide are good mod- erators but are expensive, difficult to machine, and suffer toxicity problems. Finally, graphite (carbon) is an acceptable moderator. It offers a compromise between nuclear properties, utility as a core structural material, and cost. It also has the advantage of being able to operate at very high temperatures (in the absence of oxygen). Unfortunately, the properties of graphite are markedly altered by neutron irradiation and this has to be considered in the design of graphite reactor cores. Green artifact Baked artifact Graphite isostatically pressed Baked at 800–1000 �C Impregnated to densify (petroleum pitch) Rebaked and reimpregnated artifact Graphitized 2500–2800 �C Nuclear graphite Purified Figure 2 The major processing steps in the manufacture of nuclear graphite. 4.10.2 Nuclear Graphite Manufacture The invention of an electric furnace2 capable of reach- ing temperatures approaching 3000 �C by Acheson in 1895 facilitated the development of the process for the manufacture of artificial polygranular graphite. Detailed accounts of the manufacture of polygranular graphite may be found elsewhere.2–4 Figure 2 sum- marizes the major processing steps in the manufacture of nuclear graphite. Nuclear graphite consists of two phases: a filler material and a binder phase. The pre- dominant filler material, particularly in the United States, is a petroleum coke made by the delayed coking process. European nuclear graphites are typi- cally made from a coal-tar pitch-derived coke. In the United Kingdom, Gilsonite coke, derived from natu- rally occurring bitumen found in Utah, USA, has been used. Both coke types are used for nuclear graphite production in Japan. The coke is usually calcined (thermally processed) at �1300 �C prior to being crushed and blended. Typically, the binder phase is a coal-tar pitch. The binder plasticizes the filler coke particles so that they can be formed. Forming pro- cesses include extrusion, molding, vibrational mold- ing, and isostatic pressing. The binder phase is carbonized during the subsequent baking operation (800–1000 �C). Frequently, engineering graphites are pitch impregnated to densify the carbon artifact, fol- lowed by rebaking. Useful increases in density and strength are obtained with up to six impregnations, but two or three are more typical. The final stage of the manufacturing process is graphitization (2500–3000 �C) during which, in sim- plistic terms, carbon atoms in the baked material migrate to form the thermodynamically more stable graphite lattice. Nuclear graphites require high chemical purity to minimize neutron absorption. Moreover, certain elements catalyze the oxidation of graphite and must be reduced to an acceptable level. This is achieved by selecting very pure cokes, utilizing a high graphitization temperature (>2800 �C), or by including a halogen purification stage in the manufac- ture of the cokes or graphite. Recently, comprehensive consensus specifications5,6 were developed for nuclear graphites. The electronic hybridization of carbon atoms (1s2, 2s2, 2p2) allows several types of covalent bonded structure. In graphite, we observe sp2 hybridization in a planar network inwhich the carbon atom is bound to three equidistant nearest neighbors 120� apart in a given plane to form the hexagonal graphene structure. Covalent double bonds of both s-type and p-type are present, causing a shorter bond length than in the case of the tetrahedral bonding (s-type sp3 orbital hybri- dization only) observed in diamond. Thus, in its per- fect form, the crystal structure of graphite (Figure 3) consists of tightly bonded (covalent) sheets of carbon atoms in a hexagonal lattice network.7 The sheets are 0.670 nm 0.246 nm c A B A a Figure 3 The crystal structure of graphite showing the ABAB stacking sequence of graphene planes in which the carbon atoms have threefold coordination. Reproduced from Burchell, T. D. In Carbon Materials for Advanced Technologies; Burchell, T. D., Ed.; Elsevier Science: Oxford, 1999, with permission from Elsevier. Radiation Effects in Graphite 303 weakly bound with van der Waals type bonds in an ABAB stacking sequence with a separation of 0.335 nm. The crystals in manufactured polygranular graph- ite are less than perfect, with approximately one layer plane in every six constituting a stacking fault. The graphite crystals have two distinct dimensions, the crystallite size La measured parallel to the basal plane and the dimension Lc measured perpendicular to the basal planes. In a coke-based nuclear graphite, values of La� 80 nm and Lc� 60 nm are typical.8 A combination of crystal structure bond anisotropy and texture resulting from forming imparts aniso- tropic properties to the filler coke and the manufac- tured nuclear graphite. The coke particles become preferentially aligned during forming, either with their long axis parallel to the forming axis in the case of extrusion, or with their long axis perpendicu- lar to the forming axis in the case of molding or vibrational molding. Consequently, the graphite arti- facts are often attributed with-grain and against-grain properties as in the American Society for Testing and Materials (ASTM) specifications.5,6 The degree of isotropy in manufactured graphite can be controlled through the processing route. Factors such as the nature of the filler coke, its size and size distribution, and the forming method contribute to the degree of isotropy. Nuclear graphites are typically medium or fine grain graphites (filler coke size Table 1 Currently available nuclear grade graphites Grade Manufacturer Coke type Comments IG-430 Toyo Tanso Pitch coke Isostatically molded, candidate for high-dose regions of NGNP concepts IG-110 Toyo Tanso Petroleum coke Isostatically molded, candidate for high-dose regions of NGNP concepts NBG-10 SGL Pitch coke Extruded, candidate for high-dose regions of NGNP pebble bed concepts; PBMR core graphite NBG-17 SGL Pitch coke Vibrationally molded, candidate for high-dose regions of NGNP prismatic core concepts NBG-18 SGL Pitch coke Vibrationally molded, candidate for high-dose regions of NGNP pebble bed concepts; PBMR core graphite PCEA GrafTech International Petroleum coke Extruded, candidate for high-dose regions of NGNP prismatic core concepts PGX GrafTech International Petroleum coke Large blocks for permanent structure in a prismatic core 2020 Carbone of America Petroleum coke Isostatically molded, candidate for permanent structures in a prismatic core 2191 Carbone of America Petroleum (sponge) coke Isostatically molded, candidate for permanent structures in a prismatic core 304 Radiation Effects in Graphite and the entire block is discharged from the reactor when the fuel is spent. Fuel elements of this design typically utilize ceramic (UO2 or UC2) rather that metallic fuel so as to be capable of reaching higher fuel temperatures. The ceramic fuel kernel is over coated with layers of SiC and pyrolytic carbon to pro- vide a fission product barrier and to negate the use of a metallic fuel clad (see Chapter 3.07, TRISO- Coated Particle Fuel Performance), allowing the reactor core to operate at very high temperatures (>1000 �C).1 The coated particle fuel is usually formed into fuel pucks or compacts but may be consoli- dated into fuel balls, or pebbles.1 The US designed modular high-temperature gas-cooled reactor (MHTGR) and Next Generation Nuclear Plant (NGNP), and the Japanese high-temperature test reactor (HTTR) are examples of gas-cooled reactors with high-temperature ceramic fuel. Additional vertical channels in the graphite reac- tor core house the control rods, which regulate the fission reaction by introducing neutron-adsorbing materials to the core, and thus reduce the number of neutrons available to sustain the fission process. When the control rods are withdrawn from the core, the self-sustaining fission reaction commences. Heat is generated by the moderation of the fission frag- ments in the fuel and moderation of fast neutrons in the graphite. The heat is removed from the core by a coolant, typically a gas, that flows freely through the core and over the graphite moderator. The coolant is forced through the core by a gas circulator and passes into a heat exchanger/boiler (frequently referred to as a steam generator). The primary coolant loop (the reactor coolant) is maintained at elevated pressure to improve the cool- ant’s heat transfer characteristics and thus, the core is surrounded by a pressure vessel. A secondary coolant (water) loop runs through the heat exchanger and cools the primary coolant so that it may be returned to the reactor core at reduced temperature. The secondary coolant temperature is raised to produce steam which is passed through a turbine where it gives up its energy to drive an electric generator. Some reactor designs, such as theMHTGR, are direct cycle systems inwhich the helium coolant passes directly to a turbine. The reactor core and primary coolant loop are enclosed in a concrete biological-shield, which pro- tects the reactor staff and public from g radiation and fission neutrons and also prevents the escape of radioactive contamination and fission product gasses that originate in the fuel pins/blocks. The charge face, refueling machine, control rod drives, discharge area, and cooling ponds are housed in a containment structure which similarly prevents the spread of any contamination. Additional necessary features of a fission reactor are (1) the refueling bay, where new fuel stringers or fuel elements are assembled prior to being loaded into the reactor core; (2), a discharge area and cooling ponds where spent fuel is placed while the short-lived isotopes are allowed to decay before the fuel can be reprocessed. The NGNP, a graphite-moderated, helium cooled reactor, is designed specifically to generate elec- tricity and produce process heat, which could be used for the production of hydrogen, or steam gener- ation for the recovery of tar sands or oil shale. Radiation Effects in Graphite 305 Two NGNP concepts are currently being consid- ered, a prismatic core design and a pebble bed core design. In the prismatic core concept, the TRISO fuel is compacted into sticks and supported within a graphite fuel block which has helium coolant holes running through its length.1 The graphite fuel blocks are discharged from the reactor at the end of the fuel’s lifetime. In the pebble bed core concept, the TRISO fuel is mixed with other graphite materi- als and a resin binder and formed into 6 cm diameter spheres or pebbles.1 The pebbles are loaded into the core to form a ‘pebble bed’ through which helium coolant flows. The pebble bed is constrained by a graphite moderator and reflector blocks which define the reactor core shape. The fuel pebbles migrate slowly down through the reactor core and are discharged at the bottom of the core where they are either sent to spent fuel storage or returned to the top of the pebble bed. Not all graphite-moderated reactors are gas- cooled. Several designs have utilized water cooling, with the water carried through the core in zirconium alloy tubes at elevated pressure, before being fed to a steam generator. Moreover, graphite-moderated reactors can also utilize a molten salt coolant, for example, the Molten Salt Reactor Experiment (MSRE)1 at Oak Ridge National Laboratory (ORNL). The fluid fuel in the MSRE consisted of UF4 dissolved in fluorides of beryllium and lithium, which was circu- lated through a reactor core moderated by graphite. The average temperature of the fuel salt was 650 �C (1200 �F) at the normal operating condition of 8MW, which was the maximum heat removal capacity of the air-cooled secondary heat exchanger. The graphite core was small, being only 137.2 cm (54 in.) in diameter and 162.6 cm (64 in.) in height. The fuel salt entered the reactor vessel at 632 �C (1170 �F) and flowed down around the outside of the graphite core in the annular space between the core and the vessel. The graphite corewasmade up of graphite bars 5.08 cm (2 in.) square, exposed directly to the fuel which flowed upward in passages machined into the faces of the bars. The fuel flowed out of the top of the vessel at a temperature of 654 �C (1210 �F), through the circulating pump to the primary heat exchanger, where it gave up heat to a coolant salt stream. The core graphite, grade CGB, was specially produced for the MSRE, and had to have a small pore size to prevent penetration of the fuel salt, a long irradiation lifetime, and good dimen- sional stability. Moreover, for molten salt reactor mod- erators, a low permeability (preferably 306 Radiation Effects in Graphite in two dimensions and a high proportion will recom- bine with lattice vacancies. Others will coalesce to form C2, C3, or C4 linear molecules. These in turn may form the nucleus of a dislocation loop – essentially a new graphite plane. Interstitial clusters may, on fur- ther irradiation, be destroyed by a fast neutron or carbon knock-on atom (irradiation annealing). Adja- cent lattice vacancies in the same graphitic layer are believed to collapse parallel to the layers, thereby forming sinks for other vacancies which are increas- ingly mobile above 600 �C, and hence can no longer recombine and annihilate interstitials. The migration of interstitials along the crystallographic c-axis is discussed later. Banhart11 observed typical basal plane defects in a graphite nanoparticles using high-resolution trans- mission electron microscopy (HRTEM). These defects can be understood as dislocation loops which form when displaced interstitial atoms cluster and form less mobile agglomerates. Other interstitials condense onto this agglomerate which grows into a disk, pushing the adjacent apart. Further agglomeration leads to the for- mation of a new lattice planes (Figure 4). Other deformation mechanisms have been pro- posed for irradiated graphite. Wallace18 proposed a mechanism whereby interstitial atoms could facili- tate sp3 bonds between the atomic basal planes, this mechanism allowing the stored energy (discussed in Section 4.10.5.1) to be explained. Jenkins19 argued that the magnitude of the increase in shear 1 nm Figure 4 A high-resolution electron micrograph showing the basal planes of a graphitic nanoparticle with an interstitial loop between two basal planes, the ends of the inserted plane are indicated with arrows. Reproduced from Banhart, F. Rep. Prog. Phys. 1999, 62, 1181–1221, with permission from IOP Publishing Ltd. modulus (C44) with low dose irradiation could not be explained by interstitial clusters pinning dislocations, but that a few sp3 type covalent bonds between the planes could easily account for the observed changes. More recently, Telling and Heggie,15 in their ab-initio calculations of the energy of formation of the ‘spiro- interstitial,’ advocate this mechanism to explain the stored energy characteristics of displacement dam- aged graphite, particularly the large energy release peak seen at �473K (discussed in Section 4.10.5.1). The first experimental evidence of the interlayer interstitial–vacancy (IV) pair defect with partial sp3 character in between bilayers of graphite was recently reported by Urita et al.20 in their study of double- walled carbon nanotubes (DWNTs). Jenkins19 invoked the formation of sp3 bonding to explain the c-axis growth observed as a result of displacement damage. If adjacent planes are pinned, one plane must buckle as the adjacent planes shrink due to vacancy shrinkage; buckled planes yield the c-axis expansion that cannot be explained by swelling from interstitial cluster alone. Telling andHeggie15 are very much in support of this position on the basis of their reviewof the literature and ab-initio simulations of the damage mechanisms in graphite. Their simulations showed how the spiro-interstitial (cross-link) essen- tially locked the planes together. Additionally, diva- cancies could lead to the formation of pentagons and heptagons in the basal planes causing the observed bending of graphene layers and c-axis swelling.11,21,22 The predicted c-axis crystal expansion via this mecha- nism is in closer agreement with the experimentally observed single crystal and highly oriented pyrolytic graphite (HOPG) dimensional change data. The buckling of basal planes as a consequence of irradiation damage has been observed in HRTEM studies of irradiated HOPG by Tanabe21 and Koike and Pedraza.22 In their study, Koike and Pedraza22 observed 300% expansion of thin HOPG samples subject to electron irradiation in an in-situ transmis- sion electron microscope (TEM) study. Their exper- imental temperatures ranged from 238 to 939K. They noted that the damaged microstructure showed retention of crystalline order up to 1 dpa (displace- ments per atom). At higher doses, they observed the lattice fringes break up in to segments 0.5–5 nm in length, with up to 15� rotation of the segments with respect to the original {0001} planes. The evidence in favor of the formation of bonds between basal planes involving interstitials is consid- erable. However, such bonds are not stable at high temperature. As reported by numerous authors and Radiation Effects in Graphite 307 reviewers11,15,19,20 the sp3 like bond would be ex- pected to break and recombine with lattice vacancies with increasing temperature, such that at T>500K they no longer exist. Indeed, the irradiated graphite stored energy annealing peak at �473K, and the HRTEM observations of Urita et al.20 demonstrate this clearly. Figure 5 shows a sequential series of HRTEM images illustrating the formation rates of interlayer defects at different temperatures with the same time scale (0–220 s) in DWNTs. The arrows indicate possi- ble interlayer defects. At T¼ 93K (Figure 5(a)) the electron irradiation-induced defects are numer- ous, and the nanotubes inside are quickly damaged because of complex defects. At 300K (Figure 5(b)), the nanotubes are more resistive to the damage from electron irradiation, yet defects are still viable. At 573K (Figure 5(c)), defect formation is rarely observed and the DWNTs are highly resistant to the electron beam irradiation presumably because of the ease of defect self-annihilation (annealing). In an attempt to estimate the critical temperature for the annihilation of the IV defect pairs, a system- atic HRTEM study was undertaken at elevated temperatures by Urita et al.20 The formation rate of the IV defects that showed sufficient contrast in the HRTEM is plotted in Figure 6. The reported numbers were considered to be an underestimate as single IV pairs may not have sufficient contrast to be 93 K (a) (b) 300 K0 s 110 s 140 s 220 s Figure 5 Sequential high-resolution transmission electron mic defects at different temperatures with the same time scale (0–22 nanotubes. The arrows indicate possible interlayer defects. Sca Sugai, T.; Shinohara, H.; Iijima, S. Phys. Rev. Lett. 2005, 94, 155 convincingly isolated from the noise level and thus may have been missed. However, the data was con- sidered satisfactory for indicating the formation rate as a function of temperature. The number of clusters of IV pairs found in a DWNT was averaged for several batches at every 50K and normalized by the unit area. As observed in Figure 6, the defect for- mation rate displays a constant rate decline, with a threshold appearing at �450–500K. This threshold corresponds to the stored energy release peak (dis- cussed in Section 4.10.5.1) as shown by the dotted line in Figure 6. Evidentially, the irradiation dam- age resulting from higher temperature irradiations (above �473–573K) is different in nature from that occurring at lower irradiation temperatures. Koike andPedraza22 studied the dimensional change in HOPG caused by electron-irradiation-induced dis- placement damage. They observed in situ the growth c-axis of the HOPG crystals as a function of irradiation temperature at damage doses up to�1.3 dpa. Increasing c-axis expansion with increasing dose was seen at all temperatures. The expansion rate was however signifi- cantly greater at temperatures ≲473K (their data was at 298 and 419K) compared to that at irradiation tem- peratures ≳473K (their data was at 553, 693, and 948K). This observation supports the concept that separate irradiation damage mechanisms may exist at low irradiation temperatures (�T 0 0 50 100 D ef ec t fo rm at io n ra te (b ar ns ) 150 200 200 400 Temperature (K) 600 800 1000 Figure 6 Normalized formation rates of the clusters of interstitial–vacancy pair defects per unit area of bilayer estimated in high-resolution transmission electron microscope images recorded at different temperatures. The dotted line shows the known temperature for Wigner-energy release (�473K). Reproduced fromUrita, K.; Suenaga, K.; Sugai, T.; Shinohara, H.; Iijima, S. Phys. Rev. Lett. 2005, 94, 155502, with permission from American Physical Society. 308 Radiation Effects in Graphite buckling due to sp3 bonded cross linking of the basal planes via interstitials, and at more elevated irradiation temperatures (T≳ 473K),where the buckling of planes is attributed to clustering of interstitials which induce the basal planes to bend, fragment, and then tilt. Koike and Pedraza22 also observed crystallographic a-axis shrinkage upon electron irradiation in-situ at several temperatures (419, 553, and 693K). The shrinkage increased with dose at all irradiation temperatures, and the shrinkage rate reduced with increasing irradia- tion temperature. This behavior is attributed to buck- ling and breakage of the basal planes, with the amount of tilting and buckling decreasing with increasing tem- perature due to (1) a switch in mechanism as discussed above and (2) increased mobility of lattice vacancies above �673K. Jenkins19,23 also discussed the deformation of graphite crystals in terms of a unit c-axis dislocation (prismatic dislocation), that is, one in which the Bur- gers vector, b, is in the crystallographic c-direction. The c-axis migration of interstitials can take place by unit c-axis dislocations. The formation and growth of these, and other basal plane dislocation loops undoubtedly play a major role in graphite crystal deformation during irradiation. Ouseph24 observed prismatic dislocation loops (both interstitial and vacancy) in unirradiated HOPG using scanning tunneling microscopy (STM). Their study allowed atomic resolution of the defect structures. Such defects had previously been observed as regions of intensity variations in TEM studies in the 1960s.25 Telling and Heggie’s15 first principle simulations have indicated a reduced energy of migration for a lattice vacancy compared to the previously estab- lished value. Therefore, they argue, the observed lim- ited growth of vacancy clusters at high temperatures (T>900K) indicates the presence of a barrier to further coalescence of vacancy clusters (i.e., vacancy traps). Telling and Heggie implicate a cross-planer metastable vacancy cluster in adjacent planes as the possible trap. The disk like growth of vacancy clusters within a basal plane ultimately leads to a prismatic dislocation loop. TEM observations show that these loops appear to form at the edges of interstitial loops in neighboring planes in the regions of tensile stress. The role of vacancies needs to be reexamined on the basis of the foregoing discussion. If the energy of migration is considerably lower than that previously considered, and there is a likelihood of vacancy traps, the vacancy and prismatic dislocation may well play a larger role in displacement damage induced in- crystal deformation. The diffusion of vacancy lines to the crystal edge essentially heals the damage, such that crystals can withstand massive vacancy damage and recover completely. Regardless of the exact mechanism, the result of carbon atom displacements is crystallite dimensional change. Interstitial defects will cause crystallite growth perpendicular to the layer planes (c-axis direction), and relaxation in the plane due to coalescence of vacancies will cause a shrinkage parallel to the layer plane (a-axis direction). The damage mechanism and associated dimensional changes are illustrated (in sim- plified form) in Figure 7. As discussed above, this conventional view of c-axis expansion as being caused solely by the graphite lattice accommodating small interstitial aggregates is under some doubt, and despite the enormous amount of experimental and theoretical work on irradiation-induced defects in graphite, we are far from a widely accepted understanding. It is to be hoped that the availability of high-resolution microscopes will facilitate new damage and annealing studies of graphite leading to an improved under- standing of the defect structures and of crystal defor- mation under irradiation. Interstitial (c) (a) Contraction New plane Vacancy Collapsing vacancy line Expansion Figure 7 Neutron irradiation damage mechanism illustrating the induced crystal dimensional strains. Reproduced from Burchell, T. D. In Carbon Materials for Advanced Technologies; Burchell, T. D., Ed.; Elsevier Science: Oxford, 1999, with permission from Elsevier. 3500 0 5 10 Neutron dose (dpa) 15 0 Irrad iation-ind uced shrinkage (% ) 20 20 40 60 80 100 3000 Gra phit izati on t emp erat ure (�C) 2500 2000 Figure 8 Neutron irradiation-induced a-axis shrinkage behavior of pyrolytic graphite showing the effects of graphitization temperature on the magnitude of the dimensional changes. Reproduced from Burchell, T. D. In Carbon Materials for Advanced Technologies; Burchell, T. D., Ed.; Elsevier Science: Oxford, 1999, with permission from Elsevier. Radiation Effects in Graphite 309 Dimensional changes can be very large, as demon- strated in studies on well-ordered graphite materials, such as HOPG that has frequently been used to study the neutron-irradiation-induced dimensional changes of the graphite crystallite.13,26 Price27 conducted a study of the neutron-irradiation-induced dimensional changes in pyrolytic graphite. Figure 8 shows the crystallite shrinkage in the a-direction for neutron doses up to 12 dpa for samples that were graphitized at a temperature of 2200–3300 �C prior to being irra- diated at 1300–1500 �C. The a-axis shrinkage in- creases linearly with dose for all of the samples, but the magnitude of the shrinkage at any given dose decreases with increasing graphitization temperature. Similar trends were noted for the c-axis expansion. The significant effect of graphitization temperature on irradiation-induced dimensional change accumula- tion can be attributed to thermally induced improve- ments in crystal perfection, thereby reducing the number of defect trapping sites in the lattice. Nuclear graphites possess a polycrystalline struc- ture, usually with significant texture resulting from the method of forming during manufacture. Con- sequently, structural and dimensional changes in polycrystalline graphites are a function of the crys- tallite dimensional changes and the graphite’s texture. In polycrystalline graphite, thermal shrinkage cracks that occur during manufacture and that are prefer- entially aligned in the crystallographic a-direction will initially accommodate the c-direction expansion, so mainly a-direction contraction will be observed. The graphite thus undergoes net volume shrinkage. With increasing neutron dose (displacements), the incompatibility of crystallite dimensional changes leads to the generation of new porosity, and the volume shrinkage rate falls, eventually reaching zero. The graphite now begins to swell at an increasing rate with increasing neutron dose. The graphite thus undergoes a volume change ‘turn- around’ into net growth that continues until the gen- eration of cracks and pores in the graphite, due to differential crystal strain, eventually causes total dis- integration of the graphite. Irradiation-inducedvolume anddimensional change data forH-451 are shown28 inFigures 9–11. The effect of irradiation temperature on volume change is shown in Figure 9. The ‘turn-around’ from volume shrinkage to growth occurs at a lower fluence and -10 0 1 2 3 4 -8 -6 -4 -2 0 2 4 6 8 10 Fast fluence 1026 n m-2 [E > 0.1 MeV] V ol um e ch an ge (% ) H-451 @ 600 �C H-451 @ 900 �C Figure 9 Irradiation-induced volume changes for H-451 graphite at two irradiation temperatures. From Burchell, T. D.; Snead, L. L. J. Nucl. Mater. 2007, 371, 18–27. H-451 graphite irradiated at 600 �C D im en si on al c ha ng e (% ) Perpendicular to extrusion (AG) Parallel to extrusion direction (WG) 0 1 2 3 54 -6 -5 -4 -1 -3 -2 0 1 Fast fluence 1026 n m-2 [E > 0.1 MeV] Figure 10 Dimensional change behavior of H-451 graphite at an irradiation temperature of 600 �C. From Burchell, T. D.; Snead, L. L. J. Nucl. Mater. 2007, 371, 18–27. H-451 graphite irradiated at 900 �C D im en si on al c ha ng e (% ) Perpendicular to extrusion (AG) Parallel to extrusion direction (WG) 0 0.5 1 1.5 2 2.5 3 -1 -3 -2 0 1 2 3 4 5 Fast fluence 1026 n m-2 [E > 0.1 MeV] Figure 11 Dimensional change behavior of H-451 graphite at an irradiation temperature of 900 �C. From Burchell, T. D.; Snead, L. L. J. Nucl. Mater. 2007, 371, 18–27. 310 Radiation Effects in Graphite the magnitude of the volume shrinkage is smaller at the higher irradiation temperature. This effect is attributed to the thermal closure of aligned micro- cracks in the graphite which accommodate the c-axis growth. Hence, there is less accommodating volume available at the higher temperatures and the c-axis growth dominates the a-axis shrinkage at lower doses. The irradiation-induced dimensional changes of H-451at 600 and 900 �C are shown in Figures 10 and 11, respectively. H-451 graphite is an extruded material and therefore, the filler coke particles are preferentially aligned in the extrusion axis (par- allel direction). Consequently, the crystallographic a-direction is preferentially aligned in the parallel direction and the a-direction shrinkage is more appar- ent in the parallel (to extrusion) direction, as indi- cated by the parallel direction dimensional change data in Figures 10 and 11. The dimensional and volume changes are greater at an irradiation temper- ature of 600 �C than at 900 �C; that is, both the maxi- mum shrinkage and the turnaround dose are greater at an irradiation temperature of 600 �C. This tem- perature effect can be attributed to the thermal closure of internal porosity that is aligned parallel to the a-direction that accommodates the c-direction swelling. At higher irradiation temperatures, a greater fraction of this accommodating porosity is closed and thus the shrinkage is less at the point of turnaround. A general theory of dimensional change in polygra- nular graphite due to Simmons29 has been extended by Brocklehurst and Kelly.30 For a detailed account of the treatment of dimensional changes in graphite the reader is directed to Kelly and Burchell.31 4.10.5 Neutron-Induced Property Changes 4.10.5.1 Wigner Energy The release of Wigner energy (named after the phys- icist who first postulated its existence) was histori- cally the first problem of radiation damage in graphite to manifest itself. The lattice displacement processes previously described can cause an excess of energy in the graphite crystallites. The damage may comprise Frankel pairs or at lower temperatures the sp3 type bond previously discussed and observed by Urita et al.20 When an interstitial carbon atom and a lattice vacancy recombine, or interplanar bonds are broken, their excess energy is given up as ‘stored energy.’ If sufficient damage has accumulated in the graphite, the release of this stored energy can result in a rapid rise in temperature. Stored energy accu- mulation was found to be particularly problematic in the early graphite-moderated reactors, which operated at relatively low temperatures. Figure 12 shows the rate of release of stored energy with 100 0 0.1 0.2 0.3 d S /d T (c al g– 1 �C –1 ) Annealing temperature (�C) 0.4 0.5 0.6 0.7 Exposures in MWd/at and dpa (approximately) Specific heat 5700/0.60 1075/0.10 100/0.01 200 300 400 500 Figure 12 Stored energy release curves for CSF graphite irradiated at �30 �C in the Hanford K reactor cooled test hole. Source: Nightingale, R. E.Nuclear Graphite; Academic Press: New York, 1962. From Burchell, T. D. In Carbon Materials for Advanced Technologies; Burchell, T. D., Ed.; Elsevier Science: Oxford, 1999, with permission from Elsevier. Radiation Effects in Graphite 311 temperature, as a function of temperature, for graph- ite samples irradiated at 30 �C to low doses in the Hanford K reactor.32 The release curves are character- ized by a peak occurring at�200 �C. This temperature has subsequently been associated with annealing of interplanar bonding involving interstitial atoms.20 In Figure 12, the release rate exceeds the specific heat and therefore, under adiabatic conditions, the graphite would rise sharply in temperature. For am- bient temperature irradiations it was found9 that the stored energy could attain values up to 2720 J g�1, which if released adiabatically would cause a temper- ature rise of some 1300 �C. A simple experiment,8 in which samples irradiated at 30 �C were placed in a furnace at 200 �C and their temperature monitored, showed that when the samples attained a temperature of �70 �C their temperature suddenly increased to a maximum of about 400 �C and then returned to 200 �C. In order to limit the total amount of stored energy in the early graphite reactors, it became nec- essary to periodically anneal the graphite. The gra- phite’s temperature was raised sufficiently, by nuclear heating or the use of inserted electrical heaters, to ‘trigger’ the release of stored energy. The release then self-propagated slowly through the core, raising the graphite moderator temperature and thereby par- tially annealing the graphite core. Indeed, Arnold33 reports that it was during such a reactor anneal that the Windscale (UK) reactor accident occurred in 1957. Rappeneau et al.34 report a second release peak at very high temperatures (�1400 �C). They studied the energy release up to temperatures of 1800 �C of graphites irradiated in the reactors BR2 (Mol, Belgium) and HFR (Petten, Netherlands) at doses between 1000 and 4000MWdT�1 and at tempera- tures between 70 and 250 �C. At these low irradiation temperatures, there is little or no vacancy mobility, so the resultant defect structures can only involve interstitials. On postirradiation annealing to high tem- peratures, the immobile single vacancies become increasingly mobile and perhaps their elimination and the thermal destruction of complex interstitial clusters or distorted and twisted basal planes contrib- ute to the high-temperature stored energy peak. The accumulation of stored energy in graphite is both dose and irradiation temperature dependent. With increasingly higher irradiation temperatures, the total amount of stored energy and its peak rate of release diminish, such that above an irradiation temperature of �300 �C stored energy ceases to be a problem. Accounts of stored energy in graphite can be found elsewhere.1,8,29,32 4.10.5.2 Mechanical and Physical Properties The mechanical and physical properties of several medium-grained and fine-grained nuclear grade gra- phites currently in production are given in Table 2 (see also Chapter 2.10, Graphite: Properties and Characteristics). The coke type, forming method, and potential uses of these grades are inTable 1. The most obvious difference between the four grades listed in Table 2 is the filler particle sizes. Grade IG-110 is an isostatically pressed, isotropic grade, whereas the others grades shown are near-isotropic and have properties reportedeitherwith-grainor against-grain.Asdiscussed earlier (see Section 4.10.2), the orientation of the filler coke particles is a function of the forming method. The mechanical properties of nuclear graphites are substantially altered by radiation damage. In the unirradiated condition, nuclear graphites behave in a brittle fashion and fail at relatively low strains. The stress–strain curve is nonlinear, and the fracture process occurs via the formation of subcritical cracks, which coalesce to produce a critical flaw.35,36 When graphite is irradiated, the stress–strain curve becomes more linear, the strain to failure is reduced, and the strength and elastic modulus are increased. On irra- diation, there is a rapid rise in strength, typically �50%, that is attributed to dislocation pinning at irradiation-induced lattice defect sites. This effect is largely saturated at doses >1 dpa. Above �1 dpa, a more gradual increase in strength occurs because of Table 2 Typical physical and mechanical properties of unirradiated nuclear graphites Property Graphite grade IG-110 PCEA NBG-10 NBG-18 Maximum filler particle size (mm) 10 800 1600 1600 Bulk density (g cm�3) 1.77 1.83 1.79 1.88 Tensile strength (MPa) 24.5 21.9 (WG) 20.0 (WG) 21.5 (WG) 16.9 (AG) 18.0 (AG) 20.5 (AG) Flexural strength (MPa) 39.2 32.4 (WG) 24.0 (WG) 28 (WG) 23.3 (AG) 27.0 (AG) 26 (AG) Compressive strength (MPa) 78.5 60.8 (WG) 47.0 (WG) 72.0 (WG) 67.6 (AG) 61.0 (AG) 72.5 (AG) Young’s modulus (GPa) 9.8 11.3 (WG) 9.7 (WG) 11.2 (WG) 9.9 (AG) 9.7 (AG) 11.0 (AG) Thermal conductivity (Wm�1 K�1) (measured at ambient temperature) 116 162 (WG) 148 (WG) 156 (WG) 159 (AG) 145 (AG) 150 (AG) Coefficient of thermal expansion (10�6 K�1) (over given temperature range) 4.5 (350–450 �C) 3.5 (WG) 4.1 (WG) 4.5 (WG) 3.7 (AG) 4.6 (AG) 4.7 (AG) (30–100 �C) (20–200 �C) (20–200 �C) Electrical resistivity (mOm) 11 7.3 (WG) 9.1 (WG) 8.9 (WG) 7.8 (AG) 9.3 (AG) 9.0 (AG) WG, with-grain; AG, against-grain. 0 0 10 20 30 40 50 Yo un g’ s m od ul us (G P a) 5 10 15 Fluence (dpa) 20 875 �C 600 �C 25 30 Figure 13 Neutron irradiation-induced Young’s modulus changes for GraphNOL N3M at irradiation temperatures 600 and 875 �C. From Burchell, T. D.; Eatherly, W. P. J. Nucl. Mater. 1991, 179–181, 205–208. 312 Radiation Effects in Graphite structural changes within the graphite. For nuclear graphites, the dose at which the maximum strength is attained loosely corresponds with the volume change turnaround dose, indicating the importance of pore closure and generation in controlling the high-dose strength behavior, and may be as much as twice the unirradiated value. The strain behavior of nuclear graphites subjected to an externally applied load is largely controlled by shear of the component crystallites. As with strength, irradiation-induced changes in Young’s modulus are the combined result of in-crystallite effects, due to low fluence dislocation pinning, and superimposed structural change external to the crystallite. The effects of these two mechanisms are generally consid- ered separable, and related by ðE=E0Þirradiated ¼ ðE=E0ÞpinningðE=E0Þstructure ½1� where E/E0 is the ratio of the irradiated to unirradi- ated elastic modulus. The dislocation pinning contri- bution to the modulus change is due to relatively mobile small defects and is thermally annealable at �2000 �C. The irradiation-induced elastic modulus changes for GraphNOL N3M graphite37 are shown in Figure 13. The low dose change due to dislocation pinning (dashed line) saturates at a dose 200 400 600 800 1000 12000 0 20 40 60 80 100 120 140 160 Temperature (�C) Th er m al c on d uc tiv ity (W m –1 K –1 ) 12 dpa 25 dpa 26 dpa Unirradiated Figure 14 Temperature dependence of thermal conductivity in the irradiated and unirradiated condition for typical nuclear grade graphite. Irradiation temperature ¼ 600 �C. Radiation Effects in Graphite 313 Graphite is a phonon conductor of heat. Therefore, any reduction in the intrinsic defect population causes a reduction in the degree of phonon-defect scattering, an increase in the phonon mean free path, and an increase in the thermal conductivity. In graphite, such thermally induced improvements are attributable to increases in crystal perfection and a concomitant increase in the size of the regions of coherent order- ing upon graphitization. With increasing temperature, the dominant phonon interaction becomes phonon– phonon scattering (Umklapp processes). Therefore, there is a reduction of thermal conductivity with increasing temperature.39 This decrease in the thermal conductivity with increasing temperature can clearly be seen in Figure 14. The mechanism of thermal conductivity and the degradation of thermal conductivity have been exten- sively reviewed.13,14,26,40 The increase of thermal resis- tance due to irradiation damage has been ascribed by Taylor et al.41 to the formation of (1) submicroscopic interstitial clusters, containing 4 2 carbon atoms; (2) vacant lattice sites, existing as singles, pairs, or small groups; and (3) vacancy loops, which exist in the graphite crystal basal plane and are too small to have collapsed parallel to the hexagonal axis. The contribution of collapsed lines of vacant lattice sites and interstitial loops, to the increased thermal resis- tance, is negligible. The reduction in thermal conductivity due to irradiation damage is temperature and dose sensitive. At any irradiation temperature, the decreasing thermal conductivity will reach a ‘saturation limit.’ This limit is not exceeded until the graphite undergoes gross struc- tural changes at very high doses. The ‘saturated’ value of conductivity will be attained more rapidly, and will be lower, at lower irradiation temperatures.42 In graph- ite, the neutron irradiation-induced degradation of thermal conductivity can be very large, as illustrated in Figure 14. This reduction is particularly large at low temperatures. Bell et al.43 have reported that the room temperature thermal conductivity of pile grade A (PGA) graphite is reduced by more than a factor of 70 when irradiated at 155 �C to a dose of �0.6 dpa. At an irradiation temperature of 355 �C, the room temperature thermal conductivity of PGA was reduced by less than a factor of 10 at doses twice that obtained at 155 �C. Above 600 �C, the reduc- tion of thermal conductivity is less significant. For example, Kelly8 reported the degradation of PGA at higher temperatures: at an irradiation temperature of 600 �C and a dose of �13 dpa, the thermal conductiv- ity was degraded only by a factor of �6; at irradiation temperatures of 920 and 1150 �C, the degradation was minimal (less than a factor of 4 at �7 dpa). For the fine-grained, isomolded graphite shown in Figure 14, the degradation of thermal conductivity at the irradia- tion temperature (600 �C) was only by a factor of �3, but was by a factor �6 at a measurement temperature of 100 �C. There are two principal thermal expansion coeffi- cients in the hexagonal graphite lattice; ac, the ther- mal expansion coefficient parallel to the hexagonal c-axis and aa, the thermal expansion coefficient par- allel to the basal plane (a-axis). The thermal expan- sion coefficient in any direction at an angle f to the c-axis of the crystal is aðfÞ ¼ ac cos2fþ aa sin2f ½3� The value of ac varies linearly with temperature from �25� 10�6 K�1 at 300 K to �35� 10�6 K�1 at 2500 K. In contrast, aa is much smaller and increases rapidly from �1.5� 10�6 K�1 at �300K to �1� 10�6 K�1 at 1000 K, and remains relatively constant at temperatures up to 2500 K.39 The large anisotropy in the crystal coefficient of thermal expansion (CTE) values is a direct conse- quence of the bond anisotropy and the resultant anisotropy in the crystal lattice compliances. The thermal expansion of polycrystalline graphites is con- trolled by the thermal closure of aligned internal porosity which forms as a result of thermal shrinkage strains on cooling after graphitization. Thus, the c-axis expansion of the graphite crystals is initially, partially accommodated by this internal porosity and a much lower bulk CTE value is observed. On further heating, the graphite crystals fill more of the available internal porosity and more of the c-axis expansion is observed. The bulk CTE thus increases with temper- ature (Figure 15). 0.0 0 200 400 600 800 1000 1.0 2.0 3.0 4.0 5.0 6.0 Measurement temperature (�C) A ve ra ge c oe ff ic ie nt o f t he rm al ex p an si on (1 0- 6 �C -1 ) Figure 15 Temperature dependence of the coefficient of thermal expansion for typical nuclear grade graphite. 0 2 3 4 C oe ffi ci en t of t he rm al e xp an si on a (1 0- 6 �C -1 ) 5 875 �C 600 �C 6 5 10 15 Fluence (dpa) 20 25 30 Figure 16 The irradiation-induced changes in coefficient of thermal expansion (25–500 �C) for GraphNOL N3M graphite at two irradiation temperatures. From Burchell, T. D.; Eatherly, W. P. J. Nucl. Mater. 1991, 179–181, 205–208. 0 0 500 1000 1500 2000 2500 3000 400 800 1200 1600 2000 2400 Temperature (K) S p ec ifi c he at (J K g– 1 K –1 ) Calculated value Experimental data Figure 17 The temperature dependence of the specific heat of graphite, a comparison of calculated values and literature data for POCO AXM-5Q graphite. Sources: ASTM C 781. Standard Practice for Testing Graphite and Boronated Graphite Materials for High-Temperature Gas-Cooled Nuclear Reactor Components, Annual Book of Standards. ASTM International: West Conshohocken, PA; Vol. 05.05; Hust, J. G. NBS Special Publication 260–89; US Department of Commerce, National Bureau of Standards, 1984; p 59. 314 Radiation Effects in Graphite As the CTE of polycrystalline graphite is depen- dent on the pore structure, irradiation-induced changes in the pore structure (see discussion of structural changes in Section 4.10.4) can be expected to modify the thermal expansion behavior of carbon materials. Burchell and Eatherly37 report the behav- ior of GraphNOL N3M (which is typical of many fine-textured graphites), which undergoes an initial increase in the CTE followed by a steady reduction to a value less than half the unirradiated value of 5� 10�6 �C�1 (Figure 16). Similar behavior is reported by Kelly8 for grade IM1-24 graphite. Heat energy is stored in the crystal lattice in the form of lattice vibrations. The Debye equation there- fore gives the specific heat, C, as follows: C ¼ 9R T yD � �3 ðyT 0 z4ez ðez � 1Þ2 dz ½4� where R is the gas constant (8.314 Jmol�1 K�1); T, the temperature; yD, the Debye temperature; and z¼ ho/2kTp, where o is the frequency of vibrational oscillations; k, the Boltzmann’s constant; T, the tem- perature; and h is the Plank’s constant. At low temperatures, where (T/yD) Radiation Effects in Graphite 315 (Wigner energy) will reduce the effective specific heat (see Section 4.10.4). Cp ¼ 1 11:07T�1:644 þ 0:0003688T 0:02191 J Kg �1 K�1 ½6� The electrical resistivity of graphite is also affected by radiation damage. The mean free path of the conduction electron in unirradiated graphite is rela- tively large, being limited only by crystallite bound- ary scattering. Neutron irradiation introduces (1) scattering centers, which reduce charge carrier mobility; (2) electron traps, which decrease the charge carrier density; and (3) additional spin reso- nance. The net effect of these changes is to increase the electrical resistivity on irradiation, initially very rapidly, with little or no subsequent change to rela- tively high fluence.14,37 A subsequent decrease at very high neutron doses is attributed to structural degradation. 4.10.6 Irradiation Creep 4.10.6.1 The Relevance of Creep to Reactor Design and Operation Graphite will undergo creep (inelastic strain) during neutron irradiation and under stress at temperatures where thermal creep is generally negligible. The phenomenon of irradiation creep has been widely studied because of its significance to the operation of graphite-moderated fission reactors. Indeed, if irradiation-induced stresses in graphite moderators could not relax via radiation creep, rapid core disinte- gration would result. The total strain, eTotal, in a graph- ite component under irradiation in a reactor core is given by the materials constitutive equation: eTotal ¼ ee þ et þ ed þ ec ½7� where ee is the elastic strain; et, the thermal strain; ed, the dimensional change strain; and ec is the creep strain, which is given as ec ¼ ep þ es ½8� where ep and es are the primary and secondary creep strains, respectively. Tsang and Marsden46 concluded that irradiation creep strain is particu- larly important in reactor design because without creep strain self-induced shrinkage stresses would build up to levels exceeding the graphite component failure strength. The significance of irradiation creep to reactor core design and operation has been the subject of recent work, where it has been shown how uncertainties in the assumed magnitudes of the irradiation-induced creep strains in a graphite reac- tor core component can substantially impact the predicted stress levels, and hence the predicted fail- ure probabilities of core components.47,48 Li et al.47 assumed the current UK creep law and showed that a 50% decrease in the assumed creep strain resulted in a 50% increase in the magnitude of the predicted hoop stress in a hollow cylindrical core brick. Simi- larly, a 50% increase in the assumed creep strain yielded a 30% reduction in the predicted brick hoop stress. Wang and Yu48 report the effect of varying creep strain ratio (analogous to Poisson’s ratio, but where the two perpendicular strains are creep strains) on the magnitude of the modified equivalent stress in a graphite component and the associated probability of failure, as a function of neutron dose. In addition, they examined the influence of primary creep in reducing the magnitudes of stresses and associated failure probabilities in graphite core components. Wang and Yu’s results clearly indicate that variations of the creep strain ratio resulted in considerable change in the stress distributions and the corre- sponding failure probabilities of graphite compo- nents. In addition, they showed that the primary creep appears to play the same important role as secondary creep in certain cases. Because of the significance of irradiation-induced creep to the stress levels in graphite core compo- nents, accurate models of creep have long been sought. Recently, the breakdown of the currently accepted model(s) of creep at high temperatures and doses has been reported, and possible improve- ment or alternative models have been postu- lated.49,50 Analysis of the creep behavior of H-451 at high doses indicated that further modification to the current Kelly and Burchell51 model is required to allow for the generation of new porosity at higher doses and temperatures.50 The extent to which high-dose creep strain behavior differs between the compressive load and tensile load situations is shown in Figure 18, which compares the creep behavior of ATR-2E graphite for the þ5MPa and the �5MPa loading cases;52 the dashed lines are polynomial fits to the data. A more rapidly increasing creep rate in the tensile loading case compared to the compressive case is clearly observed. Because of the importance of irradiation-induced creep to the design and operation of graphite reactor cores, the subject is treated here in considerable detail, 0 43210 0.2 0.4 0.6 0.8 1 1.2 1.4 1.6 1.8 2 Neutron dose 1022 n cm-2 [E > 50 keV] C re ep s tr ai n (% ) Tensile creep strain (% MPa) Compressive creep strain (%) Poly tensile creep strain (% MPa) Poly compressive creep strain (%) Figure 18 A comparison of the tensile (þ5MPa) and compressive creep (�5MPa) rates of ATR-2E graphite at irradiation temperature of 500–550 �C. Source: Haag, G. Properties of ATR-2E Graphite and Property Changes due to Fast Neutron Irradiation; Report No. Jul-4183; Published FZ-J, Germany, 2005; Available at http://juwel.fz-juelich.de. 316 Radiation Effects in Graphite including a review of the in-crystal creep mechanism and irradiation-induced graphite creep models. 4.10.6.2 The Irradiation-Induced Creep Mechanism (In-Crystal) A mechanism for the irradiation-induced creep of graphite was proposed by Kelly and Foreman53 which involves irradiation-induced basal plane dis- location pinning/unpinning in the graphite crystals. Pinning sites are created and destroyed by neutron irradiation (radiation annealing). Under neutron irradiation, dislocation lines in the basal planes may be completely or partially pinned depending upon the dose and temperature of irradiation. The pinning points were speculated to be interstitial atom clus- ters 4 2 atoms in size,54,55 that is, the same defects clusters assumed to contribute to the reduction in thermal conductivity. The interstitial clusters are temporary barriers as they are annealed (destroyed) by further irradiation. Thus, irradiation can release dislocation lines from their original pinning site and the crystal can flow as a result of basal plane slip at a rate determined by the rate of pinning and unpinning of dislocations. Kelly and Foreman’s the- ory assumes that polycrystalline graphite consists of a single phase of true density r0 and apparent density r. The material may be divided into elementary regions in which the stress may be considered uniform and which may be identified as monocrystalline graphite. Significantly, the model excludes porosity. It is further assumed that the only deformation mode is basal plane slip for which the strain rate is determined by e_xz ¼ kðsxzÞf ½9� and e_yz ¼ kðsyzÞf ½10� where f is the fast neutron flux; k, the steady-state creep coefficient, and s is the stress in the given direc- tion. The microscopic deformation assumes the usual relationship between the basal plane shear strain rate (e_) and the mobile dislocation density (O), and is given by e_ ¼ Obn ¼ ksf ½11� where b is the Burger’s vector and n is the dislocation velocity as a function of the pinning point concentra- tion in the basal plane as the pins are created and destroyed by neutron flux. The dislocation line flow model used the flexible line approach where the dislo- cation line is pinned/unpinned and the dislocation line bowing is a function of the line tension and pin spacing. The concentration of pinning sites increases under irradiation from the initial value (from intrinsic defects) to a steady-state concentration. The initial creep rate is high and decreases to a steady-state value as the pinning concentration saturates at a level controlled by the neutron flux and temperature. This saturation would be expected to occur over the same dose scale as the reduction of thermal conductivity to its saturation limit (see Section 4.10.5.2). Thus, a two stage model can be envisioned where the primary creep rate is initially high and falls to a secondary or ‘steady-state’ creep rate. The steady-state creep term should be the dominant termwhen the dose has reached values at which physical property changes due to dislocation pinning have saturated (see Section 4.10.5.2). Kelly and Foreman state that at higher tem- peratures the steady-state (secondary) creep rate (k) http://juwel.fz-juelich.de Radiation Effects in Graphite 317 would be expected to increase because of (1) incom- patibility of crystal strains increasing the internal stress and thus enhancing the creep rate, and (2) additional effects due to the destruction of interstitial pins by thermal diffusion of vacancies (thermal annealing as well as irradiation annealing). Kelly and Foreman53 further speculate that the nonlinearity of creep strain with stress, which is expected at higher stress levels, may also be related to the high-dose dimensional change behavior of polycrystalline graphite.56 The possibility of other dislocation and crystal deformation mechanisms being involved in irradia- tion creep must also be considered. For example, prismatic dislocations may play an enhanced role at high temperatures (>250 �C) when the graphite lat- tice is under stress, as suggested by others.57 Are there mechanisms of dislocation climb and glide that need to be explored? Can dislocation lines climb/glide past the assumed interstitial cluster barriers via a mechanism that is active only when structural rear- rangements occur during irradiation? This behavior is analogous to carbons and graphites undergoing thermal creep when they undergo structural reorga- nization, that is, during carbonization and graphitiza- tion (thermal relaxation or slumping). 4.10.6.3 Review of Prior Creep Models 4.10.6.3.1 Linear viscoelastic creep model Irradiation-induced (apparent) creep strain is con- ventionally defined as the difference between the dimensional change of a stressed specimen and an unstressed specimen irradiated under identical con- ditions. Early creep data was found to be well described by a viscoelastic creep model1,58–63 where total irradiation creep (ecÞ ¼ primary (transient) creepþ secondary (steady-state) creep. ec ¼ as E0 1� expð�bgÞ½ � þ ksg ½12� where ec is the total creep strain; s, the applied stress; E0, the initial (preirradiated) Young’s modulus; g, the fast neutron fluence; a and b are constants (a is usually¼ 1); and k is the steady-state creep coefficient in units of reciprocal neutron dose and reciprocal stress. Equation [13] thus conforms to the Kelly–Foreman theory of creep with an initially large primary creep coefficient, while the dislocation pinning sites develop to the equilibrium concentration, at which time the creep coefficient has fallen to the steady-state or secondary value. Early creep experiments in several countries showed the primary creep saturated at approximately one elastic strain (s/E0) so that the true creep may be represented as ec ¼ s E0 þ ksg ½13� This is often normalized to the initial elastic strain and written as ec ¼ 1þ kE0g ½14� in elastic strain units (esu) (esu is defined as the externally applied stress divided by the initial static Young’s modulus), or creep strain per unit initial elas- tic strain; kE is the creep coefficient in units of recip- rocal dose [United Kingdom� 0.23� 10�20 cm2 n�1 EDN up to Tirr� 500 �C]. (EDN – equivalent DIDO nickel dose, a unit of neutron fluence used in the United Kingdom and Europe.) 4.10.6.3.2 The UK creep model The UK model1,64,65 recognizes that the initial creep coefficient is modified by irradiation-induced struc- ture changes (i.e., changes to the pore structure). Hence, the total creep strain is given by ec ¼ s E0 þ dec dg � � 0 s ðg 0 S�1ðgÞdg ½15� where s is the applied stress; (dec/dg)0, the initial secondary creep rate; g, the fast neutron fluence; S(g), the structure factor, given by S(g)¼ Eg/Ep the ratio of the Young’s modulus at dose g to the Young’s modulus after the initial increase due to dislocation pinning. The structure factor, S(g), thus attempts to sepa- rate those effects due to dislocation pinning occurring within the crystallites and structural effects occur- ring ex-crystal through changes in the Young’smodulus. However, the effect of creep strain (tensile or compres- sive) on modulus is not considered when evaluating the structure term. The unstressed Young’s modulus changes are used to establish the magnitude of S(g). 4.10.6.3.3 The Kennedy model Kennedy et al.66 replaced the structure term in the UK model with a parameter based on the volume change behavior of the graphite: ec ¼ s E0 þ k0ðgÞsg ½16� where k0ðgÞ ¼ k0 ðg 0 1� m DV=V0ðDV=V0Þmax � �� � dg ½17� 318 Radiation Effects in Graphite Here, m is an empirical constant equal to 0.75 and k0 is the steady-state creep coefficient established from low dose creep experiments. Although the Kennedy et al.66 model was shown to perform well in the prediction of high-dose tensile creep data, it did not predict the compressive data nearly as well. Moreover, as with the UK model, the sign of the applied stress is not considered when evaluating the influence of structure change (as reflected in volume changes). The quotient in eqn [17] is evaluated solely from unstressed (stress- free) samples irradiation behavior. As discussed by Kelly and Burchell,51 the term (DV/Vmax) does not exist at low irradiation temperatures where graphites expand in volume. 4.10.6.3.4 The Kelly and Burchell model The Kelly and Burchell50,51 model recognizes that creep produces significant modifications to the dimen- sional change component of the stressed specimen compared to that of the control and that this must be accounted for in the correct evaluation of creep strain data. The rate of change of dimensions with respect to neutron dose g(n cm�2) in appropriate units is given by the Simmons’ theory29 for direction x in the unstressed polycrystalline graphite: gx ¼ ax � aaac � aa � � dXT dg � � þ 1 Xa dXa dg þ Fx ½18� where ax is the thermal expansion coefficient in the x-direction, and ac and aa are the thermal expansion coefficients of the graphite crystal parallel and per- pendicular to the hexagonal axis, respectively, over the same temperature range. The term Fx is a pore generation term that becomes significant at interme- diate doses when incompatibilities of irradiation- induced crystal strains cause cracking of the bulk graphite.67 For the purposes of their analysis, Kelly and Burchell ignored the term Fx . The parameters (1/Xc)(dXc/dg) and (1/Xa)(dXa/dg) are the rates of change of graphite crystallite dimensions parallel and perpendicular to the hexagonal axis, and dXT dg ¼ 1 Xc dXc dg � 1 Xa dXa dg ½19� The imposition of a creep strain is known to change the thermal expansion coefficient of a stressed speci- men, increasing it for a compressive strain and decreasing it for a tensile strain compared to an unstressed control. Thus, the dimensional change component of a stressed specimen at dose g(n cm�2) is given by g 0x ¼ a0x � aa ac � aa � � dXT dg � � þ 1 Xa dXa dg þ F 0x ½20� where a0x is the thermal expansion coefficient of the crept sample, and F 0x is the pore generation term for the crept specimen. The difference between these two equations is thus the dimensional change correc- tion that should be applied to the apparent creep strain (the pore generation terms Fx and F 0 x were neglected): g 0x � gx ¼ a0x � aa ac � aa � � dXT dg � � � ax � aa ac � aa � � dXT dg � � ¼ a 0 x � ax ac � aa � � dXT dg � � ½21� The true creep strain rate can now be expressed as de dg ¼ de 0 dg � a 0 x � ax ac � aa � � dXT dg � � ½22� where e is the true creep strain and e0 is the apparent creep strain determined experimentally in the con- ventional manner. Thus, the true creep strain (ec) parallel to the applied creep stress is given by ec ¼ e0c � ðg 0 a0x � ax ac � aa � � dXT dg � � dg ½23� where e0c is the induced apparent creep strain, ða0x � axÞ is the change in CTE as a function of dose, ðac � aaÞ is the difference of the crystal thermal expansion coefficients of the graphite parallel and perpendicular to the hexagonal axis, XT is the crystal shape change parameter given above, and g is the neutron dose. The apparent (experimental) creep strain is thus given by e0c ¼ ec þ ðg 0 a0x � ax ac � aa � � dXT dg � � dg ½24� Substituting for ec from eqn [13] gives the apparent (experimental) creep strain e0c as e0c ¼ s E0 þ ksg � � þ ðg 0 a0x � ax ac � aa � � dXT dg � � dg ½25� with the terms as defined above. The Kelly–Burchell model is unique in that it does take account of the sign of the applied stress in Radiation Effects in Graphite 319 predicting creep strain through changes in the CTE of the stressed graphite. While the model gave good agreement between the predicted H-451 graphite apparent creep strain and the experimental data at low doses and high temperatures51 (Figures 19–22), the creep model was shown to be inadequate at doses >0.5� 1022 n cm�2 [E>50 keV] (�3.4 dpa) at an irradiation temperature of 900 �C (Figure 23).50 4.10.6.3.5 The M2 model Based upon the evidence from UK and US creep experiments, Davies and Bradford49,68 suggest the following: � The strain induced change in CTE is not a func- tion of secondary creep strain, but saturates after a dose of �30� 1020 n cm�2 EDN (�3.9 dpa). � There is evidence, from both thermal and irradia- tion annealing, for a recoverable strain several -2 -1.5 -1 -0.5 0 0 0.1 0.2 0.3 0.5 Neutron dose 1022 n cm-2 C re ep s tr ai n (% ) Figure 19 Comparison of predicted apparent creep strain (fro irradiation creep at 600 �C under a compressive stress of 13.8M From Burchell, T. D. J. Nucl. Mater. 2008, 381, 46–54. -2.5 -3 -2 -1.5 -1 -0.5 0 0 0.1 0.2 0.3 0.5 1 Neutron dose 1022 n cm- C re ep s tr ai n (% ) Figure 20 Comparison of predicted apparent creep strain (fro irradiation creep at 600 �C under a compressive stress of 20.7M From Burchell, T. D. J. Nucl. Mater. 2008, 381, 46–54. times that of primary creep, and a lower associated secondary creep coefficient that has been previ- ously assumed. � The dose at which the recoverable strain saturates bears a striking similarity to that of the saturation of the CTE change. Davies and Bradford49,68 proposed a new creep model (the M2 model) without the term reflecting changes in CTE due to creep, but containing one additional term, recoverable creep: ec ¼lk1 E0 exp�k1g1 ðg1 0 s SW expk1gdg þ xk2 E0 exp�k2g1 ðg1 0 s SW expk2gdgþ b E0 ðg1 0 s SW dg ½26� 0.4 0.5 0.6 Experimental creep strain True creep strain CTE correction strain Predicted apparent creep strain [E > 50 keV] m eqn [25]) and the experimental creep strain data for Pa. The true creep strain is calculated from eqn [13]. Experimental creep strain True creep strain CTE correction strain Predicted apparent creep strain 0.4 0.5 0.6 2 [E > 50 keV] m eqn [25]) and the experimental creep strain data for Pa. The true creep strain is calculated from eqn [13]. Experimental creep strain True creep strain Dimensional change correction Predicted apparent creep strain-4 -5 -3 -2 -1 0 0 0.1 0.2 0.3 0.4 0.5 0.6 1 2 Neutron dose 1022 n cm-2 [E > 50 KeV] C re ep s tr ai n (% ) Figure 22 Comparison of predicted apparent creep strain (from eqn [25]) and the experimental creep strain data for irradiation creep at 900 �C under a compressive stress of 20.7MPa. The true creep strain is calculated from eqn [13]. From Burchell, T. D. J. Nucl. Mater. 2008, 381, 46–54. Experimental (apparent) creep True creep strain CTE change correction Predicted apparent creep strain-2.5 -3 -3.5 -2 -1.5 -1 -0.5 0 0 0.1 0.2 0.3 0.4 0.5 0.6 0.5 1.5 1 Neutron dose 1022 n cm-2 [E > 50 keV] C re ep s tr ai n (% ) Figure 21 Comparison of predicted apparent creep strain (from eqn [25]) and the experimental creep strain data for irradiation creep at 900 �C under a compressive stress of 13.8MPa. The true creep strain is calculated from eqn [13]. From Burchell, T. D. J. Nucl. Mater. 2008, 381, 46–54. Apparent (experimental) creep True creep Dimensional change correction Predicted apparent creep -2.0 -1.0 0 1 1.50.5 2 5.0 4.0 1.0 0.0 3.0 2.0 Neutron dose, 1022 n cm-2 [E > 50 KeV] C re ep s tr ai n (% ) Figure 23 Comparison of predicted apparent creep strain (from eqn [26]) and the experimental creep strain data for irradiation creep at 900 �C under a tensile stress of 6MPa. The true creep strain is calculated from eqn [13]. From Burchell, T. D. J. Nucl. Mater. 2008, 381, 46–54. 320 Radiation Effects in Graphite Radiation Effects in Graphite 321 where ec is the total creep strain; s, the applied stress; l, x, and b, are the empirical fitting parameters; k1 and k2, the primary and recoverable dose constants respectively; and W is the oxidation change factor (with respect to Young’s modulus) and is analogous to the structure factor. The terms in the eqn [26] are proportional to esu and the effects of structural changes and radiolytic oxidation (gasification of graphite by an activated species that occurs in CO2 cooled reactors) are also included. The rates of satu- ration of the primary and recoverable creep compo- nents are controlled by the dose constants k1 and k2. The first and last terms in eqn [26] are primary and secondary creep as in the prior UK creep model, with the middle term being recoverable creep. Primary creep is still fast acting, but in the AGR temperature range of 400–650 �C, appears to act on a longer fluence scale equivalent to that associated in the United Kingdom with the Young’s modulus pinning,69 k1¼ 0.1, and saturates at 1 esu (a¼ 1). The irrecoverable creep is synonymous with second- ary creep, but with a coefficient, b, derived from the irrecoverable strain postthermal anneal, as 0.15 per 1020 n cm�2 EDN (�1.3 dpa) in the AGR temperature range. The lateral creep strain ratios for primary and recoverable creep are assumed to be the Poisson’s ratio and secondary creep is assumed to occur at constant volume. Figure 24 shows the performance of the M2 models applied to some high dose ATR-2E tensile creep data52 when irradiated at 500 �C in high flux reactor (HFR), Petten. The prediction matches the observed data well up to significant fluence of �160� 1020 n cm�2 EDN (�21 dpa). Only beyond 0 0 50 100 150 200 250 0.005 0.01 0.015 0.02 0.025 EDND (1020 n cm-2) C re ep s tr ai n M2 model 500 T Figure 24 Comparison of the M2 models prediction and experimental creep strain data for ATR-2E tensile creep data, when irradiation was at 500 �C in Petten. Reproduced from Davies, M.; Bradford, M. J. Nucl. Mater. 2008, 381, 39–45. this fluence does the new model prediction deviate from the data with a delay in the increase in creep strain at high doses that is often referred to as the ‘tertiary’ creep phase. Figure 25 shows the corresponding compressive creep data,52 irradiated at 550 �C. The model over predicts the data slightly but follows the trend remark- ably well up to a significant fluence of �160� 1020 n cm�2 EDN (�21 dpa). Beyond this fluence, the compressive prediction also indicates a ‘tertiary’ creep, but the data does not extend into this region. The data52 also indicates a possible difference between tensile and compressive creep (seen more clearly in Figure 18). Saturation of CTE with creep strain as reported by Davies and Bradford49,68 is not however in agree- ment with other published data. Gray70 reported CTE behavior with creep strain (up to 3%) for three different graphites at irradiation temperatures of 550 and 800 �C. Saturation of the CTE in the manner described by Davies and Bradford49,68 for UK AGR graphite was not observed. 4.10.6.4 Deficiencies in Current Creep Models at High Neutron Doses The poor performance of the Kelly and Burchell model (eqn [25]) at predicting the high temperature (900 �C) and high dose 6MPa tensile creep data suggests that the model requires further revision.50,71 H-451 graphite irradiated at 900 �C goes through dimensional change turn-around in the dose range 1.3–1.5� 1022 n cm�2 [E>50] (�8.8–10.2 dpa). This behavior is understood to be associated with the 0 50 100 250150 2000 0.005 0.01 0.015 0.02 Dose (1020 n cm-2 EDN) C re ep s tr ai n Model 550 �C 550 �C Figure 25 Comparison of the M2 models prediction and experimental creep strain data for ATR-2E compressive creep data, when irradiation was at 550 �C in Petten. Reproduced from Davies, M.; Bradford, M. J. Nucl. Mater. 2008, 381, 39–45. 322 Radiation Effects in Graphite generation of new porosity due to the increasing mismatch of crystal strains. The Kelly–Burchell model accounts for this new porosity only to the extent to which it affects the CTE of the graphite, through changes in the aligned porosity. Gray70 observed that at 550 �C the creep rate was approximately linear. However, at 800 �C he reported a marked nonlinearity in the creep rate and the changes in CTE were significant. Indeed, for the two high density graphites (H-327 and AXF-8Q) Gray reports that the 900 �C creep strain rate reverses. Gray postulated a creep strain limit to explain this behavior, such that a back stress would develop and cause the creep rate to reduce. Other workers have shown that a back stress does not develop.62 However, Gray further argued that a creep strain limit is improbable as this cannot explain the observed reversal of creep strain rate. Note that a reversal of the creep rate is clearly seen in the 900 �C tensile creep strain data reported here for H-451 (Figure 23). Also, a creep strain limit would require that tensile stress would modify the onset of pore generation behavior in the same way as compressive stress, because the direction of the external stress should be immaterial.70 More recent data52 and the behavior reported by Burchell71 show that this is not the case. Gray70 suggests that a more plausible expla- nation of his creep data is the onset of rapid expansion accelerated by creep strain; that is, net pore generation begins earlier under the influence of a tensile applied stress. Indeed, it has been observed52 that compressive creep appears to delay the turnaround behavior and tensile creep accelerates the turnaround behavior (Figure 18). In discussing possible explanations for his creep strain and CTE observations, Gray70 noted that changes in the graphite pore structure that mani- fested themselves in changes in CTE did not appear to influence the creep strain at higher doses. The classical explanation of the changes in CTE invokes the closure of aligned porosity in the graphite crys- tallites. Further crystallite strain can be accommo- dated only by fracture. A result of this fracture is net generation of porosity resulting in a bulk expansion of the graphite. A requirement of this model is that the CTE should increase monotonically from the start of irradiation. A more marked increase in CTE would be seen when the graphite enters the expan- sion phase (i.e., all accommodating porosity filled). The observed CTE behavior, reported previously50 and in Gray’s70 work, does not display this second increase in CTE; thus, the depletion of (aligned) accommodation porosity is not responsible for the early beginning of expansion behavior. The observation by Gray70 and Kennedy63 that creep occurs at near constant volume (up to moder- ate fluence) indicates that creep is not accompanied by a net reduction of porosity compared to unstressed graphite, but this does not preclude that stress may decrease pore dimensions in the direction of the applied stress and increase them in the other, that is, a reorientation of the pore structure. Pore reorienta- tion could effectively occur as the result of a mecha- nism of pore generation where an increasing fraction of the new pores are not well-aligned with the crystal- lites basal planes (and thus they would not manifest themselves in the CTE behavior) or accompanied with the closure of pores aligned with the basal planes. Kelly and Foreman53 report that their proposed creep mechanism would be expected to break down at high doses and temperatures, and thus deviations from the linear creep law (eqn [12]) are expected. They suggest that this is due to (1) incompatibility of crystal strains causing additional internal stress and an increasing crystal creep rate, (2) destruction of interstitial pins by diffusion of vacancies (thermal annealing of vacancies in addition to irradiation an- nealing), and (3) pore generation due to incompati- bility of crystal strains. It is likely that pore generation can manifest itself in two ways: (1) changes in CTE with creep strain – thus, pores aligned parallel to the crystallite basal planes are affected by creep strain – and (2) at high doses, pore generation or perhaps pore reorientation, under the influence of applied and internal stress that must be accounted for in the prediction of high neutron dose creep behavior. Brocklehurst and Brown62 report on the annealing behavior of specimens that had been subjected to irradiation under constant stress and sustained up to 1% creep strain. They observed that the increase in creep strain with dose was identical in compression and tension up to 1% creep strain, and that the CTE was significantly affected in opposite directions by compressive and tensile creep strains. Irradiation annealing of the crept specimens caused only a small recovery in the creep strain, and therefore provided no evidence for a back stress in the creep process, which has implications for the in-crystal creep mech- anism. Thermal annealing also produced a small recovery of the creep strain at temperatures below 1600 �C, presumably because of the thermal removal of the irradiation-induced defects responsible for dislocation pinning. Higher temperature annealing Radiation Effects in Graphite 323 produced a further substantial recovery of creep strain. Most significantly, Brocklehurst and Brown62 reported the complete annealing of the creep induced changes in CTE, in contrast to the total creep strain, where a large fraction of the total creep strain is irrecoverable and has no effect on the thermal expan- sion coefficient. Brocklehurst and Brown62 discuss two interpretations of their results, but report that neither is satisfactory. One interpretation requires a distinction between changes in porosity that affect the CTE and changes in porosity affecting the elastic deformation under external loads, that is, two distinct modes of pore structure changes due to creep in broad agreement with the mechanism discussed earlier. The modified Simmons model29,30,67 for dimen- sional changes (eqn [18]) and that for dimensional changes of a crept specimen (eqn [20]) both have pore generation terms which are currently neglected. It now appears necessary to modify the current Kelly–Burchell creep model (eqn [25]) to account for this effect of creep strain on this phenomena; that is, we need to evaluate and take account of the terms Fx and F 0 x as well as include the term (F 0 x –Fx) in eqn [25]. Such a term should account for pore gener- ation and/or reorientation caused by fracture when incompatibilities in crystallite strains become exces- sive.71 Clearly, further work is needed in the area of irradiation-induced creep of graphite. 4.10.7 Outlook For more than 60 years, nuclear graphite behavior has been the subject of research and development in support of graphite-moderated reactor design and operations. The materials physics and chemistry, as well as the behavior of nuclear graphite under neutron irradiation are well characterized and understood, although new high-resolution characterization tools, such as HRTEM and STM, and other nanoscale characterization techniques, coupled with powerful computer based simulations of crystal deformation and displacement damage, are yielding new insights to the deformation mechanisms that occur in graphite throughout its life in the reactor core. Perhaps the biggest remaining challenge is to gain a fuller understanding of irradiation-induced dimensional change and irradiation creep in graphite. Currently, new creep irradiation experiments are underway at ORNL in the High Flux Isotope Reactor, and at Idaho National Laboratory in the Advanced Test Reactor. Studies of pore structure change from unirradiated reference samples, irradiated unstressed samples (controls), and irradiated stressed samples (crept samples), may advance our understanding of pore generation. Work in other countries is directed at reviewing existing creep data and assessing the observable graphite dimensional changes and creep strain in currently operating reactors. A recent Coor- dinated Research Project initiated by the Interna- tional Atomic Energy Agency (IAEA) has the goal of bringing these various strands of research together to form a single unified theory of irradiation-induced creep deformation in graphites. The knowledge gained through these many years of work, and 50 years of graphite-moderated reactor operating experience is currently being used to under- write the safety cases of graphite reactors through out the world. In the first part of the twenty-first century, more knowledge will be gained from the new graphite- moderated reactors in Japan and China that operate at higher temperatures. Several nations (within the Generation IV International Forum) are pursuing high-temperature, graphite-moderated, gas-cooled reactor projects with the goal of developing versatile and inherently safe reactor systems that can efficiently deliver both process heat and electricity. With the realization and acceptance that green- house gas emissions from fossil fueled power plants are causing global climate changes, as evidenced by the Kyoto and recent Copenhagen Agreements, the nuclear option may once again become attractive for clean electric power generation. At that time, it is to be hoped that inherently safe, graphite-moderated, gas-cooled reactors may find renewed popularity. 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Irradiation Induced Creep in Graphite at High Temperature and Dose – A Revised Model; ORNL/TM-2008/098; Oak Ridge National Laboratory: Oak Ridge, TN, Feb 2009. http://juwel.fz-juelich.de 4.11 Graphite in Gas-Cooled Reactors B. J. Marsden and G. N. Hall The University of Manchester, Manchester, UK � 2012 Elsevier Ltd. All rights reserved. 4.11.1 Introduction 327 4.11.2 Graphite Crystal Structures 327 4.11.2.1 Graphite Crystal Atomic Structure and Properties 327 4.11.2.2 Coefficient of Thermal Expansion 328 4.11.2.3 Modulus 328 4.11.2.4 Thermal Conductivity 329 4.11.2.5 Microcracking (Mrozowski Cracks) 329 4.11.3 Artificial Nuclear Graphite 329 4.11.3.1 Microstructure/Property Relationships 330 4.11.4 Graphite Core Fast Neutron Fluence, Energy Deposition, and Temperatures 332 4.11.5 Dosimetry (Graphite Damage Dose or Fluence) 333 4.11.5.1 Early Activation Measurements on Foils 334 4.11.5.2 Reactor Design and Assessment Methodology: Fuel Burnup 335 4.11.5.2.1 Calder effective dose 335 4.11.5.3 Equivalent Nickel Flux 336 4.11.5.4 Integrated Flux and Displacements per Atom 336 4.11.5.4.1 DIDO equivalent flux 337 4.11.5.5 Energy Above 0.18MeV 338 4.11.5.6 Equivalent Fission Flux (IAEA) 339 4.11.5.7 Fluence Conversion Factors 339 4.11.5.8 Irradiation Annealing and EDT 339 4.11.5.9 Summary of Fast Neutron Dose (Fluence) 339 4.11.6 Graphite ‘Energy Deposition’ (Nuclear Heating) 340 4.11.6.1 The Use of Titanium for Installed Sample Holders 341 4.11.7 Radiolytic Oxidation 341 4.11.7.1 Introduction 341 4.11.7.2 Ionizing Radiation 341 4.11.7.2.1 Energy deposition 341 4.11.7.3 Radiolytic Oxidation Mechanism 341 4.11.7.4 Inhibition 342 4.11.7.5 Internal Porosity 342 4.11.7.6 Prediction of Weight Loss in Graphite Components 343 4.11.7.7 Weight Loss Prediction in Inhibited Coolant 343 4.11.8 Graphite Temperatures 346 4.11.9 Variation of Fluence, Temperature, and Weight Loss in a Reactor Core 346 4.11.9.1 Fuel End Effects 347 4.11.9.2 Temperature and Weight Loss 347 4.11.10 Distribution of Fluence Within an Individual Moderator Brick 347 4.11.11 Fast Neutron Damage in Graphite Crystal Structures 348 4.11.11.1 Stored Energy 348 4.11.11.2 Crystal Dimensional Change 352 4.11.11.3 Coefficient of Thermal Expansion 353 4.11.11.4 Modulus 354 4.11.11.5 Thermal Conductivity 354 4.11.11.6 Raman 354 4.11.12 Property Changes in Irradiated Polycrystalline Graphite 355 325 326 Graphite in Gas-Cooled Reactors 4.11.13 Averaging Relationships 357 4.11.14 Dimensional Change 359 4.11.14.1 Pile Grade A 360 4.11.14.2 Gilsocarbon 363 4.11.14.3 Effect of Radiolytic Oxidation on Dimensional Change 364 4.11.14.4 Dimensional Change Rate 366 4.11.15 Coefficient of Thermal Expansion 366 4.11.15.1 Pile Grade A 366 4.11.15.2 Gilsocarbon 367 4.11.15.3 Methodology for Converting Between Temperature Ranges 367 4.11.15.4 Effect of Radiolytic Oxidation on CTE 368 4.11.16 Thermal Conductivity 368 4.11.16.1 Pile Grade A 369 4.11.16.2 Gilsocarbon 369 4.11.16.3 Thermal Conductivity Temperature Dependence of Irradiated Graphite 369 4.11.16.4 Predicting the Thermal Conductivity of Irradiated Graphite for Reactor Core Assessments 370 4.11.17 Young’s Modulus 371 4.11.17.1 Relationship Between Static and Dynamic Young’s Modulus 372 4.11.17.2 Pile Grade A 372 4.11.17.3 Gilsocarbon 372 4.11.17.4 Separation of Structure and Pinning Terms 374 4.11.17.5 Effect of Radiolytic Weight Loss on Dimensional Change and Young’s Modulus 374 4.11.17.6 Small Specimen Strength 375 4.11.18 Effect of Radiolytic Oxidation on Thermal Conductivity, Young’s Modulus, and Strength 375 4.11.19 The Use of the Product Rule 375 4.11.20 Irradiation Creep in Nuclear Graphite 376 4.11.20.1 Dimensional Change and Irradiation Creep Under Load 378 4.11.20.2 Types of Irradiation Creep Experiments 378 4.11.20.3 The UKAEA Creep Law 378 4.11.20.4 Observed Changes to Other Properties 380 4.11.20.4.1 Coefficient of thermal expansion 380 4.11.20.4.2 Young’s modulus 381 4.11.20.5 Lateral Changes 381 4.11.20.6 Creep Models and Theories 381 4.11.20.6.1 UKAEA creep law 383 4.11.20.6.2 German and US creep model 385 4.11.20.6.3 Further modifications to the UKAEA creep law: interaction strain 385 4.11.20.6.4 Recent nuclear industry model 387 4.11.20.7 Final Thoughts on Irradiation Creep Mechanisms 387 4.11.21 Concluding Remark 387 References 388 Abbreviations AG Against grain AGR Advanced gas-cooled reactor BAF Bacon anisotropy factor BEPO British Experimental Pile Zero CPV Closed pore volume CTE Coefficient of thermal expansion DFR Dounreay Fast Reactor DSC Differential scanning calorimeter DYM Dynamic Young’s modulus EDND Equivalent DIDO nickel dose EDNF Equivalent DIDO nickel flux Graphite in Gas-Cooled Reactors 327 EDT Equivalent DIDO temperature FWHM Full-width, half-maximum HFR High Flux Reactor HOPG Highly oriented pyrolytic graphite HRTEM High-resolution transmission electron microscopy HTR High temperature reactor IAEA International Atomic Energy Agency MTR Materials test reactor NDT Nondestructive testing OPV Open pore volume PGA Pile Grade A RBMK Reaktor Bolshoy Moshchnosti Kanalniy (there are other quoted translations) RPV Reactive pore volume SEM Scanning electron microscopy SYM Static Young’s modulus TEM Transmission electron microscopy TPV Total pore volume UKAEA United Kingdom Atomic Energy Authority WG With grain 4.11.1 Introduction Nuclear graphite has, and still continues, to act as a major component in many reactor systems. In prac- tice, nuclear graphite not only acts as a moderator but also provides major structural support which, in many cases, is expected to last the life of the reactor. The main texts on the topic were written in the 1960s and 1970s by Delle et al.,1 Nightingale,2 Reynolds,3 Simmons,4 in German, and Pacault5 Tome I and II, in French with more recent reviews on works by Kelly6,7 and Burchell.8 This text is mainly on the basis of the UK graphite reactor research and operating experi- ence, but it draws on international research where necessary. During reactor operation, fast neutron irradiation, and in the case of carbon dioxide-cooled systems radiolytic oxidation, significantly changes the graph- ite component’s dimensions and properties. These changes lead to the generation of significant graphite component shrinkage and thermal stresses. Fortu- nately, graphite also exhibits ‘irradiation creep’ which acts to relieve these stresses ensuring, with the aid of good design practice, the structural integ- rity of the reactor graphite core for many years. In order to achieve the optimum core design, it is important that the engineer has a fundamental under- standing of the influence of irradiation on graphite dimensional stability and material property changes. This chapter aims to address that need by explaining the influence of microstructure on the properties of nuclear graphite and how irradiation-induced changes to that microstructure influence the behavior of graph- ite components in reactor. Nuclear graphite is manu- factured from coke, usually a by-product of the oil or coal industry. (Some cokes are a by-product of refining naturally occurring pitch such as Gilsonite.9) Thus, nuclear graphite is a porous, polycrystalline, artificially produced material, the properties of which are de- pendent on the selection of raw materials and man- ufacturing route. In this chapter, the properties of the graphite crystal structures that make up the bulk poly- crystalline graphite product are first described and then the various methods of manufacture and resultant properties of the many grades of artificial nuclear graphite are discussed.This is followed bya description of the irradiation damage to the crystal structure, and hence the polycrystalline structure, and the implica- tion of graphite behavior. The influence of radiolytic oxidation on component behavior is also discussed as this is of interest to operators or designers of graphite- moderated, carbon dioxide-cooled reactors, many of which are still operating. 4.11.2 Graphite Crystal Structures The properties and irradiation-induced changes in graphite crystals have been studied using both ‘natu- rally occurring’ graphite crystals and an artificial product referred to as highly orientated pyrolytic graphite (HOPG), formed by depositing a carbon substrate using hydrocarbon gas6 followed by com- pression annealing at around 3000 �C. HOPG is con- sidered to be the most appropriate ‘model’ material that can be used to study the behavior of artificially produced polycrystalline nuclear graphite. It has a density value near to that of a perfect graphite crystal structure, but perhaps more appropriately, it has imperfections similar to those found in the struc- tures that make up artificial polycrystalline graphite. A detailed description of the properties of graphite can be found in Chapter 2.10, Graphite: Properties and Characteristics. 4.11.2.1 Graphite Crystal Atomic Structure and Properties In this section, the atomic structure of graphite crys- tal structures is discussed briefly, along with some of the properties relevant to the understanding of the 328 Graphite in Gas-Cooled Reactors irradiation behavior of graphite. Graphite can be arranged in an ABAB stacking arrangement termed hexagonal graphite (see Figure 1). This is the most thermodynamically stable form of graphite and has a density of 2.266 g cm�3. The a-spacing is 1.415 Å and the c-spacing is 3.35 Å. However, in both natural and artificial graphite stacking faults and dislocations abound.10 4.11.2.2 Coefficient of Thermal Expansion The coefficient of thermal expansion (CTE) as measured for natural graphite and HOPG is temper- ature dependent (Figure 2) and the data from a number of authors has been collated by Kelly.6 The room temperature values of CTE are about 27.5� 10�6 K�1 and �1.5� 10�6 K�1 in the ‘c’ and ‘a’ directions, respectively. a a c Upper layer (A) Lower layer (B) Figure 1 The crystalline structure of graphite. 0 500 2 1 0 C TE a a (1 0− 6 K −1 ) C TE a (1 0− 6 K −1 ) −1 −2 1000 Temperature (K) ‘c’ direction 1500 2000 2500 3000 Bailey and Yates Steward et al. Harrison Yates et al. Figure 2 Crystal coefficient of thermal expansion. Modified fro London, 1981. 4.11.2.3 Modulus The crystal elastic moduli6 are C11 (parallel to the basal planes)¼ 1060.0� 109Nm�2,C12¼ 180.0� 109Nm�2, C13¼ 15.0� 109Nm�2, C33 (perpendicular to the basal planes)¼ 34.6� 109Nm�2, and C44 (shear of the basal planes)¼ 4.5� 109Nm�2 as defined by the orthogonal co-ordinates given below: sxx syy szz tzx tzy txy 0 BBBBBBBBBB@ 1 CCCCCCCCCCA ¼ C11 C12 C13 0 0 0 C12 C11 C13 0 0 0 C13 C13 C33 0 0 0 0 0 0 C44 0 0 0 0 0 0 C44 0 0 0 0 0 0 12ðC11�C12Þ 0 BBBBBBBBBB@ 1 CCCCCCCCCCA exx eyy ezz ezx ezy exy 0 BBBBBBBBBB@ 1 CCCCCCCCCCA ½1� a c c 0 0 10 20 30 40 500 1000 Temperature (K) ‘a’ direction 1500 2000 2500 3000 Bailey and Yates Steward et al. Harrison Yates et al. Nelson and Riley m Kelly, B. T. Physics of Graphite; Applied Science: Graphite in Gas-Cooled Reactors 329 The strength of the crystallite is also directly related to the modulus, that is, the strength along the basal planes is higher than the strength perpendicular to the planes, and the shear strength between the basal panes is relatively weak. 4.11.2.4 Thermal Conductivity The thermal conductivity of graphite along the basal plane ‘a’ direction is much greater than the thermal conductivity in the direction perpendicular to the basal plane ‘c.’ At the temperature of interest to the nuclear reactor engineer, graphite thermal con- duction is due to phonon transport. Increasing the temperature leads to phonon–phonon or Umklapp scattering (German for turn over/down). Imperfections in the lattice will lead to scattering at the boundaries. 4.11.2.5 Microcracking (Mrozowski Cracks) During the manufacture of artificial graphite, very high temperatures (2800–3000 �C) are required in the graphitization process. On cooling from these high temperatures, thermoplastic deformation is pos- sible until a temperature of �1800 �C is reached. Below this temperature, the large difference in ther- mal expansion coefficients between the ‘c’ and ‘a’ directions leads to the formation of long, thin micro- cracks parallel to the basal planes, often referred to as ‘Mrozowski’ cracks.11 These types of cracks are even observed in HOPG (Figure 3). The high density of HOPG when compared to the large number of microcracks, a few nanometers in (a) 1µm Figure 3 Transmission electron microscopic images of highly plane, ‘c’ direction, of HOPG (reproduced from Kelly, B. T. MSc (b) Mrozowski cracks in HOPG as seen along the ‘basal’ planes, width and many micrometers in length (as seen in Figure 3(b)), appears to be counterintuitive and has led to speculation that these microcracks may contain some low-density carbonaceous structure. The pres- ence of these microcracks is very important in under- standing the properties of nuclear graphite as they provide accommodation for thermal or irradiation- induced crystal expansion in the ‘c’ direction. Therefore, two crystal structures are of interest; the ideal, ‘perfect’ structure and the nonperfect struc- tures as may be defined with reference to HOPG. It is of the latter that many of the crystal behaviors and properties have been studied. Definition: In this chapter on nuclear graphite, ‘crystal’ refers to the perfect crystal structure and ‘crystallite’ refers to the nonperfect crystal struc- tures containing Mrozowski-type microcracks (and nanocracks). 4.11.3 Artificial Nuclear Graphite The reactor designer requires a high-density, very pure graphite, with a high scattering cross-section, a low absorption cross-section, and good thermal and mechanical properties, both in the unirradiated and irradiated condition. The purity is important to ensure not only a low absorption cross-section but also that during operation the radioactivity of the graphite remains as low as possible for waste disposal purposes. Artificial graphite is manufactured from coke obtained either from the petroleum or coal industry, or in some special cases (such as Gilsocarbon, a UK grade of graphite) from a ‘graphitizable’ coke derived (b) 1µm orientated pyrolytic graphite. (a) View into the ‘basal’ thesis, University of Cardiff, Cardiff, Wales, 1966) and ‘a’ direction. Courtesy of A. Jones, University of Manchester. 330 Graphite in Gas-Cooled Reactors from naturally occurring pitch deposits.9 The raw coke is first calcined to remove volatiles and then ground or crushed for uniformity, before being blended and mixed with a pitch binder. (Crushed ‘scrap’ artifi- cial graphite may be added to help with heat removal during the subsequent baking. For nuclear graphite, this should be of the same grade as the final product.) This mixture is then formed into blocks using one of various techniques such as extrusion, pressing, hydro- static molding, or vibration molding, to produce the so-called ‘green article.’ The ‘green’ blocks are then put into large ‘pit’ or ‘intermittent’ gas or oil-fired furnaces. The blocks are usually arranged in staggers, covered by a metallurgic coke, and baked at around 800 �C in a cycle lasting about 1month to produce carbon blocks. These carbon blocks may be used for various industrial purposes such as blast furnace liners; it has even been used for neutron shielding in some nuclear reactors. (Care must be taken as the carbon blocks are not as pure as graphite and may lead to waste disposal issues at the end of the reactor life.) To improve the properties of the graphite pro- duced from the carbon block, the carbon block is often impregnated with a low-density pitch under vacuum in an autoclave. To facilitate the entry of the pitch into the body of the block, the block surface may be broken by grinding. After impregnation the blocks are then rebaked. This process of impregna- tion and rebaking may be repeated 2, 3, or 4 times. However, the improvement in the properties by this process is subject to diminishing rewards. The next process is graphitization at about 2800– 3000 �C by passing a large electrical current at low voltage through the blocks either in an ‘Acheson furnace’ or using an ‘in-line furnace.’ In both cases, the blocks are covered by a metallurgical coke to prevent oxidation. This graphitization cycle may take about 1month. If necessary, there may be a final purification step. This involves heating the graphite blocks to around 2400 �C in a halogen gas atmosphere to remove impurities. The final product can then be machined into the many intricate com- ponents required in a nuclear reactor. For quality assurance purposes, during manufac- ture the blocks are numbered at an early stage and this number follows the block through the manu- facturing process. This is clearly an expensive manufacturing process and therefore, at each stage, quality control is very important. Many samples will be taken from the blocks to ensure that the final batch (or heat) is of appropriate quality compared to previ- ous heats. It is important that the reactor operators retain this data in electronic form as it may be required to investigate any anomalous behavior as the reactor ages. Samples of ‘virgin’ unirradiated graphite blocks should also be retained for future reference. Records should include information on the batch or heat, property measurements, nonde- structive testing (NDT) results, and measurements of impurities. It is not enough just to have the ‘ash’ content after incineration and the ‘boron equivalent’ as some impurities, such as nitrogen, chlorine, and cobalt, will cause significant issues related to reactor operation and final waste disposal. It is important that the reactor operator takes responsibility for these measurements as in the past it has been found that reactor designers and graphite manufacturers close down or merge, and records are lost. Final inspection will uncover issues related to damage, imperfection, quality, etc. Therefore, a ‘con- cessions’ policy is required to determine what is acceptable and where such components can be used in reactor. Again, the reactor operator will require an electronic record of these concessions. 4.11.3.1 Microstructure/Property Relationships The microstructure of a typical nuclear graphite is described with reference to Gilsocarbon. This product was manufactured from coke obtained from a naturally occurring pitch found at Bonanza in Utah in the United States. To understand the microstructural properties, one has to start with the raw coke. The structure of Gilsonite coke is made of spherical parti- cles about 1mm in diameter as shown in Figure 4. This structure is retained throughout manufacture and into the final product. In Figure 4(b), the spher- ical shaped cracks following the contours of the spherical particles are clearly visible. This coke will be carefully crushed in order to keep the spherical structures that form the filler particles and help to give Gilsocarbon its (semi-) isotropic properties. At a larger magnification in a scanning electron microscopy (SEM), the complexity of these cracks is clearly visible, Figure 4(c), and at an even larger magnification, a ‘swirling structure’ made up of graphite platelets stacked together is discernable between the cracks. In essence, the whole structure contains a significant amount of porosity. After graphitization, the Gilsonite coke filler par- ticles are still recognizable (Figure 5(a) and 5(b)). From the polarizing colors, one can see that the main ‘a’ axis orientation of the crystallites follows the (c) (a) (b) (d) Figure 4 Gilsonite raw-coke microstructure. (a) Photograph of Gilsonite coke, (b) Scanning electron microscopy (SEM) image of polished Gilsonite coke, (c) detail in an SEM image showing the region around cracks that follow the spherical shape of the coke particles, and (d) a higher magnification SEM image showing the intricate, random arrangement of platelets. Courtesy of W. Bodel, University of Manchester. (a) (b) (c) (d) 500µm 200µm Figure 5 Polarized optical and scanning electron microscopic images of Gilsocarbon graphite. (a) Optical image, (b) optical image, (c) SEM image, (d) SEM image. Courtesy of A. Jones, University of Manchester. Graphite in Gas-Cooled Reactors 331 332 Graphite in Gas-Cooled Reactors spherical particles circumferentially, as does the ori- entation of the large calcination cracks. The crystal- lite structures in the binder phase are much more randomly oriented, and this phase contains signifi- cant amounts of gas-generated porosity. There are also what appear to be broken pieces of Gilsonite filler particles contained within the binder phase. The bulk properties of polycrystalline nuclear graphite strongly depend on the structure, distribu- tion, and orientation of the filler particles.12 The spherical Gilsonite particles and molding technique give Gilsocarbon graphite semi-isotropic properties, whereas in the case of graphite grades such as the UK pile grade A (PGA), the extrusion process used dur- ing manufacture tends to align the ‘needle’ type coke particles. Thus, the crystallite basal planes that make up the filler particles tend to align preferentially, with the ‘c’ axis parallel to the extrusion direction and the ‘a’ axis perpendicular to the extrusion direc- tion. The long microcracks are also aligned in the extrusion direction. The terms ‘with grain (WG)’ and ‘against grain (AG)’ are used to describe this phe- nomenon, that is, WG is equivalent to the parallel direction and AG is equivalent to the perpendicular direction. Thus, the highly anisotropic properties of the crystallite are reflected in the bulk properties of polycrystalline graphite (Table 1). A graphite anisotropy ratio is usually defined by the AG/WG ratio of CTE values. For needle coke graphite, this ratio can be two or more, while for a more randomly orientated structure, values in the region of 1.05 can be achieved by careful selection of material and extrusion settings. A more scientific way of defining anisotropy ratio is by use of the Bacon anisotropy factor (BAF).13 Other forming methods are usually used to pro- duce isotropic graphite grades such as the Gilsocar- bon grade described above. In this case, it was found that Gilsocarbon graphite produced by extrusion was not isotropic enough to meet the advanced gas- cooled reactor (AGR) specifications. Therefore, a Table 1 Relative properties–grain direction relationships Property With grain (WG) Against grain (AG) Coefficient of thermal expansion (CTE) Lower Higher Young’s modulus Higher Lower Strength Higher Lower Thermal conductivity Higher Lower Electrical resistivity Lower Higher ‘molding’ method where the blocks were formed by pressing in two directions was used. This had the effect of slightly aligning the grains such that the AG direction was parallel to the pressing direction and the WG was perpendicular to the pressing direc- tion. However, Gilsocarbon has proved to be one of the most isotropic graphite grades ever produced, even in its irradiated condition. Another approach is to choose an ‘isotropic coke’ crushed into fine particles and then produce blocks using ‘isostatic molding’ process. The isostatic mold- ing method involves loading the fine-grained coke binder mixture into a rubber bag which is then put under pressure in a water bath. In this way, high quality graphite can be produced mainly for use for specialist industries such as the production of elec- tronic components. This type of graphite (such as IG- 110 and IG-11) has been used for high-temperature reactor (HTR) moderator blocks, fuel matrix, and reflector blocks in both Japan and China. However, even these grades exhibit slight anisotropy. The final polycrystalline product contains many long ‘thin’ (and not so ‘thin’) microcracks within the crystallite structures that make up the coke particles. Similar, but much smaller, cracked structures are to be found in the ‘crushed filler flour’ used in the binder, and in well-graphitized parts of the binder itself. It is these microcracks that are responsible for the excellent thermal shock resistance of artificial polycrystalline graphite. They also provide ‘accom- modation,’ which further modifies the response of bulk properties to the crystal behavior in both the unirradiated and irradiated polycrystalline graphite. Typical properties of several nuclear graphite grades are given inTable 2. One can see that polycrystalline graphite has about 20% porosity by comparing the bulk density with the theoretical density for graphite crystals (2.26 g cm�3). About 10% of this is open porosity, the other 10% being closed. 4.11.4 Graphite Core Fast Neutron Fluence, Energy Deposition, and Temperatures Since the late 1940s, many journal papers, conference papers, and reports have been published on the change in properties in graphite due to fast neutron damage. Many different units have been used to define graphite damage dose (or fluence). It is impor- tant to understand the basis of these units because historic data are still being used to justify models Table 2 Typical properties of several well-known grades of nuclear graphite Property PGA CSF Gilsocarbon IG-110 H451 Production method Extruded Extruded Press-molded Iso-molded Extruded Direction WG AG WG AG WG AG WG AG WG AG Density (g cm�3) 1.74 1.66 1.81 1.77 1.76 Thermal conductivity (Wm�1 K) 200 109 155 97 131 116 158 137 CTE, 20–120 �C (10�6 K�1) 0.9 2.8 1.2 3.1 4.3 CTE, 350–450 �C (10�6 K�1) 4.5 CTE, 500 �C (10�6 K�1) 1.5 3.5 3.6 4.0 4.4 5.1 Young’s modulus (GPa) 11.7 5.4 8.0 4.8 10.9 9.8 8.51 7.38 Poisson’s ratio �0.07 0.21 0.14 0.15 Strength, tensile (MPa) 17 11 17.5 24.5 15.2 13.7 Strength, flexural (MPa) 19 12 23.0 39.2 Strength, compressive (MPa) 27 27 70.0 78.5 55.3 52.7 Graphite in Gas-Cooled Reactors 333 used in assessments for component behavior in reac- tors. Indeed, some of these historic data, for example, stored energy and strength, will also be used to sup- port decommissioning safety assessments. Early estimations of ‘graphite damage’ were based on the activation of metallic foils such as cobalt, cadmium, and nickel. Later, to account for damage in different reactors, equivalent units, such as BEPO or DIDO equivalent dose, were used where the dam- age is referred to damage at a standard position in the BEPO, Calder Hall, or DIDO reactors. The designers of plutonium production reactors preferred to use a more practical unit related to fuel burnup (megawatts per adjacent tonne of uranium, MW/Atu). Research- ers also found that the calculation of a flux unit, based on an integral of energies above a certain value, was relatively invariant to the reactor system and used the unit En> 0.18MeV and other variants of this. Today, the favored option is to calculate the flu- ence using a reactor physics code to calculate the displacements per atom (dpa). However, in the field of nuclear graphite technology historic units are still widely used in the literature. For example, reactor operators have access to individual channel burnup which, with the aid of axial ‘form factors,’ can be used to give a measure of average damage along the indi- vidual channel length. Fortunately, most, but not all, of these units can be related by simple conversion factors. However, care must be taken; for example, the unit of megawatt days per tonne of uranium (MWd t�1) is not necessarily equivalent in different reactor systems. When assessing the analysis of a particular com- ponent in a reactor, one must be aware that a single detailed calculation of a peak rated component in the center of the core may have been carried out to give spatial, and maybe temporal, distribution of that component’s fluence (and possibly temperature and weight loss). These profiles may have then been extrapolated to all of the other components in the core using the core axial and radial ‘form factors.’ In doing this, some uncertainty will be introduced and clearly, some checks and balances will be required to check the validity of such an approach. 4.11.5 Dosimetry (Graphite Damage Dose or Fluence) In a nuclear reactor, high energy, fast neutron flux leads to the displacement of carbon atoms in the graphite crystallites via a ‘cascade.’ Many of these atoms will find vacant positions, while others will form small interstitial clusters that may diffuse to form larger clusters (loops in the case of graphite) depending upon the irradiation temperature. Con- versely, vacancy loops will be formed causing the lattice structure to collapse. These vacancy loops will only become mobile at relatively high tempera- tures. The production of transmutation gas from impurities is not an issue for highly pure nuclear graphite, as the quantities of gas involved will be negligible and the graphitic structure is porous. The change in graphite properties is a function of the displacement of carbon atoms. The nature and amount of damage to graphite depends on the partic- ular reactor flux spectrum, which is dependent on the reactor design and position, as illustrated in Figure 6. It is impractical to relate a spectrum of neutron energies to a dimensional or property change at a 1800 1600 1400 1200 1000 800 Fl ux p er u ni t le th ar gy f (v ) 600 400 200 0 10 100 1000 Energy (keV) 10000 100000 TE rig in BEPO Hollow fuel element in BEPO Empty fuel channel in BEPO Empty lattice position in PLUTO Hollow fuel element in PLUTO Figure 6 Flux spectrums for various reactor positions used in graphite irradiation programs. Modified from Simmons, J. Radiation Damage in Graphite; Pergamon: London, 1965. Table 3 Relationship for BEPO equivalent flux (thermal) at a central lattice position to other positions in BEPO and 334 Graphite in Gas-Cooled Reactors single point in a material such as graphite. Therefore, an ‘integrated flux’ is used and is discussed later. other irradiation facilities Position Factor BEPO lattice 1 BEPO hollow slug 2.27 BEPO empty fuel channel 0.63 Windscale Piles 1.29 Windscale Piles thermostats 1.29 NRX fast neutron plug MWHs 1.54�1015 American data MWd/CT 5.5�1017 Source: Simmons, J. H. W. The Effects of Irradiation on Graphite; AERE R R 1954; Atomic Energy Research Establishment, 1956. 4.11.5.1 Early Activation Measurements on Foils Although one cannot directly measure the damage to graphite itself, it is possible to measure the activation of another material, because of nuclear impacts adja- cent to the position of interest. This activation may then be related to changes in graphite properties. This was done in early experiments using cobalt foils and by measuring the activation arising from the 59Co(n,g)60Co reaction. This reaction has a cross-section of 38 barns and 60Co has a half-life of 5.72 years, which need to be accounted for in the fluence calculations. Such foils were included in graphite experiments in BEPO and the Windscale Piles, and are still used today for irradiation rig validation and calibration purposes. In these early experiments, after removal from the reactor, cobalt foils were dissolved in acid, diluted, and the decay rate measured. A measure of fluence could then be calculated from knowledge of the following: � the solution concentration � the time in the reactor � the decay rate � the activation cross-section Unfortunately the 59Co(n,g)60Co reaction is mainly a measure of thermal flux and atomic displa- cements in graphite are due to fast neutrons. An improvement was the use of cobalt/cadmium foils, but this was not really satisfactory. Measurements made in this way are often given the unit, neutron velocity time (nvt). Table 3 gives an example of thermal flux deter- mined from cobalt foils defined at a standard posi- tion in the center of a lattice cell in BEPO. Graphite damage at other positions in other reac- tors could then be related to the standard position in BEPO. Graphite in Gas-Cooled Reactors 335 4.11.5.2 Reactor Design and Assessment Methodology: Fuel Burnup When designing a nuclear reactor core, a channel ‘rating’ can be related to the reactor power andweight of uranium in a particular channel. This channel ‘rating’ can be related to a rate of change in the graphite properties. The channel rating is given in MW/Atu and over time the channel burnup as mega- watt days per adjacent tonne of uranium (MWd/Atu). Note that this unit is literally the (power in a particular channel)� (number of days)� (weight of uranium in that channel). No account is taken of refueling. However, fuel burnup is a function of reactor design and therefore, the equivalence concept was used and damage was related to a standard position. In the United Kingdom, the change in graphite property was defined at a standard position in a Calder Hall reactor to give Calder equivalent dose. This was defined as the dose at a position on the wall of a fuel channel in a Calder Hall reactor. In the Calder Hall design, the lattice pitch is 8 in. The stan- dard position was chosen to be in a 3.55-in.-diameter fuel channel at a point on the shortest line between the centers of two fuel channels. The fuel is assumed to be 1.15-in.-diameter natural uranium metal rods. Calder equivalent dose was then used as a function to relate graphite property change to fuel burnup. Kinchin14 had measured the change in graphite electric resistivity as a function of distance into the BEPO reflector. By normalizing this change, he defined a ‘graphite damage function’; see also Bell et al.15 Thus, graphite damage at some position in a reac- tor core graphite component could be defined as a function of the following: � source strength � distance between position and source � attenuation in damage with distance through the intervening graphite The damage function is a measure of the last two bullet points. The source strength is related to fuel burnup. The graphite damage function df is defined as df ¼ fðRgÞ R ½2� where f(Rg) is the damage absorption curve for an equivalent distance through BEPO graphite ‘Rg’ of density 1.6 g cm�3, and R is the distance through graphite between the source and position of interest. Note that nonattenuating geometric features, that is, holes, need to be accounted for. Calder equivalent rating Pe can now be defined as Pe ¼ AdfP ACalder dfCalder ½3� where ‘ACalder’ and ‘A ’ are the uranium fuel cross- sectional areas in Calder (1.04 in.2) and in the reactor under consideration, respectively; ‘dfCalder’ and ‘df ’ are the values of the damage at the Calder standard position (1.395) and in the reactor under consider- ation, respectively, and ‘P’ is the fuel rating in the reactor under consideration. Thus, a graphite property change in a reactor under assessment can be related to the equivalent graphite property change at the Calder standard position. However, in a real reactor there is more than one fuel channel. There may also be absorbers or empty interstitial holes, the fuel rating will change with burnup, and the fuel will be replaced from time to time. Therefore, a more complex, multiple source cal- culation is required to take account of the actual chan- nel rating and the geometric features of the core. This is normally done by considering a 5� 5 lattice array: df ¼ X i Bi B fðRgÞi Ri � � ½4� where Bi and B are the accumulated fuel burnup at the ith and reference source, respectively, f(Rg)i is the damage absorption function corresponding to thickness Rg for ith source, and Rg is the distance between the ith source and target. This method was successfully used to design the Magnox reactors. However, because of the higher enriched oxide fuel and more complex fuel design in the AGRs, this approach became less satisfactory and new ‘damage functions’ that accounted for the new fuel and geometry were calculated using Monte Carlo methods (made possible by the introduction of the digital computer). This method was until recently still used in industry codes such as ‘Fairy’ (National Nuclear Company) and ‘GRAFDAM’ (UKAEA). 4.11.5.2.1 Calder effective dose When only low-dose irradiation graphite property data were available, it was assumed that irradiation damage could be obtained at one temperature and that the property change versus dose (fluence) curves could be adjusted for all other temperatures using the so called R(y) curve: Calder effective dose ¼ calder equivalent dose � RðyÞ ½5� 336 Graphite in Gas-Cooled Reactors However, the use of R(y) is valid only for very low fluence and it should no longer be used, although one may come across its use in historic papers. 4.11.5.3 Equivalent Nickel Flux Nickel foils were used to give a measure of the damage to graphite through the 58Ni(n,p)58Co reac- tion. This reaction has a mean cross-section of 0.107 barns and 58Co has a half-life of 71.5 days. The change in graphite thermal resistivity was measured in the TE10 experimental hole in BEPO and the nickel flux was also measured at the same position. It was assumed that the graphite displace- ment rate fd was equal to the nickel flux fNi at this position. For comparison, the change in graphite thermal resistivity was then measured at various other positions in BEPO, as given in Table 4 Later, the same exercise was repeated in PLUTO, the sister reactor to DIDO at Harwell, and the ratio compared to that at other positions. In this case, the ratio appeared to be largely invariant to position. Table 5 gives a few examples of the many measure- ments made.16 It was decided that the activity produced in nickel fNi could be related to the graphite damage rate by a Table 4 Ratio of graphite damage to nickel flux as measured in BEPO Position Ratio bfd/fNi Experimental hole TE10 1.0 (definition) Hollow fuel element 0.43 Empty fuel channel (at three positions) 1.0 Experimental hole E2/7 0.75 Modified from Bell, J.; Bridge, H.; Cottrell, A.; Greenough, G.; Reynolds, W.; Simmons, J. Philos. Trans. R. Soc. Lond. A Math. Phys. Sci. 1962, 254(1043), 361–395. b is a proportionality factor. Table 5 Ratio of graphite damage to nickel flux as measured in PLUTO Position Ratio fd/fNi C4 – inside fuel element stainless steel thimble 0.518 D3 – inside fuel element stainless steel thimble 0.468 C4 – inside fuel element aluminum thimble 0.507 D4 – empty fuel element 0.564 Modified from Bell, J.; Bridge, H.; Cottrell, A.; Greenough, G.; Reynolds, W.; Simmons, J. Philos. Trans. R. Soc. Lond. A Math. Phys. Sci. 1962, 254(1043), 361–395. factor. However, care was still required with respect to the choice of reactor and irradiation location. Thus, a definition of damage based on a standard position in DIDO and a calculation route for equiva- lent DIDO nickel flux (EDNF) were devised. It should be noted that there are difficulties related to a standard based on measurements made with nickel foils and the 58Ni(n,p,)58Co reaction because of the short half-life of 58Co and the interfering effect of the 58Co(n,g)59Co reaction. A method by Bell et al.15 which went back to measuring activation of cobalt foils and the 59Co(n,g)60Co reaction, and then calcu- lating the ratio fNi/fCo, was used for a short while. This method used the following relationships: 115g fuel elements fNi=fCo ¼ 0:378� 0:504b 150g fuel elements fNi=fCo ¼ 0:502� 0:530b where ‘b’ is the fuel burnup. However, this was not very satisfactory and it was clear that a validated calculation route was desirable, and is now becoming practicable through development in computer technology. 4.11.5.4 Integrated Flux and Displacements per Atom The rate of change of a material property can be related to displacement rate of carbon atoms (dpa s�1). However, it is not possible to directly measure dpa s�1 in graphite, but dpa s�1 can be related to the reactor flux. The flux depends on reactor design, and varies with position in the reactor core. Neutron flux is a measure of the neutron popula- tion and speed in a reactor. In a reactor, neutrons move at a variety of speeds in randomly orientated directions. Neutron flux is defined as the product of the number of neutrons per unit volume moving at a given speed, as given by eqn [6] below. f number cm2 s � � ¼ n number cm3 � � v cm s � � ½6� However, as there is a spectrum of neutrons, with many velocities, this is not a useful unit for the material scientist. Therefore, integrated flux is used over a range of energies E1 to E2 as given by eqn [7]. f ¼ ðE2 E1 fðEÞdE ½7� In this way, a measure of neutron damage at any position within a structural component can be defined as follows. For a material such as graphite, Graphite in Gas-Cooled Reactors 337 the damaging power (displacement rate), fd, can be expressed as an integrated flux as given in eqn [8]. fd ¼ ð1 0 cðEÞfðEÞdE ½8� where f(E) is the neutron flux with energies from E to Eþ dE and C(E) is a function to describe the ability of neutrons to displace carbon atoms. 4.11.5.4.1 DIDO equivalent flux At the standard position in a hollow fuel element, the nickel flux, fs, can be defined by eqn [9]. fs ¼ 1 s ð1 0 fsðEÞsNiðEÞdE ½9� where Ð1 0 fsðEÞsNiðEÞdE is the integral of neutron flux multiplied by the nickel cross-section at the standard position in DIDO, and s0 is the average nickel cross-section for energies >1MeV, which is equal to 0.107 barn. The value of fs at this position is 4� 1013 n cm�2 s�1. The carbon displacement rate can be calculated using eqn [10]. fd ¼ ð ð f E1ð Þs E1; E2ð Þn E2ð ÞdE1dE2 ½10� where f(E1) is the flux of neutrons with energy E1, s(E1, E2) is the cross-section for a neutron with 10 000 Thompson-Wright Norgett, Robinson, and Torrens Kinchin-Pease (Lc= 25 keV) Kinchin-Pease (Lc= 12 keV) 1000 100 N um b er o f d is p la ce d a to m s 10 1 0.1 10 102 103 104 Ene Figure 7 Comparison of various damage function models tha primary knock-on atom. energy E1 to produce a recoil atom with energy E2, and v(E2) is a ‘damage function’ giving the number of atoms displaced from their lattice site by recoil energy E2. The carbon displacement rate, fds, at a standard position in DIDO is 5.25� 10�8 dpa. The derivation of the damage function (Figure 7) is on the basis of billiard ball mechanics, energy losses to the lattice due to impacts, and to forces associated with excitation of the lattice. The early Kinchin and Pease17 form of the dam- age function was found to underestimate damage in graphite. To give greater dpa, it was recommended that ‘Lc’ was artificially increased, but this was not satisfactory. The Thompson and Wright18 damage function was used in the official definition of EDNF. However, the Norgett et al.19 damage function is used in most modern reactor physics codes and it has been recently shown that there is little difference in the calculation of graphite damage using either of these latter two functions.20,21 It is assumed that the ratio of dpa to nickel flux (fds/fs) at the standard position, which is equal to 1313� 10�24 dpa (n cm�2 s�1)�1, can be equated to the same ratio fd/fNi in the reactor of interest as given by eqn [11]: fds fs � � ¼ fd fNi � � ¼ 1313� 10�24dpaðn cm�2s�1Þ�1 ½11� This value was derived using the Thompson and Wright damage function and an early flux spectrum 105 rgy (eV) 106 107 108 t describe the number of displaced atoms versus energy of Table 6 Comparison of calculated and measured graphite damage rates using the Thompson and Wright model Location Calculated Measured/ standard DIDO hollow fuel element 1.00 1.00 PLUTO empty lattice position 0.975 1.22 DR-3 empty lattice position 0.975 0.90 BR-2, Mol, hollow fuel element 1.00 0.90 HFR-Petten core 1.02 1.0 BEPO TE-10 hole 2.31 2.04 BEPO empty fuel channel 2.36 2.04 BEPO hollow fuel channel 0.98 0.87 Windscale AGR replaced fuel stringer B 2.70 2.28 Windscale AGR replaced fuel stringer D 2.71 2.03 Windscale AGR loop stringer 2.60 2.08 Windscale AGR loop control stringer 2.60 2.51 Windscale AGR fuel element – inner ring 1.18 1.06 Windscale AGR fuel element – outer ring 1.39 1.06 Calder x-hole 2.12 2.10 Dounreay fast reactor core 0.46 0.50 Modified fromMarsden, B. J. Irradiation damage in graphite due to fast neutrons in fission and fusion systems; IAEA, IAEA TECDOC- 1154; 2000. Table 7 Energies, cross-sections, and mean number of displacement for various particles Particles Energy (eV) Cross-section (cm2) Mean number of displacements per collision Electrons 1�106 14.5�10�24 1.6 2�106 15.0�10�24 1.9 3�106 15.5�10�24 2.3 4�106 16.0�10�24 2.5 Protons 1�106 7.8�10�21 4–5.5 5�106 1.56�10�21 4–5.5 10�106 7.8�10�21 4–6 20�106 3.9�10�21 4–6 Deuterons 1�106 1.56�10�20 4–5 5�106 3.12�10�21 4–6 10�106 1.6�10�21 4–6 20�106 7.8�10�22 4–6.5 a-Particles 1�106 1.25�10�19 4–5 5�106 2.5�10�20 4–6 10�106 1.25�10�20 4–6.5 20�106 6.25�10�21 4–6.5 Neutrons 103 4.7�10�24 2.83 104 4.7�10�24 28.3 105 4.6�10�24 280 106 2.5�10�24 480 107 1.4�10�24 500 Source: Simmons, J. Radiation Damage in Graphite; Pergamon: London, 1965. Table 8 Displacements (�10�21) per unit fluence for energies above E1 for various systems Spectrum E1¼0.067MeV E1¼ 0.18MeV PEGGI 0.719 0.738 ETR(N-8) 0.697 0.810 EBR-II 0.693 0.769 DFR 0.690 0.790 HFR 0.683 0.779 Average 0.701� 2.6% 0.774�4.5% Source: Morgan, W. Nucl. Technol. 1974, 21, 50–56. 338 Graphite in Gas-Cooled Reactors for the standard position in DIDO. Hence, the EDNF or fd can be calculated at the position of interest. The equivalent DIDO nickel dose (fluence) (EDND) is derived by integrating EDNF over time, as given in eqn [12]: EDND ¼ ðt 0 fdðtÞdt ½12� Table 6 compares the calculated and measured graphite damage rates in various systems using the Thompson and Wright model. Finally, for those wishing to try and reproduce damage in graphite using ion beams, Table 7 gives the energies, cross-sections, and mean number of displacement for various particles. 4.11.5.5 Energy Above 0.18MeV Dahl and Yoshikawa22 noted that for energies above 0.065MeV, eqn [13] was reasonably independent of reactor spectrum under consideration: fðE>E1Þ ¼ Ð1 0 fðEÞsðEÞnðEÞdE Ð1 E1 fðEÞdE ½13� Equation [13] is the integral of graphite displace- ment for a position in the particular reactor of interest, divided by the integral of flux from E1 (0.065MeV in this case) to infinity at the same position. Table 8 gives this ratio for two other values of E1. Graphite in Gas-Cooled Reactors 339 4.11.5.6 Equivalent Fission Flux (IAEA) An IAEA committee recommended the use of equiv- alent fission flux23 as given by eqn [14]. fG ¼ Ð1 0 P d ðEÞfðE; tÞdE Ð1 0 P d ðEÞwðEÞdE= Ð1 0 wðEÞdE ½14� Equation [14] is essentially graphite dpa divided by a normalized fission flux. A similar unit is defined by Simmons4 in his book. However, the use of this unit was never taken up for general use. 4.11.5.7 Fluence Conversion Factors Table 9 gives the conversion factor from other units to EDND. The following should be noted: � EDND is a definition, � Calder equivalent dose and other units relating damage to fuel ratings are approximate, � BEPO equivalent dose is a thermal unit and should be avoided, � Energies above En are a good approximation, � dpa is directly proportional to EDND. 4.11.5.8 Irradiation Annealing and EDT The reasoning behind the use of equivalent DIDO temperature (EDT) is that if two specimens are irra- diated to the same fluence over two different time periods, the specimen irradiated faster will contain the most irradiation damage. The reasoning is that the spec- imen irradiated at the slower rate would have a longer time available to allow for ‘annealing’ out of defects caused by fast neutron damage as outlined below. The rate of accumulation of damage dC/dt can be described by eqn [15]. Table 9 Conversion factors to EDND Fluence unit Conversion factor EDND (n cm�2) 1.0 Equivalent fission dose (n cm�2) 0.547 Calder equivalent dose (MWdAt�1) 1.0887� 1017 BEPO equivalent dose (n cm�2) 0.123 En>0.05MeV (n cm �2) 0.5 En>0.18MeV (n cm �2) 0.67 En>1.0MeV (n cm �2) 0.9 dpa (atom/atom) 7.6162� 1020 Modified fromMarsden, B. J. Irradiation damage in graphite due to fast neutrons in fission and fusion systems; IAEA, IAEA TECDOC-1154; 2000. dC dt / f exp � E kT � � ½15� where f is the flux, E is the activation energy, T is temperature (K), and k is Boltzmann’s constant. Equating the damage rate for two identical samples at different flux levels f1 and f2 and different tem- peratures T1 and T2, f1 exp � E kT1 � � ¼ f2 exp � E kT2 � � ½16� Rearranging this gives the EDT relationship: 1 y1 � 1 y2 ¼ k E ln ðfdÞ2 ðfdÞ1 ½17� The term on the left is the difference in the recipro- cal of the temperatures in the two systems (tempera- ture has traditionally been given the symbol ‘y ’ when used in this context) and the term on the right con- tains the natural log of the damage flux (or fluence) in the two systems divided by each other. In practice, the activation energy E is an empirical constant. The use of EDT has recently been investigated24 at temperatures above 300 �C. The authors concluded that the use of EDT was inappropriate (Figure 8). However, below 300 �C, there was some evidence of the applicability,15 but at these lower temperatures there is little reliable data. Therefore, the use of the EDTconcept is not recommended for modern graph- ite moderated reactors where the graphite is usually irradiated above 300 �C. 4.11.5.9 Summary of Fast Neutron Dose (Fluence) 1. Care must be taken when interpreting graphite data because of the variety of fast neutron dose units used. Older data in particular should be treated with care. 2. ‘Graphite damage’ has been equated to activation of nickel at a standard position in DIDO. This can now be calculated and equated to dpa. 3. ‘Graphite damage’ may also be equated to channel burnup which can also be equated to dpa. 4. ‘Graphite damage’ can also be equated to En> 0.18MeV. 5. EDT is not applicable to irradiation temperatures above 300 �C; there is some evidence that it may be applicable below 300 �C. 0 −3 −2 −1 0 1 D im en si on al c ha ng e (% ) PLUTO DFR 2 50 100 150 Fluence (1020n cm−2 EDND) 200 250 300 Figure 8 Comparison of dimensional changes on Gilsocarbon graphite samples irradiated in DFR with similar samples irradiated in PLUTO. Reproduced from Eason, E. D.; Hall, G.; Heys, G. B.; et al. J. Nucl. Mater. 2008, 381, 106–113. 340 Graphite in Gas-Cooled Reactors 6. There are conversion factors between all these units but these are subject to various degrees of uncertainty. 4.11.6 Graphite ‘Energy Deposition’ (Nuclear Heating) The heat generated in the graphite (or energy depo- sition) is required for the calculation of the graphite temperatures, and in the case of CO2-cooled systems, it is required for the calculations of radiolytic weight loss. Both of these requirements are important in graphite stress analysis calculations. In the case of CO2-cooled systems it is assumed that the graphite radiolytic oxidation rate is proportional to the heat generated in the graphite. However, it is ionizing irradiation that causes the dissociation of the CO2. The energy deposition is produced by the inter- action of graphite atoms with three types of particles: � Neutron interactions with graphite atoms (�40%). � Fission g-rays (�60%). � Secondary g-rays caused by absorption by materi- als outside the moderator (e.g., steel fuel pins in AGRs) and by inelastic scattering of carbon atoms (�1% in aMagnox reactor and�10% in an AGR). The main source of gammas and neutrons arises from the fuel, mainly from prompt fission, but there are some from delayed fission. The ratios given above are for a central position in the core and for initial fuel loading. The ratio may change with position in the core and with graphite weight loss. Furthermore, in graphite material test programs, the ratio between neutron and g-heating is likely to be significantly different, because of the dif- ferent materials used to construct the various reactor cores. It is therefore important that this ratio is known and the implication of a change in this ratio onmaterial property changes, that is, the implication of the ratio between fast neutron damage versus radiolytic weight loss on graphite property changes, is understood. The gamma and neutron spectrum varies with distance from the fuel and will vary with graphite density (i.e., will change with weight loss) and fuel design. A reactor is run at constant power, and there- fore, as weight loss increases, the spectrum (gamma and neutron) will change and become harsher (higher neutron and g-flux). In the graphite, charged electrons are produced because of the following: 1. Compton scattering interaction of gamma with electrons within the carbon atoms. 2. Pair product in electrostatic field associated with carbon atoms. 3. Photoelectric absorption. Compton scattering predominates, but electrons and charged carbon ions are also produced because of the displacement of carbon atoms in the moderator, and in principle this could be calculated. Energy deposition is the energy released from the first collisions of primary gamma and neutrons.25 Energy deposition is calculated in watts per gram (W/g) and the spatial distribution can be calcu- lated using reactor physics codes such as McBend Graphite in Gas-Cooled Reactors 341 (http://www.sercoassurance.com/answers/), WIMS, and WGAM. However, a crude estimation of energy deposition can be made by assuming that �5% of the reactor power is generated in the graphite. This heat can then be proportioned to the rest of the core using interpolation and form factors, and estimates of the distribution within a moderator brick. In conclusion, energy deposition is required to calculate graphite temperatures and radiolytic oxida- tion rates. Energy deposition can be estimated but is most accurately calculated using reactor physics codes. However, care must be taken because the ratio between neutron heating and g-heating, or more appropriately a direct measure of the ionizing irradi- ation, is important. 4.11.6.1 The Use of Titanium for Installed Sample Holders During the construction of the Magnox and AGR reactors, graphite specimens were placed into ‘installed sample holders,’ the intention being that these samples could be removed at a later date to give information on the condition of the graphite core. To enhance the radiolytic weight loss of the graphite in the installed sample holders, titanium was used. Although this only slightly increased the g-heating, it did increase the number of electrons produced, because of an increase in pair production and Compton scattering caused by the higher atomic number or ‘Z-value’ of titanium compared to graphite (22 and 6, respectively). 4.11.7 Radiolytic Oxidation 4.11.7.1 Introduction In carbon dioxide (CO2)-cooled reactors, two types of oxidation can occur. The first is thermal oxidation which is purely a chemical reaction between graph- ite and CO2. This reaction is endothermic and is negligible below about 625 �C and is not impor- tant up to 675 �C. The second is radiolytic oxidation that occurs when CO2 is decomposed by ionizing radiation (radiolysis) to form CO and an active oxi- dizing species, which attacks the graphite. Radiolytic oxidation occurs predominantly within the graphite open porosity. 4.11.7.2 Ionizing Radiation Ionizing irradiation can be defined as that part of a radiation field capable of ionization (charge separation) in CO2 either directly or indirectly. This leads to the creation of reactive species, which may react with the carbon atoms at the surfaces (external and more importantly internal) of the graphite components. 4.11.7.2.1 Energy deposition Historically, ‘energy deposition’ has been used for a surrogate for ionizing irradiation, most probably because it is easy to measure using calorimetry and can be estimated from the reactor power. Energy deposition, sometimes referred as ‘dose rate,’ in the units of W/g of graphite, is a measure of the total energy absorbed in the gas in unit time from the scattering of g-radiation and fast neutrons. For a typical Magnox reactor, energy deposition is composed of approximately the following components: � 36% from the neutrons � 58% from the gamma � 6% from the interaction of graphite atoms within the moderator Of these, it is only the last two that directly con- tribute to ionization of the carbon dioxide gas, mainly through Compton scattering. These ratios will be slightly different in an AGR. An assumption is made that the dose rate received by the graphite is the same as that absorbed by carbon dioxide within the pores of the graphite and that a fraction k of the fission energy from the fuel causes heating in the moderator. For a typical Magnox reac- tor, k is �5.6% of the thermal power. The unit G�C is defined as the number of carbon atoms gasified by the oxidizing species produced by the absorption of 100 eV of energy in the CO2 contained within the graphite pores; G�C for pure CO2¼ 3. 4.11.7.3 Radiolytic Oxidation Mechanism The exact mechanism of radiolytic oxidation in a car- bon dioxide-cooled reactor is complex and has been a matter of debate for some time; the most satisfactory explanation has been given by Best et al.26 However, in its most simplistic form the mechanism can be described as follows: In the gas phase, CO2�! ionizing radiation COþO ½I� COþO ! CO2 ½II� where O* is an activated state-oxidizing species. Thus, after ionization the carbon monoxide and oxidizing species rapidly recombine back into carbon dioxide and to an uninformed observer, carbon http://www.sercoassurance.com/answers/ 0.8 0.7 0.6 0.5 0.4 0.3 0.2 Is ot he rm al r ad io ly tic o xi d at io n ra te (1 0− 8 h− 1 W kg −1 ) 0.1 0 0 100 200 300 400 Calculated CH4 concentration (vpm) 500 600 700 800 0.4–0.6% CO 1% CO 2% CO Figure 9 G�C as function of CO and CH4 concentration (41bar, 673K). 0 1 2 3 0 1 2 3 G –c CO (%) 5 vpm H2O 15 to 30 vpm H2O 200 to 400 vpm H2O (a) 342 Graphite in Gas-Cooled Reactors dioxide would appear to be stable in an irradiation field. However, in the presence of graphite which typically contains �20% porosity, 10% of which is initially accessible to the carbon dioxide gas, at the graphite pore surface (mainly internal) carbon atoms are oxidized. This can be simplistically described as O þ C ! CO ½III� The principal oxidizing species is still under debate, but the most favored candidate is the negatively charged ion, CO3 �. 0.0 0.5 1.0 1.5 G –c 4.11.7.4 Inhibition The rate of oxidation can be reduced by the addition of carbon monoxide (CO) and moisture (H2O) and can be greatly reduced by the addition of methane (CH4), as illustrated in Figures 9 and 10. As described above the radiolytic oxidation process pro- duces CO and if CH4 is added, moisture will be one of the by-products of the reaction. 0 200 400 600 800 CH4(vpm)(b) Figure 10 Inhibition in carbon dioxide due to the addition of carbon monoxide, moisture, and methane. (a) G�C as function of CO and H2O concentration and (b) G�C as function of methane (CH4) concentration. Reproduced from Best, J. V.; Stephen, W.; Wickham, A. Prog. Nucl. Energy 1985, 16(2), 127–178. 4.11.7.5 Internal Porosity As supplied, graphite components contain a significant amount of both open and closed porosity in a variety of shapes and sizes, from the nm scale to the mm scale, as illustrated in Section 4.11.3. The open pore volume (OPV) is defined as the volume of pores accessible to helium, closed pore volume (CPV) is the volume of pores not accessible to helium, and the total pore volume (TPV) is the volume of open and closed pore. The effect of pore size on the radiolytic oxidation rate was investigated by Labaton et al.27 who found the maximum range to be 2.5–5 mm. Taking this into account and referring to eqns [I]–[III] above, the oxidation process will be expected to be more effi- cient in the smaller pores than in the larger pores. Graphite in Gas-Cooled Reactors 343 This is because in the smaller pores the distance to the wall is less, making it less likely, compared to the case for larger pores, that the active species would be deactivated by collision in the gas phase. To account for this difference in oxidation rate with pore size for modeling purposes, in the case of the Magnox reactors which did not have CH4 rou- tinely added to the coolant, a pragmatic approach of defining ‘pore efficiency’ was adopted, whereas in the case of the AGRs where CH4 is routinely added, a reactive pore volume (RPV) was defined as being the volume of pores oxidizing in CH4-inhibited coolant gas. It is also clear that as the oxidation process pro- ceeds, closed porosity will be opened and the pore size distribution will change, thereby changing the oxidation rate. 4.11.7.6 Prediction of Weight Loss in Graphite Components The methodologies used to predict the oxidation rate in Magnox reactors are based on work by Standring28 as discussed below: Weight of CO2 in the pores of 1g of graphite ¼ er0 P 14:7 273 T � � ½18� where P (psi) is the gas pressure, T (K) the tem- perature, and r0 is the density of CO2 at standard conditions for temperature and pressure (STP) (g cm�3). The dose rate to the graphite can then be given in watts as follows: Dose rate to graphite ¼ eDr0 P 14:7 � 273 T � � W ½19� where D (W g�1 s�1) is the ‘energy deposition rate’ or ‘dose rate’ and e is the OPV (cm3 g�1). This reasoning can be taken further to give Percentage initial oxidation rate; g0 ¼ 145 eG�CDP T � � %per year ½20� Standring and Ashton29 measured the OPVand CPV in PGA as a function of weight loss (Figure 11). In the specimens they examined, there appeared to be a small amount of pores which opened rapidly before the pore volume increased linearly as a func- tion of weight loss over the range of the data. To account for this behavior, they modified eqn [20] by defining an effective OPV as ‘ee’: g0 ¼ 145 eeG�CDp T � � % per year ½21� Standring further developed this reasoning into a relationship for the cumulative weight loss, Ct , at a constant dose rate: Ct ¼ A exp g0t A � � � 1 h i ½22� where A ¼ 100Pe1�Peð Þ and Pe is the effective initial OPV in cm3 cm�3. However, a reactor is operated at constant power. Replacing the dose rate in eqn [21] by kPt/Wm, where Pt is the reactor thermal power, k is the fraction of the reactor power absorbed in the graphite (�5%), and Wm is the weight of the moderator, gives g0 ¼ 145eeG�C kPt Wm P T % per year ½23� From eqn [23], it is clear that the rate of oxidation will increase with loss of moderator mass. It was shown by Standring that the cumulative weight loss, Ct, for a reactor operated at constant power is given by A2 100Pe log 1þ Ct A � � � A 100 Ct ¼ g0t ½24� This equation yields higher weight loss than the constant dose rate equation. This approach was used to design the early Magnox stations. However, as higher weight loss data became available from the operating Magnox stations, it was found necessary to modify the rela- tionship to account for the pore distribution with increasing oxidation. 4.11.7.7 Weight Loss Prediction in Inhibited Coolant It had not been possible to regularly add CH4 as an inhibitor to the coolant in the Magnox reactors because of concerns regarding the metallic compo- nents in the coolant circuit. However, the higher rated AGRs were designed with this in mind by selecting denser graphite and adding CH4 gas as an oxidation inhibitor. The addition of an inhibitor causes the process of radiolytic weight loss to be more complex than that for Magnox reactors as the oxidation rate becomes a com- plex function of the coolant gas composition. This is because gas composition, and hence, graphite oxidation rate, is not uniform within the moderator bricks and keys as CH4 is destroyed by radiolysis and may thus be depleted in the brick interior. In addition, methane destruction gives rise to the formation of carbon 50 40 35 30 25 20 15 O p en p or e vo lu m e (c m − 3 p er 10 0 cm − 3 ) C lo se d p or e vo lu m e (c m − 3 p er 10 0 cm − 3 ) 10 5 0 35302520 Weight loss (%) Weight loss (%) (a) (b) 151050 7 6 5 4 3 2 1 0 0 5 10 15 20 25 30 35 45 Initial density ~ 1.68 g cm−3 Initial density ~ 1.74 g cm−3 Figure 11 (a) Open and (b) closed pore volume in pile grade A as a function of radiolytic weight loss. 344 Graphite in Gas-Cooled Reactors monoxide and moisture which may be higher in the brick interior. Graphite oxidation forms carbon mon- oxide, thereby further increasingCO levels in the brick interior. These destruction and formation processes are gas composition dependent and the flow rates of these gases within the porous structure are dependent upon graphite diffusivity and permeability values which change with graphite weight loss. The exact mechanism of radiolysis in a CH4- inhibited coolant is complex and the radicals are disputed. However, from a practical point of view the mechanisms for oxidation and inhibition can be considered as given below: In the gas phase CO2�! ionizing radiation COþ O ½IV� COþ O ! CO2 ½V� CH4 ! P ½VI� where O* is the activated oxidizing species formed by radiolysis of CO2 and P is a protective species formed from CH4 oxidation. At the graphite surface (mainly internal porosity), O þ C ! CO ½VII� O þ P ! OP ½VIII� where OP is the deactivated gaseous product of CH4 destruction. An altogether more satisfactory explanation and model for the effect of pore structure on corrosion in gas mixtures containing carbon monoxide, CH4, and Graphite in Gas-Cooled Reactors 345 water was developed by Best and Wood30 and Best et al.,26 who gave a relationship for G�C with respect to a pore structure parameter F and to P, the proba- bility of graphite gasification resulting from species which reach the pore surface: G�C ¼ 2:5FP ½25� The inhibited-coolant radiolytic oxidation rate is usually referred to as the graphite attack rate. Data on initial graphite attack rate have been obtained in experiments carried out in various materials test reactors (MTRs)31 for Gilsocarbon and to some extent other types of graphite (Figure 12). From Figure 12, it can be seen that the oxidation rate does not go on exponentially increasing as predicted by earlier low-dose work, but the increasing rate saturates at about 3 times the initial oxidation rate. The approach to predicting temporal and spatial weight loss in graphite components irradiated in inhibited coolant is to use numerical analysis to solve the diffusion equations given below: Methane concentrations rT ðD10rðC1Þ � rðn C1ÞÞ � K1 ¼ 0 Moisture concentrations rT ðD20rðC2Þ � rðn C2ÞÞ þ K1STOX ¼ 0 Carbon monoxide concentrations rT ðD30rðC3Þ � rðn C3ÞÞ þ K1STOXþ K2STOX2 ¼ 0 ½26� The basic unknowns are the CH4, C1, moisture, C2, carbon monoxide, C3, and gas concentration profiles. 0.25 0.20 0.15 Fr ac tio na l g ra p hi te w ei gh t lo ss 0.10 0.05 0.00 0 10 20 Fluence Figure 12 Typical experimental weight loss dose relationship In the CH4 part of eqn [26], the first term is the pure diffusion contribution, and D10 is the effective diffusion coefficient in graphite of CH4 in CO2. The second term is the contribution from porous flow due to permeation, and n is the velocity vector for CO2 flow through the graphite pores, and K1 is the sink term for CH4 destruction. In the moisture part of eqn [26], the first term is again the pure diffusion contribution, and D20 is the effective diffusion coefficient in graphite of moisture in CO2. The second term is the contribu- tion from porous flow. K1STOX is the source term for moisture formation from CH4 destruction in accordance with CH4 þ 3CO2 ! 4COþ 2H2O ½IX� In the carbon monoxide part of eqn [26], the first term is the pure diffusion contribution, and D30 is the effective diffusion coefficient in graphite of car- bon monoxide in CO2. The second term is the contribution from porous flow. K1STOX is defined above and K2STOX2 is the source term of carbon monoxide formation from graphite oxidation. The various terms in the diffusion equations must be updated at each time-step for changes in coolant composition, dose rate, attack rate, and all parameters controlling graphite pore structure, diffusivity, and permeability which change with oxidation. These equations can be solved numerically using finite dif- ference or finite element techniques to give point wise, temporal distributions of weight loss in a graph- ite component. 30 (MWh kg−1) 40 50 60 from materials test reactor experiments. 346 Graphite in Gas-Cooled Reactors 4.11.8 Graphite Temperatures Graphite component temperature depends on radia- tion and convection (and conduction in the case of light-water gas-cooled reactors) heat transfer from the fuel and heat generated in the graphite by neu- tron and g-heating, that is, energy deposition as dis- cussed above. Therefore, a detailed knowledge of the coolant flow is important. Thermohydraulic codes such as Panther (http:// www.sercoassurance.com/answers/) are used to cal- culate heat generated in graphite blocks. These codes estimate the following: 1. The heat generated in the fuel. 2. The coolant flow. 3. The heat transfer to the graphite. 4. The heat ‘energy deposition’ in the graphite. The calculations take account of graphite weight loss and change in thermal conductivity of the graph- ite due to fast neuron damage and radiolytic oxida- tion. The largest uncertainty is probably associated with the size of flow bypass paths and flow resistance. In an AGR, the temperature at the outside of the brick is lower than the temperature at the inside because of the interstitial flow, whereas in an Reaktor Bolshoy Moshchnosti Kanalniy (RBMK) the temper- ature is hotter at the brick outside. Using the brick ‘boundary conditions’ including energy deposition temperatures calculated by the thermohydraulic code, a standard finite element code such as ABAQUS can easily be used to calculate the spatial distribution of temperature with the graphite component. Thermal transient tempera- tures can also be calculated using a standard finite element code. Often, the temperature distribution is calculated for a central brick, and the temperatures in the bricks in the rest of the core are calcu- lated using interpolation/extrapolation, that is, form factors as described in Section 4.11.9. The calculated Reflector Ind ch Abo cha in a (a) Core Figure 13 Form factors. (a) Graphite moderator with reflector temperatures are compared with the few brick ther- mocouples that are installed in the moderator. The codes are also fine-tuned to these. In conclusion, the calculation of graphite tem- peratures is complex and involves the calculation of heat transfer flow to the fuel and flow calculations. Graphite temperature predictions should be com- pared to measurements taken from thermocouples located in most graphite cores. 4.11.9 Variation of Fluence, Temperature, and Weight Loss in a Reactor Core The flux, temperature, andweight loss will vary within each individual graphite component, for example, moderator brick. In addition, the mean component flux, temperature, and weight loss will vary through- out the core. When designing a graphite core, in order to extrapolate data from one component, which has been analyzed in detail, to the other core components, ‘form factors’ are often used, as illu- strated in Figure 13. In typical graphite-moderated reactors, the axial (vertical) flux varies approximately as a cosine with the maximum at center, whereas the radial flux is usually a flattened cosine as illustrated in Figure 13. The exact form of these profiles can be calculated using reactor physics codes. The mean core rating can be calculated from eqn [27]: Rating ¼ reactor power = weight of fuel in reactorðMWd t�1Þ ½27� and at the time of interest the mean core burnup can be calculated by eqn [28]: Core burnup ¼ reactor power � days at power = weight of fuel in reactor ðMWd t�1Þ ½28� ividual annel MHA ut 320 nnels n AGR (b) Core and (b) graphite moderator flux profile (form factors). http://www.sercoassurance.com/answers/ http://www.sercoassurance.com/answers/ Graphite in Gas-Cooled Reactors 347 Thus, a mean moderator brick burnup can be calcu- lated by multiplying the mean core burnup by the axial and radial form factor for the particular brick of interest. 4.11.9.1 Fuel End Effects The relatively small gap between fuel elements has a pronounced effect on the damage to the graphite moderator bricks. This is particularly noticeable in the brick dimensional changes, in both AGRs and RBMK reactors. In assessments, this detail needs to be accounted for and may require a three-dimensional reactor physics calculation. 4.11.9.2 Temperature and Weight Loss By using the same ‘form factors,’ the moderator brick mean weight loss can be estimated, assuming that weight loss is proportional to burnup or fluence. The gas temperature will vary roughly linearly in the axial (vertical) direction from the inlet tempera- ture T1 to the outlet temperature T2. A more detailed profile may be calculated using a thermohydraulics All values � 1011 2 2.95 33.42 3.44 3.4 4.14 4.2121.18 18.26 23.48 4.0 4.19 4.15 4.13 Figure 14 Nickel flux distribution in a quarter cell calculated f code. The radial temperature can be assumed to follow the radial flux profile. Thus, an approximate mean gas temperature for an individual moderator brick may be obtained. 4.11.10 Distribution of Fluence Within an Individual Moderator Brick Having obtained the component mean fluence, tem- perature, and weight loss, the variation of these para- meters throughout the particular component of interest is required. The fluence reduces exponentially away from the fuel in the radial direction, but is influenced by surrounding fuel sources. The exact distribution is usually calculated using a reactor physics code for a 5� 5 array pertinent to the area of interest. Figure 14 is an example for the Windscale Piles. Thus, the spatial and temporal fluence distribu- tion throughout a graphite component can be calcu- lated. The component temperature can be calculated using finite element analysis through knowledge of the surrounding gas temperature, accounting for the 1.94 1.97 2.03 2.05 2.14 2.11 2.28 2.32 2.41 2.21 2.39 2.44 2.31 .64 .04 1 2.99 2.69 2.49 2.63 2.79 3.06 3.55 3.44 3.08 2.83 2.72 2.69 2.70 2.803.053.476 2.72 2.56 2.52 2.61 or the Windscale Piles. Courtesy of A. Avery. 348 Graphite in Gas-Cooled Reactors 5% of the reactor heat which is generated within the graphite. Graphite weight loss variation within a component is more complex and is calculated by various empirical industry codes. If the axial variation in fluence, temperature, and weight loss along the brick length is deemed to be important, three-dimen- sional physics, temperature, and weight loss calcula- tions will be required. 4.11.11 Fast Neutron Damage in Graphite Crystal Structures Atomic displacements due to fast neutron irradiation modify the ‘crystallite’ dimensions and most of their material properties. Neutron energies of around 60 eV are required to permanently displace carbon atoms from the lattice. However, most damage in graphite is due to fast neutron energies >0.1MeV; a typical thermal reactor has neutron energies of up to 10MeV, with an average of 2MeV. High-energy neutrons knock an atom out of the lattice, leading to a cascade of secondary knock-ons. This process knocks atoms into interstitial positions between the basal planes, leaving vacant positions within the lat- tice. Many of the interstitial atoms will immediately find and fill these vacancies. However, others may form semistable Frenkel pairs or other small clusters or ‘semistable’ clusters. With increasing fast neutron damage, the stability, size, and number of these clus- ters will change depending on the irradiation tem- perature. The higher the irradiation temperature, the larger are the interstitial clusters or ‘loops.’ This process leads to considerable expansion in the graph- ite crystal ‘c’ axis. Conversely, vacancy loops also form and grow in size with increased irradiation temperature. It has been postulated that this process will cause the lattice to collapse leading to the ‘a’ axis shrinkage observed on irradiating graphite crystal structures. This process is illustrated in Figure 15. Thrower32 carried out an extensive review of transmission electron microscopic (TEM) studies of defects in graphite, particularly those produced by fast neutron irradiation. He demonstrated that interstitial loops and vacancy loops could be distin- guished by tilting the specimen. He was able to observe vacancy loops in graphite irradiated only at and above 650 �C, whereas interstitial loops and defects were observed at all temperatures of interest to reactor graphite. It is proposed that the dimen- sional change in bulk polycrystalline graphite may be understood by eqn [29]33: DLc Lc ffi Dc c þ r0 r1 � �2 ½29� where Lc is the crystal dimension perpendicular to the basal plane, ‘c’ is the atomic lattice parameter, and r0 and r1 are the mean defect radius and mean half separation of defects in the basal plane, respectively. However, it was noted that this does not completely explain the expansion. In order to explain basal plane contraction it is necessary to postulate that vacancy lines cause the collapse of the basal planes.34–36 More recent atomistic calculations due to Telling and Heggie37 have sought to explain the process by the ‘buckling’ of basal planes until they twist round upon themselves. This latter explanation is more satisfying as it accounts for the atomistic bonding around the edges of the interstitial loops and vacan- cies. However, more HRTEM (high-resolution trans- mission electron microscopy) observations and other techniques are required to validate these theories. Whichever mechanism is correct, empirical observations made on HOPG, and some natural crystal flakes,35 show that graphite crystal structures expand in ‘c’ axis and shrink in the ‘a’ axis, the degree of deformation being a function of fast neu- tron fluence and irradiation temperature. Crystal dimensional change is discussed in more detail later in this section. 4.11.11.1 Stored Energy It would not be appropriate to continue without some discussion on stored (or Wigner) energy. The perfect crystal configuration is the lowest energy state for the graphite lattice. However, irradiation damage will considerably alter that configuration. Wigner38 predicted that the increased lattice vibra- tion due to heating would allow carbon atoms to rearrange themselves into lower energy states, and that in doing so energy would be released in the form of heat. Early experience in operating graphite- moderated plutonium production and research reac- tors at low temperatures in the United States, Russia, France, and the United Kingdom proved that this assumption was correct. The highest value of stored energy measured was �2700 J g�1.15 If all of this were released under adiabatic conditions, the temperature rise would be 1500 �C. Fortunately, that is not the case. Furthermore, the accumulation of stored energy is insignificant above an irradiation temperature of �300 �C, it is difficult to accidentally release the stored energy above an irradiation ~0.01µm Vacant lattice site Interstitial atom 1 keV 0.001µm ~500 eV Displacement cascade Interstitial defects Vacancy defects Layer plane Submicroscopic cluster of 4±2 atoms Interstitial loop Interstitial diffusing to loop Dose increase Dose increase Layer increase Single vacancy Di vacancy Vacancy line Vacancy loop Figure 15 Formation of interstitial and vacancy loops in graphite crystals. Modified from Simmons, J. Radiation Damage in Graphite; Pergamon: London, 1965. Graphite in Gas-Cooled Reactors 349 temperature of �150 �C, and only limited self-sus- taining energy release of stored energy can be achieved in graphite irradiated below �100 �C. Thus, stored energy is now of consideration in the United Kingdom only in the decommissioning of shutdown reactors such as the Windscale Piles and BEPO and other similar overseas systems, although there are graphite ‘thermal columns’ in some research reactors that may require periodic assessment. The reason for this is the nature of the irradia- tion damage sites with respect to irradiation tempera- ture. In graphite irradiated in the early facilities, at temperatures between about ambient and 150 �C, point defects associated with Frenkel pairs and small loops can diffuse only slowly through the lattice to form larger, more stable loops because of the low irradiation temperature. However, thermal annealing at temperatures above the irradiation temperature can readily release the stored energy, and under certain circumstances, this release can be self-sustaining over certain temperature changes. (A ‘rule of thumb’ tem- perature of 50 �C above the irradiation temperature is often cited as a ‘start of release temperature.’ How- ever, this is misleading as a heat balance needs to be 900 800 700 600 500 To ta l s to re d e ne rg y (J g− 1 ) 400 300 200 100 0 0.0 0.2 0.4 0.6 Fluence (1020n cm–2nvt) 0.8 1.0 1.2 25 �C 30 �C 80–100 �C 99 �C 125 �C 140 �C 150 �C 180 �C 1.4 1.6 Figure 16 The accumulation of stored energy as a function of fluence and temperature. 350 Graphite in Gas-Cooled Reactors considered when assessing energy release rates. Thus, 50 �C above the irradiation temperature can be con- siderably overconservative.) The accumulation of stored energy, measured by burning irradiated graphite samples in a bomb calo- rimeter, is given as a function of fluence and temper- ature in Figure 16. At low fluence, stored energy quickly accumulates reaching a plateau at high flu- ence. Many measurements were made in the Wind- scale Piles, BEPO, Hanford, and Magnox reactors that clearly illustrated this behavior.15 To fully understand the thermal stability of graph- ite containing stored energy, the most appropriate measure is the rate of release of stored energy measured using a differential scanning calorimeter (DSC) as illustrated in Figure 17. A graphite sample is heated in the DSC usually at a constant rate of 2.5 �Cmin�1. In simple terms, two runs are made and the heat capacity of the samples measured in each case. When the heat capacities from the two runs are subtracted, the energy release rate is easily obtained as a function of heating temperature. This can be compared to the specific heat of graphite as given in Figure 17. When the rate of release of energy is below the specific heat, energy needs to be added to continue the process. When the rate of release is above the specific heat, the process is self-sustaining. This behavior was used to ‘anneal’ the Windscale Piles; a ‘hit and miss’ strategy that ended in damage to the fuel cartridges and eventually a ‘metal uranium fire.’ (Contrary to ‘common folklore,’ the graphite did not burn in the Windscale incident. A limited amount of graphite was oxidized leading to enlargement of fuel and control channels but it was the metal uranium that burnt. Graphite is very diffi- cult to burn and requires large amounts of heat and oxygen or air, applied to crushed graphite in a flui- dized bed or in similar form.39) The form of this rate of release curve is a func- tion of (1) the amount of stored energy in the sample, (2) the temperature the sample was irra- diated at, (3) the fluence the sample had been irradiated to, (4) the release temperature, and (5) the heating rate. Unfortunately, there are no com- prehensive datasets of these five parameters that allow a robust empirical model to be derived for assessing the stability of graphite containing stored energy. The models that usually exist take the worst- case rate of release curve and fit an Arrhenius type equation to the rate of release curve. dS dt ¼ �uS exp �EðT ; SÞ KTðtÞ � � ½30� where S is the stored energy remaining, t is time, T(t) is temperature in (K) as a linear function of time (T¼at in the case of the DSC test and is nonlinear in most practical cases), K is Boltzmann’s constant, and E(T, S) is the activation energy as a function of the stored energy remaining and temperature and u is a frequency factor usually taken as 7.5 � 1013 s�1.40 3.0 2.5 2.0 1.5 R at e of r el ea se o f s to re d e ne rg y (J g− 1 �C −1 ) 1.0 0.5 0.0 0 50 100 150 200 250 300 Temperature (�C) 350 400 450 500 Specific heat capacity Dowel 4 Dowel 11 Dowel 14 Dowel 24 Dowel 30 Dowel 36 Figure 17 Typical rate of release of stored energy. Modified from Bell, J. C.; Gray, B. S. Stored Energy Studies Made on Windscale Pile Graphite Since October 1957; TRG Report 84(W), UKAEA, 1961. Graphite in Gas-Cooled Reactors 351 It can be appreciated that the exact solution of eqn [30] requires a substantial amount of information from several rate of release curves from several samples, which is seldom available. Thus, a practical approach is usually taken, the simplest of which is to assume a single activation energy. However, this is not very satisfactory and more elegant approaches using vari- able or discrete activation energies can be found.41–44 Having derived a satisfactory model for the rate of release using a DSC, it then can be applied to a practi- cal situation using commercially available computer codes such as ‘user subroutine’ facilities.42 In assessing practical situations, it is important to use an energy balance that accounts for heat applied, heat generated by the release of stored energy, the heat capacity of the graphite itself, and heat lost to the surroundings. If the heat generation is intense and oxygen is available, the heat generated by graphite oxidation should be taken into account. However, the latter case should be unnecessary as a professional scientist or engineer would not design a system or process that would approach such conditions. It should be noted that irradiated graphite thermal conductivity and total stored energy are directly cor- related15; see eqn [31]. Therefore, the thermal con- ductivity will improve as energy is released. S ¼ 27:2 K0 K � 1 � � J g�1 ½31� The data that eqn [30] is based on is derived from rate of release curves obtained using a relatively fast heating rate. In dealing with irradiated graphite waste, much slower rates of heating are often required. Graphite samples taken from the Windscale Piles 40 years after the incident showed little change in the dS/dt curves,45 indicating that diffusion of atoms at around ambient temperature is extremely slow. Nevertheless, condi- tions relevant to any proposed encapsulation technique and repository will need to be accounted for in deter- mining if heat released from stored energy is an issue. The rate of release curves given in Figure 17 are only to a temperature of around 450 �C. It had been observed, by comparing the energy released in the DSC with the energy released on a similar sample in a bomb calorimeter, that not all of the stored energy had been released in the samples heated to a maxi- mum of 450 �C in the DSC. It was found that on increasing the temperature to around 1600 �C, a sec- ond peak could exist46,47; see Figure 18. It was observed that the ‘200 �C’ peak reduced in size and moved to a slightly higher temperature with increased irradiation, presumably as the irradiation induced defects became more stable, and the plateau between the two peaks increased in height and approached the specific heat value.15 The first of these phenomena could explain why it became more and more difficult to ‘anneal’ the Windscale Piles39 and the second had the implication that 500 450 400 350 300 250 E ne rg y (c al g− 1 �C −1 ) 200 150 100 50 0 0 200 400 600 800 1000 Temperature (�C) 1200 1400 1600 1800 2000 Graphite-specific heat Rate of release Figure 18 Schematic of the high-temperature rate of release curve. Reproduced from Rappeneau, J.; Taupin, J. L.; Grehier, J. Carbon 1966, 4(1), 115–124. 352 Graphite in Gas-Cooled Reactors eventually the rate of release curve would remain above the specific heat up to 1600 �C with the conse- quent safety implications. Fortunately, the second of these phenomena proved to be incorrect. It is interesting to note that there is a correla- tion (eqn [32]) between the height of the plateau at �400 �C and total stored energy.15 dS dT½400� ¼ S 1670 J g�1 �C�1 ½32� This equation, although not exact, was often used in reactor graphite sampling programs to avoid having to measure total stored energy. Most of the discussion on stored energy above is relevant only to low temperature reactor systems, with graphite temperatures operating from ambient to 150 �C. When graphite is irradiated at higher tem- peratures, in practice above about 100 �C, the dS/dt does not exceed the graphite specific heat. One of the operating rules for the UK Magnox reactors was that the dS/dt, as measured on surveillance samples, should always be below 80% of the specific heat, which proved to be the case. 4.11.11.2 Crystal Dimensional Change As previously discussed, graphite crystal structures, in the form of HOPG, have been observed to swell in the ‘c’ axis direction and contract in the ‘a’ axis direction for all measured fluence and irradiation temperatures. Figure 19 shows early data obtained by Kelly et al.35 It is clear from Figure 19 that the rate of swelling and shrinkage significantly changes between 200 and 250 �C, indicating that the defect population is becoming more stable above this tem- perature range. HOPG data is an important input into multiscale models of irradiation damage in graphite.48 For the purpose of understanding irradiation damage in operating reactors, data would be required ideally from 140 to 1400 �C, the maximum fluence being dependent on the irradiation temperature. Unfortu- nately, the dataset is far from complete. The data due to Brockelhurst and Kelly49 is the most complete set of HOPG irradiation data covering the fluence and part of the temperature range appropriate to AGRs (Figure 20). In the same paper, the authors show the effect of final heat treatment, between 2000 and 3000 �C, on the crystal dimensional change rate of HOPG. The data showed that the lower the heat treatment, the faster is the dimensional change rate, indicating that the dimensional change rate of a poorly graphitized component would be expected to be greater than that of majority of the components.50 Pre- viously, Kelly and Brocklehurst51 had shown that boron doping also significantly increased the dimensional 0 0 5 10 15 20 D im en si on al c ha ng e (% ) 25 30 -7 -6 -5 -3 -4 -1 -2 D im en si on al c ha ng e (% ) 0 1 5 10 15 Fluence (1020 n cm-2 EDND) c-axis a-axis 20 25 30 0 5 10 15 Fluence (1020 n cm-2 EDND) 20 25 30 150 °C 170 °C 200 °C 250 °C 350 °C 450 °C 650 °C 150 °C 170 °C 200 °C 250 °C 300 °C 350 °C 450 °C 650 °C Figure 19 Dimensional changes in highly orientated pyrolytic graphite as a function of fast neutron fluence and temperature. Modified from Kelly, B. T.; Martin, W. H.; Nettley, P. T. Philos. Trans. R. Soc. Lond. A Math. Phys. Sci. 1966, 260 (1109), 37–49. Graphite in Gas-Cooled Reactors 353 change rate in HOPG, and this was again reflected in the behavior of doped polycrystalline graphite.50 HOPG high-temperature data was mainly obtained by investigators interested in the behavior of HTR fuel coatings.52 Some of this data is for low-density pyrolytic carbons and it is not always made clear which material the data refers to. Figure 21 shows all the data known to the author, and it is clear that there is some inconsistency. 4.11.11.3 Coefficient of Thermal Expansion There are only two reasonable sets of data for the CTE in HOPG. Data for the lower temperatures (150–250 �C) is given in Figure 22. This data is associated with the dimensional change data given in Figure 19. At these low temperatures, there is a significant increase in dimensional changes with increased fluence. This could explain the increase, and the subsequent decrease, in CTE in the ‘c’ direc- tion. It is interesting that this is reminiscent of the behavior of medium grained, semi-isotropic poly- crystalline data, as discussed later. In the ‘a’ direction, there is significant scatter in the data, possibly due to the difficulty in measuring such low values of CTE. However, the behavior appears to indicate an increase to a plateau, reminiscent of the behavior of needle coke anisotropic polycrystalline data. The higher temperature HOPG CTE data given in Figure 23 appears to be invariant to the increasing 70 60 c-axis, 430 �C c-axis, 600 �C a-axis, 430 �C a-axis, 600 �C50 40 30 20 10 −10 −20 −30 0 20 40 60 80 100 Fluence (1020 n cm−2 EDND) 120 140 160 180 200 0 D im en si on al c ha ng e (% ) Figure 20 High-dose highly orientated pyrolytic graphite data. Reproduced from Brocklehurst, J. E.; Kelly, B. T. Carbon 1993, 31(1), 179–183. 354 Graphite in Gas-Cooled Reactors fluence, although the maximum fluence is limited to 30� 1020 n cm�2 EDND. The author is not aware of any data at higher irradiation temperatures. 4.11.11.4 Modulus Changes to C33 and C44 in HOPG and natural graph- ite crystals have been reported at 50, 650, and 1000 �C53 and at 150 �C.54 For HOPG, the 150 �C data indicated that C33 slightly reduced with increas- ing irradiation (Figure 24) and this was attributed to the increase in ‘c’ axis lattice spacing. However, there is no clear trend at the other temperatures (Figure 24). In the case of shear, at a very low temperature of 50 �C there was a significant increase in C44, but at higher temperatures the increase was less (Figure 25). The data for natural crystal showed similar trends but there was significantly more scatter. The trend in the increase in C44 at the lower temperature would go towards explaining the increase in modulus in polycrystalline data at low fluence. However, it is surprising that the increase is only modest at the higher temperatures, although the maximum fast neutron fluence is very low and data is required at the intermediate temperatures. Seldin and Nezbeda53 also measured the shear strength but unfortunately there is considerable scat- ter and no definite trend. 4.11.11.5 Thermal Conductivity Taylor et al.55 measured the change in thermal conductivity in HOPG with fast neutron irradiation. The thermal conductivity along the basal planes (the ‘a’ direction) is much greater than the value perpendicular to the basal planes (the ‘c’ direction). Taylor et al. also measured the change in thermal resistivity in irradiated graphite, and when this data is normalized, the data indicated that thermal resistivity temperature dependence changed with irradiation as given in Figure 26. This is the so-called ‘d’ curve that is used in the United Kingdom to predict thermal resistivity in irradiated graphite. 4.11.11.6 Raman Figure 27 gives Raman spectra for unirradiated and irradiated graphite as well as for baked carbon. In the spectral range shown, there is a prominent G-peak at 1580 cm�1 associated with the basal plane bond stretching of ‘c’ axis sp2 atoms. The D-peak at 1350 cm�1 is associated with the breathing mode of sp2 atoms and disordered carbon structure. The sec- ond D-peak at 2700 cm�1 is indicative of the crystal- line structure of the graphite. In Figure 28 the normalized positions of the G- and D-peaks, and the ratio of the peak intensities are compared for various graphites (unirradiated and irradiated). HOPG is obviously the most ordered structure followed by the PGA needle coke graphite and then the medium grained graphite grades. The most disordered materials are the baked carbon (NBG-18 baked) followed by irradiated BEPO (a UK test reactor) graphite. D im en si on al c ha ng e (% ) 120 100 80 60 D im en si on al c ha ng e (% ) 40 20 20 40 60 80 c-axis Fluence (1020 n cm−2 EDND) 100 120 140 160 0 0 0 −5 −10 −15 −20 −25 −30 20 40 60 80 a-axis Fluence (1020 n cm−2 EDND) 100 120 140 1600 Kelly (450 �C) Kelly (650 �C) Kelly (900 �C) Kelly (1200 �C) Bokros et al. (900 �C) Bokros et al. (1075 �C) Bokros et al. (1250 �C) Figure 21 High-temperature dimensional change data for highly orientated pyrolytic graphite. Graphite in Gas-Cooled Reactors 355 Figure 29 qualitatively demonstrates that there is a relationship between Raman spectra and crystal structure disorder. The higher the disorder, the higher are the D- and G-peak wave number and I(D/G) ratio. In Figure 29(a), the crystal length La has been calculated from the full width at half maximum (FWHM) using the method proposed by Tuinstra and Koenig.56 Both figures demonstrate that Raman can be used to quantify the disorder in the graphite structure, either as manufactured or due to irradiation. 4.11.12 Property Changes in Irradiated Polycrystalline Graphite Fast neutron irradiation and, in the case of car- bon dioxide-cooled reactors, radiolytic oxidation change many of the properties of graphite. The properties of interest to the nuclear engineer are the following: � Stored energy – a function of fast neutron damage and temperature, due to damage to the graphite crystallites, but not affected by radiolytic oxidation other than by a reduction in mass. � Specific heat – a function of temperature but not affected by fast neutron irradiation or radiolytic oxidation other than by stored energy, which may be considered separately. � Dimensional changes – a function of fast neutron damage and irradiation temperature. There is also some evidence of modification by radiolytic oxida- tion. It isalsomodifiedbystress (seeirradiationcreep). � CTE – a function of temperature, fast neutron damage, irradiation temperature, and stress. There is evidence that it is not modified by radio- lytic weight loss. 30 28 26 24 22 20 18 C oe ffi ci en t of t he rm al e xp an si on (1 0− 6 K −1 ) C oe ffi ci en t of t he rm al e xp an si on (1 0− 6 K −1 ) 16 14 12 10 0 2.0 1.5 1.0 0.5 0.0 −0.5 −1.0 −1.5 −2.0 −2.5 0 2 4 6 8 10 12 14 16 18 2 4 6 8 c-axis Fluence (1020 n cm−2 EDND) a-axis Fluence (1020 n cm−2 EDND) 10 12 14 16 150 �C 170 �C 200 �C 250 �C Figure 22 Low-temperature changes to coefficient of thermal expansion in highly orientated pyrolytic graphite. Modified from Kelly, B. T.; Martin, W. H.; Nettley, P. T. Philos. Trans. R. Soc. Lond. A Math. Phys. Sci. 1966, 260(1109), 37–49. 356 Graphite in Gas-Cooled Reactors � Thermal conductivity – a function of temperature, fast neutron damage, and irradiation tempera- ture. It is significantly modified by radiolytic weight loss. � Young’s modulus – a function of temperature, fast neutron damage, and irradiation temperature. It is significantly modified by radiolytic weight loss. � Strength (tensile, compressive, flexural, and frac- ture) – a function of temperature, fast neutron damage, and irradiation temperature. It is signifi- cantly modified by radiolytic weight loss. � Electrical resistivity – a function of temperature, fast neutron damage, and irradiation temperature. It is probably modified by radiolytic weight loss. � Irradiation creep – a function of fast neutron dam- age, irradiation temperature, and stress. These property changes are illustrated in Figure 30. These dimensional changes, property changes and creep mechanisms are correlated, some more strongly than others. This has been taken advantage of in various semiempirical models for irradiation damage in graphite.57–60 Indiscussing the irradiation behavior of polycrystal- line graphite, it is useful to split these changes into low, medium, and high fluence effects. At low irradiation, fluence changes in polycrystalline graphite are strongly correlated with the crystallite changes discussed else- where; seeSection4.11.11. Typicalmechanismswould 30 25 20 15 10 C oe ffi ci en t of t he rm al e xp an si on (1 0− 6 K −1 ) C oe ffi ci en t of t he rm al e xp an si on (1 0− 6 K −1 ) 5 0 0 0 −1 −2 −3 −4 −5 0 5 10 15 20 25 30 5 10 15 c-axis a-axis Fluence (1020 n cm−2 EDN) Fluence (1020 n cm−2 EDND) 20 25 350 �C 450 �C 650 �C 350 �C 300 �C 450 �C 650 �C 30 Figure 23 High-temperature changes to coefficient of thermal expansion in highly orientated pyrolytic graphite. Reproduced from Brocklehurst, J. E.; Kelly, B. T. Carbon 1993, 31(1), 179–183. Graphite in Gas-Cooled Reactors 357 be the accumulation of stored energy, pinning (leading to a rapid increase in Young’s modulus), and the rapid decrease in thermal conductivity. At medium fluence, several of the properties saturate, such as Young’s mod- ulus and thermal conductivity. At high fluence, when crystallite growth in the ‘c’ direction has taken upmuch of the accommodation provided by Mrozowski cracks,11 and larger ‘cracks,’ the polycrystalline struc- ture starts to become strained, thereby generating new cracking. At extremely high fluence, beyond that expe- rienced in a modern power production reactor, the crystallite swelling becomes so large that the polycrys- talline structure completely breaks down leading to a rapid decrease in modulus and strength. Each of the property changes is discussed in more detail below. In attempting to understand the behavior of polycrystalline graphite, reference is made to the irradiation behavior of HOPG, as previously discussed inSection4.11.11. This is becauseHOPGis considered to be a representative model material for the individual crystallite structures in polycrystalline graphite. 4.11.13 Averaging Relationships Before looking at each of the properties individually, it is first worth considering the methods developed to relate changes in the crystallites to the bulk 40 35 30 25 20 E la st ic m od ul us C 33 (G N m −2 ) 15 10 5 0 0.0 0.5 1.0 1.5 Fluence (1020 n cm−2 EDND) 2.0 2.5 50 �C 150 �C 650 �C 1000 �C Figure 24 Reduction in C33 in irradiated highly orientated pyrolytic graphite. Modified from Seldin, E. J.; Nezbeda, C. W. J. Appl. Phys. 1970, 41(8), 3389–3400; Summers, L.; Walker, D. C. B.; Kelly, B. T. Philos. Mag. 1966, 14(128), 317–323. 5 4 3 2 E la st ic m od ul us C 44 (G N m −2 ) 1 0 0 10 20 30 40 50 �C Fluence (1017 n cm−2) 50 60 70 80 E la st ic m od ul us C 44 (G N m −2 ) 650 and 1000 �C 650 �C 1000 �C Fluence (1020 n cm−2) 0.0 0.0 0.1 0.2 0.3 0.4 0.5 0.2 0.4 0.6 0.8 1.0 1.2 1.4 Figure 25 Changes in C44 with irradiation in highly orientated pyrolytic graphite. Modified from Seldin, E. J.; Nezbeda, C. W. J. Appl. Phys. 1970, 41(8), 3389–3400. 358 Graphite in Gas-Cooled Reactors properties of polycrystalline graphite. In the 1960s, Simmons derived a model based on the following assumptions: polycrystalline graphite could be con- sidered to be a single phase, porous aggregate of identical crystals with the correct graphite symmetry, and small element volumes containing only graphite could be chosen so that their internal and external stress could be considered to be uniform.61 From first principles and by applying the laws of thermodynamics for the assumptions given above, Simmons derived the following relationship for the bulk dimensional change rate: dexx dg ¼ Axdxc dg þ ð1� AxÞdxa dg ½33� where dexx/dg, dxc/dg, and dxa/dg are the linear dimensional change rate in polycrystalline graphite in some direction ‘x’ and the crystallite dimensional change rates in the ‘a’ and ‘c’ directions respectively. Simmons also derived the following relationship for the bulk CTE: axx ¼ Axac þ ð1� AxÞaa ½34� where axx, ac, and aa are the linear CTE in poly- crystalline graphite in some direction ‘x’ and the crystallite CTE in the ‘a’ and ‘c’ directions, respectively. The so-called structure factor ‘Ax’ is the summa- tion rate of change in rate of the crystallite stresses with respect to the change in bulk stress, as illustrated in the equations below. Ax ¼ X n an @s033;n @sxx � � ½35� 0.7 270 �C 300 �C 350 �C 400 �C 460 �C 510 �C 780 �C 0.6 0.5 0.4 Fr ac tio na l c ha ng e in t he rm al c on d uc tiv ity (K 0 K -1 ) 0.3 0.2 0.1 0.0 0 50 (a) 100 150 Fluence (1020 n cm−2 EDND) 200 250 300 10.0 5 4 3 2 1 0 0 100 200 300 400 500 600 700 800 1.0 0.1 C ha ng e in t he rm al re si st an ce (c m 0 K W -1 ) 0.0 0 100 200 300 400 Temperature (�K) Temperature (�K) 500 600 700 800 N or m al iz ed c ha ng e in th er m al r es is ta nc e (b) (c) 30 �C 150 �C 300 �C 450 �C 150 �C (annealed) 300 �C (annealed) 450 �C (annealed) Figure 26 Changes in the thermal conductivity of highly orientated pyrolytic graphite. (a) Change in thermal conductivity in the ‘a’ direction as a function of fluence, (b) Change in thermal resistance as a function of temperature, and (c) Normalized change in thermal resistivity as a function of irradiation and measurement temperature. Reproduced from Taylor, R.; Kelly, B. T.; Gilchrist, K. E. J. Phys. Chem. Solids 1969, 30, 2251–2267. Graphite in Gas-Cooled Reactors 359 ð1� AxÞ ¼ X n an @ @sxx ðs022;n þ s011;nÞ � For a more detailed derivation of these equations see Hall et al.61 In addition, Simmons4 provided evidence that there is a linear relationship between unirradiated CTE and initial dimensional change rate for poly- crystalline graphite (Figure 31). Brocklehurst and Bishop62 later found a similar relationship in bromi- nated graphite. Expressions to those of Simmons have been derived by Sutton and Howard12: apar ¼ K1gac þ K2baa aperp ¼ K3gac þ K4baa ½36� where K1, K2, K3, K4, g, and b are crystal accommoda- tion and Bacon13 crystal orientation factors. A similar but more complex relationship was also derived by Jenkins.63 In the discussion of the individual properties, the anisotropic graphite PGA and the semi-isotropic Gilsocarbon graphite are used as examples. 4.11.14 Dimensional Change As discussed previously (see Section 4.11.11.2), when irradiated, graphite crystallites, as simulated using HOPG, expand significantly in the ‘c’ direction and shrink in the ‘a’ direction. These dimensional changes are reflected in the behavior of polycrystal- line graphite, but the volumetric changes, although relatively large are much smaller than those seen in HOPG. The reason for this is attributed to the many microcracks, which range in size from the nano- to microscale; see Figure 32. While these cracks can accommodate the crystal growth in the ‘c,’ the shrink- age in the ‘a’ direction will directly be reflected in the polycrystalline behavior. However, as previously BEPO-20 BEPO-16 In te ns ity (a .u .) HTR 1 baked HTR 2 HTR 1 Gilsocarbon PGA 1000 2000 Wavenumber (cm−1) 3000 1000 2000 3000 HOPG Figure 27 Raman spectra for unirradiated and irradiated graphite and baked carbon. Courtesy of A. Jones, University of Manchester. 15 10 N or m al iz ed G p ea k p os iti on (c m −1 ) 5 0 Graphite grade(a) H O P G P G A G ils oc ar b on H TR 1 H TR 1 b ak ed H TR 3 b ak ed B E P O -1 B E P O -1 6 B E P O -2 0 B E P O -1 3a H TR 2 ir ra d ia te d H TR 3 ir ra d ia te d H TR 2 H TR 3 1.4 1.2 1.0 0.8 I(D )/ (G ) 0.6 0.4 0.2 0.0 (c) Grap H O P G P G A G ils oc ar b on H TR 1 H TR 1 b ak ed H TR 2 Figure 28 Relative position and ratio of I(D/G) by graphite gra (b) Normalized position of D-Peak, (c) Ratio of the D-peak and G Manchester. 360 Graphite in Gas-Cooled Reactors discussed, the crystallite dimensional change rate is much greater below �300 �C than above that temperature. 4.11.14.1 Pile Grade A The blocks of PGA were manufactured by extrud- ing needle-shaped filler particles mixed with a pitch binder. During the extrusion process, the needle-shaped filler particles tended to align with the ‘a’ axis parallel, and the ‘c’ axis perpendicular, to the extrusion direction. Thus, the final product (or block) had two orthotropic directions: parallel to the extrusion direction (WG) and perpendicular to the extrusion direction (AG). This strong ori- entation in direction is reflected not only in the unirradiated properties but also in the irradiation properties and dimensional changes. Dimensional N or m al iz ed D p ea k p os iti on (c m −1 ) 10 5 0 (b) Graphite grade H O P G P G A G ils oc ar b on H TR 1 H TR 1 b ak ed H TR 3 b ak ed B E P O -1 B E P O -1 6 B E P O -2 0 B E P O -1 3a H TR 2 ir ra d ia te d H TR 3 ir ra d ia te d H TR 2 H TR 3 hite grade H TR 3 b ak ed B E P O -1 B E P O -1 6 B E P O -2 0 B E P O -1 3a H TR 2 ir ra d ia te d H TR 3 ir ra d ia te d H TR 3 de and condition. (a) Normalized position of G-Peak, -peak intensities. Courtesy of A. Jones, University of 1.4 1.2 Irradiated Baked HTR 3 HTR 2 HTR 1 PGA Gilsocarbon BEPO-1 1.0 0.8 I(D /G ) 0.6 0.4 0.2 0.0 0 20 40 (a) (b) 60 80 Jones Knight and White Coke 120 100 80 60 FW H M (c m −1 ) 40 20 0 0 0.2 0.4 0.6 0.8 I(D/G) 1 1.2 1.4 Glassy carbon HTR 1 baked HTR 3 baked Polycrystalline graphitesHOPG Irradiated graphite La (A) 100 120 140 Figure 29 Quantitative relationship between I(D/G) ratio and crystallite length (La) and full width at half maximum. (a) I(D/G) as a function of crystal size (La) and (b) FWHM as a function of I(D/G). Courtesy of A. Jones, University of Manchester. 0.10 2 1 0 −1 −2 −3 0.05 −0.05 0.00 D im en si on al c ha ng e (% ) D im en si on al c ha ng e ra te (% p er flu en ce ) C TE (K −1 � 10 −6 ) 0 50 (a) (b) 100 Fluence (n cm−2� 1020 EDND) Fluence (n cm−2� 1020 EDND) 150 200 2 4 40 3.0 8 7 6 5 4 3 2 1 0 2.5 2.0 1.5 1.0 0.5 0.0 0 50 100 150 200 30 20e su E /E o− 1 K o/ K −1 10 0 Figure 30 Schematic of the irradiation-induced changes in Gilsocarbon graphite irradiated at 550 �C (note that there will be a similar set of curves for each irradiation temperature). (a) Dimensional change, dimensional change rate, and coefficient of thermal expansion and (b) Factorial change in Young’s modulus (E/E0�1) and thermal conductivity (K0/K�1), and irradiation creep (elastic strain units, esu). Graphite in Gas-Cooled Reactors 361 (a) (b)1µm 500µm Figure 32 Graphite microstructure showing Mrozowski cracking and larger porosity. (a) Mrozowski cracks and (b) Polarized optical image of Gilsocarbon microstructure. Courtesy of A. Jones, University of Manchester. 0.6 0.5 0.4 0.3 0.2 0.1 0.0 0 (a) (b) (c) 1 2 CTE (10−6 K−1) CTE (10−6 K−1) 3 4 0.5 0.4 0.3 0.2 D im en si on al c ha ng e ra te (% p er 10 20 n cm −2 ) D im en si on al c ha ng e ra te (% p er 10 20 n cm −2 ) D im en si on al c ha ng e ra te (% p er 10 20 n cm −2 ) 0.1 0.0 0.08 0.06 0.04 0.02 0.00 0 1 2 3 4 5 −0.02 −0.04 0 80 �C 30 �C 180 �C 1 2 3 CTE (10−6 K−1) 4 5 6 Parallel to extrusion Perpendicular to extrusion Figure 31 Correlation between initial growth rate and unirradiated coefficient of thermal expansion. Modified from Simmons, J. Radiation Damage in Graphite; Pergamon: London, 1965. 362 Graphite in Gas-Cooled Reactors change MTR data for PGA is given in Birch and Brocklehurst64 for both the parallel (WG) and per- pendicular (AG) directions. In the parallel direc- tion and below 300 �C, the large dimensional change rate in the crystallite ‘c’ axis causes the graphite to swell. Above 300 �C and parallel to the extrusion direction, the graphite shrinks in the lower fluence range. This behavior can 6 5 4 3 2 D im en si on al c ha ng e (% ) D im en si on al c ha ng e (% ) 1 0 0 (a) 0.0 0 (b) 10 20 30 40 50 60 −0.5 −1.0 −1.5 −2.0 −2.5 −3.0 10 20 30 Fluence (1020 n cm−2 EDND) Fluence (1020 n cm−2 EDND) 40 50 60 −1 −2 150 �C 200 �C 250 �C 300 �C 350 �C 450 �C 650 �C219 �C 225 �C Figure 33 Low- to medium-fluence irradiation dimensional change in pile grade A graphite. (a) perpendicular to extrusion and (b) parallel to extrusion. Graphite in Gas-Cooled Reactors 363 be compared to that of HOPG, as given in Figure 33. If PGA is irradiated to a higher fluence, the shrinkage rate reduces until the graphite begins to expand or ‘turns around,’ as illustrated in Figure 34. ‘Turnaround’ is associated with the closure of the Mrozowski cracks; see Figure 32. When all of the accommodation provided by the cracks has been taken up, the larger ‘c’ crystallite dimensional change rate would be expected to dominate the ‘a’ axis shrinkage rate. This behavior has been used, with some success, to model the dimensional change behavior in PGA, Gilsocarbon, and Russian GR-280 graphite.48,65,66 4.11.14.2 Gilsocarbon As previously discussed, Gilsocarbon graphite for the AGRs was manufactured by molding (or press- ing) the spherical filler particles and blocks, result- ing in a semi-isotropic graphite with an anisotropy ratio of �1.01 (based on the ratio of the orthotropic CTE values). Dimensional change MTR data for Gilsocarbon over a wide range of temperatures is given in Figure 35. There are two sets of data at each temperature; one WG and one AG. This illustrates how isotropic the properties of Gilsocar- bon are, even when irradiated. In Figure 35, it is also clearly illustrated that the higher the irradiation 6 5 4 3 2 1 0 D im en si on al c ha ng e (% ) 0 50 100 Fluence (1020 n cm−2 EDND) 150 200 430 �C 600 �C 900 �C 940 �C 1240 �C 1430 �C 250 −1 −2 −3 −4 Figure 35 Dimensional change in Gilsocarbon graphite. Modified from Birch, M.; Brocklehurst, J. E. A review of the behaviour of graphite under the conditions appropriate for the protection of the first wall of a fusion reactor; UKAEA, ND-R-1434(S); 1987; Brocklehurst, J. E.; Kelly, B. T. Carbon 1993, 31(1), 155–178. 4 2 0 0 20 Parallel to extrusion Perpendicular to extrusion 40 60 Fluence (1020 n cm−2 EDND) 80 100 120 140 160 D im en si on al c ha ng e (% ) −2 −4 −6 −8 −10 Figure 34 High-fluence irradiation dimensional change in pile grade A graphite irradiated at 600 �C. Reproduced from Brocklehurst, J. E.; Kelly, B. T. Carbon 1993, 31(1), 155–178. 364 Graphite in Gas-Cooled Reactors temperature, the sooner the turnaround is reached. At very low fluence, semi-isotropic graphite swells. This swelling can be quite significant as demon- strated by the irradiation of semi-isotropic NBG- 10 at 294 and 691 �C.67 This behavior has been attributed to the annealing out of residual machining stresses or shrinkage strains, but there are no micro- structural or other experimental observations to val- idate this reasoning. When graphite is irradiated past turnaround and reaches its original volume, sometimes referred to as ‘critical fluence,’ the structure of the graphite begins to break down, as illustrated in Figure 36. 4.11.14.3 Effect of Radiolytic Oxidation on Dimensional Change When designing the UK AGRs, irradiation experi- ments in a carbon dioxide atmosphere were carried out in BR-2 at Mol, Belgium. These experiments were designed to obtain high radiolytic weight loss (�35%) in a very short time, and hence, a low fast 3 2 1 0 D im en si on al c ha ng e (% ) −1 −2 −3 −4 −5 0 50 100 150 Fluence (1020 n cm−2 EDND) 200 250 300 NA NA oxidized, x 366 Graphite in Gas-Cooled Reactors graphite MTR experiments designed to carry out simultaneous radiolytic oxidation and fast neutron damage under power reactor conditions were aban- doned because of the closure of the UKAEA MTRs at Harwell in 1990. It is therefore unclear how signif- icant this behavior is for the AGRs. However, there are nowMTRexperiments being undertaken in HFR (High Flux Reactor) at Petten, the Netherlands to try and address this. 4.11.14.4 Dimensional Change Rate The constitutive models used to predict stresses in polycrystalline components often do not use dimen- sional change directly but use dimensional change rate. The dimensional change rate and dimensional change of Gilsocarbon graphite irradiated at 550 �C are com- pared in the schematic shown in Figure 38. The turn- around in dimensional change rate occurs earlier than turnaround in dimensional change. In channel-type reactors such as an AGR or Magnox reactor, it is the turnaround in rate that is associated with the peak in- bore stress. Thus, when planning a nuclear graphite MTR experiment, it is important to obtain data in the low tomedium fluence range, aswell as at high fluence. 4.11.15 Coefficient of Thermal Expansion The mechanism that drives the changes to the CTE is, at present, not well understood because of the lack of microstructural studies in this area. In the United Kingdom, CTE data is usually quoted over the range 0.5 0.0 D im en si on al c ha ng e (% ) −0.5 −1.0 −1.5 −2.0 −2.5 −3.0 0 50 100 Fluence (1020 n c Dimensional change Dimensional change rate Figure 38 Schematic of dimensional change and dimensiona 20–120 �C but, for increased accuracy, it is better to measure CTE over a much larger temperature range. For use in graphite component assessment, the CTE needs to be converted to an appropriate temperature range, that is, room temperature to irradiation tem- perature (20�Tirr). At irradiation temperatures above 300 �C, crystal- lite CTE, as measured on HOPG,49 is assumed to be invariant to fast neutron fluence. Thus, the irradia- tion changes in polycrystalline CTE are assumed to be related to closure of Mrozowski-type cracks. 4.11.15.1 Pile Grade A As previously discussed, PGA graphite is an aniso- tropic material, and this is also reflected in the CTE irradiation data. CTE measurements on irradiated PGA from an MTR program are given in Figure 39. Unfortunately, there is a considerable amount of scat- ter in this data, probably related to the difficulty in measuring CTE on irradiated graphite at the time these measurements were taken. However, there are some interesting trends illustrated in Figure 39. At the lower irradiation fluence and irradiation temperature below 300 �C, there is a rapid increase in CTE, which in the case of the perpendicular (AG) direction, is followed by a rapid reduction. This is reminiscent of the irradiation CTE behavior of HOPG in the ‘c’ direction when irradiated below 300 �C.35 At the higher irradiation temperatures (above 300 �C), the CTE first increases and then increases again. The initial increase has been associated with closure of Mrozowski cracks due to ‘c’ axis crystal growth. 0.08 0.06 0.04 0.02 0.00 −0.02 −0.04 −0.06 200 D im en si on al c ha ng e ra te (% p er 10 20 n cm −2 E D N D ) m−2 EDND) 150 l change rate in Gilsocarbon graphite irradiated at 550 �C. 0 0 2 3 4 5 6 7 10 20 30 40 50 60 0.0 0.5 1.0 1.5 C TE (1 0- 6 K -1 ) C TE (1 0- 6 K -1 ) 2.0 2.5 3.0 10 20 30 40 50 60 70 150 °C 200 °C 225 °C 250 °C 300 °C 350 °C 450 °C 650 °C 150 °C 200 °C 250 °C 300 °C 350 °C 450 °C 650 °C Fluence (10 20n cm−2 EDND) Fluence (1020n cm−2 EDND) (a) (b) Figure 39 Changes in the coefficient of thermal expansion of pile grade A graphite with irradiation. (a) parallel to extrusion and (b) perpendicular to extrusion. Modified from Birch, M.; Brocklehurst, J. E. A review of the behaviour of graphite under the conditions appropriate for the protection of the first wall of a fusion reactor; UKAEA, ND-R-1434(S); 1987. Graphite in Gas-Cooled Reactors 367 The fall in CTE in the perpendicular direction has been attributed to the generation of new cracks caused by the large growth of the crystallites. Electron micro- graphs due to Sutton and Howard12 appear to support the first assumption. Unfortunately, there is no evi- dence for the latter assumption. 4.11.15.2 Gilsocarbon Irradiation CTE data for semi-isotropic Gilsocarbon is given in Figure 40. This later CTE data has much less scatter. As with PGA, there is an initial rise in CTE, attributed to the closure of Mrozowski-type cracks, followed by a fall with increasing fluence. Whether this is the same mechanism as the perpen- dicular PGA data is unclear. However, this behavior is typical of medium-grained nuclear graphite. At very high fast neutron fluence, the CTE appears to start to saturate at a lower value than in the case of virgin Gilsocarbon, although there is insufficient data to confirm this behavior. 4.11.15.3 Methodology for Converting Between Temperature Ranges Using Simmons’s relationship,61 the instantaneous CTE at two different temperatures represented by ax and a0x can be written as ax ¼ Axac þ ð1� AxÞaa a0x ¼ Axa0c þ ð1� AxÞa0a ½37� Rearranging leads to the expression a0x ¼ ax � aa ac � aa � � ða0c � a0aÞ þ a0a ¼ Aiax þ Bi ½38� or for mean CTE að20�Ti Þ ¼Aiað20�120Þ þBi ½39� 7 430 �C 600 �C 940 and 1240 �C 6 5 4 3 C TE (1 0− 6 K −1 ) 2 1 0 0 50 100 150 Fluence (1020 n cm−2 EDND) 200 250 Figure 40 Irradiation coefficient of thermal expansion data for Gilsocarbon graphite. Modified from Birch, M.; Brocklehurst, J. E. A review of the behaviour of graphite under the conditions appropriate for the protection of the first wall of a fusion reactor; UKAEA, ND-R-1434(S); 1987; Brocklehurst, J. E.; Kelly, B. T. Carbon 1993, 31(1), 155–178. 6 5 4 C TE (1 0− 6 K −1 ) 3 2 0 10 20 30 40 Weight loss (%) 50 60 Inert Continuous oxidation Interchange (inert and oxidizing) 70 80 Figure 41 Coefficient of thermal expansion of Gilsocarbon graphite irradiated in inert and oxidizing atmospheres. 368 Graphite in Gas-Cooled Reactors Experimental data for aa and ac as a function of temperature can be used to calculate the tempera- ture dependence of A and B . The method has been compared to other empirical approaches used generally in industry and was found to give similar results.69 However, the authors have never seen the method validated for irradiated, radiolytically oxi- dized graphite. 4.11.15.4 Effect of Radiolytic Oxidation on CTE Figure 41 shows data for the combined fast neutron irradiation and radiolytic oxidation of Gilsocarbon graphite over a limited fluence range, which appears to show no effect of radiolytic oxidation on CTE. This is also supported by data on thermally oxidized (up to �60% weight loss) PGA and Gilsocarbon graphite70 (see Figure 42). 4.11.16 Thermal Conductivity Thermal conductivity in nuclear graphite is usually determined by measuring thermal diffusivity using the laser flash method at �30 �C. The mechanism for thermal conductivity in graphite over the tem- peratures of interest in nuclear reactors is lattice vibration (phonon) conductance. There is a pro- nounced reduction in thermal conductivity with increased temperature attributed to phonon– phonon scattering. At low irradiation fluence, there 9 8 7 6 5 4 C TE (1 0− 6 K −1 ) 3 2 1 0 0 10 20 30 Weight loss (%) 40 50 Gilsocarbon PGA (parallel to extrusion) PGA (perpendicular to extrusion) 60 70 Figure 42 Coefficient of thermal expansion of thermally oxidized graphite. Modified from Hacker, P. J.; Neighbour, G. B.; McEnaney, B. J. Phys. D Appl. Phys. 2000, 33, 991–998. Graphite in Gas-Cooled Reactors 369 is a significant decrease in thermal conductivity due to fast neutron irradiation, attributed to an increase in scattering in the damaged lattice. This decrease in thermal conductivity (or increase in thermal resis- tivity) saturates in the medium fluence range. At very high fluence, there is a secondary decrease attributed to microcracking due to high crystallite strains. There is also evidence of change in temper- ature dependence with irradiation. The thermal conductivity in crystallite basal plane is much larger than that perpendicular to basal plain; thus, Ka�Kc. The thermal resistivity can be described by the equation below: 1 Ka ¼ 1 KB þ 1 KU þ 1 KD ½40� � U – Umklapp scattering (German for turnover/ down) or phonon–phonon scattering (due to increase in temperature) � D – scattering due to defects (caused by irradiation) � B – boundary scattering (structural effects) Changes to these resistances will be reflected in the thermal conductivityof polycrystalline graphite. Ther- mal conductivity is significantly decreased by radio- lytic oxidation. Data is usually presented as the reciprocal of conductivity, that is, thermal resistivity. 4.11.16.1 Pile Grade A Although PGA is significantly anisotropic over the range of interest to the Magnox reactors, the change in thermal resistivity can, for practical purposes, be considered as invariant to grain direction. Changes in thermal resistivity in PGA graphite are given in Figure 43. There is a significant change in the rate of increases in thermal resistivity between 250 and 300 �C, giving a similar trend to the change in crystal growth rates between these two temperatures. The increase in thermal resistivity is significant (a factor of 100) at 150 �C, for a relatively low fluence. The low irradiation temperature data for PGA in Figure 43 do not reach a high enough fluence to saturate. However, at the higher temperatures, data is near saturation. 4.11.16.2 Gilsocarbon The irradiation-induced changes in the thermal resistivity of Gilsocarbon graphite are isotropic, as seen in Figure 44. This data extends to a much higher fluence than the PGA data, and the secondary increase in thermal conductivity, associated with cracking of the microstructure by the large swelling of the crystallites is clearly seen. This behavior occurs after dimensional change turnaround. 4.11.16.3 Thermal Conductivity Temperature Dependence of Irradiated Graphite The temperature dependence of irradiated HOPG parallel to the basal plane (the ‘a’ direction) was measured by Taylor et al.55 When the data is normal- ized to 30 �C, a single curve was produced, the so- called temperature-dependent ‘d’ curve: dðTÞ ¼ Ka;ið30Þ Ka;iðT Þ ½41� 12 10 8 6 Fr ac tio na l c ha ng e in t he rm al r es is tiv ity (K 0 /K −1 ) 4 2 0 0 50 100 Fluence (1020 n cm−2 EDND) 150 430 �C 600 �C 940 �C 1240 �C 200 250 Figure 44 Changes in thermal resistivity of Gilsocarbon graphite. Adapted from Birch, M.; Brocklehurst, J. E. A review of the behaviour of graphite under the conditions appropriate for the protection of the first wall of a fusion reactor; UKAEA, ND-R-1434(S); 1987. 120 100 80 60 Fr ac tio na l c ha ng e in t he rm al r es is tiv ity (K 0/ K −1 ) 40 20 0 0 5 10 15 20 Fluence (1020 n cm−2 EDND) 25 30 35 150 �C 250 �C 350 �C 450 �C 650 �C 300 �C 200 �C 225 �C 40 45 50 Figure 43 Changes in thermal resistivity of pile grade A graphite. Modified from Birch, M.; Brocklehurst, J. E. A review of the behaviour of graphite under the conditions appropriate for the protection of the first wall of a fusion reactor; UKAEA, ND-R-1434(S); 1987. 370 Graphite in Gas-Cooled Reactors This curve is used in the United Kingdom to predict the thermal conductivity of irradiated polycrystalline graphite at the required assessment temperature. However, above around 300 �C, ‘d’ can be taken as unity for all practical purposes. 4.11.16.4 Predicting the Thermal Conductivity of Irradiated Graphite for Reactor Core Assessments A method derived by Kelly71 is often used in the prediction of irradiated graphite thermal conductivity in the assessment of UK reactor cores. The thermal resistivity, 1/Ki(T), induced because of irradiation as a function of temperature,T, can be given by the differ- ence between the thermal resistivity due to fast neu- tron damage, and the unirradiated thermal resistivity: 1 KiðT Þ ¼ 1 KirrðT Þ � 1 K0ðT Þ ½42� However, dðTÞ ¼ Ka;ið30Þ Ka;iðT Þ ½43� Graphite in Gas-Cooled Reactors 371 and as Ka� Kc, Ka is assumed to dominate in polycrys- talline graphite, Ka,i(T)¼ Ki(T) and Ka,i(30)¼ Ki(30) ¼ Kirr(30), and hence, 1 KirrðT Þ ¼ 1 K0ðT Þ þ dðTÞ 1 Kið30Þ ½44� The irradiation-induced fractional changes in ther- mal resistivity measured at 30 �C are available for graphite irradiated at various temperatures, Tirr. Therefore, for graphite irradiated and measured at 30 �C, we can write K0ð30Þ Kið30Þ ¼ K0ð30Þ Kirrð30; 30Þ � 1 � � ¼ f ½45� Substituting in the previous equation then gives 1 KirrðT Þ ¼ 1 K0ðT Þ þ dðTÞ f K0ð30Þor K0ðT Þ KirrðTÞ ¼ 1þ dðT Þf K0ðT Þ K0ð30Þ ½46� Thus, the irradiated thermal conductivity can be pre- dicted at any temperature for graphite irradiated in an inert atmosphere. The effect of weight loss and high fast neutron fluence is dealt with by a version of the ‘product’ rule leading to 1 KirrðT Þ ¼ 1 K0ð30Þ K0ð30Þ K0ðT Þ þ dðT Þf � � Sk K0 K � � ox ½47� where Sk is the high dose reduction in thermal con- ductivity and [K0/K]ox is the reduction in thermal conductivity due to radiolytic oxidation. 0.00 0.04 Strain (%) 0.08 0.12 0 2 4 6 8 10 S tr es s (M N m -2 ) Loading Unloading (a) ( Figure 45 Stress–strain curves of unirradiated pile grade A gra H.; Orchard, J. In Fifth Conference on Carbon; Pergamon, 1962 The addition in these two factors in this way appears to be arbitrary as the author does not know of any validation of this method. Other methods do exist, but they depend on many more measurements of thermal conductivity at various irradiation and measurement temperatures.72,73 4.11.17 Young’s Modulus The unirradiated stress–strain behavior of graphite is nonlinear and exhibits hysteresis and permanent set. It is different in tension to compression (Figure 45) and graphite is also much stronger in compression than in tension. Similar curves can be found for Gilsocarbon.74 On irradiation, in an inert atmosphere, there is a rapid and significant increase in modulus attributed to pinning of dislocations in the basal plane. Also, the stress–strain behavior becomes almost linear (Figure 46). This increase soon saturates, but there is a secondary increase attributed to structure tight- ening (or closure of porosity due to high crystal strain). Finally, at very high fluence, there is a rapid fall in modulus due to the degeneration of the graphite microstructure. Brown75 also showed that the Vickers hardness of isotropic graphite was con- siderably increased by irradiation. Young’s modulus is significantly reduced by radiolytic oxidation. Graphite Young’s modulus increases with increas- ing temperature (Figure 47), which is attributed to the tightening of the structure, presumably because of the closure of microcracks; Maruyama et al.76 tested samples in vacuum to avoid thermal oxidation. 0.0 0.2 Strain (%) 0.4 0.6 0 5 10 15 20 25 Loading Unloading S tr es s (M N m -2 ) b) phite. (a) Tension and (b) Compression. Modified from Losty, ; pp 519–532. 0 2 4 6 8 10 12 14 16 0 0.02 0.04 0.06 0.08 Strain (%) Irradiated (1.0 � 1019n cm-2EDND at 50 �C) Failure stress = 47 MN m-2 Unirradiated failure stress = 23 MN m-2 S tr es s (M N m -2 ) 0.1 0.12 0.14 0.16 Figure 46 Change in the stress–strain behavior of graphite due to fast neutron irradiation. Modified from Brocklehurst, J. E. Chem. Phys. Carbon 1977, 13, 145–272. 372 Graphite in Gas-Cooled Reactors However, note the significant difference in strength at room temperature between samples tested in air and in vacuum. Several other authors have reported similar findings, attributing the difference to adsorbed moisture.77 The increase in strength with temperature is significant above 600 �C, making it of interest only for HTR reactor components. Also of interest in Figure 47 is the correlation between modulus and strength. 4.11.17.1 Relationship Between Static and Dynamic Young’s Modulus Irradiated and unirradiated Young’s modulus of nuclear graphite is usually measured either by an impulse or frequency method, giving the dynamic Young’s modulus (DYM). However, for use in com- ponent stress analysis assessments, the static Young’s modulus (SYM) is required. In addition, the UK irradiation creep modulus was originally defined using SYM, so there is a need to interconvert between the two measurements. DYM is always higher in value than SYM. Unfor- tunately, SYM has historically been defined in vari- ous ways, that is, as the chord between the origin at zero stress and half, or in some cases two-thirds the failure strength in tension or bending. There is limited amount of data on the SYM/ DYM ratio for UK nuclear graphite, as given in Figure 48.78 The data is for unirradiated graphite and for graphite irradiated in an inert atmosphere. The ratio for irradiated graphite is higher than for unirradiated graphite. At present, there are no published data on this ratio for radiolytically oxidized graphite, although safety case requirements have rekindled effort in this area in the United Kingdom. It has been shown79 that for fine-grained IG-110 nuclear graphite and PGX reflector graphite, as well as ASR-ORB baked carbon, the ratio of static to dynamic Young’s modulus strongly depends on the chord length chosen to define SYM. 4.11.17.2 Pile Grade A The orthotropic values of Young’s modulus in PGA graphite are seen in the changes due to fast neutron irradiation; see Figure 49.64 All of these measure- ments are DYM measured on small samples irra- diated in an MTR. Again, there is a distinct difference in the behavior below 300 �C compared to that above 300 �C. At low temperature (below 300 �C) and fluence, the initial increase attributed to pinning reaches a peak and then rapidly decreases. This rapid decrease has been attributed to large crys- tallite ‘c’ axis expansion, although there are no micro- structural observations to verify this postulation. There is considerable scatter in the data, which may reflect the measurement technique at that time. In some temperature ranges, the trends are defined by very few points. 4.11.17.3 Gilsocarbon Gilsocarbon graphite data extends to a much higher fluence; see Figure 50.50,64 For assessment purposes, the initial increase or pinning is normally assumed to 0.0 0.2 0.4 0.6 0.8 1.0 1.2 0 200 400 R el at iv e ch an ge in s f/ s f 0 an d E /E 0 600 800 1000 Temperature (�C) 1200 Flexural strength (b) Elastic modulus 16001400 (a) 0 10 20 30 40 50 60 70 0 200 400 Fl ex ur al s tr en gt h (M N m -2 ) 600 800 1000 Temperature (�C) 1200 1400 In air In vacuum Figure 47 The strength and modulus of IG-11 as a function of test temperature. (a) Flexural strength and (b) Flexural strength and modulus. Reproduced from Maruyama, T.; Eto, M.; Oku, T. Carbon 1987, 25(6), 723–726. 30 25 20 15 10 5 0 0 5 10 15 Dynamic Young�s modulus (GN m−2) S ta tic Y ou ng �s m od ul us (G N m −2 ) 20 25 30 Unirradiated Irradiated Figure 48 Ratio between static and dynamic Young’s modulus. Modified fromBrocklehurst, J. E.; Brown, R. G. The relation between strength, modulus, and structural changes of isotropic graphites irradiated to high fast neutron doses in DFR; UKAEA, TRG-M-5985 (AB 7/22015); 1972. Graphite in Gas-Cooled Reactors 373 0 (a) (b) 0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.6 1.8 0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.6 1.8 2.0 5 10 15 20 Fluence (1020n cm-2EDND) Fr ac tio na l c ha ng e in Y ou ng ’s m od ul us (E /E 0− 1) Fr ac tio na l c ha ng e in Y ou ng ’s m od ul us (E /E 0− 1) 25 30 35 40 0 5 10 15 20 Fluence (1020n cm-2EDND) 25 30 150 ºC 200 ºC 250 ºC 300 ºC 350 ºC 450 ºC 650 ºC 150 ºC 200 ºC 250 ºC 450 ºC 650 ºC Figure 49 Changes in Young’s modulus of pile grade A due to fast neutron irradiation. (a) Parallel to extrusion and (b) Perpendicular to extrusion. Modified from Birch, M.; Brocklehurst, J. E. A review of the behaviour of graphite under the conditions appropriate for the protection of the first wall of a fusion reactor; UKAEA, ND-R-1434(S); 1987. 374 Graphite in Gas-Cooled Reactors occur instantaneously on irradiation. After saturation, there is a significant increase in modulus with increase in fluence followed by a reduction. This behavior is often referred to as a ‘structural’ effect. At very high fluence, the decrease in modulus due to degradation of the microstructure is clearly illu- strated (Figure 50). 4.11.17.4 Separation of Structure and Pinning Terms Various proposed graphite irradiation models, includ- ing the UKAEA irradiation creep models, require the pinning and structural terms to be separated. Origi- nally, the pinning and structure terms for Gilsocarbon were separated by ‘eye,’ which is subjective. Recently, more sophisticated statistical, pattern recognition and curve fitting methods have been used.60 Having separated out the pinning term, the struc- ture term is a function of irradiation temperature and dose, whereas the pinning term is only a function of temperature. The pinning term is assumed to not affect creep rate, whereas the structure term affects it. In some models, the structure term is also assumed to be a function of radiolytic weight loss and to be correlated to dimensional change. 4.11.17.5 Effect of Radiolytic Weight Loss on Dimensional Change and Young’s Modulus Dimensional change andmodulusMTR data on small preoxidized samples, see Figure 51, appear to indi- cate that radiolytic weight loss would be expected, not only to increase dimensional change shrinkage and delay turnaround, but also to delay the structural 3.5 430 �C 600 �C 940 �C 1240 �C 1430 �C 3.0 2.5 2.0 1.5 1.0 0.5 Fr ac tio na l c ha ng e in Y ou ng ’s m od ul us (E /E 0− 1 ) 0.0 −0.5 −1.0 0 50 100 Fluence (1020 n cm−2 EDND) 150 200 250 Figure 50 Changes in Young’s modulus of Gilsocarbon due to fast neutron irradiation. Modified from Birch, M.; Brocklehurst, J. E. A review of the behaviour of graphite under the conditions appropriate for the protection of the first wall of a fusion reactor; UKAEA, ND-R-1434(S); 1987; Brocklehurst, J. E.; Kelly, B. T. Carbon 1993, 31(1), 155–178. Graphite in Gas-Cooled Reactors 375 increase in Young’s modulus. This correlation has been used as the basis of models to predict dimen- sional change in radiolytically oxidized graphite.51 4.11.17.6 Small Specimen Strength Graphite is stronger in bend than in tension, and stronger in compression than in bend. Irradiated strength tends to be correlated with Young’s modulus. Graphite strength is significantly reduced by radio- lytic oxidation. Losty and Orchard80 used the Griffith theory to try and demonstrate that the change in strength can be related to the square root of modulus as follows: Failure stress s is proportional to the square root of the product of strain energy release rate g and Young’s modulus divided by the crack length c. s2 / 2gE pc ½48� Thus, assuming that strain energy release rate, g, is not changed by irradiation and the critical crack, s s0 � � ¼ E E0 � �1=2 ½49� Figure 52 was purported to support this relationship. However, statistical scrutiny of the data given in this figure revealed that there is not enough data to support the argument in favor of the square root law, or even a relationship to another power. There is a significant amount of data indicating that the relationship may be more appropriate as a direct relationship.81 4.11.18 Effect of Radiolytic Oxidation on Thermal Conductivity, Young’s Modulus, and Strength Figure 53 illustrates that thermal conductivity, strength, and Young’s modulus are all significantly reduced by radiolytic oxidation.82 The data is usually fitted to a simple exponential decay83 of the form [P/P0]¼ exp(�lx). However, there must be a practi- cal ‘percolation limit’ to this law when all the porosity joins together and properties reduce to zero. This was recently highlighted in the statistical analysis of high weight loss PGA data by McNally et al.84 4.11.19 The Use of the Product Rule The product rule has been used for many years in the United Kingdom to combine changes due to fast neutron irradiation and radiolytic oxidation for strength, modulus, and thermal conductivity. Exam- ples are given below: Strength s s0 � � ¼ s s0 � � irr � s s0 � � oxidation ½50� Young’s modulus E E0 � � ¼ E E0 � � P � E E0 � � S � E E0 � � oxidation ½51� Thermal conductivity 1 0 0% weight loss D im en si on al c ha ng e (% ) Y ou ng ’s m od ul us s tr uc tu re t er m 5.8% weight loss 13.4% weight loss 23.1% weight loss 38.4% weight loss 0% weight loss 5.8% weight loss 13.4% weight loss 23.1% weight loss 38.4% weight loss Trend curve A Trend curve B 0 20 (a) (b) 2.0 1.8 1.6 1.4 1.2 1.0 0.8 0 20 40 60 80 100 120 140 160 40 60 80 Fluence (1020 n cm−2 EDND) Fluence (1020 n cm−2 EDND) 100 120 140 160 −1 −2 −3 −4 −5 −6 −7 Figure 51 Correlations between dimensional change and Young’s modulus structure term in Gilsocarbon. (a) Dimensional changes in pre-oxidized Gilsocarbon and (b) Young’s modulus structure terms in pre-oxidized Gilsocarbon. Modified from Schofield, P.; Brown, R. G.; Daniels, P. R. C.; Brocklehurst, J. E. Fast neutron damage in Heysham II/Torness moderator graphites (final report on 3-temperature zone rig); UKAEA, NRL-M-2176(S); 1991. 376 Graphite in Gas-Cooled Reactors 1 KirrðT Þ ¼ 1 K0ð30Þ K0ð30Þ K0ðTÞ þ dðT Þf � � Sk K0 K � � ox ½52� In recent years, it has been realized that the use of the product rule is simplistic, and most probably, only applicable for low irradiation data, up to a fluence not far beyond dimensional change turnaround and only for relatively low weight loss. Therefore, there has been a recent trend to use empirical fits to reactor or MTR data where available. 4.11.20 Irradiation Creep in Nuclear Graphite By the late 1940s, it was known that graphite compo- nents, when subjected to fast neutron irradiation, suffered significant dimensional change. It was thought that, because of the flux gradient across the brick section, these dimensional changes would gen- erate significant stresses in hollow graphite modera- tor blocks and that this would lead to significant 0 4.5 4.0 3.5 Gilsocarbon (300–400 �C) Petroleum coke (300–440 �C) Petroleum coke (560 �C) Pitch coke (300–440 �C) Unidentified coke (300–440 �C) Unidentified coke (560 �C) Gilsocarbon (560 �C) E/E0 Ö(E/E0) 3.0 2.5 2.0 1.5 Fr ac tio na l c ha ng e in s /s 0 1.0 0.5 0.0 50 100 150 Fluence (1020 n cm−2 EDND) 200 250 300 350 Figure 52 Change in strength of irradiated Gilsocarbon graphite compared to change in modulus. Modified from Brocklehurst, J. E. Chem. Phys. Carbon 1977, 13, 145–272. 1 0.5 R el at iv e ch an ge in p ro p er ty 0 0.0 0.1 0.2 Porosity 0.3 e−3.8x e−2.8x e-5.1x 0.4 Thermal conductivity (k/ko) Young’s modulus (E/Eo) Tensile strength (s/so) Compressive strength (s/so) Figure 53 Change in thermal conductivity, Young’s modulus, and strength due to radiolytic oxidation. Modified from Adam, R. W.; Brocklehurst, J. E. Mechanical tests on graphite with simulated radiolytic oxidation gradients; UKAEA, ND-R-853(S) (AB 7/26300); 1983. Graphite in Gas-Cooled Reactors 377 component failures within a few years of reactor operation. Therefore, the fuel channels in the early reactors, such as the Windscale Piles, were designed to avoid the buildup of stress. By the 1950s, it was realized that there was an irradiation-induced mechanism that was relieving stresses generated by dimensional change and the term ‘irradiation induced plasticity’85,86 was coined to describe this mechanism. Later, around 1960, the term ‘irradiation creep’87 started to be used for the difference between dimensional change in loaded and unloaded graphite specimens irradiated to the same dose. 378 Graphite in Gas-Cooled Reactors 4.11.20.1 Dimensional Change and Irradiation Creep Under Load Compressive stress increases, and tensile stress decreases, the irradiation-induced dimensional change of graphite as illustrated in Figure 54. In these experiments, two matching graphite samples, a loaded specimen and an unloaded ‘control’ specimen, were irradiated adjacent to each other in an MTR, and dimensional change in the direction of load was measured. However, as well as change in dimension in the load direction, there are also dimensional changes perpen- dicular to the load direction as shown in Figure 55. An irradiation creep curve can be simply obtained by subtraction of the unloaded dimensional change curve from the crept dimensional change curve, as illustrated in Figure 56. However, for practical use in assessments, this would require data for a range of temperatures and fast neutron fluences covering all of the expected operating conditions. Also, in the case of carbon dioxide-cooled reactors, the effect of 0 D im en si on al c ha ng e (% ) 0 50 (a) (b) 100 Fluence (1020 n cm−2 EDND) 150 200 Reference Loaded 250 −1 −2 −3 −4 0 1 D im en si on al c ha ng e (% ) 0 50 100 Fluence (1020 n cm−2 EDND) 150 200 Reference Loaded 250 −1 −2 −3 −4 Figure 54 Dimensional changes of loaded ATR-2E graphite. (a) compression and (b) tension. Modified from Haag, G. Properties of ATR-2E Graphite and Property Changes due to Fast Neutron Irradiation; FZJ, Jül-4183; 2005. the rate of radiolytic oxidation on creep rates would have to be quantified and understood. In addition, changes to the CTE and Young’s modulus with irra- diation creep have been observed, which further complicate assessment technology. 4.11.20.2 Types of Irradiation Creep Experiments There are three categories of graphite creep experi- ments. The first type was the restrained creep experiments. In these experiments a graphite specimen, usually dumbbell in shape, is restrained from shrinkage by a tube or split collar manufactured from graphite that shrinks less than the specimen of interest. In the case of anisotropic graphite, the tube or split collar could be manufactured with its longitudinal axis aligned with the more dimension- ally stable grain direction; the specimen would be manufactured with its axis perpendicular to this. These types of experiments are relatively easy to deploy but are difficult to assess as the load is not directly measured and a ‘creep law’ has to be assumed in the assessment of the results.88 Avariation on these experiments was the graphite spring tests used in Calder Hall to define primary creep.86 A second important type of experiment was the ‘out-of-pile measurements’ technique.89,90 These experiments, importantly, give ‘real-time results.’ However, this type of experiment is difficult to install in a reactor and results are obtained only for one specimen. The final type of experiment is the in-pile rig loading using a string of samples and usually taking advantage of the MTR flux gradient to obtain data on samples to various levels of fluence. There have been various designs of simple strings of specimens loaded either in tension or in compression. Tensile creep tests are vulnerable, that is, if one specimen fails, results for the whole string of specimens could be lost. However, there have been various rig designs aimed at overcoming this problem. 4.11.20.3 The UKAEA Creep Law Irradiated creep experimentswere carried out between 350 and 650 �C on both PGA and Gilsocarbon graph- ite.91 Some low fluence experiments were also carried out in Calder Hall86 which were used to define the so-called ‘primary creep.’ Creep strain data ecr was normalized to elastic strain units (esu) by dividing by the applied stress (s) and multiplying by the 2.0 1.5 1.0 0.5 C re ep s tr ai n (% ) 0.0 0 50 100 150 200 Compressive Tensile Fluence (1020 n cm−2 EDND) 250 Figure 56 Irradiation creep curves for ATR-2E at 500 �C. Modified from Haag, G. Properties of ATR-2E Graphite and Property Changes due to Fast Neutron Irradiation; FZJ, Jül-4183; 2005. 3 2 1 0 D im en si on al c ha ng e (% ) 0 50 100 Fluence (1020 n cm−2 EDND) 150 200 Reference Loaded 250 −1 −2 Figure 55 Lateral dimensional changes of loaded ATR-2E graphite. Modified from Haag, G. Properties of ATR-2E Graphite and Property Changes due to Fast Neutron Irradiation; FZJ, Jül-4183; 2005. Graphite in Gas-Cooled Reactors 379 unirradiated SYM (E0) as given below: esu ¼ E0ecr s ½53� Surprisingly, they found that by doing this, the creep data for these two types of graphite, with very differ- ent microstructures, could be fitted to a simple ‘creep equation’ of the form ecr ¼ s E0 ½expð�4gÞ� þ 0:23 s E0 g ½54� where g is the fast neutron dose. This is illustrated in Figure 57. The primary creep strain is assumed to be recov- erable on removal of the load while still under irradi- ation. Some evidence for this came from out-of-pile measurement experiments such as the FLACH experiments.90 However, if the specimens had been left unloaded for longer duration, more than 1 esu may have been recovered. In addition to this, an experiment carried out on precrept samples, that is, 1.2 1.0 0.8 0.6 C re ep s tr ai n (% ) E la st ic s tr ai n un its 0.4 0.2 0.0 0 16 14 12 10 8 6 4 2 0 0 10 20 30 40 50 60 70 10 (a) (b) 20 30 Fluence (1020 n cm−2) 40 50 BR–2 (300 – 650 �C, 900 psi) Isotropic graphites (compressive) Isotropic graphites (tensile) PGA (perpendicular) PGA (parallel) Pyrolytic graphites 60 70 Fluence (1020 n cm−2) PGA (BR-2 (300 – 650 �C, perpendicular) PGA (BR-2 (300 – 650 �C, parallel) Calder hall (140 – 350 �C) PGA graphites (PLUTO, 300 �C) Pyrolytic graphites (BR-2, 300 - 650 �C) Isotropic graphites (BR-2, 300 – 650 �C) Figure 57 UK creep data used to define the UK creep law. (a) creep strain and (b) elastic strain units. Modified from Brocklehurst, J. E. Irradiation Damage in CAGR Moderator Graphite; Northern Division, UKAEA, ND-R-1117(S); 1984. 380 Graphite in Gas-Cooled Reactors samples of PGA and Gilsocarbon irradiated in a creep experiment in the BR-2 reactor and then irra- diated with the load removed in DIDO and DFR respectively, exhibited more than 1 esu (a recovery in the region of 6 esu in the case of Gilsocarbon in DFR). 4.11.20.4 Observed Changes to Other Properties 4.11.20.4.1 Coefficient of thermal expansion Significant differences have been observed between the unstressed CTE and stressed CTE, as illustrated in Figure 58. Compressive creep strain was found to increase the CTE, and tensile creep strain to decrease the CTE. The changes in CTE caused by irradiation creep have similarities to those caused by the applica- tion of stress on unirradiated graphite. Figure 59(a) shows the changes in CTE in irradiated, crept specimens plotted as a function of creep strain92 and Figure 59(b) gives the changes in CTE in unirradi- ated graphite due to stress.93 At room temperature, the average CTE of an isotropic graphite with no porosity should be the average of the crystallite CTEs, that is, the crystallite CTEs are �27.0� 10�6 K�1 and �1.0� 10�6 K�1 in the ‘c’ and ‘a’ directions respectively, giving an average Unstressed 6 4 3 2 5 10 (a) Fluence (1020 n cm−2 EDN) C TE (1 0− 6 K −1 ) 0 20 30 40 50 60 Compressive Tensile IM1-24 (parallel) SM2-24 (parallel) VNMC (parallel) SM2-24 (perpendicular) VNMC (perpendicular) Creep strain(%) 6 1 2 2 100 20 30 40 50 60 (b) C ha ng e in C TE (1 0− 6 K −1 ) Figure 58 Additional changes to the coefficient of thermal expansion in loaded materials test reactor specimens. (a) Modified from Brocklehurst, J. E.; Brown, R. G. Carbon 1969, 7(4), 487–497 and (b) Modified from Brocklehurst, J. E.; Kelly, B. T. A review of irradiation induced creep in graphite under CAGR conditions; UKAEA, ND-R-1406(S); 1989. Graphite in Gas-Cooled Reactors 381 of 8.0� 10�6 K�1. From Figure 59(a), which is for Gilsocarbon with an unirradiated CTE of 4.0� 10�6 K�1, it is interesting to note that the increase in CTE in compression is approaching that value. 4.11.20.4.2 Young’s modulus The early BR-2 experiments showed little evidence of change in modulus with irradiation creep. How- ever, other authors94 did find evidence of a change (Figure 60). Later, higher temperature creep experi- ments carried out in the United Kingdom92 also showed a change in Young’s modulus. Figure 61(a) shows the fractional change in Young’s modulus E/E0, as a function of different compressive stresses. In Figure 61(b), the data has been normalized by dividing the crept Young’s modulus Ec by the irra- diated value of the unstressed specimen Ei. 4.11.20.5 Lateral Changes Creep in metals is purported to be related to plastic flow occurring at constant volume. However, artificial polycrystalline graphite is a porous material with a crystal structure considerably different from that of metals. It is therefore not surprising that in irradia- tion creep, graphite does not deform with a constant volume. As with much of the irradiation creep data, the quality of the transverse creep data is poor and inconsistent. However, from the data that does exist, the irradiation creep ratio does appear not to be constant with creep strain (Figure 62). 4.11.20.6 Creep Models and Theories It is unfortunate that a validated set of graphite irra- diation creep data covering the range of temperatures and fluences of interest for power producing reactors, as well as radiolytic oxidation in the case of carbon dioxide-cooled reactors, does not exist. In addition, there are no microstructural studies available to give an insight into the mechanism involved in irradiation creep in graphite. This has lead to much speculation and several model proposals. −5 −4 −3 −2 −1 −2 −1 0 0 C ha ng e in C TE (1 0− 6 K −1 ) 1 2 1 2 3 Compression (PLUTO) Compression (BR-2) Tension (BR-2) Creep strain (%) Stress (MN m−2) (a) Longitudinal stain gauge results Longitudinal stain gauge results Transverse strain gauge results −60 −50 −40 −30 −20 −10 C ha ng e in C TE (1 0− 6 K −1 ) −2 −1 0 0 10 20 1 2 3 (b) Figure 59 Synergy between changes in coefficient of thermal expansion in irradiation creep specimens and change in unirradiated, stressed graphite. (a) additional change in CTE as a function of creep strain and (b) change in CTE in unirradiated stressed samples. 2.0 1.6 1.2 0.8 0.4 0.0 0 2 4 SM1-24 (axial), irradiation temperature = 850- 920 �C Fluence (1024 n cm−2, E > 29 fJ) C ha ng e in Y ou ng ’s m od ul us (E /E 0) 6 8 10 12 14 16 0.0 MPa (Capsule 76M-18A) 3.3 MPa (Capsule 77M-10A) 4.5 MPa (Capsule 77M-10A) 6.5 MPa (Capsule 77M-10A) 0.0 MPa (Capsule 77M-10A) 6.5 MPa (Capsule 76M-18A) 4.5 MPa (Capsule 76M-18A) 3.3 MPa (Capsule 76M-18A) Figure 60 Changes in Young’s modulus in tensile crept and uncrept specimens. Reproduced from Oku, T.; Fujisaki, K.; Eto, M. J. Nucl. Mater. 1988, 152(2–3), 225–234. 382 Graphite in Gas-Cooled Reactors 0.5 0.4 0.3 0.2 0.1 0.0 0 5 (a) 10 C ha ng e in Y ou ng ’s m od ul us (E /E 0) 15 Fail 20 25 30 Compressive stress (MN m-2) 35 0.9 � 1021n cm−2 (1050 �C) 1 � 1021n cm−2 (850 �C) 2 � 1021n cm−2 (850 �C) 1.8 � 1021n cm−2 (1050 �C) 4540 50 0.8 0.9 1.0 1.1 (b) C re ep m od ul us /e la st ic m od ul us Compressive creep strain (%) 1+3(creep modulus) 1050 �C 850 �C -6 -5 -4 -3 -2 -1 0 1 2 Figure 61 Changes in Young’s modulus in irradiated-creep experiments. (a) changes to Young’s modulus as a function of stress and fast neutron fluence and (b) normalized Young’s modulus as a function of creep strain. Modified from Brocklehurst, J. E.; Kelly, B. T. A review of irradiation induced creep in graphite under CAGR conditions; UKAEA, ND-R-1406(S); 1989. Graphite in Gas-Cooled Reactors 383 4.11.20.6.1 UKAEA creep law In the United Kingdom, to extend the creep law to higher fluence and to account for radiolytic oxidation, the following approach was taken by the UKAEA. Creep strain, decr/dg, was assumed to be defined by decr dg ¼ aðTÞ s Ec ½expð�bgÞ� þ bðT Þ s Ec ½55� where a(T), b(T ), and b are temperature-dependent functions equal to 1.0, 0.23, and 4.0, respectively in the AGR and Magnox temperature ranges, and s and g are stress and irradiation fluence, respectively. The need for the temperature dependence outside this range was defined by data for HTRs obtained in the United States and Russia (Figure 63). The ‘creep modulus’ in the UKAEA model was defined as Ec ¼ E0SE½ox� ½56� where E0 is the unirradiated SYM and S is the irradi- ation temperature- and fluence-dependent struc- ture term derived from the irradiated modulus data (Section 4.11.17.4). To account for radiolytic weight loss, E[ox] is a modulus weight loss term defined as E/E0¼ exp(�lx) where l is an empirical constant equal to about 4.0 and x is the fractional weight loss. There are no rigorous observational data to underpin this model other than a few data points for preoxidized graphite given in Figure 64. It should be noted that in Figure 64, the only two Tr an sv er se c re ep s tr ai n (% ) (b) Longitudinal creep strain (%) 0.4 0.0 −0.4 −0.8 −1.2 ATR-2E (tension, 500 �C) ATR-2E (compression, 500 �C) H-337 (compression, 550 �C) H-451 (compression, 900 �C) AGOT (compression, 550 �C) AGOT (compression, 800 �C) H-337 (compression, 800 �C) −1.6 0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5 0.6 0.5 0.4 0.3 P oi ss on ’s r at io 0.2 0.1 0.0 (a) 0 1 2 3 Longitudinal creep strain (%) 4 5 SM2-24 VNMC IM1-24 6 Figure 62 Various transverse irradiation creep data. (a) UKAEA data and (b) US data. Modified from Brocklehurst, J. E.; Kelly, B. T. A review of irradiation induced creep in graphite under CAGR conditions; UKAEA, ND-R-1406(S); 1989. 1E−24 1E−25 1E−26 1E−27 C re ep c oe ffi ci en t ( k g cm −2 )− 1 (n eu tr on cm −2 )− 1 0 100 200 300 400 500 Temperature (�C) 600 700 Russian graphite American graphite (CGB) American graphite (EGCR) 800 900 1000 Figure 63 Temperature dependence of the secondary creep coefficient b(T) from US and Russian data. 384 Graphite in Gas-Cooled Reactors E la st ic s tr ai n un its (e su ) Fluence (1020 n cm-2 EDND) 12 10 8 6 4 2 0 0 5 10 25.5% 27.8% 9.8% 6% 9.5% 15 2520 30 35 40 45 50 Low density graphite (equivalent to 25% weight loss) Tensile Compressive UK creep law Figure 64 Preoxidized irradiation creep data. Modified from Brocklehurst, J. E.; Kelly, B. T. A review of irradiation induced creep in graphite under CAGR conditions; UKAEA, ND-R-1406(S); 1989. Graphite in Gas-Cooled Reactors 385 samples with a significant amount of weight loss are irradiated to a relatively low fluence. Similarly, there are some, even less convincing, data on creep sam- ples initially irradiated to high fast neuron fluence before loading.92 The main criticisms of the UKAEA creep model, for inert conditions, is that it gives a very poor fit to the high fluence creep data obtained in Germany and the United States, as discussed in the next section. 4.11.20.6.2 German and US creep model This model was devised for helium-cooled HTR applications where radiolytic oxidation was of no concern. The form of some of the US and German data is given below (Figure 65). There appears to be a difference between tension and compression at high fluence. However, this is the only data that shows this and it is not clear if it is a real effect. It was assumed that microstructural changes at medium to high fluence would modify the creep rate and account for the shape of these curves. It was assumed that this could be accounted for by modifying the secondary creep coefficient in the UK creep law by the following expression: ecrðsecondaryÞ ¼ K s E0 � � K ¼ K 0 1� m DV=V0ðDV=V0Þm � � ½57� where K0 is the UK creep coefficient, DV/V0 is the change in volume (which is a function of fluence and irradiation temperature), (DV/V0)m is the volume change at volumetric ‘turnaround,’ and m is a graphite grade and temperature-dependent variable. 4.11.20.6.3 Further modifications to the UKAEA creep law: interaction strain The theory originally developed by Simmons in the 1960s reported in detail by Hall et al.61 relat- ing the polycrystalline dimensional change rate and CTE with crystallite dimensional change rate, and CTE has been further developed95 in an attempt to explain the shape of the graphite irradiation creep behavior at high dose. The proposed theory argues that if the dimensional change rate in polycrystalline graphite can be related to the CTE, and because irradiation creep has been observed to modify CTE of the loaded specimen differently to that seen in unloaded specimens, changing the CTE by creep would be expected to change the dimensional change rate and hence, the dimensional change in the loaded specimen. This leads to the introduction of the so-called ‘interaction strain.’ The theory behind this methodology is described below. Considering two specimens (a crept specimen and an unloaded control) being irradiated under identical conditions; in the unloaded control specimen, by applying the Simmons equations, the bulk dimensional change rate gx and bulk CTE ax can be defined by gx ¼ ð1� AxÞga þ Axgc ax ¼ ð1� AxÞaa þ Axac ½58� 3.0 2.5 2.0 1.5 1.0 0.5 0.0 0 50 C re ep s tr ai n (% ) ATR-2E (stress = 5 Mpa) Tension (900 �C) Compression (500 �C) Tension (500 �C) Tension (300 �C) 100 Fluence (1020 n cm–2 EDND) 150 200 Figure 65 High-fluence German and US data. 386 Graphite in Gas-Cooled Reactors where ga and gc are crystal dimensional change rate in the a and c directions, respectively, and aa and ac are the crystal CTE in the a and c directions, respec- tively. Ax is referred to as the structure factor and by rearrangement Ax ¼ ax � aaac � aa ½59� Thus, gx ¼ ga þ gT ax � aaac � aa � � ½60� where gT is the crystal shape rate factor and is equal to gc� ga. Similarly for the loaded specimen, g 0x ¼ ga þ gT a0x � aa ac � aa � � ½61� Therefore, the difference between the dimensional change rates of the unloaded and loaded specimen is g 0x � gx ¼ gT Da ac � aa � � ½62� or g 0x ¼ gx þ gT Da ac � aa � � ½63� where Da is the change in CTE under load (a0x � ax). This leads to the following definitions: � The true dimensional change in the loaded specimen¼ the dimensional change in the controlþ the interaction term � True creep¼ dimensional change in loaded specimen � true dimensional change in loaded specimen � Apparent creep¼ dimensional change in loaded specimen� dimensional change in control Thus, the interaction term gT Da ac�aa � � is included in the finite element analysis of graphite components. The limited data that exists on irradiated HOPG indicates that the dimensional change rate of graphite crystallites increases with increasing fluence in the ‘c’ direction and decreases in the ‘a’ direction for all measured irradiation temperatures and dose range. For irradiation temperatures of 450 and 600 �C, the data indicates that ac and aa remain invariant to fluence. However, below 300 �C the crystal CTE appears to change. There are no crystal CTE data for higher temperatures. It should also be noted that Simmons equations imply that Ax ¼ ax � aaac � aa ¼ gx � ga gc � ga ½64� Close examination of typical graphite irradiation data, say for Gilsocarbon irradiated in the tempera- ture range where crystal data are available (450 and 600 �C), shows that the relationship given above does not hold. In fact, the Simmons relationship and measured data diverge at low dose. This is attributed to Simmons assuming that polycrystalline graphite can be considered as a loose collection of crystallites with no mechanical interactions. Others95 have added an extra ‘pore generation’ term to the Simmons dimensional change relationship to try and reconcile these issues, but again there is no real validation of these models. Graphite in Gas-Cooled Reactors 387 The use of this interaction term did not gain wide (international) acceptance as it appeared to be using the Simmons relationship beyond its applicability and did not explain the difference between compres- sive and tensile loading at high fluence. 4.11.20.6.4 Recent nuclear industry model Recently Davies and Bradford96 have developed a far more complex creep model as given below: ec ¼ ak1 E0 exp�k1g ðg g0¼0 s Sðg;T ÞW ðxÞexp k1g0dg0 þ b E0 ðg g0¼0 1� exp�k2g0 � � s Sðg;T ÞW ðxÞdg 0 þok3 E0 exp�k3g ðg g0¼0 s Sðg;TÞW ðxÞexp k3g0dg0 ½65� where a, 1 esu (where this is defined by s/E0); k1, 0.0857e (1630.4/T); b, 0.15 esu per 1020 n cm�2 EDN; k2, 0.0128e (1270.8/T); o, 5 esu; k3, 0.4066e (�1335.9/T); s, stress (P); E0, unirradiated SYM (Pa) appropriate to the stress applied (0.84�DYM); S(g, T), struc- ture term representing structural induced changes to creepmodulus (a function of fluence and temperature); W(x), oxidation term representing oxidation-induced changes to creep modulus (a function of weight loss, x, which is a function of fluence). The lateral strain ratio for the primary and recov- erable terms is assumed to be equal to the elastic Poisson’s ratio. The lateral strain ratio for secondary creep, nsc, is assumed to follow the relationship nsc ¼ 0:5½1� 3ScðgÞ� Sc is a structural connectivity term that the authors have used in model fits for other graphite property changes.57 This model certainly fits the available inert data better than the previous models, although it cannot be tested against radiolytically oxidized- graphite data as there is none. 4.11.20.7 Final Thoughts on Irradiation Creep Mechanisms Two main models for the mechanism of irradiation creep have been put forward but neither has any microstructural observations to support them. The first suggestion is that a model by Roberts and Cottrell97 for a-uranium may be appropriate. This model proposes that the graphite crystallite struc- tures will yield and shear because of the generation of stresses caused by dimensional change. However, it is difficult to envisage such a yield and shear mechanism in crystalline graphite. The second model98 suggests that under load, the crystallite basal planes will slide because of a pinning and unpinning mechanism during irradiation. Such a mechanism is described in detail by Was99 with rela- tion to metals and could explain primary creep and secondary linear creep. However, if irradiation creep in graphite is associated with basal plane slip due to pinning–unpinning, it is surprising that in PGA, irradiation creep is less in the WG or parallel to the basal plane direction than it is in the AG or perpen- dicular to the basal plane direction (Figure 57). Another possibility is that stress modifies the crys- tal dimensional change rate itself. In support of this are X-ray diffraction measurements100 that showed that the lattice spacing in compressive crept speci- mens is less than that in the unstressed control specimens (Figure 66). Such a mechanism would explain the PGA data and could be related to the change in CTE and the observed annealing behavior. However, the data and experimental fluence and creep range given are very limited. It is clear that changes to the lattice spacing in crept graphite would be an area worthy of further investigation in future irradiation creep programs. Irradiation creep in the graphite crystallite will be reflected in the bulk deformations observed in creep specimens and in reactor components. Changes to the bulk microstructure due to radiolytic oxidation would be expected to influence this bulk behavior, as would large crystal dimensional changes at very high fluence (past dimensional change turnaround). It would be expected that at very high fluence the behavior of graphites with differing microstructures would diverge; this appears to be the case from the limited high fluence data available. 4.11.21 Concluding Remark Nuclear grade graphite has been used, and is still used, in many reactor systems. Furthermore, it pro- vides an essential moderator and reflector material for the next-generation high-temperature gas-cooled nuclear reactors that will be capable of supplying high-temperature process heat for the hydrogen economy. Hence, nuclear graphite technology remains an important topic. Although there is a wealth of data, knowledge, and experience on the design and operation of graphite-moderated reactors, 3.56 3.54 3.52 3.50 3.48 3.46 In te rla ye r sp ac in g (Å ) 3.44 3.42 3.40 3.38 3.36 0 2 4 6 8 10 Fluence (GWd te−1) 12 14 16 18 20 Unrestrained Restrained Figure 66 The effect of stress and irradiation on the interlayer spacing of graphite. Modified from Francis, E. L. Progress Report for the JNPC-Materials Working Party: Graphite Physics Study Group; UKAEA, TRG-M-2854 (AB 7/17604); 1965. 388 Graphite in Gas-Cooled Reactors the need for present plants to extrapolate beyond current data and to predict the behavior of new graphite grades operating for longer lifetimes at higher temperatures than before means there is still a substantial amount of work for the graphite special- ist. Future understanding and validation of property/ microstructural change relationships that enable the prediction and interpolation of existing databases and the development of new graphite grades is now pos- sible using new characterization, modeling, and com- putation techniques. These allow the investigation of mechanisms and graphite behavior that were previ- ously impossible or impractical to conduct. 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All rights reserved. 4.12.1 Introduction 391 4.12.2 Vanadium Alloys for Fusion Reactors 391 4.12.3 Compositional Optimization 392 4.12.4 Fabrication Technology 393 4.12.5 Fundamental Study on Impurity Effects 396 4.12.6 Thermal Creep 396 4.12.7 Corrosion, Compatibility, and Hydrogen Effects 398 4.12.8 Radiation Effects 400 4.12.9 Tritium-Related Issues 401 4.12.10 Development of Advanced Alloys 402 4.12.11 Critical Issues 403 4.12.12 Vanadium Alloy Development for Fusion Blankets 403 4.12.13 Summary 404 References 405 Abbreviations DBTT Ductile–brittle transition temperature dpa Displacement per atom flibe Molten LiF-BeF2 salt mixture GTA Gas tungsten arc HFIR High Flux Isotope Reactor HIP Hot isostatic pressing IFMIF International Fusion Materials Irradiation Facility IP Imaging plate ITER International Thermonuclear Experimental Reactor LMFBR Liquid Metal Fast Breeder Reactor MA Mechanical alloying PWHT Postweld heat treatment RAFM Reduced activation ferritic/martensitic REDOX Reduction–oxidation reaction TBM Test Blanket Module TBR Tritium breeding ratio TEM Transmission electron microscope 4.12.1 Introduction Vanadium alloys were candidates for cladding materials of Liquid Metal Fast Breeder Reactors (LMFBR) in the 1970s.1 However, the development was suspended mainly because of an unresolved issue of corrosion with liquid sodium. Vanadium alloys attracted attention in the 1980s again for use in fusion reactors because of their ‘low activation’ properties. At present, vanadium alloys are considered as one of the three promising candidate low activation structural materials for fusion reactors with reduced activation ferritic/martensitic (RAFM) steels and SiC/SiC composites. Overviews of vanadium alloys for fusion reactor applications are available in the recent proceedings papers of ICFRM (International Conference on Fusion Reactor Materials).2–6 This chapter highlights the recent progress in the devel- opment of vanadium alloys mainly for application in fusion nuclear systems. 4.12.2 Vanadium Alloys for Fusion Reactors Various tritium breeding fusion blanket concepts have been studied with different combinations of structural materials, tritium breeding materials, and cooling materials. Vanadium alloys have been used in most cases with liquid lithium as the breeding and cooling materials (self-cooled V/Li blankets) for advanced concepts of DEMO (fusion demonstra- tion power plant) and commercial fusion reactors.7,8 Because of high atomic density of Li atoms in liquid Li relative to Li-ceramics, Li–Pb, and molten-salt 391 392 Vanadium for Nuclear Systems Flibe, V/Li systems can obtain high tritium breeding ratio (TBR) without using the neutron multiplier Be. A neutronics calculation showed that ‘tritium self sufficiency’ can be satisfied without Be both in Tokamak and Helical reactor systems.9 Without the necessity of using beryllium as a neutron multiplier, the replacement frequency of the blanket will be reduced because the blanket system is free from the periodic replacement due to the lifetime of Be, which can lead to enhanced plant efficiency. V/Li blankets can be designed with a simple structure as schematically shown in Figure 1. The blanket is composed of Li cooling channels made of vanadium alloys, reflectors, and a shielding area, which is in contrast to more complex solid breeder blankets that need a solid breeder zone, a neutron multiplier beryllium zone, cooling channels using gas or water, and tritium recovery gas flow channels in addition to reflectors and shielding. A self-cooled Li blanket using neutron multiplier beryllium was also designed in the Russian pro- gram.10 This concept can downsize the blanket area because of efficient tritium generation per zone. However, the blanket structure must be more Table 1 Breeding blanket concepts using vanadium alloys Concept V/Li V/Be/Li Breeder and coolant materials Liquid Li Liquid Li Use of neutron multiplier Be No Yes Advantages Simple structure High TBR Critical issues MHD coating, T recovery from Li MHD coati T recover D-T plasma Reflector Neutron Superconducting magnet Vanadium alloy structures Blanket Coating with W, Be, or C Flowing liquid lithium Shield Figure 1 Illustration of self-cooled Li blanket with V–4Cr–4Ti structural material. complex than V/Li and new issues need to be solved such as Li/Be compatibility. General requirements for structural materials of fusion blankets include dimensional stability, compat- ibility with breeder and coolants, high-temperature strength and low-temperature ductility during irradia- tion. For vanadium alloys, issues concerning industrial maturity such as developing large-scale manufacturing technology need to be resolved. Vanadium alloys could be a candidate structural material for molten-salt Flibe (LiF–BeF2) blankets. For this application, a concept was proposed to dis- solve WF6 or MoF6 into Flibe for corrosion protec- tion of the wall surfaces by precipitation of Wor Mo and reduction of the tritium inventory in vanadium alloys by enhancing reaction from T2 to TF, which is more highly soluble in Flibe than T2. 11 The TBR of Flibe/V blankets may be marginal, but the neutron shielding capability for the superconductor magnet systems may be superior relative to V/Li according to neutronics investigation.12 In this system, precipitates ofWorMo formed as a result of reaction fromT2 toTF needs to be recovered from the flowing Flibe. Table 1 summarizes the blanket concepts using vanadium alloys with the advantages and critical issues. 4.12.3 Compositional Optimization Vanadium alloys potentially have low-induced acti- vation characteristics, high-temperature strength, and high thermal stress factors. For the optimization of the composition, both major alloying elements and minor impurities need to be controlled. For main- taining the low activation properties, use of Nb and Mo, which used to be the candidate alloying elements for application to LMFBR, need to be avoided. Cr was known to increase the strength of vanadium at high temperature and Ti was known to enhance ductility of vanadium by absorbing interstitial impu- rities, mostly oxygen. However, excess Cr or Ti can V/Flibe Molten-salt Flibe No Small MHD pressure drop ng, Li/Be compatibility, y from Li REDOX control, recovery of W or Mo, increase in TBR Vanadium for Nuclear Systems 393 lead to loss of ductility. Hence, optimization of Cr and Ti levels for V–xCr–yTi has been investigated. It was known that with x þ y > 10%, the alloys became brittle6 as shown in Figure 2. With systematic efforts, V–4Cr–4Ti has been regarded as the leading candi- date. For low activation purposes, the level of Nb, Mo, Ag, and Al needs to be strictly controlled. Large and medium heats of V–4Cr–4Ti have been made in the United States, Japan, and Russia. 950 �C 950 �C 1000 �C 1000 �C 1050 �C 1100 �C 1100 �C 1150 �C 1150 �C Annealing temperature 5 10 15 20 25 Cr + Ti (Wt %) ±20 �C –250 –200 –150 –100 –50 0 50 100 D B TT (� C ) Figure 2 DBTT as a function of Cr þ Ti (wt%) of V–Cr–Ti alloy for various annealing temperatures. Reproduced from Zinkle, S. J.; Matsui, H.; Smith, D. L.; Rowcliffe, A. L.; van Osch, E.; Abe, K.; Kazakov, V. A. J. Nucl. Mater. 1998, 258–263, 205–214, with permission from Elsevier. 10–2 10–5 10–4 10–3 10–2 10–1 100 101 102 103 104 105 10–1 100 101 Cooling time after shut Reduced activation ferritics (F82H) Pure SiC/SiC C on ta ct d os e ra te (S v h– 1 ) V–4Cr–4Ti (NIFS-HEAT) Figure 3 Contact dose after use in first wall of a fusion comm reference ITER structural material F82H: reference reduced acti V–4Cr–4Ti alloy SiC/SiC: assumed to be impurity-free. An especially high-purity V–4Cr–4Ti ingot pro- duced by the National Institute for Fusion Science (NIFS) in collaboration with Japanese Universities (NIFS-HEAT-1 and 2) showed superior properties in manufacturing due to their reduced level of oxygen impurities.4 Figure 3 compares the contact dose rate after use in the first wall of a fusion commercial reactor for four reference alloys. The full-remote and full-hands-on recycle limits are shown to indicate the guideline for recycling and reuse.13 SS316LN-IG (the reference ITER structural material) will not reach the remote- recycling limit after cooling and hence the recycling is not feasible. F82H (reference RAFM steel) and NIFS-HEAT-2 behave similarly, but NIFS-HEAT-2 shows significantly lower dose rate before the 100- year cooling. The dose rate of F82H and NIFS- HEAT-2 reached a level almost two orders lower than the remote-recycle limit by cooling for 100 and 50 years, respectively. The dose rate of SiC/SiC com- posites (assumed to be free from impurities because of lack of reference composition) is much lower at 100 year cooling relative to F82H and NIFS-HEAT-2. 4.12.4 Fabrication Technology Figure 4 summarizes the microstructural evolution during the breakdown process of NIFS-HEAT-2 102 103 104 down (years) FFHR Li blanket first wall neutron 1.5 MW m–2 operation Full-hands-on recycling Full-remote recycling SS316 for ITER (SS316LN-IG) ercial reactor for four reference alloys. SS316LN-IG: the vation ferritic/martensitic steel NIFS-HEAT-2: reference Ingot Hot forging 1423 K Hot/cold roll 1373 K/RT Heat treatment 973 K 1273 K 1373 K 1573 K Ti-rich blocky precipitates (with N, O, C) Formation Elongation, band structure Dissolution Ti–O–C thin precipitates Formation Coarsening Dissolution V–C on GB 50 mm 50 mm 25 mm 1 mm 1 mm 1 mm 50 mm Figure 4 Microstructural evolution during the breakdown process of V–4Cr–4Ti ingots. Reproduced from Muroga, T.; Nagasaka, T.; Abe, K.; Chernov, V. M.; Matsui, H.; Smith, D. L.; Xu, Z. Y.; Zinkle, S. J. J. Nucl. Mater. 2002, 307–311, 547–554. 120 140 160 180 200 220 240 260 200 400 600 800 1000 1200 1400 1600 NIFS-HEAT-1 NIFS-HEAT-2 US-DOE 832665 V–4Cr–4Ti V ic ke rs h ar d ne ss (H v) Annealing temperature (K) Figure 5 Vickers hardness as a function of annealing temperature for NIFS-HEAT-1, NIFS-HEAT-2, and US-DOE 832665. Reproduced from Heo, N. J.; Nagasaka, T.; Muroga, T. J. Nucl. Mater. 2004, 325, 53–60. 394 Vanadium for Nuclear Systems ingots.4 Bands of small grains aligned along the rolling direction were observed at the annealing temperature below 1223K. The grains became homogeneous at �1223K. The examination showed that optimization of size and distribution of Ti-CON precipitates are crucial for good mechanical properties of the V–4Cr– 4Ti products. Two types of precipitates were observed, that is, the blocky and the thin precipitates. The blocky precipitates formed during the initial fabrication pro- cess. The precipitates aligned along the working direc- tion during the forging and the rolling processes forming band structures, and were stable to 1373K. Since clustered structures of the precipitates result in low impact properties, rolling to high reduction ratio is necessary for making a thin band structure or homo- genized distribution of the precipitates. The thin pre- cipitates were formed at �973K and disappeared at 1273–1373K. At 1373K, new precipitates, which were composed of V and C, were observed at grain bound- aries. They seem to be formed as a result of redistri- bution of C induced by the dissolution of the thin precipitates. The impact of the inhomogeneous micro- structure can influence the fracture properties.14 Figure 5 shows the hardness as a function of final heat treatment temperature for three V–4Cr–4Ti materials: NIFS-HEAT-1, NIFS-HEAT-2, and US- DOE-832665 (US reference alloy).15 The hardness has a minimum at 1073–1273K, which corresponds to the temperature range where formation of the thin precipitates is maximized. With the heat treatment higher than this temperature range, the hardness increases and the ductility decreases because the precipitates dissolve enhancing the level of C, N, and O in the matrix. Based on the evaluation of various properties in addition to the hardness as a function of heat treatment conditions, the optimum heat treat- ment temperature of 1173–1273Kwas suggested. Plates, sheets, rods, and wires were fabricated mini- mizing the impurity pickup and maintaining grain and precipitate sizes in Japanese, US, and Russian programs. Thin pipes, including those of pressurized creep tube specimens, were also successfully fabricated Vanadium for Nuclear Systems 395 in Japan maintaining the impurity level, fine grain size, and straight band precipitate distribution by maintain- ing a constant reduction ratio between the intermedi- ate heat treatments.16 The fine-scale electron beam welding technology was enhanced as a result of the efforts for fabricating the creep tubes, including plug- ging of end caps.17 In the United States, optimum vacuum level was found for eliminating the oxygen pick-up during intermediate annealing to fabricate thin-walled tubing of V–4Cr–4Ti.18 In Russia, fabrica- tion technology is in progress for construction of a Test BlanketModule (TBM) for ITER (InternationalTher- monuclear Experimental Reactor).19 Joining of V–4Cr–4Ti by gas tungsten arc (GTA) and laser welding methods was demonstrated. GTA 0 5 10 15 50 100 150 200 Test tem EU= 13 J NH 188 US/HP 183 K128 K NH1/HP 0 50 100 150 2 0 100 200 300 Oxygen in we NH1/HP US/HP DBTT = +60 K/100 wppm Plate/filler A b so rb ed e ne rg y (J ) D B TT (K ) Figure 6 Upper: Absorbed energy of Charpy impact tests of V various combinations of plates and fillers. Lower: DBTT of V–4Cr– NH1, NIFS-HEAT-2 (O: 181wppm); US, US-DOE 832665 (O: 310 (O: 36wppm). Reproduced from Nagasaka, T.; Grossbeck, M. L. is a suitable technique for joining large structural components. GTAwelding technology for vanadium alloys provided a significant progress by improving the atmospheric control. The results are summarized in Figure 6. Oxygen level in the weld metal was controlled by combined use of plates of NIFS- HEAT-1 (181 wppm O) or US-8332665 (310 wppm O) and filler wire ofNIFS-HEAT-1, US-8332665, or a high-puritymodel alloy (36wppmO).Asdemonstrated in Figure 6, ductile–brittle transition temperature (DBTT) of the joint and the oxygen level in the weldmetal had a clear positive relation. Thismotivated further purification of the alloys for improvement of the weld properties.20 Only limited data on irradia- tion effects on the weld joint are available at present. 250 300 350 400 perature (K) US/US 320 K 1/NH1 K 00 250 300 350 400 ld metal (wppm) NH1/NH1 US/US O –4Cr–4Ti weld joints as a function of test temperature for 4Ti weld joints as a function of oxygen level in the weld metal. wppm); HP, high-purity model V–4Cr–4Ti alloy ; Muroga T.; King, J. F. Fusion Technol. 2001, 39, 664–668. 396 Vanadium for Nuclear Systems The welding results in complete dissolution of Ti- CON precipitates and thus results in significant increase in the level of C, O, and N in the matrix. In such conditions, radiation could cause embrittlement. SomeTEMobservations showedenhanceddefect clus- ter density at the weld metals. However, the overall evaluation of the radiation effects remains to be per- formed. Especially, elimination of radiation-induced degradationbyapplyingappropriate conditionsofpost- weld heat treatment (PWHT) is the key issue. For the use of vanadium alloys as the blanket of fusion reactors, the plasma-facing surfaces need to be protected by armor materials such asW layers. Limited efforts are, however, available for developing the coating technology. A low pressure plasma-spraying method was used for coating Won V–4Cr–4Ti for use at the plasma-facing surfaces. The major issue for the fabrication is the degradation of the vanadium alloy substrates by oxidation during the coating processes. Figure 7 shows the result of bending tests of the coated samples. The crack was initiated within the W layer propagating parallel to the interface and followed by cracking across the interface. Thus, in this case, the quality of W coating layer is the issue rather than the property of the V–4Cr–4Ti substrate or the interface. Hardening of substrate V–4Cr–4Ti by the coating occurred but was shown to be in acceptable range.21 Figure 8 is a collection of the products from NIFS-HEAT-2. 4.12.5 Fundamental Study on Impurity Effects Effects of C, O, and N on the property of vanadium are a long-standing research subject. However, research into the effects of C,O, andNonV–4Cr–4Ti is limited. 500 µm Intergranular fracture W V–4Cr–4Ti Figure 7 Cross-section of W coating on V–4Cr–4Ti after bend Research with model V–4Cr–4Ti alloys doped with O and N provided information on the partition- ing of O and N into the precipitates and matrix. The density of the blocky precipitates and thin pre- cipitates increased with N and O levels, respectively. Figure 9 shows hardness as a function of N and O levels in V–4Cr–4Ti after melting and annealing at 1373 K for 1 h.22 Hardness after annealing at 1373K, where only the blocky precipitates were observed in the matrix, increased to a certain extent with O level (�4.5Hv/100 wppm O), but only very weakly with N level (�0.9Hv/100 wppm N). These data suggest that, after the annealing, most of the N is included in the blocky precipitates and stable to �1373K. On the other hand, O exists in the matrix, the blocky and the thin precipitates, and the partition- ing changes with the heat treatment. Thus, for the purpose of the property control of V–4Cr–4Ti, the level of N before the heat treatment is not so impor- tant but that of O is crucial. It is to be noted, however, that N contamination during the operation can influ- ence the properties of vanadium alloys seriously. Fundamental information on the impurity dis- tribution and interaction with solutes and dislocations is obtained by serrated flow in tensile deformation as shown in Figure 10. Temperature and stain rate depen- dence of the flow showed that the serrated flow above 673K is related to C andO and above 773K toN. Small serration height at 673K for NIFS-HEAT-1 (1–3MPa) relative to that of US-832665 (�9MPa) was observed and attributed to the difference in O level.23 4.12.6 Thermal Creep Thermal creep is a potential factor which can deter- mine the upper operation limit of vanadium alloys. 50 µm Crack 10 µm ing tests. Fracture started in the W coating layer. Nitrogen level (wppm) 50 100 150 200 250 300 0 200 400 600 800 1000 1200 0 200 400 600 800 1000 1200 V ic ke rs h ar d ne ss (H v) Oxygen level (wppm) V–4Cr–4Ti, as-melted V–4Cr–4Ti, 1373 K Pure V, as-melted Pure V, 1373 K Figure 9 Vicker’s hardness as a function of O and N levels for V–4Cr–4Ti after melting and annealing at 1373K for 1 h. Reproduced from Heo, N. J.; Nagasaka, T.; Muroga, T.; Matsui, H. J. Nucl. Mater. 2002, 307–311, 620–624. 26 t 1.9 t 2 d 8 d (mm) Plates, sheets, wires, and rods Laser weld joint Thin pipes W coating by plasma spraying Creep tubes 6.6 t 4.0 t 1.0 t 0.5 t 0.25 t f 4.57 � 0. 25 t � 400 mm f 10 � 0. 5 t � 100 m m 20 mm W coating NIFS-HEAT-2 (V–4Cr–4Ti) 0.5 mm 5 mm Figure 8 Collection of the V–4Cr–4Ti products manufactured by the Japanese program. Vanadium for Nuclear Systems 397 Previously, uniaxial tensile creep tests and biaxial pressurized creep tube tests were carried out in vac- uum for evaluation of the creep deformation charac- teristics. Figure 11 shows summary of the creep deformation rate as a function of applied stress.3 In this plot, the creep data were described by a power- law equation24: de=dt ¼ AðDGb=kTÞðs=GÞn where de/dt is the creep rate, s is the applied stress, D is the lattice diffusion coefficient, G is the shear modulus, b is the Burgers vector, k is the Boltzmann constant, T is the absolute temperature, and A is a constant. The exponent of the function (n) changed from 10 with the increase in the stress. A new apparatus for biaxial creep testing in Li provided opportunities for examining creep 10-3 10-5 10-6 10-7 10-8 10-9 10-10 10-11 10-12 10-2 (d e/ d t)k T/ D G b s/G Uniaxial tests 310 wppm O Biaxial tests 699 wppm O n = 4.3 n = 3.7 n = 0.84 n = 13 Figure 11 Thermal creep deformation rate of V–4Cr–4Ti as a function of applied stress for uniaxial and biaxial tests. The definition of the terms and the function from which n is extracted are indicated in the text. Reproduced from Kurtz, R. J.; Abe, K.; Chernov, V. M.; Hoelzer, D. T.; Matsui, H.; Muroga, T.; Odette, G. R. J. Nucl. Mater. 2004, 329–333, 47–55. C re ep s tr ai n ra te (1 s– 1 ) 0 10-6 10-7 10-8 10-9 10-10 2 4 6 8 10 12 Creep strain (%) In vacuum In lithium 50 MPa 70 MPa 90 Mpa Figure 12 Creep strain rate as a function of creep strain for the same batch of NIFS-HEAT-2 creep tubes in vacuum and Li environments. Modified from Li, M.; Nagasaka, T.; Hoelzer, D. T.; et al. J. Nucl. Mater. 2007, 367–370, 788–793; Fukumoto, K.; Nagasaka, T.; Muroga, T.; Nita, N.; Matsui, H. J. Nucl. Mater. 2007, 367–370, 834–838. 0 10 20 30 1073 K 973 K 873 K 773 K 673 K RT S tr es s (M P a) Strain (%) 200 Figure 10 Tensile deformation curves of V–4Cr–4Ti at various temperatures. 398 Vanadium for Nuclear Systems deformation in Li with that in vacuum.25 However, the correlation of creep data is subject to the alloy heat and manufacturing processes as well as test methods and environments. Figure 12 shows the comparison of the NIFS-HEAT-2 creep strain rate versus creep strain data for tests in vacuum and Li environments at 1073K, for the same batch of NIFS- HEAT-2 creep tubes.25,26 The figure clearly shows reduced strain rate in Li environments. A possible factor could be N pick-up from Li and the resulting surface hardening during exposure to Li. Further investigation is necessary for understanding the envi- ronmental effects on impurity redistribution and creep performance. Microstructural observations of the creep tube specimens tested at 1123K showed free dislocations and dislocation cell at 100 and 150MPa, respectively. This change of dislocation structure can cause the change in power-law creep behavior.27 4.12.7 Corrosion, Compatibility, and Hydrogen Effects In a Li/V blanket, it is believed that the interior of the wall needs to be coated with insulator ceramics for mitigating the pressure drop caused by magnetohydro- dynamic effects (see also Chapter 4.21, Ceramic Coatings as Electrical Insulators in Fusion Blan- kets). Corrosion of vanadium alloys in liquid Li might not be a concern if the entire inner wall is covered with the insulating ceramic coating. However, since the idea to cover the insulator ceramic coating againwith a thin vanadium or vanadium alloy layer was presented for the purpose of preventing liquid lithium from intruding into the cracks in the ceramics coating, the corrosion of vanadium alloys in liquid lithium again attracted attention. It is known that the corrosion of vanadium alloys in liquid lithium is highly dependent on the alloy composition and lithium chemistry. Espe- cially, the N level influences the corrosion in complex manners.28,29Figure 13 shows a summaryof theweight 10 20 30 Ti Cr Ti:Cr = 2:1 50 40 30 20 10 V 40 50 I II –47.4–8.2–2.1+7.5 +11 +0.4 +11.9 +6.7 +2.5 +1.2 –22.0–19.0 –19.7–2.4 –21.0 –26.4 –52.5 +5.8 Figure 13 The compositions of V–Ti–Cr alloys (wt%) with increase (area I) and decrease (area II) of mass (g cm�2) after holding of samples in Li at 973K, 500 h. Reproduced from Eliseeva, O. I.; Fedirko, V. N.; Chernov, V. M.; Zavialsky, L. P. J. Nucl. Mater. 2000, 283–287, 1282–1286. 0 0.02 0.04 0.06 0.08 0.1 773 V–4Cr–4Ti V–4Cr–4Ti–0.5Si V–4Cr–4Ti–0.5Al V–4Cr–4Ti–0.5Y Oxidation temperature (K) W ei gh t g ai n (m g m m –2 ) M el te d 873 973 1023 Figure 14 Weight gain of V–4Cr–4Ti with Si, Al, and Y exposed to air for 1 h. At 1023K, the weight gain was not measured for V–4Cr–4Ti because the surface oxidized layer melted. Reproduced from Fujiwara, M.; Natesan, K.; Satou, M.; Hasegawa, A.; Abe, K. J. Nucl. Mater. 2002, 307–311, 601–604. 0 10 20 To ta l e lo ng at io n (% ) 30 40 50 0 100 200 400300 600500 700 Natesan (BL-71 O:670 wppm) DiStefano (US-832665 O:310 wppm) DiStefano (Preoxidized US-832665 O:800 wppm) Chen (SWIP-Heat O:900 wppm) Chen (NIFS-HEAT-2 O:158 wppm) Hydrogen concentration (wppm) Figure 15 Total elongation as a function of hydrogen concentration for V–4Cr–4Ti alloys with different O levels. Modified from DiStefano, J. R.; Pint, B. A.; DeVan, J. H. J. Nucl. Mater. 2000, 283–287, 841–846; Chen, J. M.; Muroga, T.; Qiu, S.; Xu, Y.; Den, Y.; Xu, Z. Y. J. Nucl. Mater. 2004, 325, 79–86. Vanadium for Nuclear Systems 399 gain and loss in V–xCr–yTi systems in Li.30 High Ti alloys showed aweight increase by forming aTiN layer and high Cr alloys exhibited a weight loss as a result of the dissolution of Cr–N complexes. As the boundary of the two contradictory changes, Ti:Cr�2:1 was observed. Recently, a corrosion test using monometallic thermal convection Li loop made of V–4Cr–4Ti was conducted at 973 K for 2355 h. Because of the temperature gradient, weight loss and weight gain of V–4Cr–4Ti samples occurred at the hot leg and cold leg, respectively. However, the loss rate corresponded to only 400 Vanadium for Nuclear Systems alloys with various O levels. The loss of ductility by hydrogen charging was shown to be enhanced by impurity oxygen.35,36 4.12.8 Radiation Effects A fair amount of data is available for radiation response of vanadium alloys partly because they were candidates of cladding materials of LMFBR. For example, void swelling is known to be quite small if the alloy contains Ti. However, data are limited for V–4Cr–4Ti because this composition was decided as the reference one for fusion only recently. For this alloy, the feasibility issues of radia- tion effects are considered to be loss of ductility at lower temperature, embrittlement enhanced by trans- mutant helium at high temperature, and irradiation creep at intermediate to high temperature. The mechanism of the loss of uniform elongation of vanadium alloys at relatively low temperature ( 1024 1023 1022 1021 1020 1025 400 500 600 700 800 900 1000 1100 O:389, N:14 wppm (loop) O:28, N:27 wppm (loop) O:389, N:14 wppm (precipitates) O.28, N:27 wppm (precipitates) Temperature (K) D ef ec t d en si ty (m –3 ) Figure 18 Densities of dislocation loops and precipitates asa functionof irradiation temperature for twoV–4Cr–4Tialloys with different O and N levels (0.1dpa by Cu ion irradiation). Reproduced fromWatanabe, H.; Suda, M.; Muroga, T.; Yoshida, N. J. Nucl. Mater. 2002, 307–311, 408–411. US-832665 698 K-Li (HFIR) NIFS-HEAT-2 698 K-Li (HFIR) NIFS-HEAT-2 731 K-Na (JOYO) 0 0 50 100 Applied stress (MPa) C re ep s tr ai n (% ) 150 200 0.1 0.2 Figure 19 Creep strain as a function of applied stress for V–4Cr–4Ti (US-832665 and NIFS-HEAT-2) irradiated in Li (HFIR) and Na (JOYO) environments. The creep strain was normalized as that at two displacements per atom. Reproduced from Fukumoto, K.; Narui, M.; Matsui, H.; et al. J. Nucl. Mater. 2009, 386–388, 575–578. 0 0 2 4 6 8 10 12 100 200 300 400 Irradiation temperature (�C) 500 600 700 U ni fo rm e lo ng at io n (% ) Test temperature–Irradiation temperature Annealing conditions: 900–1125 �C, 3.6 or 7.2 ks V–4.8Ti–4.0Cr–Si, Al, Y ATR, 0.7–4.7 dpa FFTF, EBR-II, 10–54 dpa Loomis et al.,51 FFTF, 13–33dpa Zinkle et al.,52 EBR-II, 4 dpa Tsai et al.,53 ATR, 4.1–4.3 dpa Tsai et al.,54 BOR-60, 17–19 dpa Chung et al.,55 HFIR, 10 dpa Snead et al.,56 HFBR, 0.5 dpa V–3.8Ti–5.9Cr–Si, Al, Y V–(4–5)Cr–(4–5)Ti Figure 17 Uniform elongations as a function of irradiation temperature for V–(4–5)Cr–(4–5)Ti alloys and those with addition of Si, Al, and Y. Reproduced from Satou, M.; Chuto T.; Abe, K. J. Nucl. Mater. 2000, 283–287, 367–371. Vanadium for Nuclear Systems 401 conditions. DHCE is highly anticipated as a potential method to extend our understanding of the helium effects. However, for conclusive evaluation, a 14MeV neutron source is certainly necessary. The irradiation creep tests have made progress recently, partly because of the progress in fabricating high quality pressurized creep tube specimens with reduced impurity levels. Figure 19 shows the nor- malized creep strain as a function of applied stress by irradiation in HFIR and JOYO in Li and Na environments. The data also compare the perfor- mances of US and Japanese reference alloys.41 It was found that the creep strain rate exhibited a linear relationship with the effective stress up to 150MPa at �700K and the differences with the environments and the heats are small. 4.12.9 Tritium-Related Issues In the blanket, the tritium inventory is not considered to be the issue once liquid Li is used as the breeding and cooling materials owing to the high hydrogen solubility of Li. The behavior of hydrogen and its isotopes in vanadium alloys is a concern for tritium retention in the first wall. Deuterium retention of V–4Cr–4Ti was investigated by deuterium ion implantation followed by thermal desorption, in comparison with other candidate first wall materials. The study showed that the retained amount at 380 K was one and two orders of magnitude larger than graphite and tungsten, respectively. For the irra- diation at 773 K, the retained amount was almost the same as that of graphite and one order larger than tungsten.42 Surface composition was also known to influence the hydrogen transport. For example, the rate of absorption was highly influenced by prior heat treatment, inducing Ti surface segregation.43 Recent progress in detecting tritium by means of imaging plate (IP) enhanced the understanding of the tritium behavior in vanadium alloys. Figure 20 com- pares IP images of cold rolled V–4Cr–4Ti and pure V after tritium charging. Tritium is preferentially absorbed in Ti-rich precipitates that have a band structure to the rolling direction.44 Low High R ol lin g d ire ct io n V–4Cr–4Ti (NIFS-HEAT-2) 2 mm2 mm Low High Pure Vanadium Figure 20 Distribution of tritium measured by imaging plates for cold-rolled V–4Cr–4Ti and pure vanadium. The tritium was charged by gas absorption. After Homma, H.; Hatano, Y.; Daifuku, H.; et al. J. Nucl. Mater. 2007, 367–370, 887–891. 402 Vanadium for Nuclear Systems 4.12.10 Development of Advanced Alloys The performance of structural materials can strongly influence the blanket design. Especially, the opera- tion temperature window and expected lifetime are the key parameters. Increase in the upper operation temperature limit can enhance the blanket operation temperature and thus plant efficiency. Therefore, enhancing the high-temperature strength is the key issue for improving the performance of the blanket and thus the attractiveness of the fusion power sys- tems. For this purpose, efforts have been made to develop advanced vanadium alloys with potential use at higher temperature. One of the relatively simple ways to enhance the strength of the alloy is to change the thermal and mechanical treatment of the alloys. Especially, for- mation of a high density of precipitates can strengthen the alloy. Figure 21 shows microstructure and hardness of V–4Cr–4Ti as a function of the temperature of reheating for 1 h after annealing at 1373K for 1 h. The annealing at 1373K dissolves most of the thin precipitates and the reheating can form new precipitates. By choosing an appropriate reheating temperature (873–973K), the materials can be strengthened by the high density of fine precipitates. However, the strengthening by this treatment will be lost at >973K because of the coarsening of the pre- cipitates. To prevent the coarsening, cold work was applied to the specimens. Figure 22 shows the mini- mum creep rate for standard V–4Cr–4Ti and solution annealed, aged, and cold-worked V–4Cr–4Ti. Sup- pression of the creep rate occurred at 1073K but only with relatively high stresses.45 Microstructural analysis showed that the suppressive role of cold- work-induced dislocations was lost during the creep deformation by the change in the nature of the dis- locations from sessile ah100i type to gliding a/2h111i type.46 Further efforts are being made, for example, to cold-work followed by aging (strain-aging-induced strengthening). High-temperature strength of V–Cr–Ti alloys can be enhanced by increasing the Cr level. However, high Cr alloys have low ductility and fabricability issues. Recent detailed survey in V–xCr–4Ti alloys showed that the strength at high temperature increases with a small change in the DBTTwith the Cr level at �7%.47 High-strength vanadium alloys were made by addi- tion of Y, O, and N to vanadium followed by mechani- cal alloying (MA) and hot isostatic pressing (HIP). The addition of Y, O, and N was intended to enhance mechanical properties by dispersion of Y2O3 and YN and scavenging O and N from the matrix. Alloys pro- duced byoptimization of the processes had small grains and homogeneously dispersed particles and showed higher tensile strength than those of NIFS-HEATs with moderate uniform elongation, both at room tem- perature and 1073K as shown in Figure 23.48 Fine 0 Stress (MPa) M in im um c re ep r at e (s –1 ) SAACW STD 10–4 10–5 10–6 10–7 10–8 10–9 NIFS-HEAT-2 800 �C 750 �C 700 �C 100 120 140 160 180 200 220 240 260 280 300 Figure 22 Minimum creep rates as a function of applied stress for V–4Cr–4Ti with standard heat treatment (1273K for 1 h: STD) and precipitate-hardening heat treatment (1373K for 1 h, 873K for 20 h, and cold rolled: SAACW). Reproduced from Chen, J. M.; Nagasaka, T.; Muroga, T.; Qiu, S. Y.; Li, C.; Nita, N. J. Nucl. Mater. 2008, 374, 298–303. Reheat temperature (K) As solution heat treated Y ie ld s tr es s (M P a) 1373 K873 K 973 K 1073 K 1173 K 1273 K Precipitation Solution Perfectly dissolved Fine precipitates Start to dissolve Coarse precipitates (low density) Larger grain 250 873 973 1073 1173 1273 1373 300 350 400 450 200 nm Figure 21 Hardness and microstructure of V–4Cr–4Ti as a function of reheating temperature for 1 h after annealing at 1373K for 1 h. Vanadium for Nuclear Systems 403 grain and oxide dispersion increased high-temperature strength and inhibited formation of interstitial loops in the matrix by neutron irradiation because of the enhanced defect sinks. Thus, mechanically alloyed vanadium alloys have the potential to extend both low- and high-temperature operation limits. Other efforts to improve high-temperature strength of vanadium alloys include strengthening by internal oxidation.49 4.12.11 Critical Issues With the recent progress in the fabrication technology, the number of critical issues for the development of vanadium alloys for fusion reactors has been reduced. The remaining critical issues are thermal and irradia- tion creep, transmutant helium effects on high temper- ature mechanical properties, and radiation effects on fracture properties. The effect of helium, particularly, is still uncertain and can be evaluated precisely only with the use of 14MeV neutrons. This fact highly motivates the construction of a 14MeVneutron source. With the progress of the properties of vanadium alloys, the blanket concepts using the alloy become more attractive. Extension of the operation tem- perature window and lifetime of vanadium alloys contribute to the improvement of the quality of the blanket. Therefore, exploration of advanced vana- dium alloys from the current reference alloy is a valuable challenge for enhancing the expected per- formance, and then attractiveness, of fusion reactors. 4.12.12 Vanadium Alloy Development for Fusion Blankets In the fusion materials development strategy, the candidate structural materials are categorized into reference and advanced materials. As the reference First wall WC Li Be Flexible support Figure 24 A layout of the V/Be/Li test blanket module for International Thermonuclear Experimental Reactor proposed by Russia. After Kirillov, I. R.; Shatalov, G. E.; Strebkov, Y. S.; the RF TBM Team. Fusion Eng. Des. 2006, 81, 425–432. 400 nm Strain (%) S tr es s (M P a) TTest= 1073 K TTest= 298 K 0 200 400 600 800 20% V–1.0 vol% Y2O3–0.7 vol% YN NIFS-HEAT-1 Figure 23 Tensile test curves and microstructure of V–Y2O3–YN produced by mechanical alloying (MA) in comparison with V–4Cr–4Ti (NIFS-HEAT-1). After Kuwabara, T.; Kurishita, H.; Hasegawa, M. J. Nucl. Mater. 2000, 283–287, 611. 404 Vanadium for Nuclear Systems materials, RAFM steels were selected because they have the most matured industrial infrastructure. Development of the reference materials is crucial for the realization of DEMO (fusion demonstration power plant) in a timely manner. On the other hand, vanadium alloys and SiC/SiC were nominated as the advanced materials, which will contribute to increasing attractiveness of the fusion system in terms of cost of electricity and environmental benign- ness. It is recognized that the development of the advanced materials must also be enhanced now due to the long lead time necessary for their development. It should also be noted that vanadium alloys are the only nonferromagnetic and ductile materials of the three candidates. If the impact of the ferromagnetism of the RAFM on plasma operation should be unavoid- able and the brittleness of SiC/SiC should be deter- mined unaccepted by design studies, vanadium alloys could be the only candidate of low activation structural materials for fusion reactors. As shown in the summary of critical issues, a 14MeV neutron source is highly necessary for the qualification of vanadium alloys. IFMIF (Interna- tional Fusion Materials Irradiation Test Facility, a 14MeV neutron source) is under design and is recog- nized to be essential for developing structural materials for fusion reactors. The TBM to be installed in ITER is also considered to be an important milestone for technological integration. Figure 24 shows the design of the V/Li TBM in ITER proposed by Russia.50 The development of vanadium alloys is planned to pro- ceed with IFMIF for qualification of the alloy and ITER-TBM for technology integration, in addition to fundamental studies using fission reactors, etc. 4.12.13 Summary As to the application in nuclear systems, vanadium alloys were once candidate cladding materials for LMFBR, but, at present, are considered mostly as candidate low activation structural materials for fusion reactors. Vanadium for Nuclear Systems 405 Vanadium alloys, with the reference composition of V–4Cr–4Ti, are one of the three candidate low activation structural materials with RAFM and SiC/SiC. They are the only nonferromagnetic and ductile materials of the three candidates and thus are promising for advanced structural materials of fusion reactors. The self-cooled liquid lithium blanket using structural materials of vanadium alloys is an attractive concept because of the high heat transfer capability, high-temperature operation, simple structure, high tri- tium breeding capability, and low tritium leakage. Recent progress, especially in the fabrication technologies, has successfully reduced the number of critical issues enhancing the feasibility of the alloys for fusion application. Major remaining issues of vanadium alloys are thermal and irradiation creep, transmutant helium effects on mechanical properties, and radiation effects on fracture properties. For con- clusive characterization of the helium effects, the use of IFMIF is essential. Efforts are also being made to develop advanced vanadium alloys to extend the temperature window and lifetime of vanadium alloys in fusion reactor environments. References 1. Suzuki, T.; Noda, T.; Iwao, N.; Kainuma, T.; Watanabe, R. J. Nucl. Mater. 1976, 62, 205–212. 2. Kurtz, R. 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Chung, H. M.; Nowicki, L. J.; Smith, D. L. Fusion Mater. 1997, 29. DOE/ER-0313/22. 56. Snead, L. L.; Zinkle, S. J.; Alexander, D. J.; Rowcliffe, A. F.; Robertson, J. P.; Eatherly, W. S. DOE/ER-0313/23 (Dec. 31, 1997) 81. Available at http://www.ms.ornl.gov/ fusionreactor/dec97.shtml *P N 4.13 Concrete* D. J. Naus Oak Ridge National Laboratory, Oak Ridge, TN, USA � 2012 Published by Elsevier Ltd. 4.13.1 Introduction 408 4.13.2 Concrete Longevity, NPP Safety-Related Concrete Structures, Testing and In-Service Inspection Requirements, and Operating Experience 409 4.13.2.1 Historical Perspective on Concrete Longevity 409 4.13.2.2 NPP Safety-Related Concrete Structures 409 4.13.2.2.1 Boiling water reactors 412 4.13.2.2.2 Pressurized water reactors 413 4.13.2.3 Testing and In-Service Inspection Requirements 415 4.13.2.4 Operating Experience 416 4.13.3 Aging and Long-Term Durability Considerations 417 4.13.3.1 Design, Construction, Material Selection, and Maintenance Considerations 419 4.13.3.2 Materials of Construction, Degradation Mechanisms, Damage Modeling, and Long-Term Performance of Concrete Materials 419 4.13.3.2.1 Materials of construction 419 4.13.3.2.2 Degradation mechanisms 422 4.13.3.2.3 Damage modeling 422 4.13.3.2.4 Long-term performance of concrete materials 422 4.13.3.3 Assessment and Repair 424 4.13.3.3.1 Component selection 424 4.13.3.3.2 In-service inspections 425 4.13.3.3.3 Nondestructive examinations 425 4.13.3.3.4 Remedial methods 426 4.13.4 Structural Reliability Theory 427 4.13.5 Summary and Potential Research Topics 428 References 428 Abbreviations ACI repared for th o. DE-AC05-0 American Concrete Institute ANS American Nuclear Society ASME American Society of Mechanical Engineers ASTM ASTM International BWR Boiling water reactor C3A Tricalcium aluminate C2S Dicalcium silicate C3S Tricalcium silicate C4AF Tetracalcium aluminoferrite CEB-FIP International Federation for Structural Concrete CSNI Committee on the Safety of Nuclear Installations e Oak Ridge National Laboratory under Contract 0OR22725 CFR Code Federal Regulations C-S-H Calcium silicate hydrate GDC General Design Criteria GGBFS Ground granulated blast furnace slag IAGE WG Integrity of Components and Structures Working Group LWR Light-water reactor NEA Nuclear Energy Agency NPP Nuclear power plant NSSS Nuclear steam supply system PCA Portland Cement Association PWR Pressurized water reactor RG Regulatory guide RILEM International Union of Laboratories and Experts in Construction Materials, Systems and Structures RPV Reactor pressure vessel 407 408 Concrete USNRC United States Nuclear Regulatory Commission WIS University of Wisconsin 4.13.1 Introduction As concrete ages, changes in its properties will occur as a result of continuingmicrostructural changes (i.e., slow hydration, crystallization of amorphous constituents, and reactions between cement paste and aggregates) as well as environmental influences. These changes do not have to be detrimental to the point where concrete will not be able to meet its functional and performance requirements; however, concrete can suffer undesirable changes with time because of improper specifications, violation of specifications, or adverse performance of its cement paste matrix or aggregate constituents under physical or chemical attack. Additional information related to environmental effects on concrete is provided in molten core concrete interaction (Chapter 2.25, Core Concrete Interaction). Portland cement concrete durability is defined as its ability to resist weathering action, chemical attack, abrasion, or any other process or deterioration.1 A durable concrete is one that retains its original form, quality, and serviceability in the working envi- ronment during its anticipated service life. The mate- rials and mix proportions specified and used should be such as to maintain concrete’s integrity and, if applicable, to protect embedded metal from corro- sion.2 The degree of exposure anticipated for the concrete during its service life, together with other relevant factors related to mix composition, work- manship, and design, should be considered.3 Guide- lines for production of durable concrete are available in national consensus codes and standards, such as American Concrete Institute (ACI) 3184, which have been developed over the years through knowledge acquired in testing laboratories and supplemented by field experience. Serviceability of concrete has been incorporated into the codes through strength requirements and limitations on service load conditions in the structure (e.g., allowable crackwidths, limitations onmidspandeflections of beams, andmaximumservice level stresses in prestressed members). Durability gen- erally has been included through items such as specifi- cations for maximum water–cement ratios, minimum cementitious materials contents, type cementitious material, requirements for entrained air, and minimum concrete cover over reinforcement. Requirements are frequently specifiedintermsof environmentalexposure classes (e.g., chloride and aggressive ground environ- ments). Specifications in terms of service life require- ments (e.g., short 100 years) have only recently been developed, primarily through European standards.5 Water is the single most important factor con- trolling the degradation processes of concrete (i.e., the process of deterioration of concrete with time is generally dependent on the transport of a fluid through concrete), apart from mechanical deteriora- tion. The rate, extent, and effect of fluid transport are largely dependent on the concrete pore structure (i.e., size and distribution), presence of cracks, and micro- climate at the concrete surface. The primary mode of transport in uncracked concrete is through the cement paste pore structure (i.e., its permeability). The domi- nant mechanism controlling the rates of water penetra- tion into unsaturated or partially saturated concrete is absorption caused by capillary action of the concrete’s pore structure. To improve the durability of concrete, generally the capillary and pore size within the con- crete matrix should be reduced to a minimum. Although the coefficient of permeability for con- crete depends primarily on thewater–cement ratio and maximum aggregate size, it is influenced by the curing temperature, drying, cementitious materials content, and addition of chemical or mineral admixtures as well as the tortuosity of the path of flow. Concrete compressive strength has traditionally been utilized as an acceptance test for concrete, but it typically is not a good indicator of durability. Many structures have been fabricated with concretes having adequate 28-day compressive strength only to lose their func- tionality because they were facing an environment for which they had not been designed or because the concrete had not been placed or cured correctly.6 The safety-related concrete structures in nuclear power plants (NPPs) are designed to withstand load- ings from a number of low-probability external and internal events, such as earthquake, tornado, and loss- of-coolant accident. Consequently, they are robust and not subjected to high enough stresses during normal operation to cause appreciable degradation. In general, this has been the case, as the performance of reinforced concrete structures in NPPs has been good. (Operating experience is discussed in the next section.) However, as the NPPs age, degradation incidences start to occur at an increasing rate, pri- marily due to environmental-related factors. One- fourth of all containments in the United States have experienced corrosion, and nearly half of the con- crete containments have reported degradation related to either the reinforced concrete or post-tensioning Concrete 409 system.7 Although the vast majority of these struc- tures will continue to meet their functional and per- formance requirements during their initial licensing period (i.e., nominally 40 years), it is reasonable to assume that with the increasing age of the operating reactors there will be isolated examples where the structures may not exhibit the desired durability without some form of intervention. Currently, the United States has 104 NPP units licensed for commercial operation, which provide about 20% of the electricity supply. As all but one of the construction permits for existing NPPs in the United States were issued prior to 1978, the focus for the existing plants has shifted from design to condition assessment. Here, the aim is to demonstrate that struc- tural margins of the plants have not eroded or will not erode during the desired service life due to aging or environmental effects. One of the key factors to maintaining adequate structural margins to protect public health and safety in the unlikely event of an accident is implementation of effective inspection and maintenance programs. An inspection program is important for identifying and characterizing any deg- radation that may be present in a timely manner. Once degradation has been identified, or its potential to occur established, a maintenance program is imple- mented to repair the degradation and arrest (as far as possible) the mechanism(s) causing the degradation. Proper maintenance is essential to the safety of NPP structures, and a clear link exists between effective maintenance and safety. Uncertainty in condition assessment can be assessed using probabilisticmethods, which are also an essential ingredient of risk-informed management decisions concerning continued service of the NPP structures. Figure 1 Pantheon, built 119–128 AD. From Http://en. wikipedia.org/wiki/file:Pantheon_rome_2005may.jpg 4.13.2 Concrete Longevity, NPP Safety-Related Concrete Structures, Testing and In-Service Inspection Requirements, and Operating Experience 4.13.2.1 Historical Perspective on Concrete Longevity Concrete, originally based on lime that hardened by atmospheric carbonation, has been utilized as a con- struction material for several thousand years. Cement has been around for at least 12My when reactions occurred between limestone and oil shale during spon- taneous combustion in Israel to form a natural deposit of cement compounds.8 The oldest known concrete is from Yugoslavia and is about 7600 years old.9 Gypsum mortars were used by the Egyptians to fabricate the Great Pyramid at Giza about 2500BC. The Romans were the first to use hydraulic limes and discovered the benefits of pozzolans. The survival of several ancient concrete structures (e.g., Pantheon in Rome, Figure 1) attests to the durability that concrete can attain. A detailed study involving an examination of sam- ples obtained from several ancient concrete struc- tures utilizing physical and chemical techniques concluded that these structures survived primarily because of careful selection of materials and con- struction, mild climatic conditions, and the lack of steel reinforcement.9 These structures, however, were not fabricated using current ‘hydraulic Portland cement,’ as it did not exist until about 1824. Some information, however, was presented in Mallinson and Davies9 relative to samples that were obtained for testing from several structures fabricated in the mid- to late 1800s. It was concluded that the durabil- ity of these structures was not only due to high cement contents but also due to the relatively slow cement-setting times and high construction quality. These Portland cements differ somewhat from the Portland cements used to fabricate NPP concrete structures in that the formulations have changed as well as the fineness of the cement. Also, modern con- cretes have incorporated admixtures to improve work- ability, modify hardening or setting characteristics, aid in curing, and enhance the performance or durability. 4.13.2.2 NPP Safety-Related Concrete Structures All commercial NPPs in the United States contain structures whose performance and function are necessary for the protection of the safety of 410 Concrete plant-operating personnel and the general public, as well as the environment. The basic laws that regulate the design (and construction) of NPPs are contained in Title 10 of the Code of Federal Regulations (CFR),10 which is clarified by Regulatory Guides (e.g., R.G. 1.29),11 NUREG reports, Standard Review Plans (e.g., Con- crete and Steel Internal Structures of Steel or Concrete Containments),12 etc. In addition, R.G. 1.29 and Part 100 to Title 10 of the CFR state that NPP structures important to safety must be designed to withstand the effects of earthquakes without the loss of function or threat to public safety. These ‘safety-related’ structures are designed as Seismic Category I. Seismic Category I structures typically include those classified by the American Society of Mechanical Engineers (ASME) and the American Nuclear Society (ANS) as Classes 1, 2, and 3 (i.e., safety related). Initially, existing building codes such as the ACI Standard 318, Building Code Requirements for Reinforced Concrete, were used in the nuclear industry as the basis for the design and construction of concrete structural members. However, because the existing building codes did not cover the entire spectrum of design requirements and because they were not always con- sidered adequate, the United States Nuclear Regu- latory Commission (USNRC) developed its own criteria for the design of Category I structures (e.g., definitions of load combinations for both operating and accident conditions). Current requirements for nuclear safety-related concrete structures, other than concrete reactor vessels and concrete containments, are also based on ACI 318, but have incorporated modifications to accommodate the unique perfor- mance requirements of NPPs. These requirements were developed by ACI Committee 349 and first published in October 1976.13 This Code has been endorsed by the USNRC as providing an adequate basis for complying with the general design criteria for structures other than reactor vessels and contain- ments.14 USNRC15 provides additional information on the design of seismic Category I structures that are required to remain functional if the Safe Shutdown Earthquake (SSE) occurs. Current requirements for concrete reactor vessels and concrete containments were developed by ACI Committee 359 and first published in 1977.16 Supplemental load combination criteria are presented in Section 3.8.1 of the USNRC Regulatory Standard Review Plan.17 However, since all but one of the construction permits for existing NPPs have been issued prior to 1978, it is unlikely that endorsed versions of either ACI 349 or ACI 359 were used in the design of many of the concrete structures at these plants. Older plants that used early ACI codes, however, have been reviewed by the USNRC through the Systematic Evaluation Program to determine if there were any safety concerns.18 Each boiling water reactor (BWR) or pressurized water reactor (PWR) unit in the United States is located within a much larger metal or concrete con- tainment that also houses or supports the primary coolant system components. Although the shapes and configurations of the containment can vary signifi- cantly from plant to plant, leak tightness is ensured by a continuous pressure boundary consisting of non- metallic seals and gaskets and metallic components that are either welded or bolted together. There are several CFR General Design Criteria (GDC) and ASME Code sections that establish minimum requirements for the design, fabrication, construc- tion, testing, and performance of the light-water reac- tor (LWR) containment structures. The GDC serve as fundamental underpinnings for many of the most important safety commitments in licensee design and licensing bases. General Design Criterion 2, Design Bases for Protection Against Natural Phenomena, requires the containment to remain functional under the effects of postulated natural phenomena such as earthquakes, tornadoes, hurricanes, floods, tsunami, and seiches. General Design Criterion 16, Containment Design, requires the provision of reactor containment and associated systems to establish an essentially leak- tight barrier against the uncontrolled release of radio- activity into the environment and to ensure that the containment design conditions important to safety are not exceeded for as long as required for postulated accident conditions. Criterion 53, Provisions for Con- tainment Testing and Inspection, requires that the reactor containment be designed to permit (1) appropriate periodic inspection of all important areas, such as penetrations; (2) an appropriate surveillance program; and (3) periodic testing at containment design pressure of leak tightness of penetrations that have resilient seals and expansion bellows. Current LWR contain- ments are considered as a significant element of the USNRC’s safety policy, which employs a defense- in-depth approach (i.e., successive compensatory measures are exercised to prevent accidents or miti- gate damage if a malfunction, accident, or naturally caused event occurs). The defense-in-depth philoso- phy ensures that safety will not be wholly dependent on any single element of the design, construction, maintenance, or operation at a nuclear facility (e.g., the facility in question tends to be more tolerant of failures and external challenges). Concrete 411 From a safety standpoint, the containment is one of the most important components of an NPP because, independent of the fuel barrier and reactor coolant pressure boundary barrier, it serves as the final barrier to the release of fission products to the outside environment under postulated accident conditions. During normal operating conditions, the containment is subject to various operational and environmental stressors (e.g., ambient pressure fluc- tuations, temperature variations, earthquakes, and wind storms). In some containment designs, the prin- cipal leak-tight barrier is surrounded by another structure (e.g., reactor or shield building) that pro- tects the containment from external events. Ensuring Table 1 Typical safety-related concrete structures in LWR p Concrete structure Primary containment Containment dome/roof Containment foundation/basemat Slabs and walls Containment internal structures Slabs and walls Reactor vessel support structure (or pedestal) Crane support structures Reactor shield wall (biological) Ice condenser dividing wall (ice condenser plants) NSSS equipment supports/vault structures Weir and vent walls (Mark III) Pool structures (Mark III) Diaphragm floor (Mark II) Drywell/wetwell slabs and walls (Mark III) Secondary containment/reactor buildings Slabs, columns, and walls Foundation Sacrificial shield wall (metallic containments) Fuel/equipment storage pools Walls, slabs, and canals Auxiliary building Fuel storage building Control room (or building) Diesel generator building Piping or electrical cable ducts or tunnels Radioactive waste storage building Stacks Intake structures (including concrete water intake piping and canal embankments) Pumping stations Cooling towers Plant discharge structures Emergency cooling water structures Dams Water wells Turbine building Source: Hookham, C. J. In-Service Inspection Guidelines for Concrete Lockheed Martin Energy Systems, Oak Ridge National Laboratory: Oa that the structural capacity and leak-tight integrity of the containment has not deteriorated unacceptably because of aging or environmental stressor effects is essential to reliable continued service evaluations and informed aging management decisions. More detailed information on containments is available.19 In addition to the containment, a myriad of concrete-based structures are contained as a part of an LWR plant to provide foundation, support, shield- ing, and containment functions. Table 1 presents a listing of typical safety-related concrete structures that may be included as part of an LWR plant.20 Relative to general civil engineering reinforced con- crete structures, NPP concrete structures tend to be lants and their accessibility for visual examination Accessibility Internal liner/complete external Internal liner (not embedded) or top surface Internal liner/external above grade Generally accessible Typically lined or hard to access Generally accessible Typically lined Lined or hard to access Generally accessible Lined with limited access Lined Lined with limited access Internal liner/partial external access Accessible on multiple surfaces Top surface Internal lined/external accessible Internal lined/partial external Generally accessible Generally accessible Generally accessible Generally accessible Limited accessibility Generally accessible Partial internal/external above grade Internal accessible/external above grade and waterline Partially accessible Accessible above grade Internal accessible/external above grade and waterline Limited accessibility External surfaces above waterline Limited accessibility Generally accessible Structures in Nuclear Power Plants, ORNL/NRC/LTR-95/14; k Ridge, TN, 1995. 412 Concrete more massive and have increased steel reinforcement densities with more complex detailing. Information pertaining to a particular structure at a plant of interest can be obtained from sources such as the plant’s safety analysis report or docket file. Concrete structures that are considered to be ‘plant specific’ or unique have not been addressed in the discussion later, but some information provided for similar structures may be applicable. Additionally, the names of certain structures may vary from plant to plant depending on the nuclear steam supply system (NSSS) vendor, architect engineering firm, and owner preference. Typical safety-related concrete structures contained in LWR plants may be grouped into four categories: primary containments, contain- ment internal structures, secondary containment/ reactor buildings, and other structures. 4.13.2.2.1 Boiling water reactors Of the BWR plants that have been licensed for com- mercial operation in the United States, �30% utilize either reinforced or prestressed concrete primary containments. Leak tightness of each of these con- tainments is provided by a steel liner attached to the containment inside surface by studs (e.g., Nelson studs) or by structural steel members. Exposed sur- faces of the carbon steel liner are typically painted to protect against corrosion and to facilitate decontami- nation should it be required. A portion of the liner toward the bottom of the containment and over the basemat is typically embedded in concrete to protect it from damage, abrasion, etc. due to corrosive fluids and impact. A seal to prevent the ingress of fluids is provided at the interface around the circumference of the containment where the vertical portion of the liner becomes embedded in the concrete. BWR con- tainments, because of provisions for pressure sup- pression, typically have ‘normally dry’ sections (drywell) and ‘flooded’ sections (wetwell) that are interconnected via piping or vents. Requirements for BWR containments include the following: 1. Provide an ‘essentially’ leak-tight barrier against the uncontrolled release of radioactivity to the environment for all postulated design basis acci- dent conditions. 2. Accommodate the calculated pressure and tem- perature conditions resulting from a loss-of- coolant accident. 3. Withstand periodic integrated leak-rate testing at the peak-calculated accident pressure that may be at levels up to and including the containment design pressure. 4. Permit appropriate periodic inspection of all impor- tant components and surfaces, and the periodic test- ing of the leak tightness of containment penetrations. The containment vessel can also provide structural support for the NSSS and other internal equipment. The containment foundation, typically a basemat, provides the primary support and transfer of load to the earth below. Figure 2 presents a cross-section of a BWR Mark I reinforced concrete containment. Each of the three BWR primary plant types (Mark I, Mark II, and Mark III) incorporates a number of reinforced concrete containment internal structures. These structures may perform singular or several functions, including the following: 1. Radiation shielding; 2. human accessibility provisions; 3. NSSS and other equipment anchorage/support/ protection; 4. resistance to jet, pipe whip, and other loadings produced by emergency conditions; 5. boundary of wetwells and pool structures, and allow communication between drywell and wetwell (Mark II and III); 6. lateral stability for containment; 7. transfer of containment loads to underlying foun- dation; and 8. transfer of fuel to reactor (Mark III). As many of these functions are interrelated with the required containment functions, these structures are considered to be safety-related. Of the BWR plants that utilize steel primary containments, all but the pre-Mark plant type have reinforced concrete structures that serve as second- ary containments or reactor buildings and provide support and shielding functions for the primary con- tainment. Although the design parameters for the secondary containments of the Mark I and Mark II plants vary somewhat, the secondary containments are typically composed of beam, floor, and wall struc- tural elements. These structures typically are safety- related because they provide additional radiation shielding; provide resistance to environmental/opera- tional loadings; and house safety-related mechanical equipment, spent fuel, and the primary metal con- tainment. Although these structures may be massive in cross-section to meet shielding or load-bearing requirements, they generally have smaller elemental Metal liner Steel liner Vent Reactor pressure vessel support skirt Reactor pressure vessel pedestal Core Reactor pressure vessel Concrete containment vessel Equipment pool Fuel storage pool Reactor building Biological shield Drywell Drywell head Carbon steel and stainless steel (below waterline liner) Downcomer pipe Basemat Grade Pressure suppression chamber Polar crane Figure 2 Boiling water reactor Mark I reinforced concrete containment and reactor building. Concrete 413 thicknesses than primary containments because of reduced exposure under postulated accident loadings. These structures may be maintained at a slight nega- tive pressure for collection and treatment of any air- borne radioactive material that might escape during operating conditions. Other structures include such things as founda- tions, walls, slabs, and fuel/equipment storage pools. The spent- and new-fuel storage pools, and the pools for reactor internals storage, typically have a four- wall-with-bottom-slab configuration. The walls and slab are composed of reinforced concrete members lined on the interior surface with stainless steel. Cross-sections of these members are generally large because they must support a large pool of water and heavy fuel/component loads produced by high-density fuel storage considerations. The fuel storage pool in Mark III plants is located within the primary containment. 4.13.2.2.2 Pressurized water reactors Of the PWR plants that have been licensed for commercial operation in the United States, �80% utilize either reinforced or prestressed concrete pri- mary containments. In meeting the same basic func- tional and performance requirements as noted for BWR containments, the concrete containments in PWR plants are of three different functional designs: subatmospheric (reinforced concrete), ice condenser (reinforced concrete), and large/dry 414 Concrete (reinforced and prestressed concrete). The primary differences between these containment designs relate to volume requirements, provisions for accident load- ings/pressures, and containment internal structures layout. The PWR containment structure generally consists of a concrete basemat foundation, vertical cylindrical walls, and dome. The basemat may consist of a simple mat foundation on fill, natural cut, or bedrock or may be a pile/pile cap arrangement. Most of the plants have utilized the simple mat on fill or bedrock design. Interior containment surfaces are lined with a thin carbon steel liner to prevent leakage. Exposed surfaces of the carbon steel liner are typically painted to protect against corrosion and to facilitate decontamination should it be required. Depending on the functional design, the concrete containments can be on the order of 40–50m in diameter and 60–70m high, with wall and dome thicknesses from 0.9 to 1.4m and base slab Basemat Grade In-core instrument guide tubes Reactor vessel 42.7 m Seal table Prestressed reinforcing Steam generators Polar crane Ring girder (tendon anchorage) Dome Steel liner Figure 3 Pressurized water reactor large dry prestressed con thicknesses from 2.7 to 4.1m. Two of the PWR plants (Bellefonte and Ginna) have rock anchor systems to which the post-tensioning tendons are attached. Figure 3 presents a cross-section for a prestressed concrete, large, dry containment. The containment internal structures in PWR plants are typically constructed of conventionally reinforced concrete and tend to be more massive in nature than the internal structures in BWR plants, because they typically support the reactor pressure vessel, steam generators, and other large equipment and tanks. In addition, these structures provide shielding of radiation emitted by the NSSS. Some of the specific functions that these structures (typically floor slabs, walls, and columns) are required to per- form include the following: 1. provision of human accessibility; 2. support and separation of various plant equipment; Reactor cavity 2.5 m 64.7 m 8 m crete containment. Concrete 415 3. resistance to emergency loading conditions; 4. transfer of containment loads to containment foundation; 5. missile protection; and 6. channeling/routing steam and air through ice condensers (PWR ice condenser containments). PWR plants that utilize a metallic primary contain- ment (large dry and ice condenser designs) are usually contained in reinforced concrete ‘enclosures’ or ‘shield’ buildings that, in addition towithstanding environmen- tal effects, provide radiation shielding and particulate collection and ensure that the freestanding metallic primary containment is protected from the natural environment. The secondary containment consists of a vertical cylinder wall with shallow dome and is often supported by the containment basemat. Except for differences in the spent- and new-fuel storage pools, structures that fall into the other struc- tures category are essentially the same at the PWR and BWR plants. The spent- and new-fuel storage pools for PWR plants are typically located in an auxiliary building proximate to the containment. These reinforced concrete wall and slab structures are generally massive in cross-section to support a large pool of water and the fuel elements and are lined on the water side with stainless steel. The pools are connected to the reactor/refueling cavity (inside containment) via a transfer channel that is also a safety-related structure since it must provide radia- tion shielding and support for the fuel transport mechanism and fuel. 4.13.2.3 Testing and In-Service Inspection Requirements One of the conditions of all operating licenses for water-cooled power reactors in the United States is that the primary reactor containments shall meet the containment leakage test requirements set forth in Appendix J, Primary Reactor Containment Leakage Testing for Water-Cooled Power Reactors, to 10 CFR Part 50.21 These test requirements provide for preoperational and periodic verification by tests of the leak-tight integrity of the primary reactor containment as well as systems and components that penetrate contain- ment of water-cooled power reactors and establish the acceptance criteria for such tests. The purpose of these tests is to ensure that (1) leakage through the primary reactor containment and the systems and components penetrating primary reactor contain- ment shall not exceed allowable leakage-rate values as specified in the technical specifications or asso- ciated bases and (2) periodic surveillance of reactor containment penetrations and isolation valves is per- formed so that proper maintenance and repairs are made during the service life of the containment as well as systems and components that penetrate pri- mary containment. Contained in this regulation are requirements pertaining to Type A, B, and C leakage- rate tests that must be performed by each licensee as a condition of their operating license. Type A tests are intended to measure the primary reactor containment overall integrated leakage rate (1) after the contain- ment has been completed and is ready for operation and (2) at periodic intervals thereafter. Type B tests are intended to detect local leaks and to measure leakage across each pressure-containing or leakage-limiting boundary for primary reactor containment penetra- tions (e.g., penetrations that incorporate resilient seals, gaskets, or sealant compounds and air lock door seals). Type C tests are intended to measure containment isolation valve leakage rates. Requirements for system pressure testing and criteria for establishing inspection programs and pressure-test schedules are contained in Appendix J. Appendix J to 10 CFR Part 50 also requires a general inspection of the accessible interior and exterior surfaces of the containment structures and components to uncover any evidence of structural deterioration that may affect either the containment structural integrity or leak tightness. Subsection IWL of ASME Section XI addresses reinforced and post- tensioned concrete containments (Class CC). Two examination categories are provided in Subsection IWL. Examination Category L-A addresses accessi- ble concrete surfaces and examines them for evidence of damage or degradation, such as cracks. The con- crete is examined at 1, 3, and 5 years following the containment structural integrity test and every 5 years thereafter. The primary inspection method of Category L-A is visual examination (general or detailed). Examination Category L-B addresses the unbonded post-tensioning system. The unbonded post-tensioning system examination schedule is the same as for the concrete. For post-tensioned concrete containments, tendon wires are tested for yield strength, ultimate tensile strength, and elongation. Tendon corrosion protection medium is analyzed for alkalinity, water content, and soluble ion concen- trations. Prestressing forces are measured for selected sample tendons. Subsection IWL specifies accep- tance criteria, corrective actions, and expansion of the inspection scope when degradation exceeding 416 Concrete the acceptance criteria is found. Additional require- ments for inaccessible areas are specified in 10 CFR 50.55a(b)(2)(viii). The acceptability of concrete in inaccessible areas is to be evaluated when conditions that could indicate the presence or result in degrada- tion to such inaccessible areas exist in accessible areas. Information on aging management programs for masonry walls22,23 and water-control structures24 is available. Inspection requirements for steel con- tainments and liners of concrete containments are contained in Subsection IWE of ASME Section XI. Editions and addenda of the ASME Code acceptable to the USNRC are identified in 10 CFR 50.55a. 4.13.2.4 Operating Experience In general, the performance of NPP safety-related concrete structures has been very good. However, there have been several isolated incidences that, if not remedied, could challenge the capacity of the contain- ment and other safety-related structures tomeet future functional and performance requirements. Table 2 presents a summary of local degradation mechanisms Table 2 Condition survey results from several plants for NP Local degradation mechanism Plant A B C D Concrete Chemical attack b,c c b c Efflorescence and leaching b,c,d b,c Alkali–aggregate reaction Freeze–thaw cycling d a,d Thermal exposure c c Abrasion/erosion c Fatigue/vibration c Cracking c,d,f,g a,b,c,d c,d,g c,d Conventional reinforcement Corrosion b,d b,d b Prestressing system Corrosion e1 Block walls Excessive cracking c Structural steel and liners Corrosion d e c,d Soil/structure issues Differential settlement c Soil erosion (scour) d a – external structure (power block); b – subgrade structure (power blo (Intake, discharge, etc.); e – Containment vessel (power block); f – Oth block). Source: Gregor, F. E.; Hookham, C. J. Remant life preservation of LW Conference on Structural Mechanics in Reactor Technology, Stuttgart, The Netherlands, 1993; Paper DH06/2. 1Corrosion limited to exposed grease can and bearing plate surface. that have been observed by one organization during condition surveys of various concrete structures at both United States and foreignNPPs located in areas having several different climatic conditions.25 Some general observations derived from these results were that vir- tually all NPPs have experienced cracking of the con- crete structures that exceeds typical acceptance criteria for width and length; numerous NPPs had groundwater intrusion occurring through the power block or other subsurface structures; and aging con- cerns exist for subsurface concrete structures, as their physical condition cannot be easily verified. Collec- tively, it was concluded in this study that the general performance of the NPP concrete structures, has been quite favorable and proper evaluation and treatment of observed degradation at an early stage is both a cost- effective and necessary approach to long-term plant operations. Initially, degradation of NPP concrete structures in the United States occurred early in their life and has been corrected.26–28 Causes were primarily related either to improper material selection and construction/design deficiencies or environmental P structures E F G H I J c c c c b,d b,d d b,d,f a,b,c,d b,f a d f c c c,d a,b,c,d,g b,c,d,f,g b,f b,c,d,f b,c,d,f b,f,g b,d b b,d b b,f d c a c,e e g ck); c – internal structure (power block); d – water control structure er site structure (power block); g – Equipment supports (power R plant structures. In Transactions of the 12th International Germany, Aug 15–20; Elsevier Science: Amsterdam, Concrete 417 effects. Examples of some of the problems attributed to these deficiencies include low 28-day concrete com- pressive strengths; voids under the post-tensioning tendon-bearing plates resulting from improper con- crete placement; cracking of post-tensioning tendon anchor heads due to stress corrosion or embrittlement; and containment dome delaminations due to low- quality aggregate materials and absence of radial steel reinforcement or unbalanced prestressing forces.29–31 Other construction-related problems included occur- rence of excessive voids or honeycomb in the concrete, contaminated concrete, cold joints, cadweld (steel reinforcement connector) deficiencies, materials out of specification, higher than code allowable concrete temperatures, misplaced steel reinforcement, post- tensioning system button-head deficiencies, and water-contaminated corrosion inhibitors.26 Although continuing the service of a NPP past the initial operating license period is not expected to be limited by the concrete structures, several incidences of age- related degradation have been reported.28–33 Examples of some of these problems include corrosion of steel reinforcement in water intake structures, corrosion of post-tensioning tendon wires, leaching of tendon gal- lery concrete, low prestressing forces, and leakage of corrosion inhibitors from tendon sheaths. Other related problems include cracking and spalling of con- tainment dome concrete due to freeze–thaw damage, Concrete cracking outside containment wall Crease leakage outside containment wall Anchor head failure Exterior concrete wallcracks and spalling Figure 4 Examples of degradation related to nuclear power p low strengths of tendon wires, contamination of corro- sion inhibitors by chlorides, and corrosion of concrete containment liners. As the plants age, the incidences of degradation are expected to increase, primarily due to environmental effects. A listing of documented concrete problem areas by plant, type reactor, and degradation is available.34 Documented information on problem areas experienced with NPP concrete struc- tures in other countries has also been assembled.35 Figure 4 presents examples of occurrences of degrada- tion that have been observed at NPPs. Anchor head failure and containment dome delamination shown in the figure represent occurrences related to materials selection and design, respectively, with the remainder representing aging-related occurrences. 4.13.3 Aging and Long-Term Durability Considerations In the United States, the Atomic Energy Act and regulations of the USNRC limit commercial power reactor licenses to an initial 40-year period, but per- mits such licenses to be renewed. (Other countries may not have a limit set on the plant operating license period, but the utility must obtain a permanent renewal of its operating license subject to numerous and continuous justifications (e.g., periodic safety Containment dome delamination repair Water intake structure rebar corrosion Concrete wall water infiltration Corrosion of grease cap lant concrete structures. 418 Concrete reevaluations).) This original 40-year term for reactor licenses was based on economic and antitrust consid- erations – not on limitations of nuclear technology. Due to this selected period, however, some structures and components may have been engineered on the basis of an expected 40-year service life. Several nuclear power units in the United States have reached the end of their initial operating license period. To help ensure an adequate energy supply, the USNRC has established a timely license renewal process and clear requirements that are needed to ensure safe plant operation for an extended plant life. These requirements are codified in Parts 51 and 54 of Title 10, Energy, of the CFR that provides for a renewal of an operating license for an additional 20 years. In order to ensure the safe operation of NPPs, it is essential that the effects of age-related degradation of plant structures, as well as systems and components, be assessed and managed during both the current operating license period as well as subsequent license renewal periods. As these plants mature, environmental factors are going to become increasingly important. Demonstra- tion of continued safe and reliable operation of the plants will involve implementation of a program that effectively manages aging to ensure the availability of design safety functions throughout the plant ser- vice life. Examples of considerations to be addressed by such a program for the safety-related concrete structures are identified in Figure 5 and include the following: Minimum required performance I0 Imin Ir Actual performance Performance spread Repair implemented Current and future condition assessments Characterization of materials – database Remedial measures Time Example: not NPP P er fo rm an ce in d ur ab ili ty t er m s Figure 5 Components of an aging management program for n 1. What environmental stressors or aging factors are most important with respect to impacting struc- tural reliability? 2. What in-service inspection or condition assess- ment programs are most effective in demonstrat- ing structural reliability, and how often should they be applied? 3. What material sampling and testing programs should be required, if any? 4. How effective are remedial measures in enhancing the reliability of the structures and extending their usable life? 5. How have the material properties changed under the influence of aging and environmental stressors? 6. What is the residual life of the structure and how might it respond to something like a design basis event? General guidance on developing an aging manage- ment program for concrete containment buildings has been developed.35 Included in this reference is information related to practices and techniques that have been utilized by various countries for assessing the fitness for service as well as inspection, monitoring, and mitigation of aging degradation of concrete con- tainment buildings. The International Union of Laboratories and Experts in Construction Materials, Systems and Structures (RILEM) has held an interna- tional conference, prepared a report, and sponsored two workshops related to aging management of con- crete structures.36–39 Finally, the Nuclear Energy In-service inspection Material sampling/testing Example: not NPP Environmental stressors/ aging factors uclear power plant concrete structures. Concrete 419 Agency Committee on the Safety of Nuclear Installa- tions (NEA/CSNI) under its Integrity of Components and Structures Working Group (IAGE WG) has prepared several reports and held a series of work- shops that addressed various aspects of aging of NPP concrete structures.40–52 Also, there are a number of other documents that address aging of NPP concrete structures,32,53–56 as well as national programs. 4.13.3.1 Design, Construction, Material Selection, and Maintenance Considerations Design errors that can lead to subsequent deteriora- tion of concrete structures can be placed into two categories: inadequate structural design and lack of attention to details.57 Inadequate structural design occurs when the structure is exposed to a load greater than it is capable of carrying or if it sustains greater strain than its strain capacity. Inadequate considera- tions of temperature change or concrete creep and accidental impact can also result in damage. Typical symptoms of inadequate design include spalling and cracking of concrete. Poor detailing of a structure may result in localized concentration of stresses that result in cracking, which in turn can permit water or chemicals to access the concrete or ponding of water to produce saturated concrete. Poor detailing does not generally lead directly to concrete failure but can contribute to the action of one of the other specific causes of concrete failure.57 Examples of inadequate structural design include insufficient concrete cover over steel reinforcement, improper sizing and place- ment of steel reinforcement, inadequate section geometry, inadequate provision for drainage, abrupt changes in section, material incompatibility, and inadequate provision for deflection. Poor construction practices and negligence can result from not following specified procedures or from carelessness. Poor construction practices do not lead directly to failure or deterioration of concrete but can cause defects that lead to concrete cracking. Examples of concrete cracks that can result from poor construction practices include plastic shrinkage, plastic settlement, early thermal contraction, crazing, and long-term drying shrinkage. The resulting con- crete cracking then can enhance the adverse impacts of mechanisms (such as described in the next section) and lead to concrete degradation. Poor construction prac- tices and negligence are best addressed through ade- quate quality assurance/quality control in conjunction with an aggressive inspection program. Examples of poor construction practice include adding additional water to concrete to facilitate placement or finishing, improper mixing and curing, improper consolidation, and improper location of steel reinforcement. Lack of knowledge about the importance of care- ful selection and specification of materials and use of admixtures can also result in durability issues. This can include improper cement contents, use of poor quality or contaminated aggregates, incorporation of additives that can produce corrosion such as calcium chloride accelerators, and incorrect water–cement ratios. Improper or inadequate maintenance also can con- tribute to the deterioration of concrete structures. Examples of inadequatemaintenance includemoisture exposure and penetration caused by unrepaired cracks, improper application of coatings, damaged waterstops, and failure to clean drains and drain pathways. 4.13.3.2 Materials of Construction, Degradation Mechanisms, Damage Modeling, and Long-Term Performance of Concrete Materials 4.13.3.2.1 Materials of construction Nuclear safety-related concrete structures are com- posed of several constituents that, in concert, perform multiple functions (e.g., load-carrying capacity, radi- ation shielding, and leak tightness). Primarily, these constituents can include the following material sys- tems: concrete, conventional steel reinforcement, pre- stressing steel, and steel liner plate. The quality of these materials is established through regulations, qualification tests, and certification, followed by check- ing throughout construction. More detailed informa- tion on materials of construction than provided later is available elsewhere.35,58–60 Concrete is a composite material consisting of a binder (cement paste) and a filler of fine or fine and coarse aggregate particles that combine to form a synthetic conglomerate. Cement is a mixture of compounds made by grinding crushed limestone, clay, sand, and iron ore together to form a homoge- neous powder that is then heated at very high tem- peratures ranging from 1400 to 1600 �C to form a clinker.59 After the clinker cools, it is ground and mixed with a small amount of gypsum to regulate setting and facilitate placement. This produces the general-purpose Portland cement, which is mixed with water to produce cement paste that binds the aggregate particles together. (Current generation cements have higher tricalcium silicate (C3S) contents and are ground finer than previous cements. Current 420 Concrete cements attain most of their compressive strength within a 28-day period, whereas the previous cements continued to gain strength after 28 days.6,61) Portland cements are composed primarily of four chemical compounds: (C3S), dicalcium silicate (C2S), tricalcium aluminate (C3A), and tetracalcium aluminoferrite (C4AF). The type of Portland cement produced (e.g., general purpose, moderate sulfate resistance and heat of hydration, high early strength, low heat of hydration, and sulfate resistant) depends on the relative amounts of the four basic chemical compounds and fineness (high early strength). The calcium silicate hydrates (C–S–H) constitute about 75% of the mass. The C–S–H gel structure is made up of three types of groups that contribute to bonds across surfaces or in the interlayer of partly crystal- lized tobermorite material: calcium ions, siloxanes, and water molecules. Bonding of the water within the layers (gel water) with other groups via hydrogen bonds determines the strength, stiffness, and creep properties of the cement paste. There are also a number of alternative or sup- plementary cementing agents that have been used in conjunction with Portland cement, and these are pulverized fly ash, ground granulated blast furnace slag (GGBFS), and silica fume. Fly ash is collected from the exhaust flow of furnaces burning finely ground coal and reacts with calcium hydroxide in the presence of water to form cement compounds consisting of calcium silicate hydrate. GGBFS is a by-product of the iron-making process and is formed by taking the hot slag, rapidly chilling or quenching it, and grinding into a powder. When mixed with water in the presence of an alkaline environment provided by the Portland cement, GGBFS hydrates to form cementing compounds consisting of calcium silicate hydrate. Silica fume is the condensed vapor by-product of the ferrosilicon smelting process. Silica fume reacts with calcium hydroxide in the presence of water to form cementing compounds consisting of calcium silicate hydrate. High alumina cement, con- sisting mainly of calcium aluminates, has been utilized as a cementitious material because of its rapid set and rapid strength gain characteristics and resistance to acidic environments, sea water, and sulfates. However, owing to certain conditions of temperature and humidity, the cement converts over time to a different hydrate having reduced volume (i.e., increased poros- ity and reduced strength), it is recommended that calcium aluminate cements not be used for structural applications (particularly in wet or humid conditions above 27 �C).62 Selection of the proper water content of concrete is critical, since too much water reduces the concrete strength and insufficient water makes the concrete un- workable. Hardening of concrete occurs as a result of hydration, which is a chemical reaction in which the major compounds in the cement form chemical bonds with water molecules and become hydrates. The hard- ened cement paste consists mainly of calcium silicate hydrates, calcium hydroxide, and lower proportions of calcium sulfoaluminate hydrate either as ettringite or monosulfate. About 20% of the hardened cement paste volume is calcium hydroxide. The pore solution is normally a saturated solution of calcium hydroxide within which high concentrations of potassium and sodium hydroxides are present. Proper curing of the concrete during this stage is essential, as it affects the concrete’s durability, strength, water-tightness, abrasion resistance, volume stability, and resistance to freezing and thawing. Since cement is the most expensive ingredient in concrete, it is desirable to utilize the minimum amount necessary to produce the desired properties and characteristics. Aggregate typically occupies 60–75% of the volume of concrete and therefore its characteristics strongly influence the chemical, physical, and thermal properties of concrete, its mix proportions, and economy. (The balance of the concrete mix generally consists of 10–15% cement, 15–20% water, and air (5–8% if entrained).) Aggre- gates thus are important with respect to the concrete durability. The aggregates come in various shapes, sizes, and material types ranging from fine sand par- ticles to large coarse rocks. Selection of the aggregate material is determined in part by the desired char- acteristics of the concrete. Aggregate materials are available ranging from ultra-lightweight (e.g., ver- miculite and perite) to lightweight (e.g., expanded clay shale or slate-crushed brick), normal weight (e.g., crushed limestone or river gravel), and heavy- weight (e.g., steel or iron shot). Sometimes chemical or mineral admixtures are added during the mixing process to enhance durability (air entrainment), improve workability (enhanced placement and com- paction), modify hardening and setting characteris- tics, aid in curing, reduce heat evolution, or provide other property improvements.63 The concrete typically used in nuclear safety- related structures consists of Type II Portland cement,59 fine aggregates (e.g., sand), water, various minerals, or chemical admixtures for improving properties or performance of the concrete and either normal-weight or heavy-weight coarse aggregate. Concrete 421 American Society of Testing and Materials (ASTM) C 150,64 Type II Portland cement, typically has been used because of its improved sulfate resistance and reduced heat of hydration relative to the general- purpose Type I Portland cement. Both the water and fine and coarse aggregates are normally acquired from local sources and subjected to material charac- terization testing prior to use. Coarse aggregate can consist of gravel, crushed gravel, or crushed stone. Chemical (e.g., air-entraining or water-reducing) or mineral (e.g., fly ash or ground granulated blast fur- nace slag) admixtures have been utilized in many of the mixes to impart improved characteristics or per- formance. For those concrete structures in NPPs that provide primary (biological) radiation shielding, heavy-weight or dense aggregate materials, such as barites, limonites, magnetites, and ilmenites, may have been used to reduce the section thickness and meet attenuation requirements. The constituents are proportioned and mixed to develop Portland cement concrete that has specific properties. Depending on the characteristics of the specific structure, the concrete mix may be adjusted to provide increased strength, higher durability, or better workability for placement. The hardened concrete typically provides the compressive load- carrying capacity for the structure. Specified con- crete unconfined compressive strengths typically have ranged from 13 to 55MPa, with 35MPa being a typical value achieved at 28 days age. Concrete tensile strength is about one-tenth to one-fifth of its compressive strength, so concrete cannot be relied upon to withstand very high tensile stresses. This limitation is overcome by embedding steel reinforcement in the concrete so that the concrete and steel work in concert. In addition to resisting tensile loads, the bonded steel reinforce- ment is used to control the extent andwidth of cracks, especially where it is desirable to reduce member cross-sections. Steel reinforcement is also used in compression members to safeguard against the effects of unanticipated bending moments that could crack or even fail the member. The effectiveness of rein- forced concrete as a structural material depends on the interfacial bond between the steel and concrete so that it acts as a composite material, the passivating effect of the highly alkaline concrete environment to inhibit steel corrosion (see next section), and the similar coefficients of thermal expansion of the con- crete and steel. Most of the mild, or conventional, reinforcing steels used in NPPs to provide primary tensile and shear load resistance/transfer consist of plain carbon steel bar stock with deformations (lugs or protrusions) on the surface. These bars typically conform to ASTM A61565 or A 70666 specifications. The minimum yield strength for the steel reinforce- ment ranges from 280 to 520MPa, with the 420MPa strength material being most common. Post-tensioning is a method of reinforcing (or strengthening) concrete with high-strength tendons to resist tensile loadings and to apply compressive forces to the concrete to provide increased resistance to concrete cracking. A number of NPP concrete containment structures utilize post-tensioned steel tendons that are designed to have (1) consistently high strength and strain at failure, (2) serviceability throughout their lifetime, (3) reliable and safe pre- stressing procedures, and (4) ability to be retensioned and replaced (nongrouted systems). The tendons are installed within preplaced ducts in the containment structure and post-tensioned from one or both ends after the concrete has achieved sufficient strength. After tensioning, the tendons are anchored by button-heads, wedges, or nuts. Corrosion protection is provided by filling the ducts with wax or corrosion- inhibiting grease (unbonded) or portland cement grout (bonded). (Although bonded post-tensioning tendons are less vulnerable to local damage than ungrouted tendons, ungrouted tendons have been primarily used in the United States because the grouted tendon systems cannot be visually inspected, mechanically tested, or retensioned in the event of a larger than anticipated loss of prestressing force.) Supplemental conventional reinforcing is also used to minimize shrinkage or temperature effects and to provide local load-carrying capacity or load transfer. Three major categories of post-tensioning system exist depending on the type of material utilized to fabricate the tendons: wire, strand, or bar that con- form to ASTM specifications A 421,67 A 416,68 and A 722,69 respectively. Minimum tensile strengths range from 1620 to 1725MPa for the A 421 material and 1725 to 1860MPa for the A 416 material. The A 722 material has a minimum tensile strength of 1035MPa. Typical NPP tendon systems group sufficient numbers of wires, strands, or bars to have minimum ultimate strengths ranging from 2000 to 10 000 kN. The trend has been to increase the strength of the tendons to reduce the total number (e.g., in the early 1970s, the typical tendon had a capacity of 3000 kN and since then has progressed to capacities of 10 300 and 15 300 kN).19 With the exception of Robinsion 2 (bar tendons) and Three Mile Island 2 (strand tendons), plants that have post-tensioned 422 Concrete containments utilize unbonded tendons so that the tendons can be inspected and replaced (if necessary). Bellefonte and Ginna each has grouted tendons (rock anchors) to which tendons are attached. Leak tightness of reinforced and post-tensioned concrete containment vessels is provided by a steel liner plate. A typical liner is composed of steel plate stock Concrete 423 stressors. This information also has application to establishing limits on hostile environmental exposure for these structures and to developing inspection and maintenance programs that will prolong com- ponent service life and improve the probability of the component surviving an extreme event such as a loss-of-coolant accident. Prior reviews of research conducted on concrete materials and structures indicate that only limited data are available on the long-term (40–80 years) properties of reinforced concrete materials.26 Where concrete properties have been reported for condi- tions that have been well documented, the results were generally for concretes having ages 40 years. Additional applications of a concrete material sampling activity would be for assessment of construction quality, development of improved damage models, assess- ment and validation of nondestructive testing meth- ods, and evaluation of the performance of repair activities. Reference 28-day value Age (days) NPP data 100 101 102 103 104 105 M-CC(640) Sizewell Torness Hartlepool HEY SHAM I HEY SHAM II WYLFA TROJAN Mix E-2 M-AUX(634.5) M-CC(659) NL-HD ETR-HW ETR-BS RII-SB RII-BS NL-FC ained from the literature and by testing nuclear power plant- 424 Concrete With respect to post-tensioning systems, current examination programs such as ASME Section XI Subsection IWL80 are adequate for determining the condition of the post-tensioning system materials and evaluating the effects of conventional degradation. Isolated incidences of wire failure due to corrosion have occurred. Leakage of tendon sheathing filler (ungrouted tendons) has occurred at a few plants but, except for the potential loss of corrosion protec- tion, the problem appears to be primarily aesthetic.81 Tendon forces at most plants are acceptable by a significant margin, but larger than anticipated loss of force has occurred at a few older plants. The hypothetical effect of reduced prestressing force and degradation of prestressing tendons (e.g., broken wires) has been investigated for a typical PWR post- tensioned concrete containment during a loss-of- coolant accident using finite-element analysis.56 (Results for the scenario investigated indicated that loss of prestressing force leads to increased concrete cracking at lower pressurization levels, complete fail- ure of selected hoop tendons can have a significant impact on the containment ultimate capacity, and failure of selected vertical tendons does not have a significant impact on ultimate capacity.) With the potential use of grouted tendon systems in some of the new reactor designs proposed for construction in the United States, improved guidance on in-service inspection of grouted tendons is desired. Other potential research topics related to post-tensioning systems include development of an improved relation- ship between the end-anchorage forcemeasured by the lift-off test and change in mean force along the tendon length for unbonded tendons, as well as an assessment of the validity of using estimates of time-dependent loss of prestressing force based on limited-duration relaxation tests (e.g., 1000 h) and concrete creep results (e.g., 6months): at a plant 60-year old, this involves application of time factors of 500 and 120, respectively. 4.13.3.3 Assessment and Repair Operating experience has demonstrated that periodic inspection, maintenance, and repair are the essential elements of an overall program to maintain an acceptable level of reliability for structures over their service life. Assessment and management of aging in NPP concrete structures requires a more systematic approach than simple reliance on existing code margins of safety.82 What is required is the integration of structural component function, poten- tial degradation mechanisms, and appropriate control programs into a quantitative evaluation procedure. A methodology for demonstrating the continued reli- able and safe performance of these structures should include (1) identification of structures important to public health and safety; (2) identification of environ- mental stressors, aging mechanisms and their signifi- cance, and likely sites for occurrence; (3) a monitoring- or in-service-inspection-based methodol- ogy that includes criteria for resolution of existing conditions; and (4) a remedial measures program. 4.13.3.3.1 Component selection The most effective structural condition assessment programs are those that focus on the components most important to safety and at risk due to environ- mental stressor effects. Aging assessment methodolo- gies have been developed to provide a logical basis for identifying the critical concrete structural ele- ments and degradation factors that can potentially impact the performance of these structures.83 An evaluation of the impact on plant risk due to structural aging can also be used in the selection of structural components for evaluation.84 Probabilistic risk assessments conducted to date indicate that the structural systems generally play a passive role in miti- gating design basis (or larger) internal initiating events: a notable exception being the pressure-retaining func- tion of the containment following a degraded core incident involving failure of the reactor pressure vessel. The structural components play an essential role in mitigating extreme events initiated by earthquake, wind, and other extreme influences, and their failure probabilities due to external events can be higher. Moreover, failure of major structural components may impact the operation of a number of mechanical and electrical systems and lead to so-called common cause failures. Thus, deterioration of structural com- ponents and systems due to aging and other aggressive environmental influences may be more serious in terms of overall plant risk than might be evident from a cursory examination of their role in accident mitiga- tion. The significance of structural aging and deterio- ration to plant risk can be evaluated by considering the impact that they have on risk associated with external initiating events, especially earthquakes. It is in miti- gating the effects of strong ground motion due to earthquakes that structural systems play a particularly significant role. The apparent impact of structural aging can be investigated using a margins analysis to assess suitability for continued service. Sensitivity analysis can help to identify the structures of impor- tance that should warrant particular attention. Concrete 425 A third approach involves the combination of finite-element analysis and nondestructive testing methods for evaluation of aging and degradation of concrete containments.85 The CONMOD Project objective was to find a practical means to determine the condition of a containment structure as well as how this condition can be expected to change with time under the influence of various loading condi- tions, including aging. Applications of the approach developed include (1) identification of the critical parts of a structure for nondestructive evaluation including critical parameters, (2) updated structural analyses using input from nondestructive evaluations, (3) prediction of nondestructive responses for a known condition at a given time using finite-element method modeling techniques, and (4) prediction of the nondestructive evaluation responses using finite- element modeling techniques based on a known con- dition and how this will change because of aging processes. One of the conclusions of this study was that development of new containment designs should focus on establishing rules, designs, and novel ideas on how to significantly improve the accessibility of the concrete structures for diagnostic investigations. 4.13.3.3.2 In-service inspections In-service inspection programs have the primary goal of ensuring that the NPP structures have sufficient structural margins to continue to perform in a reli- able and safe manner.86,87 A secondary goal is to identify environmental stressors or aging factor effects before they reach sufficient intensity to poten- tially degrade structural components. Routine obser- vation, general visual inspections, leakage-rate tests, and destructive and nondestructive examinations are techniques available to identify areas of NPPs that have experienced degradation. Determination of the existing performance char- acteristics and extent and causes of any observed distress is accomplished through a structural condi- tion assessment that routinely initiates with a general visual inspection to identify suspect areas followed by application of destructive or nondestructive exami- nations to quantify the extent and significance of any observed degradation. Basic components of a condi- tion assessment include (1) a review of ‘as-built’ drawings and other information pertaining to the original design and construction so that information, such as accessibility and position and orientation of embedded steel reinforcing and plates in concrete, is known prior to the site visit; (2) detailed visual exam- ination of structure to document easily obtained information on instances that can result from or lead to structural distress (e.g., crack mapping); (3) determination of the need for additional surveys or application of destructive or nondestructive test- ing methods; (4) analysis of results; and (5) prepara- tion of a report presenting conclusions and recommendations. More detailed information on guidelines on conduct of surveys of existing civil engineering buildings is available.88–91 Some general guidance on assessment of NPP deg- radation is also available.92–95 However, NPP rein- forced concrete structures present special challenges for development of acceptance criteria because of their massive size, limited accessibility in certain areas, sto- chastic nature of past and future loads, randomness in strength, uncertainty in material changes due to aging and possibly degradation, and somewhat qualitative nature of some nondestructive evaluation techniques. Improved guidelines and criteria to aid in the inter- pretation of condition assessment results, including development of probability-based degradation accep- tance limits, are required. (Some information on probability-based crack acceptance limits for beams and shear walls considering loss of steel area and concrete spalling is available.96) 4.13.3.3.3 Nondestructive examinations Application of nondestructive examination methods to NPP reinforced concrete structures presents chal- lenges: wall thicknesses can be in excess of 1m; struc- tures often have increased steel reinforcement density with complex detailing; there can be a number of penetrations or cast-in-place items; accessibility may be limited because of the presence of liners or other components, harsh environments, or structures located below ground; experience with nondestruc- tive examinations of NPP concrete structures is somewhat limited; and methods utilized for the NPP structures are often based on equipment developed for other materials or technologies. Available methods are relatively good at identifying cracking, voids, and delaminations as well as indicating the relative quality of concrete. Methods for determining concrete prop- erties, however, generally are somewhat more quali- tative than quantitative because they tend to be indirect in that they often require the development of correlation curves for relating a measured parame- ter (e.g., ultrasonic velocity or rebound number) to a property (e.g., concrete compressive strength). Infor- mation on identification and description of methods for determining the strength of concrete and evalua- tion of concrete structures is available.97–100 A practical 426 Concrete guide related to nondestructive examination of con- crete, which not only identifies and describes the cap- abilities, limitations, and applications of the various methods that are available but also presents results from a number of examples, has been developed.101 The status of nondestructive examination meth- ods and priorities for its development with respect to examination and instrumentation and monitoring of concrete structures in nuclear plants was addressed by prior NEA/CSNI IAGE workshops.41,45,51 It was noted that although nondestructive examination techniques have been successfully used on a variety of reinforced and post-tensioned concrete structures, there has been somewhat limited experience in their use to evaluate typical NPP safety-related structures. With respect to these structures, three conditions exist where performing inspections or conduct of nondestructive examinations is not straightforward and requires development – inspection of thick- walled, heavily reinforced concrete sections, base- mats or foundations, and inaccessible portions of a containment metallic pressure boundary. Information summarizing the activities conducted addressing these conditions has been documented.102 Noninvasive techniques for characterization, inspection, and monitoring of thick-walled, heavily reinforced concrete sections to provide additional assurances of their continued structural integrity are desirable (e.g., as-built or current structural features determination, flaw detection and characterization, identification of honeycomb areas and embedded items, and location of voids adjacent to the liner). Methods that can be used to inspect the basemat without the requirement for removal of material and techniques that can detect and assess corrosion are of particular interest. Acoustic (e.g., ultrasonic pulse velocity, spectral analysis of surface waves, impact echo, and acoustic tomography), radar, and radiography appear to have potential for application to thick-walled, heavily reinforced concrete struc- tures in NPPs; however, additional development is required. The most commonly used type of founda- tion for both concrete and steel NPP containments is a mat foundation, which is a flat, thick slab support- ing the containment, its interior structures, and any shield building surrounding the containment.103 As such, the concrete foundation elements of NPPs are typically either partially or totally inaccessible for inspection unless adjacent soil, coatings, waterproof materials, or portions of neighboring components or structures are removed. As a result, indirect methods related to environmental qualification are often utilized to indicate the potential for degradation of the NPP concrete foundations.20 This is generally done through an evaluation of the surrounding medium (e.g., air, soil, humidity, groundwater, or cool- ing water). Methods employed are based primarily on chemical evaluations to assess the presence and con- centration of potentially aggressive ions (e.g., sulfates or chlorides). In addition to an assessment of the aggres- siveness of the surrounding environment, the CFR requires a completedescriptionof theeffects of ground- water levels and other hydrodynamic effects on the design bases of the plant foundations and other struc- tures, systems, and components important to safety.104 Inspection of inaccessible portions of metallic pressure boundary components of NPP containments (e.g., fully embedded or inaccessible containment shell or liner portions, the sand pocket region in Mark I and II drywells, and portions of the shell obscured by obstacles such as platforms or floors) requires special attention. Embedded metallic por- tions of the containment pressure boundary may be subjected to corrosion resulting from groundwater permeation through the concrete; a breakdown of the sealant at the concrete–containment shell inter- face that permits entry of corrosive fluids from spills, leakage, or condensation; or in areas adjacent to floors where the gap contains a filler material that can retain fluids. NPP inspections have identified corrosion of the steel containment shell in the dry- well sand cushion region, shell corrosion in ice con- denser plants, corrosion of the torus of the steel containment shell, and concrete containment liner corrosion. Corrosion incidences such as these may challenge the containment integrity and, if through- wall, can provide a leak path to the outside environ- ment. Several techniques have been investigated that exhibit potential for performing inspections of inaccessible portions of NPP metallic pressure boundaries (i.e., ultrasonics, electromagnetic acoustic transducers, half-cell potential measurements, mag- netostrictive sensors, and multimode guide waves).102 However, these techniques tend to be time consum- ing and costly because they tend to examine only a small area at a time. A technique that can be applied remotely to perform global inspections and determines the overall condition of the contain- ment metallic pressure boundary in a cost- and performance-effective manner is desired. 4.13.3.3.4 Remedial methods Deterioration of reinforced concrete generally will result in cracking, spalling, or delamination of the Concrete 427 cover concrete. Whenever damage is detected, cor- rective actions are taken to identify and eliminate the source of the problem, thereby halting the degrada- tion process. The first step in any repair activity is a thorough assessment of the damaged structure or component including evaluation of (1) cause of dete- rioration, (2) extent of deterioration, and (3) effect of deterioration on the functional and performance requirements of the structure or component. From this information, a remedial measures strategy is for- mulated based on the consequence of damage (e.g., effect of degradation on structural safety), time requirements for implementation (e.g., shutdown requirements, immediate or future safety concern), economic aspects (e.g., partial or complete repair), and residual service life requirements (e.g., desired residual service life will influence action taken).105 Basic remedial measures options include (1) no active intervention; (2) more frequent inspections or con- ducting specific studies; (3) if safety margins are presently acceptable, taking action to prevent deteri- oration from getting worse; (4) carrying out repairs to restore deteriorated or damaged part of the structure to a satisfactory condition; and (5) demolishing and rebuilding all or part of the structure. Basic guidance on the repair of degraded structures is available,27,106 and a workshop has been held addressing repair of NPP concrete structures.46 Results of the workshop indicate that improved guidance is required on the assessment of defects (e.g., cracks), and information is desired on the performance and effectiveness of subsequent repairs to concrete structures in NPPs (e.g., durability of repair materials). Information on past performance and current practices for repair materials and systems for general civil engineering structures is being assembled (http://projects.bre.co. uk/conrepnet/pages/contents.htm). 4.13.4 Structural Reliability Theory If properly designed and constructed, the concrete structures in NPPs generally have substantial safety margins; however, additional information for quanti- fying the available margins of degraded structures is desired. In addition, how age-related degradation may affect dynamic properties (e.g., stiffness, frequency, and dampening), structural response, structural resistance/ capacity, failure mode, and location of failure initiation is not well understood. A better knowledge of the effects of aging degradation on structures and passive components is necessary to help ensure that the current licensing basis is maintained under all loading conditions.96 Decisions as to whether to invest in maintenance and rehabilitation of structures, systems, and compo- nents as a condition for continued service and risk mitigation, and the appropriate level of investment, should consider the nature and level of uncertainties in their current condition and in future demands.107,108 Recent advances in structural reliability analysis, uncertainty quantification, and probabilistic risk assessment make it possible to perform such evalua- tions and to devise uniform risk-based criteria by which existing facilities can be evaluated to achieve a desired performance level when subjected to uncertain demands.109,110 Consideration of in situ conditions, redundancy, and uncertainties in important engineer- ing parameters often can lead to significant economic benefits when assessing the condition of an existing structure in a (possibly) degraded condition, and the maintenance or rehabilitation strategies that might be required as a condition for future service. Reliability-based approaches have been applied to the NPP concrete structures111,112 and in evaluation of the prestress level in concrete containments with unbonded tendons.113 Degradation effects can be quantified with fragil- ity curves developed for both undegraded and degraded components.114 Fragility analysis is a tech- nique for assessing, in probabilistic terms in the pres- ence of uncertainties, the capability of an engineered system to withstand a specified event. Fragility mod- eling requires a focus on the behavior of the system as a whole and specifically, on things that can go wrong with the system. The fragility modeling process leads to a median-centered (or likely) estimate of system performance, coupled with an estimate of the varia- bility or uncertainty in performance. The fragility concept has found widespread usage in the nuclear industry, where it has been used in seismic prob- abilistic safety and/or margin assessments of safety- related plant systems.115 The fragility modeling procedures applied to degraded concrete members can be used to assess the effects of degradation on plant risk and can lead to the development of probability-based degradation acceptance limits. This approach has been applied to a limited extent to degraded flexural members and shear walls.96 Additional work is desired in this area for the purpose of refining and applying the time-dependent reliabil- ity methodology for optimizing in-service inspec- tion/maintenance strategies and for developing and evaluating improved quantitative models for http://projects.bre.co.uk/conrepnet/pages/contents.htm http://projects.bre.co.uk/conrepnet/pages/contents.htm 428 Concrete predicting future performance (or failure probability) of a degraded concrete structure, either at present or at some future point in time. 4.13.5 Summary and Potential Research Topics As concrete ages, changes in its properties will occur as a result of continuing microstructural changes (i.e., slow hydration, crystallization of amorphous constitu- ents, and reactions between cement paste and aggre- gates), as well as environmental influences. These changes do not have to be detrimental to the point that concrete will not be able to meet its performance requirements. Concrete, however, can suffer unde- sirable changes with time because of improper spec- ifications, a violation of specifications, or adverse performance of its cement paste matrix or aggregate constituents under either physical or chemical attack. Concrete durability and the relationship between durability and performance, a review of the historical perspective related to concrete and longevity, a description of the basic materials that comprise rein- forced concrete, and information on the environmental factors that can affect the performance of NPP con- crete structures have been provided. Primary aspects related to management of aging of NPP concrete structures have been noted: degradation mechanisms, damage models, and material performance; assessment and repair (i.e., component selection, in-service inspection, nondestructive examinations, and remedial actions); and estimation of performance at present or some future point in time (i.e., application of structural reliability theory to the design and optimization of in- service inspection/maintenance strategies, and deter- mination of the effects of degradation on plant risk). Several areas have been identified where addi- tional research would be of benefit to aging manage- ment of NPP concrete structures: (1) compilation of material property data for long-term performance and trending, evaluation of environmental effects, and assessment and validation of nondestructive eval- uation methods; (2) evaluation of long-term effects of elevated temperature and radiation on concrete behavior; (3) improved damage models and accep- tance criteria for use in assessments of the current condition as well as estimation of the future condition of the structures; (4) improved constitutive models and analytical methods for use in determination of nonlinear structural response (e.g., accident conditions); (5) nonintrusive methods for inspection of thick-walled, heavily reinforced concrete structures and basemats; (6) global inspection methods for metal- lic pressure boundary components (i.e., steel contain- ments and liners of concrete containments) including inaccessible areas and backside of liner; (7) data on application and performance (e.g., durability) of repair materials and techniques; (8) utilization of structural reliability theory incorporating uncertainties to address time-dependent changes to structures to ensure that minimum accepted performance require- ments are exceeded and to estimate on-going compo- nent degradation to estimate end of life; and (9) application of probabilistic modeling of component performance to provide risk-based criteria to evaluate how aging affects structural capacity. 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All rights reserved. 4.14.1 Introduction 434 4.14.1.1 Historical Review of Fracture Toughness Determination for Ferritic Steels 434 4.14.1.2 Theoretical Background Leading to Use of EPFM and Data Distributions for Ferritic Steels 436 4.14.1.3 Master Curve Methodology as Developed by Wallin 438 4.14.1.3.1 History 438 4.14.1.3.2 Basics 439 4.14.1.3.3 Methodology 440 4.14.2 Analysis of Fracture Toughness Test Data for Master Curve Application 442 4.14.2.1 Standard Test and Analysis Procedure (ASTM E1921) 442 4.14.2.1.1 Test specimens used 442 4.14.2.1.2 Determination of reference temperature T0 443 4.14.2.1.3 Data qualification 443 4.14.2.1.4 Determination of lower bound curves 444 4.14.2.1.5 Limits of applicability 444 4.14.2.2 Aspects of Applying Small and Miniature Specimens 445 4.14.2.3 Effect of Constraint 447 4.14.2.4 Effect of Ductile Crack Growth 449 4.14.2.5 Inhomogeneous Materials 450 4.14.3 Application to Integrity Assessments 451 4.14.3.1 Transferability of Test Data 451 4.14.3.2 Accomplished Analyses for Specific RPV Integrity Assessments 455 4.14.3.2.1 PTS test by Framatome 455 4.14.3.2.2 PTS test by ORNL 455 4.14.4 Application to Lifetime Assessment 456 4.14.4.1 Assessment of Irradiation Embrittlement Changes 456 4.14.4.2 Definition of Reference Curves and Their Use 457 4.14.5 On Correlations Between T0 and Other Related Parameters 458 4.14.5.1 Crack Arrest Reference Temperature TKIa 459 4.14.5.2 Dynamic Versus Static Fracture Toughness 459 4.14.5.3 T0 Versus Charpy V-notch Transition Temperatures 460 4.14.6 Summary and Conclusions 460 References 460 Abbreviations bcc Body-centered cubic COD Crack opening displacement C(T) Compact tension specimen DC(T) Disk-shaped compact tension specimen EPFM Elastic–plastic fracture mechanics LEFM Linear-elastic fracture mechanics NDT Nil-ductility transition temperature NPP Nuclear power plant RPV Reactor pressure vessel SE(B) Single-edge bend specimen Symbols a Surface crack depth a0 Original (physical) crack depth/length 433 434 Fracture Toughness Master Curve of bcc Steels b0 Specimen initial ligament (W � a0) B Specimen thickness B0 Reference thickness Bnet Net specimen thickness c Half length of surface crack d Particle diameter d̄ Mean particle diameter dc Critical particle diameter dN Particle diameter scale factor E Modulus of elasticity J J-integral Jc J at fracture instability Je Elastic component of J Jp Plastic component of J K, KI Stress-intensity factor K0 Fracture toughness scaling factor corresponding to 63.2% probability K0Tref K0 corresponding to specific reference temperature K0F Local K0 value at crack tip KIc Linear-elastic fracture toughness KI eff Effective stress-intensity factor KI F Local stress intensity at crack tip KJc Elastic–plastic fracture toughness KJc(0.xx) Fracture toughness for specific probability KJc(limit) KJc capacity limit for specimen KJc(med) Average median KJc of test data Kmin Theoretical minimum fracture toughness M Constraint coefficient m Weibull exponent or material-dependent constant N Number of initiators in volume element P Specimen load Pf Cumulative failure probability Pfr Probability of particle fracture Pmax Maximum load in KIc test PQ Limit load in KIc test Pr{A/I} Conditional probability of arrest Pr{I} Probability of cleavage initiation Pr{I/O} Conditional probability of cleavage initiation Pr{O} Probability of ‘no event’ Pr{P/I} Conditional probability of propagation Pr{V} Probability of void initiation Pr{V/O} Conditional probability of void initiation r Number of valid test data rpl Radius of crack tip plastic zone RTNDT Reference temperature by ASME Code RTT0 Reference temperature by ASME Code Case s Distance dimension along crack front T Temperature T0 Master Curve reference temperature 100 MPa √m T0deep T0 measured for deep crack case T0(margin) T0 including margin for uncertainty T28J 28 J Charpy-V transition temperature T41J 41 J Charpy-V transition temperature Ti Single test temperature TKIa Crack arrest reference temperature 100 MPa √m Tref Reference temperature Tstress T-stress parameter W Specimen width Z Coefficient for selected confidence level b Sample size uncertainty factor gp Particle surface energy di Censoring parameter Da Stable crack growth G Gamma function n Particle size distribution shape factor or Poisson’s ratio s Stress s0 Stress scaling factor spart Particle stress sys Yield stress syy Normal stress V Material constant 4.14.1 Introduction 4.14.1.1 Historical Review of Fracture Toughness Determination for Ferritic Steels Fracture mechanics is an engineering discipline which concerns the behavior of crack-like defects in struc- tures or components and their effect on integrity. Initially conceived by Griffith during World War I, early applications were limited to the study of frac- ture of highly brittle materials (e.g., glass).1 Interest in the discipline languished until World War II, when �25% of the all-welded US Liberty ships experi- enced brittle fracture, exposing the urgent need to understand failure in ferritic structural steels and weldments. The earliest development in fracture mechanics of metals was focused on linear-elastic theory for understanding the fracture behavior of primarily high-strength steels and aluminum alloys. Application to brittle cleavage fracture in structures Ductile fracture Transitional region No statistical size effect Lower shelf J–R curve KJc KIc Temperature Statistical size effect Ductile tearing KJc KJc KJc KJc KIc KIc Figure 1 Schematic description of the fracture toughness transition region and parameters used to characterize fracture toughness in the lower shelf, over the transition region, and in and near the upper shelf where ductile cracking gradually becomes the predominant fracture mode. Fracture Toughness Master Curve of bcc Steels 435 made of welded ferritic steels of low-moderate strength evolved later. One of the basic quantities of fracture mechanics is the stress-intensity factor, K, which is used to describe the loading condition of a cracked structure as a function of crack depth, a, and the applied stress, s. In the simplest case, a wide plate containing a central crack (2a), the loading condition can be expressed using the stress-intensity factor in the form: K ¼ s ffiffiffiffiffipap ½1� Increasing the stress eventually results in a situation where the crack starts to propagate. Depending on the material and loading condition, this can occur by a ductile, cleavage, intergranular, or some mixed- mode mechanism. Once the crack starts to propagate, the critical stress intensity (i.e., the fracture tough- ness) of the material has been exceeded. The fracture toughness is thus a material property, but it may be strongly affected by many environmental factors like temperature and air humidity. Using fatigue pre- cracked test specimens, the fracture toughness of a material can be determined experimentally. This frac- ture toughness quantity can be used to evaluate the integrity of real structures with real or postulated flaws. The fracture mechanism can be different depending on the material, and for ferritic steels, different fracture modes are possible at different tem- peratures due to the ductile–brittle transition. There- fore, different approaches for the characterization of fracture mechanics are needed. The fracture toughness in body-centered cubic (bcc) ferritic steels exhibits a temperature depen- dence characterized by: (1) a low toughness cleavage initiation shelf at low temperatures; (2) an increasing transitional rise in toughness (going from cleavage to a mixture of cleavage and ductile-tearing fracture) with increasing temperature defined arbitrarily as a ductile–brittle transition temperature; and (3) an upper shelf characterized by fully ductile initiation fracture (Figure 1). The lower shelf and the region around the ductile–brittle transition temperature can be characterized using linear-elastic fracture mechanics (LEFM). LEFM considers material defects (flaws and cracks) and the effects of those defects on brittle cleavage crack behavior. LEFM is based on elastic stress analysis of the stress–strain field in the vicinity of the crack tip and a singularity called the stress-intensity factor, K. The linear-elastic theory was soon followed by elastic–plastic fracture mechanics (EPFM), which involved a different type of singular- ity parameter called the J-integral. Determination of material J-integral fracture resistance (J–R) curves expanded the scope of application to also include stable crack growth characterized by ductile tearing. Regardless of which theory is being used, it is neces- sary to know the material resistance to fracture, that is, the fracture toughness of the material being eval- uated. Standardized test methods for determining material fracture toughness properties have been developed. In LEFM, fracture toughness is charac- terized by the parameter KIc; in EPFM, the initiation toughness parameter Jc (often converted to an approximate equivalent K value termed KJc) is used to characterize the onset of unstable crack growth under significant crack-tip plastic deformation con- ditions. (The statistical size effect and the elastic– plastic parameter KJc are associated with the Master Curve methodology discussed in Section 4.14.1.3.3. The linear-elastic parameter KIc is not presently recommended to characterize the transition region, as shown in Figure 1, due to the inherently large scatter of data in this region.) The J–R curve deter- mination and parameters for the onset of stable crack growth are described in separate standards or sec- tions of standards (not discussed here in detail).2 Historically, LEFM concepts for determining the fracture toughness of ferritic, bcc steels have been used, often together with conventional Charpy V-notch impact tests, to characterize the lower shelf and the transition fracture toughness region. There have been few alternatives to the LEFMmeth- odology when combined with Charpy V-notch tran- sition temperature results, as this is the current 436 Fracture Toughness Master Curve of bcc Steels methodology applied for irradiated reactor pressure vessel (RPV) integrity following the ASME Boiler and Pressure Vessel Code. The LEFM approach itself is simple, because only the load record and specimen dimensions are needed for KIc determination, that is, due to the qualification requirements, the test tends to be invalid if there is any significant plastic area under the load versus displacement record. This restriction imposes a major disadvantage in that the amount of test material needed is often large, even if only one large specimen is tested. In this respect, the J-integral EPFM concepts using the parameter JIc are more applicable, as they make possible testing with smaller specimens due to less severe size requirements. Although current KIc and crack opening displace- ment (COD) testing standards better correspond to the latest fracture mechanics understanding (e.g., size restrictions relative to test specimen ligament and thickness dimensions), these standards generally are no longer recommended for characterizing the tran- sition behavior of ferritic steels; they are more appli- cable to cases where the fracture mode is known to be ductile or possibly quasicleavage, and the material shows predominantly elastic behavior. The reason is that these older standards do not account for the statistical nature of the brittle fracture process in ferritic steels. More recently, a statistical assessment methodology, called the Master Curve procedure, has been developed as an improved method for char- acterizing the material fracture toughness (both LEFM and EPFM) of ferritic bcc materials, and for characterizing the temperature dependence of the transition temperature fracture toughness curve. It is the purpose of this chapter to provide a summary review of the Master Curve methodology. The following summary review of the Master Curve fracture toughness approach provides the basis and general framework for the methodology, but it also focuses on some key technical details that are often misunderstood. The reader should consult several of the references for greater detail regarding the various considerations needed in applying the Master Curve methodology for structural integrity assessments. In this chapter, the discussion is first devoted to LEFM involving the standard methods for experimentally determining the value of KIc; then, the more advanced approach based on EPFM and the approximate equiv- alent KJc is reviewed. Finally, the Master Curve proce- dure is discussed in depth and is the primary focus of this chapter. 4.14.1.2 Theoretical Background Leading to Use of EPFM and Data Distributions for Ferritic Steels The fracture toughness of ferritic steels has been characterized by numerous different parameters. It is not the purpose here to discuss history, so the main parameter is linear-elastic KIc and its use with respect to the elastic–plastic KJc. The KIc parameter has, in the past, been one of the most commonly used parameters, also for structural steels, but its limita- tions for describing the transition behavior controlled by both cleavage and ductile cracking are widely recognized (discussed later in detail) today. Many high-alloyed quenched and tempered steels, which exhibit practically no plastic deformation, still have moderate fracture toughness, KIc, and can be used and no additional benefit is achieved by using an elastic–plastic parameter like KJc. For these steels, the measurement of fracture toughness at one or a few temperatures is all that is necessary. For low- alloyed structural steels, which typically exhibit a pronounced ductile–brittle transition and may be loaded in a wide temperature range, the situation is different. In this case, an elastic–plastic parameter is needed. An example of such an application is an irradiated RPV where safety and performance have to be demonstrated in accident and abnormal service conditions which are more severe (i.e., at lower tem- peratures) than in normal operation. The integrity analyses must be based on the material fracture toughness covering the entire transition tempera- ture range. In the mid and lower temperature por- tions of the transition curve, cleavage fracture has to be explicitly considered, and the parameter KJc and Master Curve analysis of data are the preferred methods to determine fracture toughness. It should be noted that application of the methodology is not just limited to ferritic steels, but the fracture is generally cleavage or a stress-controlled type of mechanism. When a cracked material is loaded, a plastic zone will develop at the crack tip. The size of this plastic zone depends on the crack tip loading (stress inten- sity) and the material yield strength. The radius of the plastic zone (rpl) can be expressed for mode I loading (tension loading perpendicular to crack plane) in a simple form as follows: rpl ¼ 1 2p KI sys � �2 ½2� Fracture Toughness Master Curve of bcc Steels 437 where KI is the stress-intensity factor and sys is mate- rial yield strength. The plastic zone size is thus a measure of plasticity at the crack tip and can be used to assess the applica- bility of different fracture toughness parameters. In predominantly elastic cases, the plastic zone size can be very small, and the material may be analyzed using LEFM. The parameter KIc is generally determined without correction term for plasticity, when the plas- tic zone size is small in relation to the specimen dimensions. Otherwise, an elastic–plastic parameter such as the J-integral should be used. For a compact specimen [C(T)] geometry and load- ing, the value of KI can be calculated directly from the load (Pi) and the original crack length (a0/W): KðiÞ ¼ Piffiffiffiffiffiffiffiffiffiffiffiffiffiffi WBBnet p f a0 W � � ½3� where W is specimen width, B is thickness, Bnet is net thickness, and function f (a0/W ) is defined dependent on specimen crack depth. For elastic–plastic conditions where a large plastic zone has developed and some stable crack growth can occur, one normally cannot use the LEFM parameter KIc and eqn [3] without at least some correction term(s). The use of EPFM is more practical to deter- mine directly the J-integral which takes into account the plastic component of the work done to the speci- men or component. The fracture toughness equiva- lence KI, denoted as KJc, can then be converted from the J-integral, which is first divided into elastic and plastic components: Jc ¼ Je þ Jp ½4� where the elastic component of J is calculated from the elastic K (Ke) as follows: Je ¼ ðKeÞ 2ð1� v2Þ E ½5� where E is the elastic modulus and n is Poisson’s ratio. The plastic component of J (Jp) is calculated from the total and elastic work using the measured load versus load line displacement or crack opening data. The value of Jc is converted to KJc as follows: KJc ¼ ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi Jc E 1� n2 r ½6� Despite the limitations of LEFM as discussed earlier, one of the basic standards applied in the past in fracture toughness testing has been American Society for Testing and Materials (ASTM) E 399 which defines the methodology for KIc determination. This standard has recently been revised and issued as part of ASTM E1820-08 (Annex 5),2 but the basic approach is essentially the same as in the previous versions of ASTM E399. Because these standards, as well as a corresponding linear-elastic European stan- dard, BS 7448, are still being used to assess some ferritic steels, it is important to know how they differ from the Master Curve approach used in ASTM E1921-08.3 In particular, the scope of the application requires discussion since it affects how KIc data should be analyzed with respect to KJc. A common feature of LEFM KIc standards (such as the previous ASTM E399) is that they express the fracture toughness as a single value, which should be a material property characterizing the resistance of a material to fracture. The value of KIc should be insensitive to specimen size, if the measured value fulfills the specified size criteria. When these condi- tions are met, the stress–strain condition at the crack tip has been thought to be predominantly plane strain, thus ensuring sufficient constraint to produce a minimum fracture toughness value for the material. The qualification to obtain a valid KIc measurement requires that relatively large test specimens have to be tested. The value of KIc is supposed to represent a lower limiting value of fracture toughness, but the method (ASTM E1820, Annex 5) does not ade- quately cover ferritic steels where fracture by a cleav- age mechanism in the transition or lower shelf region has known statistical characteristics different from ductile initiation fracture toughness.2 Based on the current knowledge of fracture mechanics, confirmed by numerical finite element model analyses of local stress–strain fields, many of the arguments for the LEFM parameter KIc are not valid, especially for ferritic steels.4 First, for crack tip constraint, the KIc size requirement has been shown to be overly conservative, leading to highly oversized specimens. Also, it has been shown that KIc is not a size-insensitive parameter; a size adjustment, similar to that made for Master Curve specimens (discussed later in Section 4.14.1.3), should also be made to the values of KIc to correct for size effects and to make the data comparable with KJc. Justification of this size effect argument has been demonstrated in many com- parisons made between the KJc and the older KIc data measured with different size specimens.4 It is impor- tant to note that increasing the specimen size (both ligament and thickness) gradually diminishes the effect of the size adjustment, which means a reduced, -80 -60 -40 -20 20 40 60 800 0 50 100 150 200 K Jc (M P a Öm ) T-T0 (�C) Master Curve KJc= 30 + 70 exp [0.019 (T-T0)] Figure 2 Definition of Master Curve which is defined as a representative mean fracture toughness curve for most structural bcc steels with moderate strength. 438 Fracture Toughness Master Curve of bcc Steels but still existing, dependence on specimen size even with large dimensions (discussed later in Section 4.14.1.3.3). One should also remember that the Mas- ter Curve size adjustment is valid only in the transi- tion region where cleavage fracture initiation is expected to occur, even after some stable crack exten- sion before cleavage fracture. Due to the size effect, the fracture toughness decreases in the transition region with increasing specimen size whenever cleav- age fracture is encountered. Examples of applying the statistical size adjustment are given in Section 4.14.3. As mentioned previously, ASTME1820 (Annex 5) actually invalidates the KIc determination if the value exhibits transition behavior indicative of some cleav- age fracture. If KIc values characterize only the upper shelf behavior of ferritic steels, no size effect typical of the transition behavior should exist even if later cleav- age initiation occurs. Thus, KIc determinations per- formed as per the previous ASTM standard versions do not necessarily fulfill the requirement or follow the recommendations now in ASTM E1820. This issue is significant, because it concerns an essential qualifica- tion criterion of KIc determination. Another related issue is that ASTM E1820-08 does not specify selec- tion of the test temperature to make sure that there is no cleavage fracture in the test. No posttest measures are required or recommended to confirm that the test data are not affected by cleavage initiation. The ques- tion arises: which of the reported KIc values should be size-adjusted consistently with KJc data and which should not be size corrected? Another important aspect is the 95% secant requirement (Pmax/PQ� 1.1), which limits stable crack growth to very small amounts for typical struc- tural steels exhibiting a rising tearing resistance curve.5 The methodology for KIc determination according to the ASTM E1820 standard is thus not applicable for high-tearing resistance or high-toughness steels, which is also mentioned in Annex 5. The size require- ment for KIc is essential, since it means that the speci- men (ligament) size is the only factor which can be affected in pursuing a valid test result, if the test has to be performed at a specific temperature. If the secant requirement cannot be met, it is possible that no value of KIc can be determined at that temperature. For all of these reasons, fracture toughness testing in the tran- sition region is recommended to be made following standard ASTM E1921, which takes into account the statistical nature of cleavage fracture. Direct fracture toughness determination for the reactor vessel surveillance programs of nuclear power plants (NPPs) was one of the first applications of the Master Curve methodology. Compared to the conven- tional Charpy V-notch methods, the Master Curve concept represents a new approach which makes pos- sible direct fracture toughness determinationwith only a few relatively small specimens, which is a more efficient use of limited test material. Using the Master Curve method allows statistical confidence to be applied to the directly measured data. The traditional practice of estimating fracture toughness from Charpy data (using correlations) ismore unreliable due to large uncertainties associated with the correlations and the subsequent safety margins to meet regulatory require- ments. However, the Charpy-based methods are still in use andwill continue to be used until a large amount of surveillance capsule Master Curve data is available. Existing data from correlations is discussed in Section 4.14.5, and material reference curves are discussed in Section 4.14.4.2. Use of these Charpy-based methods may remain to be the only way of estimating static fracture toughness in some cases, but is not the pre- ferred approach for future surveillance programs. 4.14.1.3 Master Curve Methodology as Developed by Wallin 4.14.1.3.1 History Master Curve methodology is based on the observa- tion that the fracture toughness transition curve for any ferritic steel has the same shape, no matter the steel (see Figure 2). Thus, a single ‘Master Curve’ can be used for all ferritic steels; the curve is simply shifted along the temperature axis to match a mean fracture toughness value, which is established from measured fracture toughness data for the particular steel being evaluated. The primary material fracture Measurement of elastic– plastic KJc data Statistical estimation to determine reference temperature T0 Construction of KJc–T reference curves Figure 3 Master Curve fracture toughness determination according to ASTM E 1921. Fracture Toughness Master Curve of bcc Steels 439 toughness parameter is the transition temperature, T0, which characterizes the Master Curve position, and is defined as the temperature where the median fracture toughness is 100MPa√m (Figure 2). The advantages of Master Curve technology over past methods for estimating the fracture toughness of materials (particularly irradiated materials) are (1) it is based on direct measurements of the property of interest (e.g., fracture toughness); (2) it provides a direct method of establishing the transition curve for irradiated materials (instead of inferring a shift in an assumed baseline bounding curve using Charpy data); and (3) it can be used for materials even with a limited availability of archival materials. The development of Master Curve methodology was started in the 1980s by Wallin and his coworkers by the introduction of a mathematical model to describe the probability of cleavage fracture initiation in a material containing a distribution of potential fracture initiators (flaws). The model was completed by including the temperature dependence of KJc, which was estimated empirically from a dataset including various ferritic structural steels. The scat- ter definition, the size adjustment, and the definition of the temperature dependence are the basic ele- ments of the Master Curve methodology described in ASTM E1921.6–8 The approach has been verified in several round-robin and research programs.9 The first version of the Master Curve standard comprised a procedure for analyzing only single- temperature test data; in later versions, the approach was extended to consider multitemperature test data. The multitemperature approach requires finding a maximum likelihood solution for the value of the transition temperature, T0, from data measured over a range of temperatures, rather than at a single tem- perature. The first version of ASTM E1921 was approved in 1997 and issued in 1998 (ASTM E1921- 97). The multitemperature approach was included in the second revision, after which several other revisions with some minor changes have been released. The present revision, ASTM E1921-08, describes proce- dures for the experimental determination of the elas- tic–plastic fracture toughness, KJc, estimation of the reference temperature, T0, and principles for the lower bound curve definition of fracture toughness (Figure 3). Further detailed descriptions on the meth- odology and applications are given in McCabe et al.10 and Sattari-Far and Wallin11. The model12 has also been validated numerically to more accurately describe the true fracture behavior and the stress–strain distribution of bcc steels on a micromechanical scale.13 The advanced numerical assessment capabilities presently available for multi- scale modeling of materials have made it possible to validate the main elements of the model. Despite several further developments primarily related to material inhomogeneity (in the basic model material macroscopic homogeneity is assumed), the basic approach being applied today is essentially the same as that developed over 20 years ago. No other defi- ciencies or assumptions requiring readjustment have been identified. Further verification of the approach and the validity of the empirically determined tem- perature dependence have been conducted. Some aspects, such as the lower shelf definition (Kmin), are practically impossible to verify only using experimen- tal methods, so that numerical modeling studies have been very instructive. The Master Curve methodology is currently being used in both structural integrity and lifetime assess- ments. Typical areas of application are pressure vessels and piping, nuclear RPV surveillance programs, other energy production structures, off-shore structures, and various welded components and bimetallic joints. 4.14.1.3.2 Basics The applied fracture model12 defines a conditional probability of fracture assuming that the distribution of flaws in the material follows a statistical distribu- tion. The conditional probability term takes into account the possibility of void formation (blunting of a crack initiator) and the crack arrest–propagation event (Figure 4). The fracture event is controlled in the model by assessing the criticality of a single crack initiator from the weakest link principle. The basic elements of the methodology – the scatter definition and the specimen size correction6,7 – are based on this cleavage fracture model. It is assumed that the material has uniform macroscopic properties and Stress applied to material element No initiation Void initiation Cleavage initiation PropagationArrest Pr{O} Pr{A/I} Failure s s = Stress V = Volume Pr{I}, N Cleavage initiator distribution { Pr{I} = Probability of cleavage initiation Pr{V} = Probability of void initiation Pr{O} = Probability of ‘no event’ Pr{I/O} = Conditional probability of cleavage initiation (no prior void initiation) Pr{V/O} = Conditional probability of void initiation (no prior cleavage initiation) Pr{P/I} = Conditional probability of propagation (in the event of cleavage initiation) Pr{A/I} = Conditional probability of arrest (in the event of cleavage initiation) Pr{V/O} Pr{I/O} Pr{P/I} Figure 4 Definition of the conditional probability of cleavage fracture. Reproduced from Wallin, K.; Laukkanen, A. Eng. Fract. Mech. 2008, 75(11), 3367–3377. 440 Fracture Toughness Master Curve of bcc Steels that no global interaction exists between the crack initiators. An overview (basic equations) of the cleav- age fracture model (also called the Wallin, Saario, Törrönen (WST) model) is described next. The conditional probability of cleavage initiation, Pr{I/O}, is expressed as a product of the probability of having a cleavage initiation and that of not having a void initiation as follows: Pr I=Of g ¼ Pr If gð1� PrfV=OgÞ ½7� Equation [7] can be approximated as: Pr I=Of g � Pr If gð1� PrfIgÞ ½8� TheWSTmodel expresses Pr{I/O} as the product of the particle fracture and nonfracture probabilities so as to take into account the previously broken parti- cles which do not contribute to the cleavage process as follows: PrfI=Og � ð1 dc Pfrð1� PfrÞPfdg@d ½9� where the critical particle size dc is defined by a Griffith type expression as follows: dc ¼ 2pE0gp s2yy ½10� where syy is the tensile stress ahead of the crack tip, E 0 is the plane strain modulus of elasticity, and gp is the particle surface energy. The probability of particle fracture is described by a Weibull-type dependence accounting for particle size (d) and particle stress (spart) as follows: Pfr ¼ 1� exp � d dN � �3 spart s0 � �m ! ½11� where dN and s0 are scaling factors. The particle size distribution, P{d}, is given by eqn [12], which describes the size distribution with two parameters, the average particle size (d ) and the shape factor n. For pressure vessel steels, the shape factor appears generally to be in the range 4–6. Pfdg ¼ ðn� 2Þ ðn�1Þ Gðn� 1Þ d d � ��n exp �n� 2 d=d � � ½12� A detailed description of the WSTmodel is given in Wallin et al.12 and Wallin and Laukkanen.13 The recent numerical validation of the model is presented in Wallin and Laukkanen.13 4.14.1.3.3 Methodology The methodology of determining the Master Curve value of T0 is described next and follows that given in ASTM E1921-08.3 ASTM E1921-08 also covers the testing procedure and specimen preparation, which are not described here. The reference temperature, T0, is defined as the temperature at which the mean fracture toughness for a 1-in. (25.4mm) thick fracture toughness specimen equals 100MPa√m. It is the main parameter used to Fracture Toughness Master Curve of bcc Steels 441 define the curves for the mean and lower bound fracture toughness. If the dataset covers several test temperatures, T0 is found as a solution of equations giving the maximum likelihood estimate to this value. Other than the LEFM testing standards for fracture characterization, there is no specified limit for the minimum specimen size. However, the number of specimens to be tested has to be at least six and generally increases when the specimen size is decreased in order to produce a statistically accept- able confidence level for the estimate. The method also includes a censoring procedure that allows the use of adjusted invalid test data, which contain statis- tically useful information. The probability of cleavage fracture initiation is described as a three-parameter Weibull distribution, which defines the relationship between the cumula- tive failure probability (Pf) and the fracture toughness level before or at KI as follows: Pf ðKJc � KIÞ ¼ 1� exp � KI � Kmin K0 � Kmin � �4" # ½13� where Kmin is the theoretical lower bound fracture toughness, set normally to 20MPa√m for steels with yield strength from 275 to 825MPa, and K0 is the scale parameter corresponding to Pf¼ 63.2%. The Weibull exponent is assumed to have a constant value equal to 4 based on theoretical and experimental arguments.7 The measuring capacity (maximum KJc) of a spec- imen depends on its dimensions (ligament size) and the material yield strength as follows: KJcðlimitÞ ¼ ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi Eb0sys Mð1� n2Þ s ½14� where E is the modulus of elasticity, n is Poisson’s coefficient, b0 is the initial specimen ligament length (specimen width minus the initial crack depth, W�a0), M is the constraint value (usually set equal Small specimens No size adjustment Large specimens 95% 95% 5% 5% T K JC Figure 5 Schematic presentation of Master Curve size adjustm to 30), and sys is the material yield strength at the test temperature. Toughness values exceeding this mea- suring capacity require censoring and are set at the level of maximum KJc. The expression for predicting the specimen size effect is based on the cleavage fracture model. The fracture toughness (KJc(x)) corresponding to the desired specimen thickness or crack front length (Bx) is obtained from the values of KJc(0) and B0, respectively, as follows: KJcðxÞ ¼ Kmin þ ðKJcð0Þ � KminÞ B0 Bx � �1=4 ½15� When analyzing measured data according to ASTM E1921, the size adjustment is normally made to 1 in. (Bx¼ 1 in. or 25.4mm). Note that the influence of side grooves on the specimen thickness is ignored. After the size adjustment is made for each KJc mea- surement, the data for different size specimens can be described as one population following the same Master Curve form and scatter as shown in Figure 5. The aforementioned formula applies in the transi- tion range, and it is not necessary to perform the size conversion at temperatures below (T0� 50 �C), because the size effect diminishes (see the fracture toughness curve 50MPa√m shown in Figure 6). The procedure for estimating the maximum like- lihood solution for T0 from data measured at various temperatures was published in 199514 and added to the second and later revisions of ASTM E 1921. The value of T0 is solved iteratively from the fol- lowing equation, which includes a factor d for data censoring: Xn i¼1 di exp 0:019 Ti�T0½ �f g 11þ77exp 0:019 Ti�T0½ �f g � Xn i¼1 ðK JcðiÞ�KminÞ4exp 0:019 Ti�T0½ �f g 11þ77exp 0:019 Ti�T0½ �f gð Þ5 ¼0 ½16� Small specimens Statistical size adjustment Large specimens 95% 5% T K JC ent. 300 250 200 150 100 50 0 350 K Jc ( M P a Öm ) KB = 25= (KJc-Kmin)*(B/25 mm) 0.25+ KminJc 20 40 60 80 100 120 140 160 180 200 B (mm) KJc= 100 MPa Öm KJc= 150 MPa Öm KJc= 200 MPa Öm KJc= 50 MPa Öm B = 25 Figure 6 Effect of the statistical size adjustment at different levels of KJc. -60-50 -40 -30 -20 -10 0 10 20 30 40 50 60 0.00 0.03 0.06 0.09 0.12 0.15 0.18 0.21 0.24 0.27 60 80 100 120 140 160 180 200 220 min. 6 specimens min. 8 specimens T-T0 (�C) n KJc(med) (MPa Öm, B = 25 mm) min. 7 specimens Figure 7 The number of specimens required by ASTM E 1921 for T0 determination (n ¼ 1/number of specimens) and how it depends on temperature and median KJc. 442 Fracture Toughness Master Curve of bcc Steels where KJc(i) is the size-adjusted fracture toughness measured at temperature Ti and di is the censoring coefficient: di ¼ 1 if the KJc(i) datum is valid (less than the limit determined from eqn [14]) or di ¼ 0 if the KJc(i) datum is not valid but may be included as censored data. When the value of T0 has been solved, the fracture toughness curves for specific levels of fracture prob- ability are obtained from the following equations for probability level 0.xx: KJcð0:xxÞ ¼ 20þ ln 1 1� 0:xx � �� �1=4 11þ 77 exp 0:019 T � T0ð Þ½ �f g ½17� The mean curve (50% failure probability) for 25.4mm specimen thickness is obtained as a function of temperature from KJcðmeanÞ1T ¼ 30þ 70 exp 0:019ðT � T0Þ½ � ½18� 4.14.2 Analysis of Fracture Toughness Test Data for Master Curve Application 4.14.2.1 Standard Test and Analysis Procedure (ASTM E1921) 4.14.2.1.1 Test specimens used ASTM E1921 specifies three specimen types for testing: standard compact tension (C(T)), disk- shaped compact tension (DC(T)), and single-edge notched bend (SE(B)). The type shall be selected based on the form and dimensions of the end-product or component (plate, forging, bar, etc.). The C(T) specimen is commonly used, however, in the nuclear power industry, the SE(B) specimen with a square cross-section is preferred since that is a form that can be used directly from the Charpy V-notch test speci- mens contained in most current RPV surveillance programs. The standard test method does not directly specify the minimum specimen size, but only an upper limit for the measuring capacity (KJc(limit)) depending on the size of the ligament (b0¼W�a0) (see eqn [14]). A value exceeding this limit may be included in the analysis as censored data, even if the test did not occur by cleavage fracture initiation. Those values shall be lowered to the limit value and included as censored (with reduced weight as indicated in eqn [16]) data. It should be noted that censored values shall not be included in determining the minimum number of specimens required for the T0 determination (discussed in Section 4.14.2.1.3). When very small-size specimens are used, more specimens need to be tested for a valid analysis. The effect of test temperature (T�T0) or the corresponding median KJc on the required minimum number of valid test values is given in Figure 7, where parameter n is a weight factor (ASTM E1921) representing the inverse of the number of specimens needed for overall validity. When using subsize specimens, the expected number of speci- mens required for a valid analysis is, however, larger. The estimated numbers of specimens needed with different size single-edge notched bend specimens are given in Table 1.15 Two SE(B) specimen configurations are specified: square (W¼ B) or rectangular (W¼ 2B) cross-sections. Table1 Expected number of specimens needed for a valid Master Curve analysis with different size single-edge notched bend specimens Specimen type Expected number of specimens 10�10 7 5� 10 7 5� 5 12 3� 4 28 3.3�3.3 40 Source: Wallin, K.; Planman, T.; Valo, M.; Rintamaa, R. Eng. Fract. Mech. 2001, 68, 1265–1296. 55 45� 0. 5 1 B = 5 W = 10 60� 4.5 b = 5 a 0 = 5 Figure 8 Example of a 5� 10mm single-edge notched bend specimen design for fracture toughness testing according to ASTM E 1921. Fracture Toughness Master Curve of bcc Steels 443 Due to the same ligament size (b0), theW¼ 2B geom- etry has the same measuring capacity as the C(T) specimen with the same thickness. Commonly used configurations are, for example, in RPV surveillance programs, the Charpy-size (10mm square cross- section) and the half Charpy-size (5� 10mm rect- angular cross-section), which also have the same measuring capacity (equal b0). For typical medium strength (quenched and tempered) structural steels, these specimens have been found to give almost identical results.15 Side grooving of specimens is optional (total side-grooved depth may not exceed 0.25B), but recommended to increase stress triaxial- ity near the specimen surfaces.15 The side-grooved half Charpy-size design, with dimensions, is shown in Figure 8. It should be noted that the total speci- men thickness (not net thickness) shall always be used in the Master Curve analysis independent of the side-grooving. Testing with miniature specimens are discussed in Section 4.14.2.2. 4.14.2.1.2 Determination of reference temperature T0 There are two methods of determining the value of T0: the single-temperature method, which is used when the tests have been conducted at a single tem- perature; and the multitemperature method, used if testing is performed at more than one temperature. In the first case, the analysis is simpler and can be performed in analytical form, but the results do not provide further insight on temperature dependence or the lower shelf behavior. The multitemperature analysis can be performed only iteratively, but it gives more comparative information on temperature dependence, scatter, and the lower shelf behavior. As for their effects on T0, both methods give statisti- cally equal confidence levels if the numbers of the valid test data are equal and the measurements have been made close enough to the final T0 (see Figure 7). If sufficient specimens are available, the multitempera- ture method is generally preferred since there is less risk in exceeding the limits of T0 50 �C (discussed next) and the maximum KJc from eqn [14]. 4.14.2.1.3 Data qualification Data qualification is described in detail in ASTM E1921 and is not repeated in detail here. In addition to actions associated with the test procedure and equipment, there are measures which have to be undertaken after each test or test series to ensure that only valid data will be included in the final T0 analysis. The optimum test temperature range for the T0 determination is generally selected iteratively by taking into account the already measured data. For example, three or four tests may be conducted at a selected temperature and then a preliminary T0 is determined from the data; subsequent test tempera- tures are then based on that preliminary T0. As discussed previously, the Master Curve cleav- age fracture model is accurate only in the transition area, where stress state and cleavage crack initiation are the main controlling factors for the fracture event. The data used for T0 determination should therefore be populated in the mid-transition area rather than near the lower or upper shelf. The shape of the transition curve causes the uncertainty of the T0 determination to increase when data measured near the lower shelf are used. There can be an optimum range within the ASTM E1921 validity range of T0 50 �C (Figure 9) depending upon test speci- men size. The resultant data outside of the validity range are excluded from the analysis but still can be compared to the Master Curve determined for the KJc(limit) 95% 5% 50250-25-50 0 50 100 150 200 -75 75 K JC (M P a Öm ) T-T0 (�C) Figure 9 Example of a Master Curve analysis showing the validity window for material with yield strength 455MPa and specimen size 10�10mm SE(B). 444 Fracture Toughness Master Curve of bcc Steels valid temperature range data. The closer the test temperature is to T0, the fewer the number of test specimens that are needed. However, when testing small specimens, the maximum KJc limit is closer to the 100MPa√m level at T0, and the test temperature generally has to be moved to temperatures below T0 with more specimens than the minimum number needed for a valid analysis. 4.14.2.1.4 Determination of lower bound curves A lower bound curve can be constructed to corre- spond to a lower limiting fracture probability, which normally is set to 5% or 2%, taking into account the uncertainty of the T0 determination. The uncertainty of determining T0 depends on the number of speci- mens used to establish T0. The uncertainty (DT0) is defined from a normal distribution with two vari- ables, the test temperature, and the number of speci- mens used for the T0 determination, as follows: DT0 ¼ bffiffiffi r p Z ½19� where b¼ 18–20 �C, depending on the value of T�T0, r is the number of valid test data used to determine T0, and Z is the confidence level (e.g., Z85%¼ 1.44). The median KJc is used to determine the value of b and the uncertainty of T0 according to ASTM E1921. When KJc(med) is equal to or greater than 83MPa√m, b¼ 18 �C. Alternatively, b¼ 20 can be used for all values of KJc(med) not less than the mini- mum of 58MPa√m. The lower and upper tolerance bound (KJc(0.xx)) for the estimated fracture toughness in KJc¼ f (T) is calculated from a revised T0 (T0(margin)) as follows: T0 ðmarginÞ ¼ T0 þ DT0 ½20� KJcð0:xxÞ ¼ 20þ ln 1 1� 0:xx � �� �1=4 11þ 77 exp 0:019 T � T0 ðmarginÞ � � � ½21� where 0.xx is the selected cumulative failure probability. For 1, 2, and 5% cumulative failure probability, the bounds are as follows: KJcð0:01Þ ¼ 23:5þ 24:4 exp 0:019 T � T0ðmarginÞ � ½22� KJcð0:02Þ ¼ 24:1þ 29:0 exp 0:019 T � T0ðmarginÞ � ½23� KJcð0:05Þ ¼ 25:2þ 36:6 exp 0:019 T � T0ðmarginÞ � ½24� When the dataset consists of several test tempera- tures, the median KJc is obtained from the basic rela- tionship (eqn [18]) as follows: K eq JcðmedÞ ¼ 1 r Xr i¼1 30þ 70 exp 0:019 Ti � T0ð Þ½ �f g ½25� where r is the number of valid test data. The 2% lower bound curve is sometimes used as a criterion to determine if the material should be analyzed by taking into account possible material inhomogeneity. This further analysis can be done using the SINTAP procedure (discussed in Section 4.14.2.5), which ensures that a conservative lower bound definition is obtained regardless of possible low fracture toughness values.16,17 If there are values below the 2% curve, the data preferably should be analyzed using the not yet standardized multimodal procedure.17 An example of the 2% curve and the effect of the basic SINTAP analysis are shown in Figure 10. 4.14.2.1.5 Limits of applicability The applied cleavage fracture model strictly applies only to the transition region of the material, although the model also includes a term to take into account the conditional probability of crack propagation. This term is needed because in and near the lower shelf, the fracture event is mainly controlled by crack Fracture Toughness Master Curve of bcc Steels 445 propagation and therefore cannot be correctly described by only the crack initiation term. The situation is complicated by the fact that the lower shelf is not always close to the assumed, theoretically estimated, constant value of 20MPa√m (although it usually is). It is also assumed (Section 4.14.1.3.2) that the crack initiators are randomly distributed and that no global interaction exists between the crack initiators. The material is also assumed to be macro- scopically homogeneous. The method applies for transgranular, cleavage fracture events, although it may be used, with caution, for intergranular-type fracture especially when the fracture event is pre- dominantly stress controlled (as is typical in the lower transition area). Figure 11 shows scanning electron microscope views of both transgranular cleavage and intergranular fracture mode surfaces. The scope of applicability and the limitations of the method are also discussed in Sections 4.14.2.5 and 4.14.3.1 and the application to structures in Section 4.14.3.2. 0 0 50 100 150 200 250 (a) (b 2% 5% 95% T (�C) Cleavage Excessive DCG T0= 69 �C B0= 25 mm 20 40 60 80 100 120 140 M = 30 K JC (M P a Öm ) Figure 10 Example of the 2% lower bound definition for a da same dataset analyzedwith the SINTAPprocedure showing in this 10�10mm single-edge bend specimens (data with excessive d Figure 11 Examples of fracture surfaces of ferritic steels: the and the right one as mostly intergranular mode. 4.14.2.2 Aspects of Applying Small and Miniature Specimens An advantage of the Master Curve method is that it makes possible the use of Charpy-size and even smaller specimens for a valid determination of frac- ture toughness and T0. The number of tests always has to be determined so that the minimum required confidence level is achieved for the estimation. ASTM E1921 describes a special weighting system to ensure a sufficient confidence level for the analysis. The final check can be made only after the value of T0 is rather well known, when the ade- quacy of tests can be determined from the conditionP ni ri 1, where ri is the number of valid data in the valid temperature range i and ni is the weight factor of this range. Normally, the required number can be tentatively determined only after some tests have been conducted, because the optimal test tempera- ture range is not known beforehand. With very small specimens, the final number of tests needed for a ) 0 50 100 150 200 250 M = 30 SINTAP 5% 95% Cleavage Excessive DCG T0= 74 �C B0= 25 mm T (�C) 0 20 40 60 80 100 120 140 K JC (M P a Öm ) taset with DT0 ¼ 9 �C, assuming b ¼ 18 �C (a), and the case no inhomogeneity (b). Material: irradiated A508Cl. 2 steel, uctile crack growth are indicated). left one appearing as pure cleavage (crack initiation shown) Charpy size specimen 55 10 5 10 5 � 10 Charpy specimen 3 � 4 (KLST) specimen 5 � 5 specimen 27 3 4 55 27 5 5 10 Figure 12 Possible single-edge bend specimen geometries for testing using ASTM E 1921; dimensions are in millimeter. 446 Fracture Toughness Master Curve of bcc Steels valid estimate will likely be larger than the given minimum of six because of the smaller validity win- dow, which is reduced with decreasing specimen size and the material yield strength (see Figure 7). Exam- ples of possible small and miniature SE(B) specimen geometries for fracture toughness testing are com- pared in Figure 12. As mentioned previously, a commonly used speci- men configuration for irradiated RPV steels is the full Charpy-size geometry with 10� 10mm cross- section. For most applications, this geometry provides a sufficiently large validity window (Figure 9), because as few as six specimens may be sufficient for a valid estimate. From the present experience for irradiated steels with different size specimens, Charpy square or half Charpy rectangular SE(B) or 0.4T or 0.5T C(T) geometries are generally optimum when the amount or form of the test material is limited. In some cases, even smaller test specimens may be required due to very small amounts of test material. A comparison made between T0 and scatter estimates from test results measured with miniature specimens (i.e., smaller than the Charpy-size) shows that the definitions of scatter and the measuring capacity (specimen size) criterion (eqn [14]) apply even to miniature specimens that are of 3� 4mm and 3.3� 3.3mm cross-section SE(B).15 The results indi- cate no bias between the T0 estimates measured with the miniature specimens compared to the overall mean values of T0, which is shown in Figure 13 for the Charpy-size specimen and three subsize SE(B) specimens. The comparison demonstrates, in all respects, applicability of the miniature size specimens for the fracture toughness estimation using the Mas- ter Curve approach. In some datasets (e.g., on grade 15Kh 2MFA and on the International Atomic Energy Agency (IAEA) reference material JRQ), the smallest specimens (3� 4 and 5� 5mm) produced some low T0 values, which most likely were caused by macroscopic inhomogeneity encountered with such small specimen dimensions. An example of miniature specimen results for the A508 Cl. 3 steel FFA (a French steel grade) is presented in Figure 14, demonstrating nearly consistent fracture toughness data independent of the specimen size. Another benefit of using the Master Curve approach is that the uncertainty associated with the T0 estimation can be determined and taken into account for assessing a conservative, realistic estimate for the lower bound fracture toughness. As discussed previously, ASTM E1921 does not set limits for the specimen size or configuration; however, the mini- mum number of test results, dependent on test tem- perature relative to T0, is predefined to ensure an acceptable minimum confidence level for the esti- mate. If a larger uncertainty for the T0 estimation is accepted, the minimum number of specimens required can be reduced from that given in the stan- dard. Correspondingly, having less than the minimum -150 -100 -50 -50 0 50 100 -100 -150 0 10 � 10 mm ASTM E 1921 ASTM E 1921 5 � 10 mm T0 All (�C) T 0 10 � 10 (� C ) 50 100 -150 -100 -50 -50 0 50 100 -100 -150 0 T0 All (�C) T 0 5 � 10 (� C ) 50 100 -150 -100 -50 -50 0 50 100 -100 -150 0 5� 5 mm 3� 4 mm ASTM E 1921 ASTM E 1921 T0 All (�C) T 0 5 � 5 (� C ) 50 100 -150 -100 -50 -50 0 50 100 -100 -150 0 T0 All (�C) T 0 3 � 4 (� C ) 50 100 Figure 13 Comparison of the values of T0 measured with normal and subsize Charpy specimens relative to the mean of all specimens (10� 10, 5�10, 5�5, and 3� 4mm single-edge bend specimens). Reproduced from Wallin, K.; Planman, T.; Valo, M.; Rintamaa, R. Eng. Fract. Mech. 2001, 68, 1265–1296. Fracture Toughness Master Curve of bcc Steels 447 number of valid data in the test series does not necessarily invalidate the dataset, but it does result in a lower confidence level of the estimate. It is essential that the most realistic confidence level is estimated and taken into account in final integrity assessments. Very small specimens (around 3� 3mm) tend to give slightly (1–3 �C) higher values of T0 compared to 10� 10mm specimens.15 This trend is likely due to the censoring procedure, which screens out pro- portionally more data from the upper tail of the dataset than from the lower tail. This screening affects both the scatter caused by possible material inhomogeneity and that from statistical outliers. The optimal test temperature range for miniature specimens has been proposed to be �50 �C� T –T0��20 �C. Even though more specimens are needed when smaller specimens are tested, the consumption of test material becomes smaller, even if more than the minimum number of specimens were tested for the T0 estimate. In this respect, the 5� 5mm specimen is the least material consuming SE(B) size (�12 specimens are needed for the standard estimate if sys¼ 500MPa).15 Using the 5� 5 or 3� 4mm specimen geometry, it is possible to prepare up to 8 or 12 subsize specimens from the halves of one tested 10� 10mm specimen (see Figure 15). When selecting specimen configuration, it should be noted that with deeply cracked specimens having the liga- ment size equal to or less than the thickness, the ligament is the primary dimension limiting the specimen-measuring capacity, not thickness. Also, with slim (reduced thickness) and very small speci- mens (like the 3� 4mm cross-section), it is recom- mended to side-groove the specimens to increase stress triaxiality near the surfaces. 4.14.2.3 Effect of Constraint Loading a cracked structure creates a local stress concentration ahead of the crack tip. A situation 55 27 3 10 10 4 KLST-size specimenCharpy-size specimen Figure 15 Preparation of 3�4mm miniature specimens from a 10�10mm specimen. 100 50 0 150 200 250 300 350 K JC (M P a Öm ) 10�10 5�10 5�5 3�4 T0= –114 �C B0= 25 mm M = 30 censoring A508 Cl. 3 (FFA) sY= 434 MPa B = 3–10 mm W = 4–10 mm -160 -140 -120 -100 -80 -60 -40 T (�C) 5% 95% Figure 14 Master Curve analysis of steel FFA (A508 Cl. 3) KJc datameasured with Charpy-size and different subsize or miniature specimens (10�10, 5� 10, 5�5, and 3�4mm single-edge bend specimens). Reproduced from Wallin, K.; Planman, T.; Valo, M.; Rintamaa, R. Eng. Fract. Mech. 2001, 68, 1265–1296. 448 Fracture Toughness Master Curve of bcc Steels where brittle fracture initiation is likely only in this restricted area around the maximum stress corre- sponds to small-scale yielding conditions. The high stress area is localized ahead of the crack tip extending to the border of the crack tip plastic zone or three to five times the distance from the crack tip to the stress maximum.18 In such a situation, the stress distribution ahead of the crack can be described correctly with the J-integral. When loading exceeds small-scale yielding, large-scale yielding is involved and the J-integral can no longer describe the crack tip stress distribution correctly. At that point, the measuring capacity of the specimen has been exceeded. The limit for measuring capacity is normally given as a function of the material yield strength and specimen ligament in the form of eqn [14]. The value ofM has been proposed from a variety of finite element calculations to be from 50 up to 200.18 Based on present understanding and examination of many experimental studies, a value of M¼ 30 can be regarded as a realistic estimate for cleavage fracture with most structural steels. Exceeding the load level given by eqn [14] will lead to gradually increasing amounts of ductile tearing before cleavage fracture initiation. The basic analysis procedure described in ASTM E1921 is intended for specimens and structures for which at least a moderate level of constraint (triaxial stress state) is achieved. With deep-cracked speci- mens or thick-wall structures including internal cracks, the constraint is typically high enough for the standard T0 analysis. The situation is different with low-constraint geometries like surface cracks and thin-sheet structures where the small-scale yielding condition may not prevail (note that the limit for this condition also depends on material strength properties). In principle, the basic Master Curve approach can also be applied for such low- constraint conditions, but the estimation may be overly conservative due to high plastic deformation which is outside the applied fracture model assump- tions. In ASTM E1921 test conditions, sufficient constraint is assured by defining valid tests only as those that exhibit brittle fracture initiation below or at the capacity limit value (eqn [14]) and by limiting the ductile crack extension prior to brittle fracture initiation. The probability of cleavage initiation is controlled by the narrow zone ahead of the crack tip where small- scale yielding condition prevails. Several approaches like the Tstress, the Q-parameter, and small-scale yield- ing corrections have been developed to account for the effect of plastic deformation due to low constraint.19 The Tstress analysis is relatively straightforward since a simple elastic stress analysis can be used instead of a numerical large-scale yielding model. Another advan- tage of using the Tstress is that it may be performed assuming no change in temperature dependence which allows the constraint effect to be described only as a shift inT0. Consequently, theTstress yields conservative estimates compared to the more complex Q-parameter approach. On the other hand, the Q-parameter is more accurate for very low-constraint situations. On the basis of a simplified linear-elastic analysis, the correc- tion to T0 due to low-constraint SE(B) geometry (a0/W 0.1) with a negative Tstress has been expressed in the form19: T0 � T0 deep þ Tstress 10MPa �C�1 for Tstress < 0 ðfor SEðBÞ specimens onlyÞ ½26� 20 0 -20 -40 -60 -80 -100 -800 -600 -400 -200 0 200 Tstress (MPa) T 0 -T 0 d ee p (� C ) 290–350 Mpa 490–680 Mpa 720–1380 Mpa Best fit Best fit Best fit T0 ≈ T0 deep+Tstress/10 MPa / �C -1 Tstress< 0 Figure 16 Effect of Tstress on the value of T0 based on tests made with single-edge bend specimens having different crack depths. Reproduced from Wallin, K. Eng. Fract. Mech. 2001, 68(3), 303–328. Fracture Toughness Master Curve of bcc Steels 449 whereT0 is the corrected value andT0 deep is the value measured for a deep crack case. Equation [26] is an empirical result from data consisting of only SE(B) specimens (Figure 16). Additional work has been conducted to refine the expression by including test data from C(T) speci- mens and by a comparison of solutions based on Tstress and the Q-parameter. Thus, a more accurate formula for estimating the T0 from Tstress and the T0 of a deep crack case has been proposed in the form20: T0 � T0 deep þ Tstress 12MPa �C�1 for Tstress 450 Fracture Toughness Master Curve of bcc Steels low-strain hardening capacity or for steels exhibiting low ductile-tearing resistance. An example of analyses corrected for ductile crack growth is shown in Figure 17. The material is a ther- mally embrittled pressure vessel steel (A508 Cl. 3), which therefore has a high T0 (þ69 �C). In this case, the correction lowers the fracture toughness in the upper transition region, but has a negligible effect on the value of T0 and the behavior near the lower shelf. 4.14.2.5 Inhomogeneous Materials In the Master Curve brittle fracture model, it is assumed that the steel is macroscopically homoge- neous. In real wrought steels, the macro- and micro- structures are seldom fully homogeneous due to various effects occurring during production and cool- ing, such as segregation of elements and development of inclusions. The resultant inhomogeneity may man- ifest itself as excessive data scatter. However, in welds or dissimilar metal joints, the material in the asso- ciated heat-affected zone (HAZ) may exhibit strongly distributed ‘inhomogeneity’ so that two or more sub- populations can be distinguished in the test data. Often in the final application of the fracture toughness data, it is only essential to determine a statistically sound conservative estimate for the lower bound frac- ture toughness, which can then be used to construct a conservative reference curve for the material. In most cases, even those where materials show slight inhomogeneity, the standard Master Curve 250 200 150 100 50 0 -100 -50 0 50 100 150 T (�C) Cleavage DCG corrected Ductile T0= 69 �C B0= 25 mm 95% 5% K JC (M P a Öm ) Figure 17 Effect of ductile-tearing correction on a pressure vessel steel (embrittled A508 Cl. 3 with sy ¼ 676MPa, and specimen thickness 10–20mm). The vertical lines indicate the ASTM E 1921 temperature validity area. From Wallin, K.; Planman, T., Eds. In Use and Applications of the Master Curve for Determining Fracture Toughness, Proceedings of the Workshop MASC 2002, Helsinki/Stockholm, June 12–14, 2002; VTT Industrial Systems: Espoo. estimation is sufficient, without further assessments. For testing the data population for possible inhomoge- neity, several approaches have been developed. A com- mon feature for all of these approaches is that they are based on the standard Master Curve, but are extended to properly take into account the inhomoge- neity. The following extended approaches have been proposed for analyzing inhomogeneous materials: 1. Simplified SINTAP procedure for determining a conservative lower bound estimate of fracture toughness when the inhomogeneity is of randomly distributed type.16,20 2. Bimodal or multimodal models for analyzing data- sets showing distributed inhomogeneity (typically welds and HAZ materials).17,20 3. Model for analyzing randomly inhomogeneous materials, including, for example, macroscopic segregations.17,20 Of the three methods, SINTAP is the simplest and the one recommended for initial assessment of the quality of data; it can generally be applied in conservative structural integrity analysis whenmaterial inhomoge- neity is suspected. An advantage of the SINTAP procedure is that it can be performed for a small dataset, unlike the bimodal or multimodal analyses, which require 15 or more data points for a valid assessment. It should be noted that the SINTAP analysis does not produce a statistically representa- tive description of the whole dataset, but its primary purpose is to determine whether the dataset is homogeneous and secondly, to develop a generally conservative lower bound when possible material inhomogeneity is present. The model for analyzing randomly distributed inhomogeneities gives a statis- tically correct description of the fracture toughness data, but requires a large dataset and is more com- plicated to perform than the SINTAP analysis. When there are a sufficient number of fracture toughness data points, the bimodal/multimodal and the randomly distributed inhomogeneous models should be used to better identify a more statistically correct description for a modified Master Curve. The simplified SINTAP procedure is described next. The SINTAP analysis is based on the maximum likelihood method like the basic Master Curve pro- cedure and it consists of three steps as follows. Step 1 is the standard estimation giving the first estimate of T0 and the median fracture toughness according to ASTM E1921. The resulting T0 is used as an input value for Step 2. Fracture Toughness Master Curve of bcc Steels 451 Step 2 is a lower tail maximum likelihood estima- tion and is performed so that all data exceeding the median fracture toughness are censored by substituting d¼ 0 and by reducing the corresponding KJc values to the median curve. If the resulting new T0 (T0–2) is lower than the T0 from Step 1 (T0–1), then T0-SINTAP¼T0–1; oth- erwise, Step 2 shall be repeated using the last T0 estimate (T0–2) as a new input value until a con- stant T0 (T0–2 adj) is obtained. If the number of valid data is ten or more, T0-SINTAP¼T0–2 adj, otherwise Step 3 shall be performed. Step 3 gives the minimum value estimate of T0 and is performed only if the number of valid test data is less than 10. In Step 3, an additional safety factor is incorporated for cases where the number of tests is small. Here, the value of T0 is estimated for each single data point to find the maximum value of T0 (T0max). If the result- ing T0max>T0–2 adjþ 8 �C, T0-SINTAP¼T0max, otherwise T0-SINTAP¼T0–2 adj. The steps and the data censoring in Steps 1 and 2 are described schematically in Figure 18. The flow chart of the iterative procedure is shown in Figure 19. Data us MML es of T0 Censoring P = 50% Specime capacityKJc (MPa Öm) T (�C) Data used for MML estimate of T0 Censoring P = 50% Lower tail MML estimate KJc T (�C) Step 2 St use w n Figure 18 Data censoring in SINTAP Steps 1 and 2 and the T 4.14.3 Application to Integrity Assessments 4.14.3.1 Transferability of Test Data The statistical size adjustment enables one to extend the fracture toughness estimation to specimens and structures with different crack front lengths. A longer crack front means that a larger volume of material is exposed to tensile stress ahead of the crack tip, which increases the probability that a crack exceed- ing the critical size will exist in this volume leading, according to the weakest link theory, to brittle frac- ture. This size effect is addressed in a conservative way by the size adjustment formula (eqn [15]). Generally, the size adjustment is made for all test specimens so that a single reference thickness (nor- mally 1 in.) is used for the Master Curve. As described earlier, when approaching the lower shelf, the size effect diminishes to zero for both EPFM KJc data and LEFM KIc data; note that KIc data are often characterized as being size independent, but as pre- viously described in the transition region, a size cor- rection appears to be applicable. An example is presented in Figure 20, which shows LEFM fracture toughness values (including both KIc and KQ) from a 30(1−ν2) Eb0sysKJc (limit)= ed for timate n measuring limit P = 50% Minimum value estimate KJc T (�C) ep 3 d only hen < 10 Single data used to estimate T0 0 determination in Step 3. Step 1: ASTM E 1921 MML estimation Step 2: Lower tail MML estimation Step 3: Minimum value estimation Repeat Step 2 until constant T0–2 using the last T0–2 in estimation n < 10 No No No Yes Yes Yes T0 max T0= T0–2 T0 max– T0–2> 8 �C T0= T0–1T0–2> T0–1 T0–2 (adj.) T0–2 T0–1 T0= T0 max T0= T0–2 (n = total number of specimens) Figure 19 Flow chart for the SINTAP procedure. −150 −100 150 250 200 150 100 B0 = 25 mm 50 0 100 50 0 −50 0 1 & 2 T 4 & 6 T 1 & 2 & 4 & 6 T T (�C) −150 −100 −50 0 T (�C) A533B Cl. 1 (HSST 02) sY= 480 MPa center A533B Cl. 1 (HSST 02) sY= 480 MPa center K IC (M P a Öm ) 1 T KIC KQ 2 T 4 T 6 T 1 T KIC KQ 2 T 4 T 6 T K IC (M P a Öm ) Figure 20 Fracture toughness data (KIc and KQ) from the heavy-section steel technology program measured with different size specimens showing the data before (left) and after (right) the size adjustment. Material: A533B Cl.1 steel of HSST 02 (sys ¼ 480MPa). Modified from Sattari-Far, I.; Wallin, K. Application of Master Curve Methodology for Structural Integrity Assessments of Nuclear Components; SKI Report 2005:55; Swedish Nuclear Power Inspectorate: Stockholm, 2005; p 178. 452 Fracture Toughness Master Curve of bcc Steels large dataset measured in the heavy-section steel technology (HSST) program with different size spe- cimens ranging from 25.4 up to 152mm (6 in.) thick- ness.23 The data from different specimen sizes follow the same Master Curve prediction after the size adjustment as shown in the second plot in Figure 20. Another example of applying the size adjustment is presented in Figure 21, showing LEFM KIc data measured by MPA (Materialprüfungsanstalt Universität Stuttgart, Germany) with different size specimens, including two very large ones (B¼ 500 mm).24 A distinct size effect is visible in the uncor- rected data (Figure 21, plot on the left), but not in the data size adjusted to the 25.4mm specimen thickness (plot on the right in Figure 21). The above transferability principle also applies for a structure when the size adjustment is made for a real or postulated crack front length, if the material and load conditions remain constant over the whole crack length. If they do not, their variation has to −100 −50 0 −100 −50 0 350 300 250 200 150 95% 5% 100 50 0 50 20 MnMoNi 5 5 (KS 15) sY= 602 MPa B = 25–500 mm 20 MnMoNi 5 5 (KS 15) sY= 602 MPa B = 25–500 mm 200 150 100 50 K IC (M P a √m ) 0 T (�C) T (�C) 50 B = 25 mm B = 50 mm B = 500 mm B = 25 mm B = 50 mm B = 500 mm T0= – 19 �C B0= 25 mm K IC (M P a √m ) Figure 21 Fracture toughness data measured on steel 20MnMoNi55 (KS 15, sys ¼ 602MPa) by MPA without size adjustment (left) and size adjusted to specimen thickness 25mm (right). Modified from Sattari-Far, I.; Wallin, K. Application of Master Curve Methodology for Structural Integrity Assessments of Nuclear Components; SKI Report 2005:55; Swedish Nuclear Power Inspectorate: Stockholm, 2005; p 178. t a c f S Figure 22 Definition of quantities along the surface crack front. Fracture Toughness Master Curve of bcc Steels 453 be taken into account. This situation is typical in a thick-wall structure, where the crack front is not usually straight and the temperature and loading may significantly vary along the crack length. By combining the three-parameter Weibull distri- bution (eqn [13]) and the formula for the statistical size adjustment (eqn [15]), one can derive a formula for the cumulative failure probability of the form: Pf ¼ 1� exp � B B0 KI � Kmin K0 � Kmin � �4( ) ½31� where B is the specimen thickness, B0 is the normaliza- tion thickness (usually 1 in. or 25mm), Kmin is the mini- mum fracture toughness (20MPa√m), KI is the stress- intensity factor, and K0 is the scaling fracture toughness corresponding to 63.2% failure probability. In this formula,KI represents the crack driving force, whereas K0 is specific for the material, and B for the specimen or geometry of the structure or component. In test configurations with simple and small specimen geometry, both KI and K0 can be taken to be constants over the whole crack length with sufficient accuracy. In a real structure, where a real three-dimensional surface crack is possible, both KI and K0 may vary depending on the location (F) along the crack front and should be treated as variables (Figure 22). Note in Figure 22 the designation of the surface crack depth (a), the crack length on the surface (2c), the total surface crack front length (s), the angle f along the crack front, and the wall thickness (t). A more general expression for the cumulative failure proba- bility can thus be expressed in the form25: Pf ¼ 1� exp � ðs 0 KIF � Kmin K0F � Kmin � �4 ds B0 8< : 9= ; ½32� By defining an effective stress-intensity factor (KI eff) corresponding to a specific reference temperature (Tref), which can be the minimum temperature along the crack front, eqns [31] and [32] can be combined to determine the effective stress intensity corresponding to the same failure probability as eqn [32]: KIeffTref ¼ ðs 0 KIF � Kmin K0F � Kmin � �4 ds B0 8< : 9= ; 1=4 K0Tref � Kminð Þ þ Kmin ½33� KIF is obtained from a stress analysis as a function of location (F). K0Tref is the standard, high constraint Tmin-T0 deep (�C) K Ie ff , K IC (M P a Öm ) Kl eff KlC, N-629 -50 0 KIC 5% MC 250 200 150 100 50 0 50 100 150 200 Figure 23 Comparison of Master Curve analysis of real flaws (5% curve) and the ASME Code Case N-629 – curve. Reproduced from Sattari-Far, I.; Wallin, K. Application of Master Curve Methodology for Structural Integrity Assessments of Nuclear Components; SKI Report 2005:55; Swedish Nuclear Power Inspectorate: Stockholm, 2005; p 178. 454 Fracture Toughness Master Curve of bcc Steels Master Curve K0, corresponding to the reference temperature (Tref) along the crack front and has the form for 63.2% failure probability of: K0Tref ¼ 31þ 77 exp 0:019ðTref � T0Þf g ½34� K0F is the local K0 value, based on local temperature and constraint and can be expressed in the form: K0 F ¼ K0T ;Tstress ¼ 31þ 77 exp 0:019 T � T0 deep � Tstress 10 MPa �C�1 � �� � ½35� Equation [35] is directly applicable with the ASME Code Case N-629 fracture toughness reference curve,26 since it is written in terms of the standard deep specimen T0. 11(ASME Code Case N-629 and N-631 allow the determination of RTT0 when T0 is measured, see Section 4.14.4.2) Equations [33]–[35] give the effective crack driving force, normalized to represent a standard Master Curve 25.4mm crack front (B0) and the minimum temperature along the crack front. It should be noted that the area of applicability of the constraint correction based on the Tstress has not yet been fully established.19 The Tstress actually is a LEFM concept that does not work when excessive plasticity is present. In this case, a more advanced concept should be used, such as the Q-parameter or a local approach.27 The Tstress equation (eqn [26] as applied in eqn [35]) works well for the negative values of Tstress, but the effect saturates at higher values. However, in actual components, theTstress is generally negative. The new Tstress equation [27] is expected to be valid up to the Tstress value of 300MPa. 25 The fracture toughness can be expressed either with the 5% lower bound Master Curve, which can be expressed in the form: K5%;Tref ¼ 25:2þ 36:6 exp 0:019 Tref � T0 deep � ½36� or by using the fracture toughness reference curve from ASME Code Case N-629 or N-631.26,28 Details on these Code Cases are presented in Section 4.14.4.2. The following expressions are derived: KIC�ASME;Tref ¼ 36:5þ 11:4 exp 0:036 Tref � T0 deep � ½37� or KIC�ASME;Tref ¼ 36:5þ 3:083 exp 0:036 Tref � RTT0 þ 56 �Cð Þf g ½38� The curves are compared in Figure 23. Note that the fracture toughness curve is not directly com- pared to the crack driving force estimated from stress analysis. Instead, the fracture toughness is compared to an effective driving force, which accounts for the local stress and constraint state and temperature along the crack front, as well as the crack front length. In this way, it is possible to combine the classical fracture mechanics and Master Curve analyses, and to present the comparison in a conventional format. One should remember, how- ever, that postulated flaws often contain unrealisti- cally long crack fronts. An assumed quarter thickness elliptical surface flaw (a/t¼ 1/4; c/a¼ 3), as used in the ASME Code for pressure–temperature operating curves for RPVs, may be used from a conservative deterministic driving force perspective (as was the intention), but from a statistical size adjustment point of view, this assumption is too conservative. If such postulated flaws are analyzed using KI eff, an additional size adjustment to the Master Curve is recommended. Note that s for a 200-mm thick RPV would correspond to an a of 50mm, 2c of 300mm, and an s of about 400mm. A more realistic maxi- mum crack front length (s) is 150mm or less. This value of s is also consistent with much of the original KIc data for the ASME Code KIc curve and therefore justifiable in terms of the functional equivalence principle. The form of KI eff for an excessively large postulated flaw (s 150mm) becomes11: -100 200 150 100 50 K IC (M P a Öm ) 0 -50 0 T (�C) 50 100 5% 95%3T-CT B = 75 mm 4T-CT B = 100 mm TSE 2c = 1000 mm Filled point 1st initiation T0= -35 �C (25 mm) B0= 1000 mm Fracture Toughness Master Curve of bcc Steels 455 KIeffTref ¼ ðs 0 KIF � Kmin K0F � Kmin � �4 150ds B0s 8< : 9= ; 1=4 K0Tref � Kminð Þ þ Kmin ½39� where s is the postulated crack front length. Applied nondestructive evaluation (NDE) techniques and other evidence suggest that a value of s less than 150mm generally can be justified. If warm prestress (WPS) transients need to be analyzed, they can be assessed based on the maximum KI along the crack front and eqns [36]–[38]. Figure 24 Data from the thermal shock test performed by Framatome29 showing the first initiation (filled symbol) and two other initiations (open symbols) together with fracture toughness data (B0 ¼ 1000mm). The Master Curve analysis was performed by Wallin.30 4.14.3.2 Accomplished Analyses for Specific RPV Integrity Assessments A typical pressurized thermal shock (PTS) experi- ment can be conducted by loading a thick-wall pres- sure vessel with an embedded crack by an external load and a severe thermal transient. The aim of a PTS experiment is to load the test vessel so that unstable crack initiation occurs. In connection with these large-scale PTS tests, a material test program is typically performed to produce the necessary mate- rial property data. If the thermal transient includes large temperature changes, any crack initiation after the peak temperature may be affected by the WPS effect. WPS involves the material loaded at the peak temperature during the thermal transient to a stress-intensity level that is higher than the critical fracture toughness at lower temperatures during the transient. Because the material was preloaded to a higher toughness level, the critical fracture toughness at a lower temperature of the transient is higher than without the preloading. The WPS effect is complex and is related to the crack tip stress state and is expected to be most pronounced in short transients where strain ageing will not diminish the effect. 4.14.3.2.1 PTS test by Framatome In the 1980s, Framatome performed a thermal shock pressure test for a thick-wall pressure vessel, includ- ing a very long but shallow crack.29 The crack length (2c) in the 230-mm-thick cylinder was 1000mm and the depth (a) was 17mm. The vessel material (A508 Cl. 3) was characterized by KIc tests conducted with large, 75 and 100-mm thick, C(T) specimens. The aim was to characterize the material directly ahead of the crack front. The data from these characterization tests have been reanalyzed using the Master Curve method.30 The analysis was performed by making the size adjustment of the material test data to the 1000mm crack length, corresponding to that of the vessel. After this correction, the material characteri- zation data and the crack tip load data for the test vessel for the observed first, second, and third initia- tions should fall on the same Master Curve. The comparison in Figure 24 shows that the vessel initia- tions occurred within the 5% and 95% fracture probability bounds estimated for the material, that is, two initiations fell almost on the mean Master Curve and the third close to the 95% probability level. In this case, the result indicates no WPS effect between the initiations, although some effect can be expected due to some warm prestressing during the decreasing temperature of the transient. The C(T) fracture toughness and the PTS test data generally coincide for all three initiations. 4.14.3.2.2 PTS test by ORNL Oak Ridge National Laboratory (ORNL) performed several PTS tests as part of the HSST program. One test (TSE-7) was performed in the 1980s on a pres- sure vessel (steel A508 Cl. 2) with 152-mm wall thick- ness.31 The initial crack shape was semielliptic, but the form changed in the test to something very irreg- ular, which complicated the analysis of this test. The material was characterized by KJc tests made in the transition region with 25.4mm C(T) specimens. In the PTS test, neither KI nor temperature was constant along the crack length. Due to subsequent inaccura- cies in the original KI data, the Master Curve reanal- ysis30 was performed using the apparent maximum 456 Fracture Toughness Master Curve of bcc Steels KI and the local temperature of the crack tip area. The crack length 2c¼ 37mm was used as the effec- tive crack length. Due to the crack shape changes in the test, only the first initiation was included in the Master Curve analysis. The estimated stress intensity in the first initiation is shown together with the fracture toughness data adjusted to a crack length of 37mm in Figure 25. Despite the inaccuracies, the result of the test is also consistent with the measured fracture toughness data. 4.14.4 Application to Lifetime Assessment 4.14.4.1 Assessment of Irradiation Embrittlement Changes The Master Curve methodology was originally developed for applications like RPV surveillance programs, for which the conventional methods are generally less accurate (based on Charpy V-notch energy shifts) or not suitable due to specimen size requirements (LEFM KIc and KIa tests). As a direct measurement approach, the Master Curve approach is preferred over the correlative and indirect meth- ods, based mostly on the Charpy test, used in the past to assess irradiated RPV integrity. It is therefore reasonable to expect that the future determination of plant-operating limits will be based on the Master Curve and related methods rather than on the 0 −100 −50 0 K JC , K IC (M P a Öm ) M = 30 5% 95% 50 T (�C) 50 100 150 200 250 300 350 1T-CT B = 25 mm TSE-7 2c = 37 mm T0= –32 �C (25 mm) B0= 37 mm Figure 25 Data from the heavy-section steel technology (TSE-7) test showing the first initiation (filled symbol) and the fracture toughness results (B0 ¼ 37 mm). Reproduced from Cheverton, R. D.; Ball, D. G.; Bolt, S. E.; Iskander, S. K.; Nanstad, R. K. Pressure vessel fracture studies pertaining to the PWR thermal-shock issue: Experiment TSE-7; NUREG: CR-4304 (ORNL-6177); Oak Ridge National Laboratory: Oak Ridge, TN, 1985, p 133. The Master Curve analysis was performed by Wallin.30 indirect methods or the trend curves based on the chemical composition of materials and their expected neutron fluence. The advantages achievable with the ASTM E1921 methodology especially for RPV applications are as follows: � Direct fracture toughness estimation of the RPV using irradiated small SE(B) or C(T) type speci- mens and determination of a statistically correct mean behavior for ageing assessment and a realistic lower bound curve definition for integrity assess- ments.15,32–34 � Expansion of the analysis to cover issues related to material inhomogeneity and the quality of measured data are possible utilizing the proposed statistical methods. � Expansion of the maximum likelihood estimation to take into account specimen-specific fluence data is possible when needed. � Different data and those measured with different size and type specimens can be included in the analysis (including LEFM KIc with caution). (Note that C(T) and SE(B) specimens may show a (usu- ally about 8 oC) bias due to different geometry so that C(T) specimens yield higher T0.) � Utilization of correlations between the parameters characterizing different loading conditions like crack arrest and dynamic loading. A typical situation especially for older RPVs is that the existing material data consists of very miscella- neous information on the properties of materials, such as test results measured with numerous different material conditions, test standards, specimen types, equipment, etc. In such a situation, a method for handling and analyzing all available data in a syner- gistic way can provide significant savings especially if there are no archive test materials available for addi- tional testing. These aspects of characterization are described schematically in Figure 26. It is important to note that, based on present knowledge, the shift in Charpy transition temperature (e.g., DT41 J) due to neutron irradiation on average is close to or less than the transition temperature shift in fracture toughness (DT0 from the Master Curve method); that is, the shift in Charpy data is generally unconservative in respect of the cor- responding shift in T0. However, there is large scatter in the relationship between these two shifts, and cau- tion is needed when assessing equivalence. Although there are no specific requirements for Master Curve testing in RPV surveillance programs E (T) USE LE (T) KIc KIR KIa Statistically defined fracture toughness of irradiated RPV steels Fluence CVN Neutron dosimetry Instrumented CVN-tests Correlations Reconstitution Prefatiguing Data evaluation using the Master Curve approach ASTM E 1921 test Dynamic KJc test Surveillance specimens 3-PB Prefatiguing ASTM E 1921 test KJc (T) KJc (T) KJd (T) & Figure 26 Characterization of irradiated reactor pressure vessel materials and related parameters. Fracture Toughness Master Curve of bcc Steels 457 in the current Codes and Regulations, the methodol- ogy has been applied in national RPV surveillance programs, and numerous retroactive analyses have been made using data measured in the past in the surveillance and material characterization programs of NPPs. Today, there are also many publications and guidelines which can be used as guidance for using the Master Curve and the related methodologies effec- tively and in a proper manner. Applications for nuclear grade pressure vessel materials and irradiated materials are addressed in the IAEA publication Technical Report Series No. 429.33 Chapter 4.05, Radiation Damage of Reactor Pressure Vessel Steels, provides additional information and mecha- nistic details of ‘Radiation Damage of Reactor Pres- sure Vessel Steels.’ 4.14.4.2 Definition of Reference Curves and Their Use The use of a Master fracture toughness Curve is not a new concept. The ASME Boiler and Pressure Vessel Code, Section III, Appendix G, has used a lower bound static fracture toughness curve that is a ‘Mas- ter Curve’ indexed using the reference temperature RTNDT. 35 Currently, the ASME KIc and KIR curves, indexed to the RTNDT of the material, describe the fracture toughness of the RPV and its lower bound variance with temperature. These curves were adopted in the early 1970s as a lower bound repre- sentation to a relatively small set of linear-elastic fracture toughness (KIc) and linear-elastic arrest toughness (KIa) values for 11 heats of RPV steel. 36 The use of RTNDT to normalize temperature was intended to account for the heat-to-heat differences in fracture toughness transition temperature, thereby collapsing the fracture toughness data onto a single curve. However, RTNDT is not always successful in this regard, often providing a very conservative characterization of fracture toughness. RTNDT is the material/heat-specific ASME Code-defined temperature per Section III, NB-2300,37 based on a combination of drop-weight nil-ductility transition temperature (NDT) and Charpy V-notch tests (>68 J) for nonirradiated materials; or for irradiated materials, RTNDT is the nonirradiated RTNDT (IRT) plus the shift in the 41 J Charpy V-notch temperature to account for irradiation (assumed to be DRTNDT). Appendix A to Section XI of the ASME Code38 uses the same lower bound Master Curve as Section III, 458 Fracture Toughness Master Curve of bcc Steels Appendix G, for crack arrest toughness, and another Master Curve (again indexed using RTNDT) for static initiation fracture toughness. This approach has been used now for over 30 years in the US nuclear industry and has been shown to be very reliable in that there have been no vessel failures, albeit very conservative for most materials. Kim Wallin’s direct fracture toughness Master Curve provides a more complete representation of the material fracture toughness. The Master Curve reference temperatureT0 can be used in an analogous manner as RTNDT to index the position of the Master Curve. The obvious advantages of this approach are: � The index temperature itself is based on measured fracture toughness rather than Charpy V-notch and drop-weight tests. � The Wallin Master Curve has a well-described statistical shape that allows for better-defined direct use in either deterministic or probabilistic analyses. � Direct measurement of irradiated fracture tough- ness eliminates the need to add a shift to an initial value for many applications and it provides, to some extent, the possibility to extrapolate outside the already characterized fluence area. ASME Code Cases N-62926 and N-63128 were pub- lished in 1998 and utilize the ASTM E1921 test method for determining T0. These Code Cases per- mit the use of a Master Curve -based index tempera- ture (RTT0 ¼T0þ 19.4 �C) as an alternative to RTNDT. Code Case N-629 is for Section XI applica- tions for both irradiated and nonirradiated RPV steels; Code Case N-631 is essentially the same Code Case, but it is for Section III design applica- tions for only nonirradiated RPV steels. These Code Cases allow the determination of RTT0 when T0 is measured. Application to RPV integrity requires the knowledge of uncertainties associatedwith the use of a measured RTT0 in place of RTNDT. The use of Master Curve in the United States has been limited to a few examples for which the Nuclear Regulatory Commis- sion (NRC) has written a safety evaluation (SE): � The first use was indirect in that Master Curve data were used to justify a lower value of non- irradiated RTNDT for some Linde 80 welds rather than defining a value of RTT0 39; this application for the Zion RPVs was actually submitted and approved before ASTM E1921-97 or the ASME Code Cases were finalized. � The key approved use of Master Curve was for the Kewaunee RPV. The NRC did not accept the utility submittal approach,40 but modified it to reflect their interpretation of approximating the current Charpy shift-based regulatory approach.41 This interpretation resulted in the use of a deter- ministic Margin term that was larger than that used for the Charpy data application. However, there was still enough beneficial gain using the Master Curve approach for determining an irra- diated RTT0 (over the current regulatory approach) to allow the utility to move forward to replace steam generators and pursue license renewal. � The most recent SE was issued for the plants that have vessels containing Linde 80 welds. The Bab- cock and Wilcox (B&W) fabricated welds were used in all of the B&W design vessels and some Westinghouse design vessels fabricated by B&W. The initial RTNDT for these weld metals has always been uncertain since these weld metals tend to have low upper shelf levels that often fall below 68 J after irradiation. Since 68 J is part of the transition temperature definition for RTNDT, these welds may be unduly affected by the Charpy 68 J temperature requirements. Therefore, the B&WOwners Group developed a program to bet- ter define the initial nonirradiated RTNDT using the Master Curve and the RTT0 approach in Code Cases N-629 and N-631. Their approach utilized Charpy V-notch testing to get the 41 J transition temperature change for assessing the effects of radiation embrittlement in the same manner as currently used for RTNDT. Irradiated surveillance program materials were evaluated using the Mas- ter Curve to compare with the predictive method of initial RTT0 þDT41 J. The methodology was accepted by the NRC but requires explicit margins to be applied.42 4.14.5 On Correlations Between T0 and Other Related Parameters The Master Curve transition temperature T0 for quasistatic loading conditions is statistically precise and measurable following testing standards such as ASTM E1921. Extension of this same type of approach to dynamic loading (KId and KJd) and crack arrest (KIa) situations seems logical, and empir- ical studies with a large variety of structural steels have confirmed a similar relationship. In both cases, the material property (KJd and KIa) can generally be Fracture Toughness Master Curve of bcc Steels 459 described with the same temperature dependence as the quasistatic initiation fracture toughness, but with an increased value of T0. The correlations between T0 and other related toughness parameters are discussed next. 4.14.5.1 Crack Arrest Reference Temperature TKIa For an integrity assessment of real structures, it is often necessary to have information not only on the initiation fracture toughness but also on the crack arrest toughness, KIa. A definition of the reference curve for crack arrest toughness is given in the ASME Code, Appendix A of Section XI. The para- meters describing crack initiation, including the Master Curve T0 and the related parameter RTT0, cannot be used to directly describe the crack arrest toughness. Associated with the development of the Master Curve concept, studies have concluded that it is possible to develop correlations describing the rela- tionship between the crack initiation and arrest toughness.43 These studies have focused on clarify- ing which elements of the Master Curve approach should be modified for assessing crack arrest; and finding possible correlations between initiation and arrest parameters. Due to different mechanisms and differences in factors controlling fracture initiation and arrest events (e.g., the local properties are crucial for crack initiation, but not so critical for crack arrest), the weakest link theory applied in the Master Curve approach is not directly suitable for crack arrest. The analyses of nine well-defined crack arrest datasets, consisting of various pressure vessel base and weld metals (including those used to construct the ASME reference curve), confirm that43: � No statistical size adjustment should be made to crack arrest data. � The scatter seems to be material independent, but lower than the scatter for crack initiation. � KIa data follow the same Master Curve tempera- ture dependence as KJc. Crack arrest data can thus, in general, be described following the same Master Curve approach, using the reference temperature TKIa to characterize the tem- perature corresponding to the crack arrest toughness at a level of 100MPa√m, consistent with the crack initiation transition temperature T0. To clarify if a reasonable correlation exists between T0 and TKIa , a total of 54 datasets, for which both crack initiation and arrest data were available, have been analyzed with the Master Curve concept.43 The result shows an exponentially decreasing trend for TKIa �T0 as a function of T0, but the standard deviation of this correlation is high. Taking into account the observed scatter, a rough estimate for the maximum TKIa can be obtained from the value of T0 at a confidence level of 85% from: TKIa ¼ T0 þ DT þ 19 �C ½40� where DT is obtained from the quasistatic T0 as follows: DT ¼ exp 5� T0 þ 273 136:3 �C � þ sys 683:3MPa � ½41� where sys is the material yield strength. Equation [41] is recommended only for steels where the nickel content is less than 1%. The crack arrest toughness (TKIa ) can also be assessed from instrumented Charpy V-notch data using the correlation developed between TKIa and the temperature corresponding to the crack arrest force of 4 kN. The method has been used for assessing the crack arrest toughness of irradiated RPV steels from the existing Charpy V-notch data. This correla- tion and its application are described in Wallin.43 4.14.5.2 Dynamic Versus Static Fracture Toughness The Master Curve concept, developed originally for quasistatic loading condition, has proven to be useful for dynamic KJd tests conducted at a high loading rate. Assuming that only the value of T0 is rising (along with the material yield strength) with increasing loading rate, dK/dt, the dynamic fracture toughness, KJd and associated T0 should, in principle, be esti- mated from the value of T0 measured by static tests. This dependence was empirically evaluated in 1997 by Wallin using a large dataset consisting of dynamic and static fracture toughness data measured for vari- ous structural steels (yield strength ranging from about 200 to nearly 1000MPa). The Master Curve method was applied to both the static and dynamic loading rates.44 Based on these results, as well as an IAEA round-robin exercise (report to be published in the IAEA Report Series within the framework of the Technical Working Group on Life Management of Nuclear Power Plants), the Master Curve approach 460 Fracture Toughness Master Curve of bcc Steels appears to be fully applicable to dynamic fracture toughness measurements conducted in the ductile- to-brittle transition region. 4.14.5.3 T0 Versus Charpy V-notch Transition Temperatures In the case of dynamic loading with notched speci- mens, the correspondence with the fracture tough- ness test and T0 is complicated due to several uncertainties associated with the Charpy V-notch impact test. First, the loading situation is very differ- ent in the dynamic loading of a notched specimen compared to the quasistatic loading of a fatigue pre- cracked specimen. Due to the differences in the load- ing conditions, the measured Charpy energy includes a significant proportion of both crack initiation and propagation, and often some energy associated with crack arrest; whereas, the quasistatic SE(B) test in the transition region characterizes mainly crack initiation conditions. Additionally, the inherent data scatter and curve fitting required to obtain Charpy V-notch parameters increase the uncertainty of estimating and correlating the transition temperatures. Correlations between the Charpy V-notch tem- peratures T28 J and T41 J versus T0 are presented in Sattari-Far and Wallin.11 The correlations, which are based on data from over 200 pressure vessel steels, are currently being reassessed in more detail, but in applications where the direct estimation of T0 is not possible, the correlations can be used as indicated below (s is standard deviation): T0 ¼ T28J � 19 �C ðs ¼ 22 �CÞ ½42� T0 ¼ T41J � 26 �C ðs ¼ 25 �CÞ ½43� 4.14.6 Summary and Conclusions The Master Curve methodology has been described as an advanced, direct technique of determining the fracture toughness of ferritic structural steels. The application of the methodology has increased during the last decade and spread worldwide extending beyond the initial applications associated with NPP surveillance and integrity assessment programs. Today, the methodology is well known and increas- ingly accepted by safety authorities as a standar- dized method for application in safety assessments. Methods based on conventional approaches, such as the Charpy V-notch test, are still widely used, and probably will be used in parallel into the foreseeable future. Once a sufficient amount of reference data have been measured using the Master Curve method, it will gain even further acceptance. Also, further understanding of the limits of applicability for differ- ent steels will be obtained. Over this transfer period, the correlations developed between the different methods should play a significant role, providing support for properly analyzing data and encouraging the use of the more advanced methods. Note that some correlations, like those proposed for estimating crack arrest toughness from Charpy V-notch tests have brought new applications for the instrumented Charpy test. 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PVRC Recommendations on Toughness Requirements for Ferritic Materials; Welding Research Council Bulletin No. 175; Aug 1972. 37. ASME NB-2331. ASME Boiler and Pressure Vessel Code, Rules for Construction of Nuclear Power Plants, Division 1, Section III, SubSection NB, Class 1 Components, 2002. 38. ASME Boiler and Pressure Vessel Code. Section XI, Appendix G, 2002. 39. Yoon, K. K J. Press. Vess. Technol. 1995, 117, 378–382. 40. Lott, R. G.; Kirk, M. T.; Kim, C. C. Master Curve Strategies for RPV Assessment; WCAP-15075; Nov 1998. 41. Nuclear Regulatory Commission. Safety evaluation by the Office of Nuclear Reactor Regulation to include the use of a Master Curve-based methodology for reactor pressure vessel integrity assessment; Docket No. 50–305; Kewaunee Nuclear Power Plant, May 2001. 42. Nuclear Regulatory Commission. Initial RTNDT of Linde 80 WeldMaterials; Safety Evaluation for Topical Report BAW- 2308, Revision 1; 2005. 43. Wallin, K. InFatigue andFractureMechanics; Chona, R., Ed.; ASTM: Philadelphia, PA, 2001; Vol. 32, 17–34, ASTM STP 1406. 44. Wallin, K. Effect of strain rate on the fracture toughness reference temperature T0 for ferritic steels. Recent Advances in Fracture, Orlando, FL, Feb 10–13, 1997; pp 171–182. 4.15 Ceramic Breeder Materials J. G. van der Laan, A. V. Fedorov, and S. van Til Nuclear Research and Consultancy Group, Petten, The Netherlands J. Reimann Karlsruhe Institute of Technology, Karlsruhe, Germany � 2012 Elsevier Ltd. All rights reserved. 4.15.1 Introduction 465 4.15.1.1 Tritium Breeding 465 4.15.1.2 Breeding Blankets 465 4.15.2 Ceramic Breeder Blankets 466 4.15.2.1 Pellets/Pins/Blocks 466 4.15.2.2 Pebble-Bed Concepts 467 4.15.2.3 Blanket Design Parameters 469 4.15.2.4 Testing of Blanket Modules in ITER 469 4.15.2.5 Ceramic Breeder Requirements 470 4.15.3 Ceramic Breeder Fabrication 470 4.15.3.1 Base Properties 472 4.15.3.2 Fabrication of Shapes 473 4.15.3.2.1 Pellets or blocks 473 4.15.3.2.2 Pebbles 473 4.15.4 Pebble and Pebble-Bed Thermomechanics 475 4.15.4.1 Introduction 475 4.15.4.2 Single Pebble Testing 476 4.15.4.3 Properties of Pebble Beds 476 4.15.4.3.1 Pebble-bed density and packing factor 477 4.15.4.3.2 Mechanical behavior of pebble beds 479 4.15.4.3.3 Thermal creep 479 4.15.4.3.4 Cyclic loading 481 4.15.4.4 Heat Transfer Properties 481 4.15.4.4.1 Thermal conductivity 481 4.15.4.4.2 Heat transfer 483 4.15.4.5 Pebble-Bed Modeling 484 4.15.4.5.1 Continuum models 484 4.15.4.5.2 Discrete-element modeling 485 4.15.4.6 In-Pile Behavior 486 4.15.5 Tritium Production and Release 491 4.15.5.1 Tritium Release 491 4.15.5.1.1 Out-of-pile 492 4.15.5.1.2 In-pile testing 495 4.15.6 Irradiation Parameters 503 4.15.6.1 Irradiation Damage 503 4.15.7 Activation and Waste Issues 504 4.15.8 Summary and Outlook 504 4.15.8.1 Microstructure 505 4.15.8.2 High Burnup 505 4.15.8.3 High Fluence 506 4.15.8.4 High Temperature 506 4.15.8.5 Effects of Transients and Off-Normal Conditions 506 4.15.8.6 Accident Behavior (Safety and Investment Integrity) 506 4.15.8.7 Development of Tools 506 463 464 Ceramic Breeder Materials 4.15.8.8 Compatibility with Structure 506 4.15.8.9 Waste Management and Reuse/Recycling 506 References 507 Abbreviations AECL Atomic Energy of Canada Limited BIT Breeder-in-tube BOT Breeder-out-of-tube BU Lithium Burn-Up CB Ceramic breeder CEA Commissariat à l’Energie Atomique DEM Discrete-element modeling DEMO DEMOnstration reactor for power generation DPC model Drucker Prager Cap model (from soil mechanics) D–T fusion Deuterium–tritium fusion EBR Experimental breeder reactor (US) EOL End of life EXOTIC Extraction of tritium in ceramics FEM Finite-element modeling FFTF Fast Flux Test Facility FPR Fusion power reactor FPY Full power year FW First wall FZK Forschungszentrum Karlsruhe, Germany, changed into KIT 2009 HCCB Helium-cooled ceramic breeder HCPB Helium-cooled pebble bed HFR High Flux Reactor at Petten, the Netherlands INTOR Precursor of ITER project ITER is Latin for ‘‘The Way’’ (before that, ITER was an acronym for International Thermonuclear Experimental Reactor) JAEA Japan Atomic Energy Agency (former JAERI) JAERI Japan Atomic Energy Research Institute (now JAEA) JRR Japan Research Reactor KIT Karlsruhe Institute of Technology, Germany, formerly FZK Li4SiO4 Lithium orthosilicate Li2TiO3 Lithium metatitanate Li2ZrO3 Lithium metazirconate MAPI Mitsubishi Atomic Power Industry NET Next European Torus NRG Nuclear Research & consultancy Group PIE Postirradiation examinations R&D Research & development SBZ Schlünder Bauer Zehner model SCW Super-critical water TBM Test blanket module TBR Tritium breeding ratio TPD Temperature programmed desorption UCT Uniaxial compression test WCCB Water-cooled ceramic breeder WCSB Water-cooled solid breeder Symbols d Diameter d1, d2 Pebble diameters in binary bed D Pebble-bed diameter E Elastic modulus G Tritium production rate h Heat transfer coefficient H Pebble-bed height I Tritium inventory kpb Pebble-bed thermal conductivity ks Thermal conductivity of pebble material Ø Diameter Pp Total pebble porosity¼ average fraction of porosity within a single pebble R Tritium release rate t Time T Temperature Tew Temperature extrapolated from bulk pebble- bed conductivity TM Melting temperature Tw Undisturbed wall temperature e Strain ecr Creep strain g Packing factor¼ ratio of pebble volume to pebble-bed volume rk Contact surface ratio rp Pebble density rpb Pebble-bed density¼ ratio of pebble-bed mass mpb to pebble-bed volume Vpb, also called ‘tap density’ or ‘apparent density’ rs Density of the pebble base material s Stress t Tritium residence time Ceramic Breeder Materials 465 4.15.1 Introduction 4.15.1.1 Tritium Breeding The fusion reaction of tritium and deuterium is con- sidered one of the most suitable options for near- term large-scale fusion power generation, through Dþ T ! 4Heð3:56MeVÞ þ nð14:03MeVÞ Deuterium is a hydrogen isotope with an abundance of 1 out of 6500 atoms in seawater, implying virtually boundless resources. Tritium is the next hydrogen isotope, and it is radioactive with a half-life of 12.3 years under emission of a b-particle; it cannot be obtained from natural resources. Therefore, the D–T fuel cycle requires the breeding of tritium from lithium using one of the following reactions: nþ 6Li ! Tþ 4Heþ 4:78MeV nþ 7Li ! Tþ 4He� 2:47MeVþ n The neutron supplied by the D–T fusion reaction shown above is also the one that provides useful energy. The reaction with 6Li is exothermic, providing a small energy gain; on the other hand, the reaction with 7Li is endothermic but does not consume the neutron, though a more thermalized neutron is released. Natural lithium contains 7.42% 6Li and 92.58% 7Li. In fact, lithium has been identified as the only viable element to breed tritium. 6Li has a very high cross-section to capture a neutron (see Figure 1), and through the use of isotope enrichment, the effective 6Li(n,α)t 7Li(n,α)t X -s ec tio n (b ar n) Energy (eV) 100 10–2 10–1 100 101 102 103 102 104 106 Figure 1 Tritium production nuclear cross-sections for Lithium-6 and Lithium-7 isotopes. Calculated with JANDL- 4.0.(http://www.oecd-nea.org/janis/). 6Li density can be raised from the natural 7.42% to about any desired value. 4.15.1.2 Breeding Blankets In the development of magnetic fusion power plants, the tritium breeding function is effectively integrated in the blanket, which also serves as the main thermal power conversion system and an effective shield for the adjacent reactor components from neutrons and g-rays. An integrated tritium breeding blanket acts as a shield, as a heat exchanger, and as a breeding zone of tritium fuel, as pictured in Figure 2. The blanket has a primary or first wall that faces the plasma, and this component is in direct contact with the edge of the fusion plasma. The latter is typically designed to remove the plasma radiative power and part of the nuclear heating, either with or without a cooling circuit separate from the blanket system. In order to obtain a closed D–T fuel cycle for a fusion power plant, it is mandatory that the tritium production rate is, at least effectively, equal to its consumption rate and accounts for decay and losses at scheduled or unscheduled plant outages; this prin- ciple is usually called ‘tritium self-sufficiency.’ These conditions will not be achieved in near-term fusion devices, where tritium resources available from fis- sion plants can be used and where production from a so-called ‘driver’ blanket is an additional or alterna- tive source for fuel supply. Effective tritium production requires that the lith- ium compounds are located in such a way that the Blanket Shield Vacuum vessel Magnets Breeding zone First wall Coolant Neutrons Radiation DT plasma Figure 2 Schematic of tritium breeding blanket configuration: the breeding zone is between the plasma facing wall and the neutron shield protecting the vessel and magnets. http://www.iter.org/mach/tritiumbreeding 466 Ceramic Breeder Materials maximum capture of D–T neutrons is obtained in the so-called tritium breeding blanket. As most fusion devices require partial use of the plasma- facing area for plasma heating, plasma diagnostics, plasma control, and fuel exhaust, the effective cap- ture of D–T neutrons for breeding tritium requires the use of neutron multipliers in the blanket. Net tritium breeding ratios (TBRs) foreseen for power plants should be about 1.05–1.1. The breeder material used in blanket designs that have attractive thermal efficiency for magnetic con- finement power plants should conform to certain requirements. It should 1. breed tritium in a relatively small volume with a high production rate 2. release tritium in a manner that allows fast proces- sing into plasma fueling 3. possess physical and chemical stability at high temperature 4. display compatibility with adjacent structures and other blanket components 5. exhibit adequate irradiation behavior 6. not pose specific safety risks under off-normal and accidental conditions 7. have activation characteristics allowing recycling or treatment as low-active waste. Lithium-based ceramics are recognized as attrac- tive tritium breeding materials for the first generation of fusion power plants, due to their inherent thermal stability and chemical inertness. This chapter describes the development of ceramic breeder (CB) blankets and material production routes applied or investigated and summarizes the properties and R&D results for a number of lithium-based ceramic materials.1 In this chapter the chemical for- mulas are used, though the actual composition are very often non-stoichiometric, which is more evident when a larger fraction of lithium has been burned. Most of the work presented in this chapter is subject to rapid evolutions in local, national or international programs. The authors like to stress that any of the activities, e.g. those concerning ITER can be quite different in their evolution. 4.15.2 Ceramic Breeder Blankets When ceramic breeder research was initiated in the 1970s, relevant data for lithium-based ceramics were scarce or nonexistent. Initial screening of candidates was mainly based on examination of the physical and chemical characteristics and neutronic behavior. An extensive R&D effort focused on determining the properties of unirradiated materials and on designing irradiation experiments to understand and quantify the effect of neutron irradiation onmaterial properties and on recovery of generated tritium. With the publication of these data, the relative merits of the candidates became known and interest changed accordingly. With the evolution of the INternational TOkamak Reactor (INTOR) and International Thermonuclear Experimental Reactor (ITER) projects, much of the international R&D effort was focused on the opportu- nities for implementing breeding units or driver blan- kets in such a device and its role in the technological development of power reactor blanket systems. Several breeder blanket design options had been developed such as high-temperature water-cooled and helium-cooled concepts for DEMO and power reactors, and low-temperature water-cooled concepts for ITER.2–21 The ceramics under consideration exhibit different characteristics, which can make one ceramic more adaptable to a specific blanket concept. In this chapter, the issues being addressed in R&D in support of current blanket design studies are high- lighted. In this chapter no reference is made to R&D for other fusion systems like inertial confinement. Two major types of ceramic breeder material con- figurations have been developed based on pressed and sintered pins or pellets or as a collection of packed spheres or pebble beds. The actual arrange- ment of pebbles though may be in tube, which may lead to some confusion when ‘breeder-in-tube’ (BIT) is mentioned. The paper by Ihli et al.21 provides an overview of blanket design developments and refer- ences ongoing work by various parties. Until 2010, pebble-bed concepts were the preferred options for all parties involved in ceramic breeder test blanketmodule (TBM) programs for ITERas reported by Giancarli and coworkers.23,24 Typically, inert gas with hydrogen addition, whose characteristics are dis- cussed later in this chapter, is used for the extraction of the tritium produced from lithium ceramics.25–28 4.15.2.1 Pellets/Pins/Blocks Early developments of ceramic breeder blanket con- cepts were based on fission reactor technologies, and the R&D activities concentrated on pellet or pin type specimens. An example is the European BIT project, using machined annular pellets or pins.2–5 Figure 3 shows a schematic view of the BIT concept, based on stacking annular pellets of Central dummy rodSpacer Purge gas channel Breeder ceramic annular pellet Beryllium Flow separator (baffle) Beryllium cladding Pressure vessel Figure 3 Breeder inside tube (BIT) concept based on lithium ceramic pellets. Reproduced from Dalle Donne, M.; Anzidei, L. Fusion Eng. Des. 1995, 27, 319–336. Purge gas Be/Li4SiO4 pebble bed breeder zone Coolant systems Diffusion-welded first wall Stiffening plate Inlet Inlet Outlet Figure 4 Breeder-out-of-tube (BOT) with a mixed bed of Li4SiO4 and beryllium pebbles. Reproduced from Dalle Donne, M.; Anzidei, L. Fusion Eng. Des. 1995, 27, 319–336. Ceramic Breeder Materials 467 LiAlO2, enriched up to 90% 6Li, operating at tem- peratures between 400 and 600 �C.5 Li2ZrO3 was considered as an alternative material, which has been further developed at the Commissariat à l’Energie Atomique (CEA).29 Other breeding concepts using machined shapes have been developed, for example, in the Russian Federation,30 and recently Sharafat et al.31 proposed shaping the breeder ceramic as foams. In this chapter, only a few R&D results with pel- lets or pins are mentioned. 0 1 2 Figure 5 Mixed-bed test configuration of Li4SiO4 pebbles with small and large beryllium pebbles packed to high density used in EXOTIC-7 irradiation experiment. Reproduced from van der Laan, J. G.; Kwast, H.; Stijkel, M.; et al. J. Nucl. Mater. 1996, 233–237, 1446–1451. 4.15.2.2 Pebble-Bed Concepts In the 1970s, alternative fission fuel technologies had been developed based on packing of spheres to reduce the problems associated with excessive swelling and fragmentation of pellets.32 One of the early pebble-bed blanket designs was developed by Dalle Donne and coworkers6,7,9 at Forschungszentrum Karlsruhe (FZK), now called Karlsruhe Institute of Technology (KIT), Germany. In this concept, breeder ceramic and neutron multiplier were both shaped as small spheres or pebbles and arranged in a so-called mixed bed (see Figure 4). The concept was based on small (0.1–0.2mm diam- eter) pebbles of Li4SiO4 and a binary mixture of beryllium pebbles (0.1–0.2mm and 1.5–2.3mm diam- eter) (see Figure 5), taken from the extraction of tritium in ceramics 7 (EXOTIC-7) irradiation proj- ect.33 It was found that the compatibility of Li4SiO4 and beryllium was drastically reduced under neutron- irradiation conditions.34 This initiated the separation of breeder and neutron multiplier in different pebble beds in further blanket design evolution. Extensive R&D on breeder pebbles has also been performed by the Canadian Atomic Energy of Canada Limited (AECL), where in particular Li2ZrO3 and Li2TiO3 were developed. 9,35 With growing insight into thermodynamics and the experimental results obtained from neutron- irradiation testing, the European breeder out of tube (BOT) concept evolved into the helium-cooled pebble-bed (HCPB) concept,7 in view of preparing a test module program for ITER. This concept evolved further in Europe within the scope of the Power Plant Conceptual Study.15,16 The key features of this early HCPB concept are given in Figure 6. It consists of First wall Breedin g zone Berylliu m pebb le bed BZ HTS LTS Man ifoldC eramic breede r pebb le bed 40 m m 10 m m Figure 6 The Helium Cooled Pebble Bed (HCPB) blanket concept from the European Power Plant Conceptual Study (PPCS), model B. Reproduced from EFDA, A. Conceptual Study of Commercial Fusion Power Plants; Final Report of the European Fusion Power Plant Conceptual Study (PPCS); Report EFDA-RP-RE-5.0; 2005. To FW From caps and grid To caps and grid Bypass He inlet From FW MF 2 Be Ceramic breeder Purge gas He coolant MF 3 MF 4 Manifold (MF) 1 Figure 7 Evolved helium-cooled pebble-bed concept. Reproduced from Poitevin. Y; et al. Fusion Eng. Des. 2010, 85, 2340–2347. 468 Ceramic Breeder Materials alternating beds of ceramic breeder and beryllium pebbles, perpendicular to the plasma-facing wall, between flat coolant plates of Eurofer-97, a so-called reduced activation steel based on conventional 9Cr steels.36,37 In this study, the blanket box is considered a con- sumable component, with (1) the maximum irradia- tion damage of primary wall structures set at 150 dpa (about 5 FPY (full power year)) and probably the limiting factor of the box lifetime; and (2) the burnup of the ceramic breeder and swelling of the beryllium neutron multiplier depending on the design.38 Further evolution of the HCPB line in Europe concentrated on the strategies for the ITER TBM, as explained by Poitevin and coworkers.17,39–41 The internal structure of the blanket is given in Figure 7. All structures contain dense patterns of cooling channels, with beds of Be and ceramic breeder in the form of near-spherical particles (Ø 0.25–0.63mm for Li4SiO4, Ø 1mm for alternative breeder Li2TiO3, and Ø 1mm for beryllium) separated by cooled steel plates and bed heights sufficiently low (about 10mm) to conduct heat to the cooling plates without exceeding material temperature limits. Tritium is removed from the pebble beds by a slow purge flow of helium at near- atmospheric pressure, with hydrogen (typically 0.1 vol%) and defined levels of other constituents such as H2O, and so on for optimized integral performance. Some alternative concepts were explored such as those using a 9-Cr steel variant with higher First wall Module box (F82H) beryllide Neutron multiplier pebble (Be, beryllide) Tritium breeder pebble (Li2TiO3) Cooling panel ~500mm Figure 8 Schematic of water-cooled ceramic breeder concept developed in Japan. Reproduced from Akiba, M.; Enoeda, M.; Tsuru, D.; et al. Fusion Eng. Des. 2009, 84, 329–332. Beryllide He out He turn 2 He in pol. rad. tor. Ceramics Tube with thermal insulation Figure 9 Breeder inside tube concept with ceramic pebbles. Reproduced from Ihli, T.; Basu, T. K.; Giancarli, L. M.; et al. Fusion Eng. Des. 2008, 83, 912–919. Ceramic Breeder Materials 469 temperature resistance, see Hermsmeyer et al.,18 or SiC-based composite structure.21 The Japanese designed water-cooled solid breeder blanket consists of two submodules (see Figure 8).20 One submodule consists of a module box, tritium breeding pebbles, neutron multiplier pebbles, and cooling panels. The tritium breeding pebbles and neutron multiplier pebbles are separated by the cool- ing panels, and the beds are oriented parallel to the plasma-facing wall. The module box and the cooling panels are made of reduced activation ferritic– martensitic steel, F82H.36,42 For the tritium breeding pebbles, Li2TiO3 is selected as the primary candidate material. These pebbles are about 0.2–2mm in diam- eter, with either a monodisperse or binary size distribution. Among the pebble-bed concepts, there are also designs for a low-pressure water-cooled ITER driver blanket (see, e.g., Lorenzetto et al.10) and the ITER 1998 design document by Ioki and coworkers.11,13 Nardi et al.13 developed ideas for a driver blanket for the reduced size ITER-FEAT. Recent work reported by Ihli et al.21 also included mixed bed options for DEMO blankets, as shown in Figure 9. 4.15.2.3 Blanket Design Parameters Table 1 provides a non-exhaustive summary of ceramic breeder blanket designs, and key parameters, compiled from the references in this chapter. Most of the concepts have a ferritic–martensitic steel as the structural material. Typical values for the expected blanket neutron wall load are about 3MWm�2, and its lifetime is mostly considered to be limited by about 150 dpa for the structural material. For a DEMO reactor, which is not intended to be a power plant with high availability, typical lithium burnups are 11% (Li4SiO4) or 17% (Li2TiO3), and fast neutron damage in the reduced activation ferritic/ martensitic (RAFM) steels is about 70 dpa.47,48 A typical lifetime for a power reactor blanket is esti- mated to be of the order of 5 years, implying about ten replacements during a 60-year reactor life. The blanket structure should, therefore, be designed with low-activation materials, enabling it to be recycled typically in a period up to 100 years. Such activation requirements have led to Li4SiO4 and Li2TiO3 as the preferred ceramic breeder systems for the European HCPB concept. A case study for Li4SiO4 has been elaborated by Fischer and Tsige-Tamirat.49 4.15.2.4 Testing of Blanket Modules in ITER Breeding blanket and associated systems in fusion power plants have to ensure tritium breeding self- sufficiency, show a sufficient power conversion effi- ciency, and withstand high neutron fluences.50 TBMs for ITER should be representative for such blanket modules (see Giancarli et al.23 and Chuyanov et al.24). 470 Ceramic Breeder Materials Among the technical objectives of ITER, it is specifi- cally stated that ‘‘ITER should test tritium breeding module concepts that would lead in a future reactor to tritium self-sufficiency and to the extraction of high-grade heat and electricity production.’’51 The main testing objectives shall be � Validation of structural integrity theoretical pre- dictions under combined and relevant thermal, mechanical, and electromagnetic loads � Validation of tritium breeding predictions � Validation of tritium recovery process efficiency and T-inventories in blanket materials � Validation of thermal predictions for strongly heterogeneous breeding blanket concepts with volumetric heat sources � Demonstration of the integral performance of the blanket systems. Table 1 also provides values for the ITER TBM loading parameters and tentative requirements for DEMO as a next step. It is seen that the neutron wall load in ITER is relatively small, which requires specific measures to make meaningful nuclear tests with TBMs. Four versions of a TBM are considered with specific objectives as follows: 1. H-phase and H–He-phase: focus on electromag- netic behavior; 2. D-phase: focus on thermal and neutronic behavior; 3. First D–T-phase: DEMO-relevant data acquisi- tion on neutronics, tritium production and man- agement, and thermomechanics; 4. Second D–T-phase: DEMO-relevant data acquisi- tionwithhighduty, longpulseswithan integralTBM. All ITER parties consider helium-cooled ceramic breeder (HCCB) blankets. This type of blanket requires beryllium as the neutron multiplier and ferritic–martensitic steel as the structural material. The ceramic breeder is a Li-based compound, either Li2TiO3 or Li4SiO4, and is used in pebble beds. A water-cooled ceramic breeder (WCCB) blanket is proposed for the Japanese TBM. ITER prepares three horizontal ports for TBM testing, and two TBMs can be installed in one port. Though TBMs will be installed and tested as part of the ITER activities, each TBM will be developed under the responsibility of the distinct ITER party. Four of the ITER parties, China, European Union (EU), Japan, and Russian Federation (RF), have made TBM design proposals with solid breeder materials, while the United States and Korea propose to test submodules integrated into one of them. All ceramic breeder-based TBMs use pebble beds and ferritic– martensitic steel structures and He coolant at 8MPa with inlet temperature of 300 �C and outlet up to 500 �C depending on the operating conditions. Only the Japanese party proposes a water-cooled concept in addition. Figure 10 shows the typical arrangement of a port cell in ITER to accommodate TBMs. 4.15.2.5 Ceramic Breeder Requirements The development of any ceramic breeder blanket concept toward demonstration and realization of fusion power must ensure that the ceramic breeder material meets the following specific requirements: � Though the ceramic breeder has no structural function in the blanket, the pebbles or pellets must withstand the stresses induced under reactor operating conditions (pressures, temperatures, temperature gradients and thermal shocks, irradiation-induced swelling, creep) without an excessive fragmentation, which might result in degradation of the heat transfer parameters and purge gas flow, up to end of life (EOL) peak burnup and displacement damage. � Stability of the ceramic at the maximum operating temperature with regard to lithium transport (e.g., by evaporation or redistribution). � Compatibility between the ceramic and the structural material in the reference purge gas con- ditions under neutron irradiation. Compatibility is one of the criteria defining the maximum inter- face temperature between ceramic breeder and the structural material. � Sufficiently low tritium residence time to mini- mize the tritium inventory in blanket and auxiliary systems that determine source term in off-normal and accidental conditions. � Activation as low as possible under neutron irradi- ation, including activation from impurities, so as to reduce the D–T fuel cycle back-end issues (includes the materials’ recycling aspects). 4.15.3 Ceramic Breeder Fabrication Because significant quantities of ceramics will be needed in the near future for the fabrication of ITERTBMs and for a potential ITER driver blanket, various efforts have been initiated to evaluate fabri- cation process development. One of the fabrication issues is the hygroscopic nature of several candidate Table 1 Some design and loading parameters for a number of fusion reactor blanket concepts utilizing ceramic breeder materials. For comparison values for the EU ITER Test Blanket Modules are given also. In addition to the references mentioned in the table, reader can turn to references in the text like e.g. 16, 21, 24 and 38. Design BIT BOT HCPB WCSB SSTR A-SSTR2 PPCS-B (HCPB) HCPB-TBM in ITER EU EU EU Japan Japan Japan EU EU References 6,7,8 2,3,4,8 15 20 217 218 16 17–19 Ceramic breeder (CB) LiAlO2 Li4SiO4 Li4SiO4 or Li2TiO3 Li2TiO3 or other Li2O Li2TiO3 Li4SiO4 Li4SiO4 Li-6 enrichment % 90 25 30% (Li4SiO4), 60% (Li2TiO3) 90 30-40 90 alternative ceramic breeder (Li2ZrO3; Li2TiO3) Li2TiO3 Li2TiO3 Li2TiO3 Shape Pellet Pebble Bed Pebble Bed Pebble Bed Pebble Bed Pebble Bed Pebble Bed Pebble Bed Coolant water He He water H2O or SCW He He He Neutron multiplier Be-pellets Be-pebbles Be-pebbles Be-pebbles, or Be12Ti Be-pebbles Be-pebbles Be-pebbles Be-pebbles Minimum ceramic breeder operation temperature C 410 300 400 400 450 700 400 250 Maximum ceramic breeder operation temperature C 590 660 890 900 750 800 920 900 Structural material 316 SS MANET Eurofer F82H F82H (RAFS) SiCf/SiC Eurofer Eurofer Surface heat flux MW/m2 0.5 1 1 10 10 7–8 0.1 C e ra m ic B re e d e r M a te ria ls 4 7 1 Cryostat plug Vacuum vessel port extension FW TBM frame and shield plug Vacuum vessel plug Cryostat extension Drain pipe Breeder concentric pipe Bio-shield plug Transporter Figure 10 View of a typical test blanket module port cell arrangement in ITER. Reproduced from http://www.iter.org/mach/ tritiumbreeding. 472 Ceramic Breeder Materials lithium ceramics. Sensitivity to moisture increases as the lithium oxide content increases and as the specific surface area increases. The research activity initially involved g-LiAlO2, Li2O, Li2SiO3, and Li2ZrO3; see, for example, Johnson et al.25 and Roux et al.52 Later work concerned Li4SiO4, Li8ZrO6, and Li2TiO3. 26,33,53–55,185,189 Cur- rently, most blanket concepts are based on Li4SiO4 or Li2TiO3, though recently, work on other systems such as Li3TaO4 56 and Li8PbO6 57 as well as compo- sites of Li2TiO3 with Li2O or Li4TiO4 additives 58 has been reported.57 Breeder development has also started in Korea and India.59–61 4.15.3.1 Base Properties Table 2 summarizes lithium compounds considered as breeding material and lists some key properties. Table 2 Overview of most relevant basic ceramic lithium co Material Lithium- Composition Li density (g cc�1) Theoretical density (g cc�1) Meltin temp (�C) Oxide Li2O 0.93 2.01 1432 Octo-zirconate Li8ZrO6 0.69 3.01 1295 Ortho-silicate Li4SiO4 0.54 2.39 1255 Ortho-tantalate Li3TaO4 0.46 5.87 1679 Meta-titanate Li2TiO3 0.44 3.55 1270 Hexa-zirconate Li6Zr2O7 0.43 3.43 1535 Meta-silicate Li2SiO3 0.39 2.53 1201 Meta-zirconate Li2ZrO3 0.38 4.15 1600 g-Aluminate LiAlO2 0.27 2.55 1610 Octo-plumbate Li8PbO6 0.66 4.24 Lithiumoxidewas favored in early blanket concepts, in particular because of its very high lithium density and good thermal conductivity. Its biggest disadvantage is its strong sensitivity to moisture and the evidence it shows of significant swelling under irradiation. The silicates studied most widely are Li2SiO3 and Li4SiO4. In practice, traces of Li2SiO5 or Li6Si2O7 can be found as well. A variety of lithium zirconates have also been stud- ied, such as metazirconate Li2ZrO3 and Li8ZrO6 octa- zirconate by Roux and coworkers.26,33,35,52,53,62,63,156 Because of activation issues for Zr in the fusion blanket spectra, these compounds became less attractive. Zirconates have been shaped in pellets as well as pebbles, and their irradiation performance at high burnups was considered promising in terms of pellet integrity and tritium release characteristics.33 mpounds and their major characteristics g Specific heat at 400 �C (JgK�1) Thermal conductivity at 400 �C (Wm�1 K�1) Linear expansion at 400 �C, relat. to 25 �C (%) 2.5 6 1 1.5 1.5 0.7 1.9 2.5 0.9 1.4 1.7 0.5 1.3 1.7 0.6 1.6 2.4 0.6 1 1.4 0.4 1.3 2.6 0.4 http://www.iter.org/mach/tritiumbreeding http://www.iter.org/mach/tritiumbreeding Ceramic Breeder Materials 473 Lithium titanates were initially introduced at AECL by Kopasz and coworkers,64,65 and early tri- tium release experiments have shown promising results. Li2TiO3 has been studied by Japanese and EU parties. Hoshino et al.66–70,179–181,190 found that not only oxygen-deficient but also lithium oxide-deficient defects form on changing the atmo- sphere from hydrogen to oxygen. Thus, the doubly nonstoichiometric composition, Li2–xTiO3–y, has been confirmed. Further, it has been shown by ther- mal diffusivity measurement that 95Li2TiO3 has a higher thermal conductivity than 100Li2TiO3. 66 Lithium aluminates have been studied mostly in the form of g-LiAlO2. Data accumulated have been summarized by Billone and Grayhack.71 Due to its low lithium density and modest tritium release per- formance, the material has gradually begun to receive less consideration.62 More recently, Zhu et al.56 in China started inves- tigations of lithium tantalates such as Li3TaO4, which has a reasonably high lithium density. Sedano and coworkers57 report on the develop- ment of lithium plumbates such as Li8PbO6, revisit- ing earlier work of Hayashi et al.72 (a) 3 2 1 5 4 6 (b) Figure 11 Illustration of the melt drop and jet spraying processes developed for production of Li4SiO4 pebbles at KIT. Reproduced from Kolb, M. H. H.; Knitter, R.; Kaufmann, U.; Mundt, D. Fusion Eng. Des. 2011, doi: 10.1016/j.fusengdes.2011.01.104. 4.15.3.2 Fabrication of Shapes The first ceramic breeder blanket concepts used blocks or plates; later, pellet designs emerged. The larger the breeder component, the greater the threat to its integrity because of cracking or fragmentation due to thermal stresses and irradiation-induced swelling and embrittlement. This is the major reason for favoring pebble-bed concepts. 4.15.3.2.1 Pellets or blocks Pellet or block fabrication makes use of proven technologies in the ceramic industry. Pressing and sintering of ceramic powders is a widely used and cost-effective industrial process. Pellets and rectan- gular blocks can be manufactured up to some centi- meters in size with excellent material homogeneity and controlled density. Thus, LiAlO2, Li2ZrO3, and Li2TiO3 pellets meeting dimensional, microstruc- tural, and purity characteristics were produced by Pechiney in collaboration with CEA.52 Similar results were obtained by ENEA, SCK/CEN, UKAEA- Springfields, and US laboratories.65,73–75 Kapychev et al.30 fabricated pellets of Li4SiO4, metasilicate (Li2SiO3), and aluminate (LiAlO2), with a diameter of about 10mm and heights of 5, 10, and 14mm. 4.15.3.2.2 Pebbles For TBR considerations, the density of the pebbles should be high and enable a dense packing to achieve a high lithium density. Further comments on pebble shapes are given in a later section. The presently used or developed processes are as follows: 1. A melting–spraying process was used at KIT (formerly FZK), in collaboration with Schott Glas- werke, for the production of 0.25–0.63mm Li4SiO4 and Li4SiO4–SiO2 pebbles 76 (see Figure 11). After annealing, spherical pebbles of 95–96% of theoretical density (TD) exhibiting satisfactory mechanical strength were obtained. Long-term annealing experiments on various candidates ceramic breeder materials were performed by Piazza et al.77 An alternative route avoiding use of carbonate and using hydroxide was developed by Knitter et al.,78 with slightly lower density. The reference composition is Li4SiO4þ 2.5wt SiO2, resulting in a two-phase Li4SiO4þLi6Si2O7 struc- ture in as-melted condition, and Li4SiO4þLi2SiO3 after heat treatment (see Figure 12). A melting-dropping process was investigated by Tsuchiya et al.79 in collaboration with Mitsubishi to produce 1-mm Li2O spheres. 2. Sol–gel type processes were developed at Japan Atomic Energy Agency (JAEA), with Nuclear Fuel Industries, to produce 1mm Li2O and 1.6mm Li2TiO3 pebbles 80 (see Figure 13). Similarly, Muis and coworkers81 at Energy Research Centre of The Netherlands (ECN) produced 0.5–1.0mm Li2TiO3 Fabrication process Calcination sintering Drying Dropping Dissolving (mixing) Gelation solvent In air Li2TiO3 solvent (H2O2 etc.) (binder) Heating Figure 13 Manufacturing process for Li2TiO3 by sol–gel meth Tsuchiya, K.; Kawamura, H.; Uchida, M.; Casadio, S.; Alvani, C. (a) (b)100 mm 50 mm Figure 12 Cross-section micrograph of pebble fabricated by a melt spray process featuring large domains of dendritically grown crystals composed of Li4SiO4 (light) and Li6Si2O7 (dark): (A) overview and (B) detail. Reproduced from Kolb, M. H. H.; Knitter, R.; Kaufmann, U.; Mundt, D. Fusion Eng. Des. 2011, doi: 10.1016/j. fusengdes.2011.01.104. ECN SEI 5.0 kV �50 100 �m WD15 mm Figure 14 Scanning electron micrographs of Li2TiO3 pebbles 474 Ceramic Breeder Materials pebbles. In these cases, the pebble densities were ECN SEI 15.0 kV �50 100 �m WD36 mm ECN SEI 15.0 kV �500 10 �m WD36 mm Figure 15 Scanning electron micrographs of Li2TiO3 pebbles produced by an extrusion-spheronization method (see text). Ceramic Breeder Materials 475 for producing 1mm Li2O, Li4SiO4, and Li2ZrO3 pebbles. Pebble densities in the 90% TD range were obtained.85 This process has also been inves- tigated at CEA for producing 1mm Li2TiO3 peb- bles. Pebble density of 90% TD and good mechanical strength were obtained.84 5. Zhu et al. developed a wet process for fabrication of Li3TaO4 ceramic pebbles. Typical pebble dia- meters are about 0.7–1.0 mm, and the density achieved is over 90% TD, with crush loads more than 40N.56 X-ray diffraction (XRD) patterns showed 99% of b-Li3TaO4 and traces of LiTaO3. The necessity to recycle ceramic breeders after service imposes specific requirements on pebble manufacturing technologies. This reprocessing aspect may become a significant driver in fusion power economics on a longer term. See Sections 4.15.7 and 4.15.8.8 for further discussion. 4.15.4 Pebble and Pebble-Bed Thermomechanics 4.15.4.1 Introduction One of the issues that need to be addressed is the thermal and mechanical behavior of constrained peb- ble beds under (cyclic) nuclear loading conditions experienced in a tritium breeding blanket. The way a pebble bed responds to a thermal load depends primarily on the thermal transport proper- ties of the bed, such as the packing factor of the bed, and the thermal conductivities of the pebble material and the surrounding purge gas. In addition, due to differences in the thermal expansion coefficients between the pebble-bed material and its surrounding structure, stresses will be induced in both the pebble bed and the structural material, and the contact areas between pebbles and pebbles and walls become an important parameter. Also, during longer term operation in the neutron-irradiation environment, swelling of the breeder material will generate stresses. Furthermore, the pebble-bed thermal transfer prop- erties may deteriorate with irradiation dose and lithium burnup. These induced stresses may directly or indirectly affect the functional operation if the mechanical integrity of the blanket element is endangered or if heat or tritium removal is significantly deterio- rated due to pebble fracture, sintering, or melting. Various creep phenomena will affect the actual evo- lution of stresses, like thermal creep and irradiation creep. In addition, chemical reactions in between pebbles and between structures and pebbles may be enhanced under high contact pressures and the compositional changes arising from long-term oper- ation by lithium burnup and other transmutation reactions. When the (constrained) pebbles experience stres- ses, by either compression due to thermal expansion or irradiation-induced swelling, these stresses will be (partially) relieved through thermal creep, that is, irreversible deformation of the pebbles. Under high stresses and at high temperatures, typically 0.6–0.8 times the melting temperature TM, these relaxation processes include sintering of the material. This irre- versible strain and deformation in the material can lead to embrittlement and, ultimately, to fragmenta- tion of the pebbles. The fracture of pebbles results in an inhomoge- neous pebble-bed density and contact area and would lead to an inhomogeneous temperature distribu- tion that is hard to model or predict. The lack of predictability of the temperature distribution can be a major safety issue. A mechanically stable pebble bed with a high packing factor is desired. 476 Ceramic Breeder Materials 4.15.4.2 Single Pebble Testing The simplest way to determine the mechanical strength of pebbles is by performing crush tests with individual pebbles. These tests are very useful for optimization of pebble properties and for quality control during pebble production. In crush tests, pebbles are arranged between two plane plates. In a pebble bed, the pebbles in contact with the cavity wall experience similar conditions because there is only one contact at the corresponding pebble hemisphere. Pebbles that are in contact only with other pebbles have on an average six contacts86,87 and the contact forces are in general smaller than on the pebble-wall contact. The crush loads of 0.5-mm diameter Li4SiO4 and 1-mm diameter Li2TiO3 pebbles used for the European HCPB project are in the range of 4–5 and 35–50N, respectively, as shown by Knitter et al.191 and Roux et al.192 (see Figure 16).90 The crush loads of the Li2TiO3 pebbles were independent of 6Li enrichment. Assuming coverage of the wall by pebbles of about 0.7, the lower values of 4 and 21N for the 0.5 and 1.0mm diameter pebbles correspond to pres- sures on the wall pebbles of about 12 and 15MPa. These values are significantly above the assessed maximum pressures of about 6MPa in the HCPB ceramic breeder pebble bed.91 However, there are several effects that decrease this margin: � In pebble beds, more generally, in granular mate- rials, an inhomogeneous branching of forces occurs (see Jaeger et al.92), with the effect that some 10 8 6 4 2 Initial material Cond. material Annealing time (days) C ru sh lo ad (N ) C ru sh lo ad (N ) 3 24 48 96 Li4SiO4 01/3-3 (ex carbonate) Li4SiO4 01/3-4 (ex hydroxide) Figure 16 Typical results of single pebble crush tests as a fun Dr. R. Knitter, KIT. pebbles experience much larger forces than the average value. � Thermal annealing results in a significant grain size growth in Li2TiO3 pebbles; a slight decrease in crush load during aging was found for the Li4SiO4 and Li2TiO3 pebbles. 77 � Neutron irradiation will have the strongest effect to reduce the pebble strength, as mentioned here. Fortunately, there is one mechanism, as outlined below, that is expected to alleviate the problem con- siderably: thermal creep. A general remark to be made here is that care is to be taken with respect to statistics: postirradiation tests from the EXOTIC-7 high burnup experiment showed fragmented pebbles as well as intact pebbles maintaining average strength levels.33,93 This statis- tical aspect is to be addressed more rigorously in optimization of pebble-bed technologies serving pre- dictable and reliable operation of breeding blankets. 4.15.4.3 Properties of Pebble Beds As the ceramic breeder material is used in the form of pebble beds, the macroscopic properties of the gran- ular material are of particular interest. Most of these properties (mechanical, thermal, etc.) cannot be deduced directly or precisely from the properties of the base material or the single pebble; therefore, dedicated experiments are required. In these experi- ments, particular attention is paid to reproduce rep- resentative conditions in terms of, for example, pebble bed typical dimensions, packing factor, reduc- ing atmosphere, and so on, as envisaged for the blan- ket component. 65 55 45 35 25 Initial material Cond. material Annealing time (days) 3 24 48 96 Ti 1100 CEA CTI 13 B0 ction of annealing time. Published by courtesy of Ceramic Breeder Materials 477 This section mostly follows the approach devel- oped for the European blanket program and the work Reimann and coworkers94–106,110,115 have performed from the late 1990s onwards. 4.15.4.3.1 Pebble-bed density and packing factor The following definitions are used: � rs¼ density of the pebble base material � rp¼ pebble density � Pp¼ total pebble porosity¼ average fraction of porosity within (macroscopic, single) pebble volume � rpb¼ pebble-bed density¼ ratio of pebble-bed mass mpb to pebble-bed volume Vpb � g¼ packing factor¼ ratio of pebble volume to pebble-bed volume The quantities are connected by rpb ¼ mpb=Vpb ¼ ð1� PpÞrsg ½1� In literature, the term ‘fraction of theoretical density’ is also used. This term is identical to (1�Pp). The pebble-bed density rpb is often named ‘apparent’ or ‘tap’ density. The pebble-bed density rpb is an impor- tant quantity for nuclear calculations as these require the specific density of lithium and other constituents as inputs. In pebble-bed engineering, the packing factor g (generally given in percentage) is the charac- teristic quantity. The packing factor is influenced by the following parameters: filling procedure, pebble shape, surface roughness diameter d, diameter distribution, and con- tainer (cavity) dimensions: height H, diameter D. Filling procedure: The selection of an optimum pro- cedure is not trivial, especially for larger mock-ups: filling should be assisted by vibration in such a way that a particle flow within the cavity is prevented or, in case of too small vibration energies, friction forces between pebbles and walls can be overcome. Pebble shape and surface roughness: ceramic breeder pebble shapes are almost spherical, partly with in- dentations (melting, spraying processes) or egg-like (extrusion, spheronization, sintering processes). The pebbles originating from melting have a smoother surface than the others. The packing factor is influ- enced by the surface roughness but negligibly by the pebble shape. Diameter distribution: The diameter d is of influence with respect to the cavity dimensions, discussed later. Ideally, the packing factor increases with increasing diameter distribution; see, for example, McGeary.87 However, sedimentation during the filling is a critical issue. For blanket pebble beds, two groups are of inter- est: (1) Pebble beds with a relatively small d variation: for example, Li2TiO3 pebbles with a nominal diame- ter of 1mm ranging from 0.8 to 1.2mm and Li4SiO4 pebbles ranging from 0.25 to 0.63mm. Production costs generally favor a wider spread, which is not detrimental in respect to the packing factor, at least for the diameter ranges cited earlier. Packing factors in the range of 63–66% are reached; (2) Binary beds: here, large pebbles, dl, and small pebbles, ds, are used, with ds� 0.1dl. First, the large pebbles are vibrated in the cavity. Then the pebbles are fixed in their posi- tion, for example, by a sieve in the filling pipe, to avoid sedimentation during the following pouring in of the small pebbles, which fill the spaces between the large pebbles. Packing factors of about 82% were obtained for beryllium beds with dl� 2mm, ds� 0.2 mm.94 Critical issues are (1) the stability of such pebble packing structure during thermal cycles in the blanket and (2) the buildup of large stresses due to irradiation-induced swelling causing pebble deg- radation, leading to fracture. Diameter d, bed height H, diameter D: For finite cavities, there exist two characteristic ratios: H/d and D/d. The reason why g depends on H/d and D/d is the existence of two different characteris- tic pebble packing structures as demonstrated by tomography experiments101,105,106: in the pebble bulk, there is no preferential direction of contacts between pebbles, whereas at the wall, a zone with a thickness of about 4d exists, where regular packing structures are observed and contact zones are no longer homogeneously distributed. Figure 17 shows tomography results106 for a cylindrical container (D¼ 49mm, H¼ 50mm) filled with spherical pebbles (d¼ 2.3mm). Distinct pebble layers close to the walls can be clearly seen in Figure 17(a) where the positions of the pebble cen- ters are plotted in a horizontal plane (left) and a vertical plane (right). Figure 17(a) shows axial void fraction distribu- tions. The degree of regularity is largest for the pebbles in contact with the wall and decreases with increasing wall distance. The regularity is most expressed for pebbles in contact with plane walls (D¼1); here, large isles with a dense hexagonal pebble arrangement are observed; see Reimann et al.105 The average packing factor of a plane wall zone is about equal to the bulk packing factor if the bed height is sufficiently large for the two wall zones Horizontal position of sphere centers Vertical position of sphere centers 0.025 0.000 y( m ) f ro m t he R O I c en te r R O I h ei gh t (m ) –0.025 –0.025 (a) (b) 50 1.0 0.9 0.8 0.7 0.6 0.5 0.4 0.3 0.2 0.1 0 2.5 Radial wall distance (mm)Axial void fraction (1) 5.0 7.5 10.0 12.5 15.0 17.5 20.0 22.5 25.0 45 40 35 30 25 20 Ve rt ic al d is ta nc e (m m ) R ad ia l v oi d fr ac tio n (1 ) 15 10 5 0 0.2 0.3 0.4 0.5 0.6 0.7 E0: p = 0 MPa E0: p = 0 MPa E2: p = 9.4 MPa E4: p = 9.2 MPa E2: p = 9.4 MPa 0.8 0.9 1.0 0.000 0.000 0.04 0.03 0.02 0.01 0.00 0.005 ROI radius (m)x(m) from the ROI center 0.010 0.015 0.0200.025 Figure 17 Visualization of pebble-wall layers by tomography: (a) Positions of sphere centers in region of interest (ROI) and b) Void fraction distributions. Reproduced from Pieritz, R. A.; Reimann, J.; Ferrero, C. Adv. Eng. Mater. 2011,13, Nr.3, 145–155. Best location for filling Y Z X Figure 18 Pebble-bed filling experiments relevant for ITER Test Blanket Modules. Reproduced from Abou-Sena, A.; Neuberger, H.; Ihli, T. Fusion Eng. Des. 2009, 84, 355–358. 478 Ceramic Breeder Materials to interfere. This is an important result because it demonstrates that the often cited statement that wall zones are characterized by reduced packing factors is not generally valid. However, with decreasing D, both the regularity of the packing and the packing factor decrease, as seen clearly in Figure 17(b). Corners in pebble-bed cavities and inserts, for example, thermocouples, generally decrease the packing factor. This is also a reason why in small experimental setups the achieved packing factor is often distinctively below the maximum value. Pebble beds in HCPB blanket concepts are typi- cally ‘thin’ pebble beds, that is, one dimension; the bed height H, is small compared with the other dimensions. For the ceramic breeder pebble beds, H� 10mm; for beryllium, H� 30mm. For a height of H� 10mm and a pebble diameter of �1mm, the maximum packing factor is already difficult to achieve. TE Displacement transmitter (in total 4) Furnace Piston Pebble bed Container Al2O3-disc Figure 19 Set-up for Uniaxial Compression Testing of pebble-beds. Reproduced from Reimann, J.; Harsch, H. In CBBI-12, 12th International Workshop on Ceramic Breeder Blanket Interactions, Karlsruhe, Germany, 2004; FZKA 7078. Ceramic Breeder Materials 479 A disadvantage of a large pebble diameter is also the smaller heat transfer coefficient, as outlined below. The blanket modules will probably be filled with pebbles through small pipes, as schematically shown in Figure 18. For simple cubic cavities, a homogeneous filling was achieved only by carefully vibrating the container in different tilted positions.97,100 The filling of a TBM mock-up using glass balls for the beryllium pebble beds was investigated by Abou-Sena et al.,107 by varying the number and posi- tions of the filling pipes. The best results were obtained by using a single filling pipe close to the box corner and, keeping the filling pipe at the highest position, tilting the box around two axes. Here, a packing factor of 63.6% was achieved, which is considered to be very close to the maximum obtain- able value. In the HELICHETTA experiment (see Dell’Orco et al.108,109) where elongated pebble-bed cavities (10� 100� 480mm3) were filled through a pipe at the top (largest dimension orientated vertically), packing factors of 60% and 62% were achieved for Li2TiO3 pebbles with diameters between 0.8 and 1.2mm, and 0.6 and 0.8mm, respectively. For Li4SiO4 pebbles with d between 0.25 and 0.63, the packing factor was more than 64%. 4.15.4.3.2 Mechanical behavior of pebble beds Reimann et al.94,96,110 obtained typical load displace- ment data for the ceramic pebble beds introduced above. The typical setup used is given in Figure 19.110 The pebble beds, contained in cylindrical cavities, are compressed by a piston and the pressure (equal to the uniaxial stress s) and the bed deformation (strain e) are measured, and the modulus E of the pebble bed is derived. These uniaxial compression tests (UCTs) should not be performed with bed height H to diameter D ratios larger than 1 so as to avoid a significant influence of friction effects at the cylindrical wall. Figure 20110 shows a typical result in terms of stress and strain. Key features of the pebble-bed deformation under monotonous isothermal mechan- ical loading are as follows: � Nonlinear elasticity: the pebble-bed stiffens with higher degrees of deformation. � Irreversible deformation is observed after initial unloading due to pebble relocation and plastic deformation at pebble–pebble and pebble–wall contacts (which disappears after multiple loading- unloading cycles). � Thermal creep at constant load with further irreversible deformation. � Further stiffening observed at loading and unload- ing cycles. � In addition, friction forces may provide some hysteresis at any loading-unloading cycle. 4.15.4.3.3 Thermal creep Characteristic differences of thermal creep effects between homogeneous materials and pebble beds are: � In pebble beds, the contact surfaces between the pebbles increase with time. Blanket relevant creep time periods of greatest interest are in the order of less than a day because stress relaxation effects are expected to be quite fast.104 � Pebble beds are loaded first with relatively small stress gradients ds/dt ; therefore, it is not differen- tiated between instantaneous plastic deformations and conventional thermal creep effects (the ther- mal creep correlations given below include both effects). Thermal creep strains are measured by UCTs by keeping the uniaxial stress s constant at a given temperature T. Figure 21110 shows creep strains for Li4SiO4 pebble beds for different values of s and T; the thermal creep strain occurring during the stress increase period has been taken into account in this representation. 5 4 3 2 1 0 0 1000 2000 3000 Time (min) C re ep s tr ai n (% ) 4000 5000 s = 2 MPa 750 �C: Li4SiO4 ex hyd 750 �C: Li4SiO4 ex silic 750 �C: Li2TiO3 900 �C: Li4SiO4 ex hydr 900 �C: Li4SiO4 ex silic 900 �C: Li2TiO3 Figure 21 Thermal creep strain as a function of time for Li4SiO4 & Li2TiO3 pebble-beds using UCT. Reproduced from Reimann, J.; Wörner, G. In Proceedings of the 12th International Workshop on Ceramic Breeder Blanket Interactions, Karlsruhe, Germany, FZKA; 2005; p 7078. 0.001 0.0001 0.88 0.92 0.96 1 1.04 1000/T(K) C /s 0 CEA-titanate C1/s 0.85= 0.67exp(-7576/T(K)) FZK-silicate p = 8.6 MPa p = 4.3 MPa p = 2. 2 MPa p = 6.4 MPa C1/s 0.85= 12.12exp(-10 220/T(K)) JAERI-titanate C1/s 0.85= 0.37exp(-6947/T(K)) 1.08 1.12 1.16 Figure 22 Pebble-bed creep: temperature dependence for a constant value of the stress exponent. Reproduced from Reimann, J.; Wörner, G. In Proceedings of the 12th International Workshop on Ceramic Breeder Blanket Interactions, Karlsruhe, Germany, FZKA; 2005; p 7078. Uniaxial strain e (%) Uniaxial strain e (%) U ni ax ia l s tr es s s (M P a) 0.0 0.0 1.0 2.0 3.0 4.0 5.0 1.0 2.0 3.0 4.0 5.0 U ni ax ia l s tr es s s (M P a) Li4SiO4 ex hydroxide T(°C) 25 850 (a) (b) 0 0 0.2 0.4 0.6 0.8 1 1 2 3 4 5 1st cycle Li2TiO3-M 2nd cycle 3rd cycle(b) (a) (a) (b) Figure 20 Example of uniaxial compression test results for Li4SiO4 pebble beds. (a) Li4SiO4 pebble beds; (b) Li2TiO3 pebble beds. Reproduced from Reimann, J.; Wörner, G. In Proceedings of the 12th International Workshop on Ceramic Breeder Blanket Interactions, Karlsruhe, Germany, FZKA; 2005; p 7078. 480 Ceramic Breeder Materials The data are well fitted by a constant exponent n for the time dependence. A constant exponent n was also found for most of the Li2TiO3 pebble beds. 95 For some batches, during the first several hours, the same exponent n was observed; however, there was a subsequent increase in the creep rates. The Li2TiO3 batches that showed this behavior were in general characterized by low sintering temperatures, large porosities, and small grain sizes, (for details, see Reimann et al.98). Eventually, impurities could play a role, too. The data are fairly well plotted in Arrhenius graphs by straight curves, as shown in Figure 22.95 For a selection of ceramic breeder material candidates, the influence of pebble-bed thermal conductivity as a function of, for example, packing factor and pebble-bed compression was studied by Ceramic Breeder Materials 481 means of UCTs. These were performed for the ceramics Li4SiO4, Li2TiO3, Li2ZrO3, and Li2O with temperatures up to 480 �C and pressures up to 8MPa, with packing factors varying between 56% and 63.5%. Creep strains in the pebble beds are identified to be functions of temperature, stress, and time and are found to be of the form ecr ¼ AðTÞsmt n where m and n are independent of temperature and need to be deduced experimentally.97 This research was extended for Li4SiO4 pebble beds with temperatures up to 850 �C and pressures up to 9MPa and concludes that thermal creep effects are negligible at temperatures below 600 �C.96,97 Creep behavior is also determined by the pebble properties: as mentioned earlier, lower creep strains were found for Li2TiO3 with small grain sizes (�5mm) and high sintering temperatures.98 4.15.4.3.4 Cyclic loading An ITER TBM will experience not a constant heat load but a cyclic heat load behavior due to burn pulses of the plasma. Dalle Donne et al.111 per- formed exploratory thermal cycling experiments with densified Li4SiO4 pebble beds in a horizontal tube (D¼ 20mm, L¼ 110mm). After about 500 cycles (350–600 �C), fractions of 1.8% and 3.5 wt% of broken pebbles in two measurement campaigns were found. The pressure drop in the helium purge flow was found to saturate after some hundred cycles, indicating that no further pebble fragmentation occurred. Thermal shock is not considered an issue for ceramic breeder pebbles; the experiments showed that pebbles only fail for dT/dt>60 �Cs�1, that is, values much larger than expected in the blanket. Cycles of UCTs up to 4MPa at ambient tempera- tures lead to negligible plastic deformations in the Li4SiO4 and Li2TiO3 pebble beds. The residual com- paction is likely to be due to pebble relocation or pebble cracking. Multiple 6MPa compression cycles at 750 �C led to irreversible strains of 3.5% and 4% for Li2TiO3 and Li4SiO4 pebble beds, respectively. When 4MPa compression cycles are performed on an Li4SiO4 pebble bed at 840 �C, the creep strains reach values of 6%. In these last tests, the influence of thermal creep is clearly visible.110 Similar cyclic compression tests on Japanese Li2TiO3 pebble beds reveal a compression ‘equilib- rium’ of 1.5% only after several cycles of 10MPa at 400 �C.112 Performing these tests at 600 �C gives lower strain values in the pebble bed, likely due to the increased compaction.113 This is in agreement with the above-mentioned results.110 Several HCPBmock-up experiments have been per- formed in the HE-FUS3 facility at ENEA, Brasimone by Dell’Orco et al.108,109,114 In the HELICHETTA and HELICA experiments, the thermomechanical behavior of Li4SiO4 or Li2TiO3 pebble beds under thermal cycling loads was investigated. The following larger HEXCALIBUR type experiments were char- acterized by two beryllium pebble beds and one ceramic breeder pebble bed in between. Nuclear heating was simulated by using electrically heated plates within the pebble beds. In the HELICA mock-up tests, the thermal cycling of the pebble beds showed a saturation of the pebble-bed strains after 30 cycles. The bed heights were then reduced by �4% and �1.5% for the Li4SiO4 and Li2TiO3 pebble beds, respec- tively. For constant power levels, the system did not change thermally with increasing cycle number. During demounting, submicron pulverized material was released, especially from the Li4SiO4 pebble beds, and pebble fragmentation was observed. An important objective of these experiments was the validation of pebble-bed thermomechanical models as outlined below. 4.15.4.4 Heat Transfer Properties 4.15.4.4.1 Thermal conductivity There are several advanced models to describe the thermal conductivity of pebble beds, which take into account the relevant parameters such as the thermal conductivity of the pebble material ks as a function of temperature T, the thermal conductivities of the surrounding gas kg as a function of temperature T and pressure p, the pebble diameter d, packing factor g, contact surface ratio rk 2, and several other second- order effect parameters. In the following section the Schlünder Bauer Zehner model (SBZ model)114 is used to demonstrate the influence of some para- meters. Figure 23 shows ks as a function of tempera- ture for both Li4SiO4 and Li2TiO3 and kg for helium and other gases at 0.1MPa.99 For most ceramic breeder materials, ks decreases first with increasing T, reaches at high temperatures a plateau, or increases again slightly; for details, see Abou-Sena et al.116 The helium conductivity increases strongly with increasing T. Figure 24(a) and 24(b) from Reimann et al.99 shows the influence of T and rk 2 on the pebble-bed conduc- tivity k of an Li4SiO4 pebble bed with mean diameter of d¼ 0.4mm and g¼ 64%. For a noncompressed bed, rk 2¼ 0, k increases moderately with T because kg increases with T. For a moderately compressed bed, 4 0.5 0.4 0.3 0.2 0.1 0 3.5 2.5 2 1.5 1 0.5 0 0 200 400 600 Temperature T(�C) B re ed er m at er ia l c on d uc tiv iti es (W m –1 K -1 ) G as c on d uc tiv iti es (W m –1 K -1 ) 800 1000 3 Orthosilicate Helium Metatitanate Air Argon Figure 23 Thermal conductivity of ceramic breeder pebble beds in helium and air. Reproduced fromReimann, J.; Hermsmeyer, S. Fusion Eng. Des. 2002, 61/62, 345–351. 7 6 5 4 3 2 1 0 0 0.26 0.26 0.26 0.26 0.24 0.27 0.27 0.28 0.28 0.29 Orthosilicate pebble bed T = 25 �C 0.3 0.3 k = 0.29 W mK-1 0.4 U ni ax ia l s tr es s (M P a) (a) (b) 0.8 Uniaxial strain (%) 1.2 1.6 2 7 6 5 4 4 3 3 2 2 1 1 0 0 0.55 0.59 0.59 0.64 0.63 Stress = constant for 1500 min 0.65 0.67 k = 0.60 W mK-1 Orthosilicate pebble bed T = 800 �C U ni ax ia l s tr es s (M P a) Uniaxial strain (%) Figure 24 Stress–strain dependence of Li4SiO4 pebble beds at (a) 25 and (b) 800 �C and evolution of thermal conductivity. Reproduced from Reimann, J.; Hermsmeyer, S. Fusion Eng. Des. 2002, 61/62, 345–351. 482 Ceramic Breeder Materials rk 2¼ 0.02, k is larger than that for rk2¼ 0, but the temperature dependence is quite small. The thermal conductivity of noncompressed Li4SiO4 and Li2TiO3 pebble beds was measured by several authors. Figure 25 from Enoeda et al.86 shows that there is a good agreement among different authors. The Li4SiO4 pebble-bed data are best fitted with a correlation established by Dalle Donne204; the Li2TiO3 data were well predicted by the SBZ model using a value rk 2¼ 0.0049.117 Li2TiO3 pebble-bed data for 2mm pebbles from Abou-Sena116 are char- acterized by the tendency of a decrease in k with increasing T. Results for noncompressed beds including further ceramic breeder pebble materials were summarized by Abou-Sena et al.116: k should increase with increas- ing ks in the sequence Li2ZrO3, Li4SiO4, Li2TiO3, Li2O for equal values of the other parameters. Because the other parameters differed, this tendency was masked. For compressed pebble beds, the SBZ model con- tains the parameter rk 2, which is a priori not known. If the pebble-bed strain is measured, it is easier to use this quantity as the relevant parameter. Figures 26(a) and 26(b) from Reimann and Hermsmeyer99 show results for Li4SiO4 and Li2TiO3 pebble beds, differ- ent gas conditions, and strain values e. For noncom- pressed beds, rk 2¼ 0, the measured data agree fairly well with SBZ model predictions. With increasing strain, k increases; however, only very moderately compared with beryllium pebble bed. Even for a strain of about 4%, obtained in air at 800 �C because of significant thermal creep, k is only increased by about 20% for both types of pebble beds. For helium atmosphere, this difference becomes even smaller. Figure 26(b) also contains some results for binary Japanese Li2TiO3 pebble beds (0.2 and 2mm peb- bles, g¼ 81.5%) in air atmosphere at ambient tem- perature. Compared with the monodisperse pebble bed with d¼ 2mm pebbles and g¼ 64.3%, k is increased by a factor of 2. For blanket relevant con- ditions, this factor reduces to 1.3 for T¼ 600 �C and helium atmosphere. 2 1 0 0 200 400 Temperature (�C) 600 800 1000 He pressure = 101-103 kPa Contact area fraction = 1. Ce-0 Accommodation coeff. = 0.2 4 3 2 1 0 Contact area fraction = 4.9 × 10−3 Accommodation coeff. = 0.2 0 200 400 Temperature (�C) 600 800 1000 Th er m al c on d uc tiv ity (W m K -1 ) Th er m al c on d uc tiv ity (W m K -1 )FZK Li4SiO4+SiO2 0.25–0.63 mm 60.13%PF, hot wire method FZK Li4SiO4 0.5 mm 52%PF by M.D.D. & G.Sondon, fusion tech. volt 7 697(1990) Dr. Plazza FZK Li4SiO4+TeO2 66%PF SZB mod HM mod Li2O data CEA Li2TiO3 60.8%PF CEA Li2TiO3 50.4%PF CEA Li2ZrO3 54.42%PF SZB mod HM mod Figure 25 Comparison of pebble bed thermal conductivity data after Enoeda, M.; Ohara, Y.; Roux, N.; Ying, A.; Malang, S. In Proceedings of the CBBI-8, Colorado Springs, CO, Oct 6–8, 1999. 1.1 1.0 0.9 0.8 0.7 0.6 0.5 0.4 0.3 0.2 0.1 0.0 0 1 P eb b le b ed c on d uc tiv ity k (W m K −1 ) 2 3 4 Helium; 475 25 25 750 800 Helium; Air; Gas Orthosilicate pebble beds; T(�C) Air; Air; 0 k( W m K −1 ) 1 2 e(%)Strain e(%) 3 Metatitanate pebble beds Type Ti-J Air 25 25 25 25 25 25 750 800 Air Air Air Air Air Helium Argon Ti-J-bin Ti-D Ti-D Ti-D Ti-E Ti-D Ti-D Gas T(�C) 4 5 0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1 1.1 Figure 26 Evolution of thermal conductivity of compressed Li4SiO4 and Li2TiO3 pebble-beds as a function of imposed strain. Reproduced from Reimann, J.; Hermsmeyer, S. Fusion Eng. Des. 2002, 61/62, 345–351. Ceramic Breeder Materials 483 Experiments with compressed Li2TiO3 pebble beds (d¼ 2mm, g¼ 65–67%) were also performed by Tanigawa et al.113 For a strain of about 1%, T¼ 600 �C, and helium at 0.1MPa, k increased only by 3% compared with the noncompressed pebble bed. After annealing the pebble bed at 700 �C without com- pression for 1 day, larger bed strains (factor 2) were obtained in the subsequent cycles and with this an additional increase in conductivity. 4.15.4.4.2 Heat transfer Close to the wall, the pebble packing differs signifi- cantly from that in the bulk, as demonstrated in Figure 17. For noncompressed pebbles, the void fraction at the wall surface is close to 100%. Heat transfer characteristics in the wall zone are, therefore, different from those in the bulk. This fact is taken into account by using the heat transfer coefficient h, which is defined with the temperature difference (Tew�Tw), where Tew is obtained by extrapolating the bulk pebble-bed thermal conductivity k up to the wall, and Tw is the undisturbed wall temperature. In measurements, (Tew�Tw) is very small and is sensitively dependent on extrapolated temperature profiles. Accurate measurements are, therefore, extremely difficult, and the discrepancies in experi- mental data are significant. The smaller the pebble diameter, the thinner the wall zone, and with this the difficulty to obtain accurate data increases. Again, different models exist to predict h exist, which have been validated in better suited experiments. Figure 27 shows h¼ f (T ) for Li4SiO4 and Li2TiO3 pebble beds for noncompressed beds calcu- lated with the model from Yagi and Kunii118: the strong increase of h with decreasing pebble diameter d is obvious. With progressing compression, the wall Yagi and Kunii model 1000 300 400 500 600 700 800 900 2000 3000 4000 5000 6000 Temperature T(°C) H ea t tr an sf er c oe ff ic ie nt h (W m –2 K ) Li 4 SiO 4 d = 0.4 mm Li 2 TiO 3 d = 0.6 mm Li 2 TiO 3 d = 0.8 mm Li 2 TiO 3 d = 1.0 mm Figure 27 Heat transfer coefficients h from Yagi and Kunii model for different Li4SiO4 and Li2TiO3 pebble diameters. p-hor (MPa) 0.12 0.4 0.08 0.06 Vibrated bed; symbols: ABAQUS Horizontal strain (%) 0 0 5 5 Ve rt ic al p re ss ur e (M P a) Figure 28 Bi-axial pebble-bed deformation tests: experimental results and calculations with ABAQUS code system. Reproduced from Hermsmeyer, S.; Reimann, J. Fusion Eng. Des., 2002, 61–61, 367–373. 200 SCATOLA calculations for 3.0 mm plates Exp.) 484 Ceramic Breeder Materials contact surfaces increase and with this h. Again, this increase is expected to be much smaller than that in beryllium pebble beds because of the small thermal conductivity of ceramic breeder materials. At present, mechanistic models are developed taking into account the pebble arrangements close to the wall as determined by tomography101,105,106,199 or by discrete-element modeling (DEM), outlined later. A typical example is the work of Gan et al.133 175 150 A2 A1 B2 B1 125 100 75 50 P la te s ax ia l d is p la ce m en t (m m 25 0 0 100 200 300 400 Temperature (�C) 500 600 700 Calc. A1 Calc. A2 Calc. B1 Calc. B2 Exp., 1st heating Figure 29 SCATOLA bench-mark calculation with a continuum model.127 Reproduced from van der Laan, J.; et al. Proceedings of the 8th International Workshop on Ceramic Breeder Blanket Interactions (CBBI-8), Colorado Springs, CO, 1999; Ying, A., Ed.; UCLA Report. 4.15.4.5 Pebble-Bed Modeling There are different types of models to describe the behavior of pebble beds (see, e.g., Reimann et al.97): finite-element models based on continuum mechanics, also called finite-element modeling (FEM), and the so-called discrete-element models, DEM, for description of mechanics at micromecha- nics level (i.e., individual pebble–pebble interactions). The development of computational tools at these two different length scales (macro- and microscales) allows for a better description of the thermomecha- nics of the pebble beds. 4.15.4.5.1 Continuum models The macroscopic behavior of pebble beds is de- scribed by constitutive equations commonly used in soil mechanics, considering the granular material as a continuum material that can undergo reversible elastic deformations, inelastic volume compaction (consolidation), and pressure-dependent shear fail- ure. To account for these properties, different models have been developed, which are implemented in structural computational programs.91,115,117,120–125 Pebble-bed data (see sections above) had to be implemented, and user-defined subroutines had to be written, for example, for the thermal creep laws.115 The codes were first validated with fairly simple experiments.104,123,126,127 Figures 28 and 29 show examples of results from the biaxial experiment126 and the SCATOLA experiment for calculational benchmarking127 (Figure 29). These codes were later validated with more com- plex mock-up experiments such as HELICA and HEXCALIBUR91,108,114,128,129 and, as outlined below, aimed to be validated by the pebble bed assembly (PBA) experiment.130 Figures 30(a)117 and 30(b)91 showa comparison betweenmeasured and calculated HELICA temperatures. 800 700 600 500 400 300 200 1 2 3 4 5 TC_01 Cassette Heater Time (h) (a) 55 mm Te m p er at ur e (� C ) TC_02 TC_03 TC_04 6 7 8 9 10 11 12 13 14 15 16 17 (b) 900 800 700 600 500 400 300 200 (a) 0 60 120 180 Time (min) Te m p er at ur e (� C ) 240 300 360 Thermocouple TC 1 Thermocouple TC 2 Thermocouple TC 3 Thermocouple TC 4 Figure 30 (a) Measured temperatures in the HELICA experiment, reprinted fromDell’Orco, G.; di Maio, P. A.; Giamusso, R.; Tincani, A.; Vella, H. Fusion Eng. Des., 2007, 82, 2366–2374. (b) Calculated temperatures for the HELICA experiment, reprinted with courtesy from Gan, Y. Ph.D. Thesis, Universität Karlsruhe, 2008. 6 5 1st cycle (PD = 68.29%) 2nd cycle (PD = 60.95%) 3rd cycle (PD = 61.10%) A xi al lo ad in g (M P a) A xi al lo ad in g (M P a) 4 3 2 0 0.5 Axial strain (%)Axial strain (%)(a) (b) Loading process Loading process Unloading process Unloading process 1 1.51.510.50 2 3 4 5 7 30 3030 25 2525 20 2020 15 1515 w L H 10 1010 5 55 0 00 6 2 Figure 31 Stress–strain behavior of granular materials in (a) a rectangular box under uniaxial compaction and (b) packing density effect. Reproduced from An, Z.; Ying, A.; Abdou, M. Fusion Eng. Des. 2007, 82, 2233–2238. Ceramic Breeder Materials 485 The final goal of these codes is to determine the thermomechanical behavior of pebble beds in TBMs for ITER and DEMO blanket modules. At present, the codes are set up for this task. Up to now, only small portions of a DEMO blanket have been modeled. In these calculations, maximum stresses (very loca- lized) of 6 and 2MPa were obtained for the ceramic breeder and beryllium pebble beds, respectively. Because of thermal creep, these values were re- duced after 2 h to 75% and 25%, respectively, of the initial value. 4.15.4.5.2 Discrete-element modeling DEM is used to study the interparticle force distri- bution and translate the microscopic information into macroscopic information such as stress–strain re- sponse. This method is important for estimating the overall properties of pebble beds, such as yield strength and crush probability. The use of DEM for fusion reactor blanket analyses was started at the Uni- versity of California, Los Angeles (UCLA)124,131,205 and has been continued by KIT.91,119,132,133 A standard experiment, used by both groups, to validate DEM is UCTs; see corresponding section above. Figure 31 from An et al.205 shows three cycles of a pebble bed with an initial packing factor of 60.3%. With increasing cycle number, the pebble bed becomes stiffer. The strong dependence of the stress–strain curves on the packing factor was also stated by Gan and Kamlah.119 DEM analyses were also compared with the SCA- TOLA experiments.124 Some features were well described (see Figure 32), but thermal creep was not satisfactorily predicted. With increasing load on the pebble bed, the number of contacts between pebbles, Nc, increases, as shown in Figure 33.91 An important quantity to assess the frac- tion of crushed pebbles during blanket operation is the Simply supported boundary Fixed boundary 250 200 150 100 50 0 0 100 200 300 Temperature (�C) D ef or m at io n (m m ) 400 500 600 Experimental data (increasing T ) Numerical estimations (increasing T ) Numerical estimations (increasing T ) Numerical estimations (decreasing T ) Numerical estimations (decreasing T ) Experimental data (decreasing T ) Figure 32 Calculational results for SCATOLA experiments. Reproduced from Ying, A.; Huang, H.; Abdou, M.; Lu, Z. In: Proceedings of the 9th International Workshop on Ceramic Breeder Blanket Interactions, Toki, Japan, Oct. 27–29, 2000. 6.7 6.6 6.5 6.4 6.3 6.2 6.1 6.0 5.9 5.8 0 1 2 3 Hydrostatic pressure (MPa) C oo rd in at io n nu m b er 4 5 nc= 5.80 � 0.490 p 0.288 6 7 8 H-1 H-2 H-3 M-1 M-2 M-3 Figure 33 The number of contacts between pebbles in a pebble-bed increases with load, reproduced with courtesy from Gan, Y. Ph.D. Thesis, Universität Karlsruhe, 2008. 0.15 All 0.1 0.05 0 1 2 3 Normalized force P ro b ab ili ty d is tr ib ut io n 4 5 Loading = 1.0 Loading = 3.0 Loading = 5.0 Loading = 7.0 Loading = 9.0 Figure 34 The probability of maximum contact forces between pebbles in a pebble-bed increases with load. Reproduced from An, Z.; Ying A.; Abdou, M. Fusion Eng. Des. 2007, 82, 2233–2238. 486 Ceramic Breeder Materials evaluation of the probability distribution of the maxi- mum contact forces for different loading conditions (see Figure 34).205 A comparison of results obtained byDEM and tomography is shown in Figure 35133: the packing factor distributions agree quite well, and the radial and vertical positions of particles show the same structure as shown in Figure 17(b). Other approaches are found like those of Aquaro and co- workers.123,199,202,203 Although the present DEM codes have proven to be very helpful for the understanding of the interaction between pebbles, there is still consider- able development work required until quantitative results for small thermally loaded pebble-bed geo- metries can be expected. The method will certainly not be applicable for large components in the near future because of computational costs, but the improved understanding of the micromechanism will be beneficial for the improvement of the con- tinuum codes. 4.15.4.6 In-Pile Behavior Ceramic breeder materials have been tested under various conditions to determine their irradiation response in terms of tritium production and release and their microstructural, thermal, chemical, and mechanical stability. The nuclear loading of ceramic breeder affects their performance in various ways, of which the most important are the following: � Lithium burnup and other neutron transmutation reactions gradually change the material composi- tion and affect the chemical and physicochemical interactions. � Transmutation reactions of major and minor con- stituents gradually change the radioactivity levels relevant for in-tokamak operation, hot-cell opera- tions, and management of waste, including recy- cling options. 201510 -10 10 20 20 20 10 10 (a) (b) 0.9 0.7 0.6 0.5 0.4 0.3 5 10 15 20 0.8 0.90.70.60.50.40.3 0.8 Distance to the center (mm) Packing factor D is ta nc e to t he c en te r p la ne (m m ) P ac ki ng fa ct or -10 -20 -20 -10 (c) -20 5 25 (3) (2) (1) -10 10 0 20 (d) -20 Figure 35 Comparison of discrete-element modeling calculation with tomography results. Reproduced from Gan, Y.; Kamlah, M.; Reimann, J. Fusion Eng. Des. 2010, 85, 1782–1787. Ceramic Breeder Materials 487 � Atomic displacements are induced by neutron impact, most effectively by fast neutrons, causing significant lattice defects, damage, swelling, and so on. � Generated tritium affects the physicochemical and mechanical behavior, and the inert, 3He decay product stays within the material if it is not desorbed earlier into the purge gas flow. � Generated helium resides in the material, and through formation of clusters of bubbles will generate stresses leading to macroscopic effects such as swelling, altered thermal transport, and/or fracture. Numerous irradiation experiments have been per- formed on ceramic breeder pebbles using material test reactors with thermal, mixed, or irradiation phe- nomena fast neutron spectrum. As the 6Li cross- section, in particular, is much higher for thermal neutrons, many of the irradiation phenomena can be effectively studied in thermal or mixed-spectrum reac- tors, that is, without using 14MeV neutrons. This section concentrates on the thermal–mechanical behavior, while tritium production and transport are dealt with in the following section. High fluence and high lithium burnup were achieved in fast reactor irradiations at experimental breeder reactor II (EBR-II) and fast flux test facility (FFTF) facilities, as reported byHollenberg and cow- orkers.28,134 The bulk of these experiments concerned Li2O, LiAlO2, and LiZrO2, with only a few data on Li4SiO4 shaped as annular pellets and LiZrO3 peb- bles. Good tritium release behavior of Li2O and LiZrO3 has been reported, even for temperatures higher than 1000 �C; pellet thermal conductivity of Li2O and LiAlO2 was decreased at lower irradiation temperatures but appeared fairly unaffected when operated over 400 �C.28,134,163,165,166,186–188 AECL tests at NRU and JAEA tests at JMTR addressed the impact of neutron irradiation on pebble-bed properties, such as conductivity. In these cases, the constraints were modest: either higher 488 Ceramic Breeder Materials burnup with few pebbles in the heat flow direction or low volumetric heat loads and low lithium burnup.135–139 Verrall et al.138 seem to the first to report on bubble formation in a ceramic breeder. They observed this in Li2O irradiated in NRU up to 1% lithium burn-up up, see Figure 45. The effective thermal diffusivity of a Li2TiO3 pebble bed was studied in an in-pile irradiation experiment by Kawamura et al.140 at the JMTR test reactor. The cylindrical assembly of the Li2TiO3 pebble-bed was instrumented with a number thermo- couples to determine the radial temperature profile. The derived thermal diffusivity as function of temperature is shown in Figure 36. The dimension of the pebble bed was 20mm in diameter and 260mm in length. The effective thermal diffusivity of the Li2TiO3 pebble bed was found to decrease with increasing irradiation temperature. This tendency remained up to a thermal neutron fluence of 1� 1024 nm�2, while the thermal conductivity at given temperatures also remained constant. None of these experiments addressed the pebble- bed deformation behavior at the strong temperature gradients envisaged for breeding blankets. Out- of-pile testing is less representative as it requires the use of heater plates, with their specific impact on reduced bed thickness, and running a ceramic- heater interface at the highest temperature. Also, when material is irradiated in a stressed state, there is an additional phenomenon of irradiation- induced creep. Reactor power 6 4 2 0 400 450 500 Temperature of point C (�C) E ffe ct iv e th er m al d iff us iv ity (� 10 -3 cm 2 s– 1 ) 550 600 650 Sweep gas flow rate Temp. raising rate : 50 MW : 4 �C min : 200 cm3min–1 (�1023n m–2) : Fth � 8 : Fth � 10 Figure 36 The effective pebble-bed conductivity derived from by Kawamura and coworkers.140 Reproduced from Kawamura, 2003, 69, 263–267. As a major step in the preparation of the European HCPB TBM program in ITER, an in-pile test of pebble-bed assemblies was defined. This experiment was designed to address the neutron-irradiation effects on the thermal–mechanical behavior of a breeder pebble bed at HCPB DEMO representative levels of temperature and defined thermal–mechanical loads.121,130,141 A schematic is given in Figure 37. The core of each test element is a horizontal cylindri- cal bed of ceramic breeder pebbles, either Li4SiO4 (OSi) or Li2TiO3 (MTi), with an outer diameter of about 45mm and bed thickness of about 10mm, sand- wiched between two beryllium pebble beds. The breeder and beryllium pebble beds are separated by Eurofer-97 steel plates. The heat flow is managed so as to have a radial temperature distribution in the ceramic breeder pebble bed as flat as reasonably possible. The test element design and test matrix required extensive pretesting, improved pebble-bed modeling, design curves for the material character- istics, and performance analyses allowing in-reactor operation in High Flux Reactor (HFR) Petten121,130,141 (Figure 38). A specific pretest compaction procedure by press- ing and heating has been developed, in particular to increase the thermal conductivity of the beryllium pebble beds and provide conditions that would result in limited changes during in-pile operation.130,141 The compaction procedure consisted of a subsequent loading of the pressure plate of the total assembly to 3MPa. The X-ray pictures combined with the Reactor power : 50 MW 30 25 20 15 10 5 00 2 4 6 8 10 12 Thermal neutron fluence (�1023n m–2) TA : Temperature of point A (400 �C) TB : Temperature of point B ΔT = TA-TB Te m p er at ur e d iff er en ce (� C ) in-pile Li2TiO3 pebble-bed irradiation in JMTR, as performed H.; Kikukawa, A.; Tsuchiya, K.; et al. Fusion Eng. Des. HCPB pebble-bed assembly Purge gas lines 2nd containment: AISI-316L structure 1st containment: Eurofer-97 structure Thermal barrier 2: Eurofer-97 Thermal barrier 1: Inconel718 Thermocouple tubes: AISI 321+Pt-clad in breeder zone Beryllium pebble-bed Beryllium pebble-bed Self-powered neutron detector Neutron dosimeters Floating plate: Eurofer-97 Ceramic breeder pebble-bed AI filler AI filler Sealing plate: all Eurofer-97 Threaded ring, Pressure plate, Figure 37 Schematic of HCPB pebble-bed assembly (PBA) test-element for in-pile testing in the High Flux Reactor at Petten, The Netherlands. Each test-element has a cylindrical ceramic breeder section, either Li4SiO4 or Li2TiO3, in between two cylindrical shaped beryllium pebble beds and separated by Eurofer plates. Figure 38 Picture of PBA test element during assembly: Ceramic breeder section (Li2TiO3 pebbles from CEA) with penetrating thermocouple tubes; note pebble alignment at circumference. Ceramic Breeder Materials 489 dimensional measurements taken during assembly allowed the determination of actual bed size and compaction values. An example of a postassembly X-ray picture is shown in Figure 39. The design and safety requirements for in-pile operation required the development of a full-coupled thermomechanical model in theMARC finite-element code. In this way, the pressure buildup and stress relax- ation in the pebble beds could be simulated in detail to guide the required reactor startup profile.121,130,141 The two test elements with Li4SiO4 pebbles were irradiated at nominal temperatures of 600 and 800 �C in the breeder bed, to see any effects of thermal creep. The other two test elements contained Li2TiO3 peb- bles with different grain sizes and were irradiated at the same temperature, the nominal temperature of 800 �C. The pebble beds were typically purged with a helium–hydrogen mixture of reference composition (0.1% H2). The gas purge entered at the lower beryl- lium bed and exited at the upper beryllium beds. The PBA has been irradiated for 294 FPD (full power days), achieving lithium burnups of 1.5–2.2% for Li4SiO4 and 2.8–2.9% for Li2TiO3, without enriching in 6Li. The dose in the Eurofer-97 parts ranged from 2 to 3 dpa.142 Extensive analyses of the in-pile data using FEM calculations showed that the maximum Purge line splitter After inspection of assembled test- element: Stepped precompaction, 1, 2, 3 MPa plus 24 h at 350 �C Last inspections prior and after sealing. Large pebble: low density near wall Pebbles well aligned with plate Pt clad on TC- tubes Figure 39 X-ray picture of PBA as assembled. Figure 40 Neutron radiograph of PBA taken after few irradiation cycles. 490 Ceramic Breeder Materials temperatures in the Li4SiO4 pebble beds of the top and bottom test-elements are about 600 and 800 �C, respectively. The average temperatures in the Li4SiO4 beds are about 550 and 740 �C, respectively.142 In postirradiation examinations of both Li4SiO4 samples, a little sintering and a significant amount of cracking or fragmentation have been observed. No significant difference between the lower and higher temperature case was found. Most of the evidence of cracking and fragmentation in the Li4SiO4 pebbles is observed toward the middle of the bed (highest temperature, highest deformation). This is visible in scanning electron microscopy images (Figure 42). There is some evidence of grain growth. Reactions of Li4SiO4 pebbles with Eurofer were found to be very small. The maximum temperatures in the Li2TiO3 pebble beds of both the two test-elements in the middle are about 780 �C. The average temperatures in the Li2TiO3 beds are about 690 and 720 �C, respectively.142 In both test-elements, Li2TiO3 pebbles showed a significant amount of sintering and necking, which was found most significant in the test-element #3. The average temperature of this test-element was higher by about 35 �C. Almost no fracture or frag- mentation was seen. There appeared to be a small reaction layer, distributed uniformly along the Euro- fer (Figures 43–44). In the high burnup irradiation EXOTIC-7 (see details in Section 4.15.5.1.2), the pellet stacks and pebble beds were found to be essentially intact by neutron radiography analyses after irradiation, and except for one capsule containing Li2ZrO3 pellets, three out of five were found intact after unloading. Fragmentation of the 0.1–0.2mm Li4SiO4 Ceramic Breeder Materials 491 pebbles was also observed but was very difficult to quantify33,143–145 (Figure 46). The pores observed in the images are related to the pebble fabrication method rather than to neutron irradiation. Chikhray et al. irradiated Li2TiO3þ 5mol% TiO2 in the Kazakhstan water water research reactor (WWRK) reactor to lithium burnups of about 20% at 760 and 920K; see details in the next section and Chikhray et al.82 Pebble crush tests showed reduc- tion of strength, whereas microhardness tests also revealed ingrowth of soft phases. X-ray diffraction measurements showed traces of LiTi2O4, LiTiO2, and Li4Ti5O12; see Chikhray et al. 82 for more details. (a) ( Figure 42 Optical micrographs of cracked Li4SiO4 from PBA Temperature 827 727 627 527 427 327 227 127 Test module #4 (�C) Figure 41 Temperature fields in the PBA as calculated with NRG’s fully coupled (cylindrical) pebble-bed thermo-mechanical model. In an IEA-framed international collaboration, European Li4SiO4 and Li2TiO3 and Japanese Li2TiO3, reference pebble materials were tested in a high fluence irradiation project at the HFR in Petten, named high neutron fluence irradiation of pebble stacks for fusion (HICU).146–148,194 The neutronic analyses as reported by Fischer and coworkers149,150 demonstrated that relevant nuclear irradiation para- meters such as the displacement damage accumula- tion, the lithium burnup, and the damage production function W(T) are met with the selected neutron shielding and 6Li enrichments chosen. This project is conceived to irradiate ceramic breeder pebble stacks at high temperatures under blanket prototypical ratios of fast neutron damage (dpa) and lithium burnup. Compared with the PBA experiment, the pebble stacks are smaller, but capsule dimensions are up to about 10mm; X-ray tomography was used for detailed mapping of pebble location prior to irradiation.148,151 4.15.5 Tritium Production and Release 4.15.5.1 Tritium Release Awide range of mechanisms play a role in the tritium transport and release processes of the lithium- containing ceramics, of which an impression is given in Figure 48.152 Tritium generated from neutron cap- ture is first transported to the grain boundary by bulk diffusion. The bulk diffusion and trapping inside the grains are affected by the neutron radiation-induced defects. Via the intergranular diffusion, the tritium is then delivered to the grain surfaces, which are exposed to open and closed porosity. The closed porosity frac- tion provides another means to build up inventory in b) postirradiation examinations. (b)(a) Figure 43 Optical micrographs of sintered Li2TiO3 from PBA postirradiation examinations. Figure 44 Optical micrograph on Li2TiO3 interaction with structure from PBA postirradiation examinations. 1 mm Figure 45 Scanning electron micrograph showing small bubbles, randomly distributed, amid larger bubbles in Li2O after irradiation in NRU.138 492 Ceramic Breeder Materials the material. At the surface isotope exchange with hydrogen (H2) and water (H2O) lead to desorption of tritium in molecular forms of HT and HTO, respec- tively. Further, the tritium in molecular form is trans- ported through the interconnected pores and enters the flow of the purge gas. In order to assess the tritium retention in the candidate ceramic breeder material, one needs to know which of the steps are rate deter- mining and which operation parameters are the most relevant for facilitation of the tritium release (Table 3). The tritium release characteristics of lithium cera- mics are typically studied in two parameter ranges: 1. Out-of-pile: Tritium production through exposure to neutron irradiation, followed by out-of-pile tri- tium desorption through stepwise isothermal or ramp annealing tests in laboratory setups, also known as temperature programmed desorption (TPD). If irradiation doses are very low, such activity is typically called ‘tritium doping,’ and tritium transport parameters reflect beginning of life (BOL) conditions only because irradiation damage and lithium burnup remain negligible. 2. In-pile: In case of in-pile experiments, typically steady-state tritium production and release condi- tions. In general, such parameters are closer to breeding blanket conditions, as they allow the application of a wide range of temperatures and purge gas conditions, and the study of long-term performance issues such as irradiation damage and lithium burnup. At present, such data are limited in terms of fast neutron damage doses (thermal and mixed spectra materials test reactor (MTR) only). 4.15.5.1.1 Out-of-pile In the majority of cases, lithium ceramic samples are irradiated in research reactors with thermal or mixed neutron spectra. Extensive studies on tritium retention in neutron irradiated lithium ceramics using the TPD Ceramic Breeder Materials 493 method have been reported in the literature.152–178,201 The chemical form of released tritium from Li4SiO4 (from FZK), LiAlO2 (from JAERI), Li2TiO3 (from CEA), and Li2ZrO3 (from MAPI) was studied in the out-of-pile tritium release experiment under various purge gas conditions. The pebbles were irradiated for a few minutes in the fluxes 4� 1017m�2 s�1 at Japan Research Reactor 4 ( JRR-4) or 1.65–2.75� 1017m�2 s�1 at the Kyoto University Reactor. The tritium was Figure 46 High burn-up Li4SiO4 pebble irradiated to 11% lithium burn-up, with fracture features, and large pores that originate from the manufacturing process. Near DEMO blanket 1 FPY 0 0 10 20 30 40 50 2 4 6 Total lithium bu D is p la ce m en t d am ag e in L i 4 S iO 4 (d p a) 0.06 % unshielded 7.5% shielded Figure 47 Irradiation parameter space projected to be covere burn-up, for Li4SiO4 with 7.5 and 20% Li-6, in a shielded case, a strong effect of reduced burn-up on the dpa/burn-up ratio (see released in the purge gas of dry nitrogen, nitrogen with 0.1% of helium, and nitrogen with 0.1% water vapor. Even if hydrogen was added to the purge gas, a considerable fraction of tritium was released in the molecular form of water, HTO. Addition of water vapor to the purge gas greatly enhanced the release rate of tritium. It was concluded that the isotope exchange of tritium with water at the exposed surfaces of the grains is much faster than the isotope exchange of tritium with hydrogen. The water is adsorbed at the grain surfaces from the water vapor present in the purge gas. Even small traces of water of 30 ppm in dry purge gas can be enough to promote tritium release in the HTO molecular form. Another source of water at the surfaces of the grains is the reduction reaction that takes place in theH2 reducing atmosphere. Among the studied materials, Li2TiO3 showed the largest water formation capacity.161 Figure 49 shows out- of-pile annealing tests for Li4SiO4 with 0.1% H2/N2 sweep gas and 0.1% H2O/N2 sweep gas, and tritium doped for 2min at a neutron flux of 2.75� 1013 cm2 s�1.158 Later work on palladium deposited as catalyst on orthosilicate indicated that almost all tritium was released as tritiated water vapor from lithium ortho- silicate pebbles and tritium at higher temperatures remains slow.160 In contrast, it was also found that a considerably larger amount of tritium was released as the molecular form (HT) from the lithium first wall Near back plate 7.5% unshielded EXOTIC 7, 8 rnup in Li4SiO4 (%) 8 10 12 20% shielded DEMO blanket 20 000 h d by HICU experiment: dpa accumulation versus lithium nd 0.06% and 7.5 Li-6 in the unshielded case, revealing the also section 4.15.6.1 on irradiation damage). 494 Ceramic Breeder Materials orthosilicate pebbles already deposited with palla- dium at lower temperatures (see Figure 50). Alvani and coworkers172,173 correlated TPD after short irradiations in the Casaccia reactor ( 2� 1021m�2) with those from long-term irradia- tion in the HFR; Petten (thermal 0.5� 1025m�2) revealed two peaks, at 770 and 941K (b¼ 5Kmin�1) (see Figure 51). It was proposed that the second peak is related to tritium trapping at the oxygen vacan- cies located along the grain boundary interface. The concentration of these trapping sites is signifi- cantly increased by the reduction effect of the R-gas, which results in a shift in the release peak to higher temperatures for the pretreated pebbles. The observed effect is more pronounced for the pebbles with finer Table 3 Comparison of surface reactions on ceramic breed Nishikawa, M.; Kinjyo, T.; Nishida, Y. J. Nucl. Mater. 2004, 325, Dry purge gas Purge 573K Adsorption/desorption Adsor 573–473K Adsorption/desorption Adsor Water Isotop Isotop 773–973K Adsorption/desorption Adsor Water Isotop Isotop Surfac 973K Adsorption/desorption Adsor Water Isotop Isotop Surfac Heat Defects L i2O + T2O ) He LiOT LiOT LiOH LiOT H2O H2 Vapor purge Neutron 7Li(n, n*a)T 6Li(n, a)T HT T T T T2O HTO HTO H2O Solid breeder grainBreeder pebble Dry He purge Hydrogen purge (LiOT + H2O LiOH + HTO) (LiOT + H2 LiOH + HT) (2LiOT (LiOT + LiOH Li2O+ HTO) Figure 48 Schematic of tritium transfer phenomena through the ceramic breeder structure. Reproduced from Nishikawa, M.; Kinjyo, T.; Nishida, Y. J. Nucl. Mater. 2004, 325, 87–93. grains. It was also found that a thermal pretreatment at 473K for 2 h removes only the environmental H2O and CO2 contamination from the surface of the pebbles without affecting the H2O desorption at higher tem- peratures. The observed H2O release above 1173K is ascribed to the reduction of Li2TiO3 to Li2TiO3�x, with x reaching a steady-state value of xeq¼ 0.01. Later work by Casadio et al.174 concerned a batch of Li2TiO3 pebbles (ENEA code FN5) prepared following the ‘citrate’ route. Analysis of TPD spectra gave the correct order of magnitude of the time constants characterizing the main desorption sites, in rough agreement with the residence times obtained by the in-pile step-perturbation methods performed during EXOTIC-8/9 experiment (see Figure 52).175 Pure helium purge increases the tritium inventory; during the last cycle of this irradiation experiment, variations of the H2 concentration in the He purge showed an increase in tritium release rate from Li2TiO3 pebbles that was found to be proportional to P0:34H2 at 473 �C, Figure 53. The effect of open and closed porosity on tritium release behavior in Li2O single crystal and sintered pellets was studied by Tanifuji et al.176–179 The pellets had densities in the range of 70–98% and grain sizes from 10 to 60mm. Irradiations were performed by ther- mal neutrons in JRR-4, JAERI, up to 2� 1023m�2. The porosity dependence of tritium release behavior from the Li2O sintered pellets has been investigated through isothermal heating tests, and the results are shown in Figure 54. For 88% TD specimens irradiated up to neutron fluences of 2� 1022 and 2� 1023 nm�2, no er grains associated with tritium release. Reproduced from 87–93. gas with hydrogen Purge gas with water vapor ption/desorption Isotope exchange 2 Adsorption/desorption ption/desorption Isotope exchange 2 formation Adsorption/desorption e exchange 2 e exchange I ption/desorption Isotope exchange 2 formation Adsorption/desorption e exchange 2 e exchange I e condition change ption/desorption Isotope exchange 2 formation Adsorption/desorption e exchange I e exchange 2 e condition change Ceramic Breeder Materials 495 irradiation effects on the tritium residence time have been observed. 4.15.5.1.2 In-pile testing In-pile experiments have the strong advantage that the tritium release characteristics can be studied, as a function of neutron damage and lithium burnup in combination with thermal–mechanical behavior under neutron irradiation. Such in-pile experiments allow steady-state tritium production and release con- ditions to be achieved. In general, such parameters are closer to breeding blanket conditions, as theyallow the application of a wide range of temperatures and purge gas conditions and the studyof long-termperformance issues such as irradiation damage and lithium burnup. At present, such data are limited in terms of fast 0.04 0.03 Breeder:Li4SiO4 Irradiation:12.0 min Sweep gas:0.1% H2/N2 Flow rate:100 ml min–1 (0.05% H 2 +0.5 % H 2 O)/N 20.02 Released HTO + HT Released HT0.01 0.00 0 1 2 3 4 Time (h)(a) Tr iti um c on ce nt ra tio n (m C ic m -3 ) Te m p er at ur e (� C ) 5 6 7 0 200 400 600 800 1000 (b Figure 50 Out-of-pile annealing experiment of (a) Li4SiO4 and Shinozaki, T.; Inoue, K.; et al. J. Nucl. Mater. 2009, 386–388, 10 Water vapor Background Time (h) 0 0.000 0.001 0.002 0.003 Tr iti um c on ce nt ra tio n (m C ic m -3 ) 0.004 0.005 1 2 3 4 0 200 Te m p er at ur e (� C ) 400 600 800 1000 Temperature 0.1% H2O/N2 0.1% H2/N2 Figure 49 Out-of-pile annealing tests for Li4SiO4 with 0.1% H2/N2 sweep gas and 0.1% H2O/N2 sweep gas (amount of breeder, 0.3 g); flow rate, 100mlmin�1; irradiation time, 2min; neutron flux, 2.75� 1013 cm2 s�1. Reproduced from Munakata, K.; Yokoyama, Y.; Baba, A.; Penzhorn, R. D.; Oyaidzu, M.; Okuno, K. Fusion Eng. Des. 2005, 75–79, 673–678. neutron damage doses (thermal and mixed spectra MTR only). Several irradiation programs have been executed around the world, involving thermal, mixed- spectrum, and fast reactors [as in Dido (Germany), Siloe (France), EBR-2 (US), FFTF (US), HFR (The Netherlands), JMTR (Japan), and WWRK (Kazakhstan)]. The European irradiation projects under the acronym EXOTIC, for extraction of tritium in ceramics, commenced during the mid 1980s at the HFR in Petten26,33,63,144,175,180–183,195 (see Table 4). The initial series concerned both closed capsule and vented capsule operation. The materials investigated were mainly Li2O, Li2SiO3, LiAlO2, LiZrO3, Li8ZrO6, and Li4SiO4 in the series EXOTIC-1 to-6. 26 The objective of the EXOTIC-7 experiment has been to irradiate candidate ceramic breeder materials in the HFR to a high lithium burnup (target 10%) and to determine the effects on the mechanical integ- rity of pellets and pebble-bed configurations, and those on tritium-inventory and-release characteris- tics.33 The experiment concerned 8 capsules, and during 11 HFR cycles (261 FPD), lithium burnups of 6–18% were achieved. The test matrix comprised pellets of Li2ZrO3, Li8ZrO6, and LiAIO2 and pebbles of Li2ZrO3 and Li4SiO4. Two capsules contained a mixture of Li4SiO4 and beryllium pebbles. To obtain a high lithium burnup within a reasonable irradiation time, the target materials were enriched with 6Li to about 50%. The tritium residence time of Li2ZrO3 pellets a higher temperatures was not changed during the irradiation. After a lithium burnup of about 5%, the residence time at lower temperatures decreases significantly with increasing burnup. Breeder:0.2%Pd/Li4SiO4 Irradiation:12.0 min Sweep gas:0.1% H2/N2 Flow rate:100 ml min–1 (500 ppmH 2 +5000 ppmH 2 O)/N 2 Released HTO + HT Released HT 0.04 0.03 0.02 0.01 0.00 ) Time (h) Tr iti um c on ce nt ra tio n (m C ic m -3 ) Te m p er at ur e (� C ) 0 1 2 3 4 5 6 7 0 200 400 600 800 1000 (b) Pd-coated Li4SiO4. Reproduced from Munakata, K.; 91–1094. He 100 10 1.2 1T rit iu m r es id en ce t im e (h ) 1.3 1000/T (K−1) 1.4 1.5 1.6 Figure 52 Arrhenius plot of the tritium residence time for Li2TiO3 (FN5) pebbles irradiated in EXOTIC-8/9 experiment under reducing atmosphere; the black point was obtained in pure helium purge gas174. Reproduced from Casadio, S.; van der Laan, J. G.; Alvani, C.; Magielsen, A. J.; Stijkel, M. P. J. Nucl. Mater. 2004, 329–333, 1252–1255. 1 1.0 ceo_clean 0.8 0.6 (a .u .) 0.4 0.2 0.0 400 600 800 Temperature (K) 1000 TRIGA casaccia lazy susan capsule Tirr about 330 K HFR petten EXOTIC 8.1 capsule Tirr about 800 K0 400 (a) (b) 600 800 1000 Temperature (K) N or m al iz ed d es or p tio n ra te (C n) 1200 Figure 51 (a) Normalized temperature programmed desorption spectra of tritium release from CEA pebbles after tritium doping in TRIGA (circles) and medium-term irradiation HFR-EXOTIC-8/1 (squares), 5 Kmin�1 in both cases, and (b) image of the deconvolution-fit analysis performed on the spectrum in TRIGA reactor. Reproduced from Alvani, C.; Casadio, S. T.; Casadio, S. Fusion Eng. Des. 2003, 69, 275–280. 100 10 1 10 100 1000 H2 concentration in purge gas (vpm) Tr iti um r es id en ce t im e (h ) 10 000 Figure 53 Tritium residence time in Li2TiO3 (FN5) pebbles at 473 C as a function of H2 concentration in the helium purge. The black point results from the best fitting line obtained for the reference composition176. Reproduced from Casadio, S.; van der Laan, J. G.; Alvani, C.; Magielsen, A. J.; Stijkel, M. P. J. Nucl. Mater. 2004, 329–333, 1252–1255. 496 Ceramic Breeder Materials After selection of the HCPB as the single solid breeder concept in the European Blanket Project, the EXOTIC-8 and-9 series concentrated on Li4SiO4 and Li2ZrO3 pebbles and a range of Li2TiO3 pro- ducts.175 The irradiation test program concentrated on two types of experiments: 1. Tritium release to low or moderate lithium burnups 2. High lithium burnup and mechanical integrity The typical designs for these experiments are given in Figure 55, with the general layout and an example of a cross-section from postirradiation testing. Figure 56 shows a sample temperature and tritium quantities in the purge gas for a typical irradiation cycle. Tritium residence time (t): t ¼ I G 1000/T ( K-1 ) 1.31.2 0.1 1 10 (c) (d)(b)(a) 100 1000 104 2 � 1023 n m -2 Ea = 138 kJ mol-1 Ea = 184 kJ mol-1 Ea = 160 kJ mol-1 2 � 1022 n m -2 4 � 1020 n m -2 4 � 1020 n m -2 105 1.4 88% TD As-irrad. A ve ra ge r es id en ce t im e (m in ) As-irrad.Preadsorbed HTO 81% TD 81% TD 1.5 1.6 1.7 1.8 1.9 2 Figure 54 Arrhenius plots of the average residence time t. Reproduced from Tanifuji, T.; Yamaki, D.; Jitsukawa, S. Fusion Eng. Des. 2006, 81, 595–600. Ceramic Breeder Materials 497 I, tritium inventory (Bq); G, tritium production rate (Bqmin�1). At steady state : Tritium release rate (R)¼Tritium generation rate (G) Temperature transients are performed: DT Difference in tritium inventory (area): DI¼ I2�I1 Difference in residence time: Dt¼DI/G Data set is processed to obtain: t(T) See Figure 57. In the EXOTIC-8 program, the tritium release characteristics and mechanical stability of the refer- ence breeder materials for the European HCPB project have been studied, including the effect of long-term neutron irradiation and high lithium burnup.175 The EXOTIC-8 program started in 1997 and ended in 2002. It consisted of ten experi- ments, named EXOTIC-8/1 to EXOTIC-8/10. The irradiations were carried out in the HFR in Petten in peripheral core positions with the typical neutron fluence rate of about 9� 1017m�2 s�2 (fast, En> 0.1MeV) and 5� 1017m�2 s�1 (thermal). The following materials were used: Li2TiO3 peb- bles produced by agglomeration – sintering and extrusion – sintering, provided by CEA; Li2TiO3 pellets produced by cold pressing, provided by CEA; Li2TiO3 pebbles produced by wet processing and sintering, provided by ENEA; Li2ZrO3 pebbles produced by extrusion – sintering process, provided by CEA; and Li4SiO4 pebbles produced by melt spray process, provided by FZK Karlsruhe. Two types of experiments were targeted: 1. Major focus on tritium release characteristics by determining differential tritium inventories by thermal transients and achieving low to medium lithium burnups of 1–3%. 2. Major focus on high lithium burnup experiments using pebbles with 50% 6Li enrichment and achieving 11% lithium burnup for Li4SiO4 and 17% for Li2TiO3, at relatively constant tempera- tures, and only few data on release, mostly from postirradiation annealing tests. Tritium release characteristics are measured both in situ by applying the temperature transients and after irradiation in the TPD setup. The tritium release is characterized by the tritium residence time t with the Arrhenius temperature behavior, as depicted in the summary graph; Figure 58. A correlation has been established between the pebble density and the measured residence time. The obtained results confirm the understanding that open porosity and small grain size are favorable for faster tritium release. The corresponding activation energies derived from the temperature transients and from the TPD measurements are in fair agreement. For Li2TiO3 pebbles, Q¼ 82–93 kJmol�1 and for Li4SiO4 peb- bles, Q¼ 112–123 kJmol�1. The activation energy for Li2ZrO3 pebbles is derived only from the temperature transients: Q¼ 84 kJmol�1. In TPD experiments, it was shown that the tritium release can involve multiple release processes. Moreover, in some in- stances the release can be limited by recombination at the grain surface, which introduces uncertainties in the measured values of the activation energy. The results of in-pile tritium behavior during nor- mal operation and during transients in temperature and gas chemistry measured in the latest irradiation experiment from the EXOTIC series, EXOTIC-9/1, are reported in Peeters et al.180 The Li2TiO3 pebbles produced by extrusion–spheronization sintering at CEA were irradiated in the HFR in Petten (thermal 0.5� 1018m�2 s�1) for 301 FPD and achieved a burnup of 3.8–4.1%. The temperature varied between 613 and 853K. Based upon the in-pile tritium release measurements and the analysis of the tritium resi- dence time, it was concluded that tritium release in the new batch of the high-density Li2TiO3 pebbles (93.0%TD) is rather slow comparedwith the ceramics irradiated in the EXOTIC-8 irradiation campaign.174 Thus, the tritium residence time measured at 773K in the EXOTIC-9/1 experiment was 30 h, whereas Table 4 Overview of ceramic breeder irradiation experiments performed in fast or mixed spectrum fission reactors to relatively high levels of thermal and/or fast neutron fluence. Values for the experimental parameters quoted were taken from the mentioned references. Readers are also referred to the text throughout the whole chapter. Experiment ID Materials Shape Li-6 (%) Li-6 burnup (%) Total Li burn up (%) Fluence thermal (1020 cm�2) Fluence E > 1 MeV (1022 cm�2) EFPD (Days) Temp (�C) Reference FUBR-IA Li2O, LiAlO2, Li2ZrO3, Li4SiO4 pellets 56, 95 1.1–1.2 1.4–1.45 96.7, [186] 2.0–2.2 3.9–4.1 177.6, 3.1–3.5 2.5–2.7 274.3 FUBR-IB Li2O, LiAlO2, Li2ZrO3, Li4SiO4 pellets 0.07,7.5,56,95 3.9–4.3 4.7–5.1 341.5 450–1225 [186] BEATRIX-I Li2O, LiAlO2, Li2ZrO3, Li4SiO4 pellets 0.07,7.5,56,95 6.7–7.4 8.2–8.9 599.3 [186] 10.3–11.3 13–14 940.8 10.3–11.3 13–14 936.9 MOTA-2A (BEATRIX-II) Li2O, LiAlO2, Li4SiO4, Li2ZrO3 rings, pebbles 0.07, 0.2, 1.7, 7.6–9.9 0.2, 7.4–8.8 299.7 550–640 [186] 7.5, 440–1000 56, 95 MOTA-2B (BEATRIX-II) Li2O, LiAlO2, Li4SiO4, Li2ZrO3 rings, pebbles 0.2 1.27, 5.3–7.9 0.2, 7.7–7.9 203.3 530–640 [186] 85 440–1100 95 CRITIC-3 Li2TiO3 pebbles 1.85 0.9 0.57 334 200–900 [208] COMPLIMENT (ELIMA-2, DELICE-3) LiAlO2, Li2SiO3, Li4SiO4, Li2O, Li2ZrO3 pellets, pebbles 7.5 0.25 178 (HFR) 400–450 [209] 1 75 (OSIRIS) 650–700 In-Pile Pebble- Bed Li2TiO3 pebbles 0.01 JMTR 400–610 [140] EXOTIC-1 (BEATRIX) LiAlO2, Li2SiO3, Li2O pellet 0.06 0.004 0.01 0.013–0.025 25 350–700 [26],[210] 0.6 0.035 7.5 0.28 EXOTIC-2 LiAlO2, Li2SiO3, Li2O pellet 0.55–0.6 0,1 0.02–04 0.025–0.05 50 350–725 [26],[210] EXOTIC-3 Li2ZrO3, Li2O, Li2SiO3 pellet, solid 0.55–0.6 0.12–0.13 0.03–0.06 0.03–0.08 75 385–650 [26] EXOTIC-4 Li2ZrO3, Li6Zr2O7, Li8ZrO6, Li2O, Li2SiO3 pellet, solid 0.55–0.6, 7.7 0.13–0.15, 0.8 0.04–0.07 0.05–0.1 97 310–680 [26] EXOTIC-5 LiAlO2, Li2ZrO3, Li4SiO4, Li6Zr2O7, Li8ZrO6 annular pellets and pebbles 7.5 1.8–2.1 0.07–0.13 0.08–0.17 166 325–630 [26] 4 9 8 C e ra m ic B re e d e r M a te ria ls EXOTIC-6 LiAlO2, Li2ZrO3, Li6Zr2O7, Li8ZrO6, Li4SiO4 annular pellets and pebbles 7.5 3.0–3.2 0.08–0.15 0.1–0.2 199 315–520 [26] EXOTIC-7 Li2ZrO3, Li8ZrO6, Li4SiO4, LiAlO2 annular pellets and pebbles 38–50 5.8–18.1 0.27 0.13 261 410–745 [33] EXOTIC-8/1 Li2TiO3 pebbles 7.5 1.86 0.046 201 310–570 [175] EXOTIC-8/2 Li2TiO3 annular pellets 7.5 1.34 0.039 173 420–640 [175] EXOTIC-8/3 Li4SiO4 + 2% TeO2, Be pebbles 50 2.6–4 0.003–0.013 201 200–700 [175] EXOTIC-8/4 Li4SiO4 + 2% TeO2 pebbles 7.5 1.38 0.033 201 280–590 [175] EXOTIC-8/5 Li2TiO3 pebbles 7.5 2.1 0.06 200 500–660 [175] EXOTIC-8/6 Li2ZrO3 pebbles 7.5 2.5 0.071 299 400–640 [175] EXOTIC-8/7 Li4SiO4, Be pebbles 51.5 8–17.5 0.025–0.064 648 400–690 [175] EXOTIC-8/8 Li2TiO3, Be pebbles 50 4.1–10.9 0.11–0.38 450 400–700 [175] EXOTIC-8/9 Li2TiO3 pebbles 7.5 3.51 0.1 448 420–560 [175] EXOTIC-8/10 Li4SiO4 pebbles 7.5 0.09 448 520–695 [175] PBA Li4SiO4, Li2TiO3 pebbles 7.5 1.5–2.9 (for natural 6Li 0.04–0.07 0.11–0.17 294 550–800 [142] EXOTIC-9 Li2TiO3 pebbles 7.5, 30.5 3.1 (for natural 6Li) 0.15 0.09 300 340–580 [143] Li4SiO4 7.5, 20 K-578 (WWRK) Li2TiO3 pebbles, pellets 96 18–23 220 500–900 [212], [82] HICU Li4SiO4 pebbles 0.06, 7.5, 20 0.7–13 1.5 0.7 403 600–900 [211] Li2TiO3 0.06, 11, 30 C e ra m ic B re e d e r M a te ria ls 4 9 9 Four channel irradiation rig External fluence detector Purge gas tube Second containment gas gap Filler element Tritium breeder Central tube First containment Second containment Third cont. gas supply 24.1 mm 37.2 mm (a) Thermocouple (b) Figure 55 Cross section of the EXOTIC series of ceramic breeder irradiation tests: a) schematic drawing, and b) macrograph from postirradiation microscopic investigations. 30 0 0.5 1 1.5 2 Time (days) Tr iti um r el ea se (m C im in –1 ) 0 5 10 15 20 25 0 200 400 600 800 Te m p er at ur e (� C ) ¬ 1 ¬ 2 ¬ 3 ¬ 4 ¬ 5 ¬ 6 ¬ 7 ¬ 8 Figure 56 Examples of tritium release from the EXOTIC-8 experiment series with a number of programmed temperature transients (including reactor power scram at day 19+20). 500 Ceramic Breeder Materials the same characteristic measured on the Li2TiO3 pebbles obtained from CEA and ENEA in the EXOTIC-8 campaign was 1.3 and 5.2 h, respectively (Figure 59). Changes in the tritium inventory resulting from the variation of the H2 concentration in the purge gas (from 0.1% to 1.0%) appeared to be much smaller than those resulting from the temperature transients. From this observation, it was concluded that the tritium inventory was determined by the thermally activated processes taking place in the bulk of the material (dissociation from traps, diffusion) rather than by recombination and isotope exchange with hydrogen at the surface. 104 350 400 450 500 550 600 103 102 101 100 10-1 10-2 104 103 102 101 100 10-1 10-2 1.1 1.2 1.3 1.4 1000/T (K) t (h ) 1.5 1.6 1.7 E-8/1: Li2TiO3 E-8/2: Li2TiO3 E-8/5: Li2TiO3 E-8/9: Li2TiO3 E-8/10: Li4SiO4 E-8/6: Li2ZrO3 E-8/4: Li4SiO4 Figure 58 Summary graph of tritium residence time obtained from the temperature transient experiments for different pebble materials used in the EXOTIC-8 program. Difference in inventory Steady- state release 0.0 0.2 0.4 0.6 0.8 Time (h) Tr iti um r el ea se r at e (m C im in -1 ) 1.0 1.2 1.4 1.6 0.15 0.20 0.25 0.30 0.35 0.40 T1 T2 Steady- state release Figure 57 Determination of differential tritium inventory from a small negative temperature transient in a steady state tritium production/release experiment. Ceramic Breeder Materials 501 Tritium recovery characteristics of a binary bed containing 0.3 and 2mm diameter Li2TiO3 pebbles were studied under continuous (20 h) and pulsed (200, 400, and 800 s) neutron irradiation in the JMTR.185 From the temperature transients from 573 to 623K, the tritium residence time was estimated as 3 h (63.2% of the steady-state value). The complete recovery of the steady-state conditions was achieved after 20 h. The tritium recovery behavior under the pulsed operation was almost the same as under contin- uous operation, except for the modulations introduced by the pulse operation, which did not exceed 20% of the total signal variation (Figures 60–62). Effects of irradiation temperature, purge gas flow rate, and hydrogen content in the purge gas on the tritium release characteristics of the Li2TiO3 pebbles were studied in the in-pile irradiation experiment in the JMTR.185 The Li2TiO3 pebbles were fabricated 0 107 108 109 50 723 773 773 673 773 673 [K]723 100 Elapsed time (h) 150 To ta l t rit iu m r el ea se (B q m in -1 ) 200 : 200 cm3min-1Sweep gas flow rate Moisture concentration Hydrogen content in sweep gas : Ceramic Breeder Materials 503 of ceramic samples were examined simultaneously using a system for in-pile tritium monitoring: one (pebbles) – under constant temperature of 650 �C, and two (pebbles and pellets) – within temperature change ranges from 500 to 900 �C. Lithium burnup reached 23% for the active ampoule pebbles, 20% for the passive ampoule pebbles, and 18% for the pellets. The tritium measurement system permitted the 100 90 80 70 60 50 40 30 20 10 0 0 5 Lithium b Li2O (15 at.%) HCPB DEMO blanket : d p a 1 FPY 2 FPY 4 FPY Figure 63 Displacement per atom (dpa) accumulation versus Li2TiO3 in a helium-cooled pebble bed DEMO blanket, as analyz breeder zone as a function of distance from the plasma-facing w Hydrogen partial pressure (Pa) O ve ra ll ra te c on st an t of t rit iu m d es or p tio n (B q m in -2 ) 100 105 106 107 108 101 102 103 104 : 1E+1 Recycling limit Hands-on-limit 1E+0 Time after shutdown (years) C on ta ct d os e ra te (S v h- 1 ) 1E+2 1E+4 1E-1 1E-3 1E-5 1E-7 1E-4 1E-2 Former Li4SiO4 Li4SiO4 pebbles 6Li Li4SiO4 Pure Li4SiO4 Figure 64 Contact dose rate versus time after shutdown of pure Li4SiO4 without impurities, and the main contributors. Reproduced from Knitter, R.; Fischer, U.; Herber, S.; Adelhelm, C. J. Nucl. Mater. 2009, 386–388, 1071–1073. 504 Ceramic Breeder Materials preparation for an IEA-framed international program on high fluence breeder irradiation, Fischer et al.150 performed analyses of lithium burnup and displace- ment damage in Li2O, Li4SiO4, and Li2TO3 for a HCPB-type DEMO blanket design (Figure 63). 4.15.7 Activation and Waste Issues During neutron irradiation of ceramic breeders, not only does the transmutation of Li to T and He take place but other isotopes are formed as well, and the impurities contained will all contribute to the induc- tion of radioactivity from exposure to the neutron irradiation environment. This enables assessment of the feasibility of recycling ceramic breeder material from blanket components having reached EOL con- ditions. Recycling options for ceramic breeder material not only avoid large waste volumes requiring long-term storage but also contribute to resource effi- ciency of valuable constituents such as lithium. From a ceramic breeder perspective, Li2O, Li2TiO3, and Li4SiO4 are more attractive than Li2ZrO3 due to their long-term activation characteristics. Knitter et al.187 assessed contact dose rates of different Li4SiO4 materials for high radiation levels expected from a fusion power reactor for 1 FPY. As an example, Figure 64 displays the contact dose rate versus time after shutdown for pure Li4SiO4, which specifically results from the production of 28Al, 24Na, 7Be, and 26Al. Due to the activity of 7Be, the recycling limit for remote handling (10mSv h�1) and the hands-on limit (10muSv h�1) are reached after Fabrication process/science Grain size Pebble density Pebble size Open porosity Sphericity/roundness % of pebble breakage, i.e., 0.1 mg m–2 i.e., 4–8% Economic, recycling, radioisotope removal Material characterization Out-of-pile pebble bed TM performance and chemical stability In-pile T release at high burnup In-pile PBA tests at 2–3% BU ITER testing DEMO application Figure 65 Outline of a ceramic breeder material development roadmap, prior to Iter testing, as proposed by Ying et al.200 The road connects actual experiments like PBA (pebble-bed assembly), HICU (small pebble stacks), using parameters as fluence (dpa), lithium burn-up/BU, and temperature in order to cover the thermo-mechanical (TM) loadings of a pebble-bed anticipated in a DEMO power plant. Ceramic Breeder Materials 505 blanket component prior to a DEMO, the questions of the choice of materials for the ITERTBM and the definition of a set of requirements (and the related qualification program) to ensure safety, reliability, and test performances become particularly impor- tant. Accordingly, Ying et al.190 proposed a roadmap outlining the necessary development steps for quali- fying and accepting the pebbles for ITER and fusion applications (see Figure 65). For each development step, a set of criteria is presented as a means for initial screening before proceeding to the next evaluation tests to reduce development costs. However, it is important to rec- ognize that ITER conditions (neutron fluence about 1.5�2 orders of magnitude lower) are far from suffi- cient to qualify any specific breeder material to be used in DEMO. Thus, parallel with ITER and subsequent to ITER testing, tests such as HICU or in fusion relevant neutron sources such as International Fusion Materials Irradiation Facility (IFMIF) for any candidate ceramic breeders under typical reactor blanket conditions with relevant nuclear environment are necessary for this purpose. Though a significant R&D effort on ceramic breeder development has already been made and a vast amount of data on material performance have been obtained, the knowledge to date on the limiting factors in blanket designs for long-term operation is still modest. These limitations are addressed here: 4.15.8.1 Microstructure Only a few types of microstructures have been studied until today, preventing clear conclusions regarding their suitability for reactor application. Both for pebble-bed and pellet/block type designs, significant opportunities still exist to tailor stoi- chiometric composition, grain size, and porosity shape and size as well as surface treatments, coatings, and other pre-conditioning treatments. 4.15.8.2 High Burnup Though some high lithium burnup samples have been achieved in experimental programs, the find- ings are not yet conclusive as to whether the 506 Ceramic Breeder Materials ceramic still functions adequately under normal operation, and whether interaction with the blanket structure satisfies, for example, future power plant safety requirements. 4.15.8.3 High Fluence Few data were obtained in fast breeder experiments; the influence of fast neutron damage on the integrity and functionality of the breeder ceramics is yet to be addressed, in particular for the newly developed peb- ble types and compositions. 4.15.8.4 High Temperature The effect of long-term exposure in a high radiation environment, causing damage, induced swelling, with increasing burnup, may also affect the vaporization behavior and even induce lithium transport within the breeder area. 4.15.8.5 Effects of Transients and Off-Normal Conditions Though work has been done already on thermal cycling of pebble beds, there is uncertainty with regard to the effects of disruption and electromag- netic (EM) loads on, for example, pebble reloca- tion, fragment relocation, and ratcheting in case of inclined or vertical beds. Specific attention is required on the change of heat transfer properties and free flow of purge gas. 4.15.8.6 Accident Behavior (Safety and Investment Integrity) In addition to the integrity and safety requirements according to which a blanket component has to be designed in order to be licensed, there is another matter of economics: in case of transient events or accidents not affecting safety, they may prohibit fur- ther operation of the blanket and require replace- ment. Those blanket concepts that are more tolerant to design base and other accidents will be preferred by utilities. In this context, the use of beryllium-based neutron multipliers, the type of coolant, and the tritium issues for the purge and coolant processing units (isotope separation, purification, steam generator operation, toxicity at accidental conditions) also appear to be very important aspects associated with the breeder material use. 4.15.8.7 Development of Tools Thermomechanical codes to describe the interaction of pebble beds with the structural material have achieved significant progress in recent years. For complete TBMs and, even more, DEMO blanket modules, computational codes based on continuum mechanics will be the first choice in the near future. These codes should be quickly developed to enable fairly good predictions for the blanket behavior at the BOL of a DEMO plant or power reactor. The database for irradiated material is still very small and the amount of relevant engineering data in the near future will also be very limited, as material development is still ongoing. Pebble-bed experiments that require a considerable amount of material will be costly and hardly possible for some time. This is the significant challenge for discrete-element methods (DEM). These codes must be improved in order to describe more realistically the interaction between nonirradiated pebbles (taking into account pebble shape, surface condition, and material properties, including thermal creep). If this goal is achieved, it should be possible to do this for irradiated materials, because a small pebble mass is required to make corresponding experiments. The results obtained then with DEM codes should be fed into formal correlations in the continuum codes in order to assess the blanket behavior toward EOL conditions. 4.15.8.8 Compatibility with Structure Though experimental evidence is now accumulating for the irradiation behavior of current candidates for ceramic breeder and structure, there is no clear insight into the extent of chemical/physicochemical interactions taking place in long-term operations under a reducing atmosphere in a DEMO or power reactor. While these may affect the breeder proper- ties, they may also affect the blanket components structural integrity. 4.15.8.9 Waste Management and Reuse/Recycling The large volume of ceramic breeder and multiplier required for breeding blankets in future power reac- tors necessitate ecological and sound economic solu- tions for intermediate storage and back end. 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All rights reserved. 4.16.1 Introduction 511 4.16.2 Background 513 4.16.2.1 Equation of State of Gases 513 4.16.2.2 Diffusivity 513 4.16.2.3 Solubility 514 4.16.2.4 Trapping 514 4.16.2.5 Permeability 516 4.16.2.6 Recombination 517 4.16.2.7 Irradiation and Implantation 518 4.16.3 Fusion Reactor Materials 518 4.16.3.1 Plasma-Facing Materials 518 4.16.3.1.1 Carbon 518 4.16.3.1.2 Tungsten 521 4.16.3.1.3 Beryllium 524 4.16.3.2 Structural Materials 527 4.16.3.2.1 Austenitic stainless steels 527 4.16.3.2.2 Ferritic/martensitic steels 528 4.16.3.2.3 V–Cr–Ti alloys 532 4.16.3.2.4 Zirconium alloys 534 4.16.3.2.5 Other structural metals 536 4.16.3.3 Barrier Materials 536 4.16.3.3.1 Oxides 536 4.16.3.3.2 Aluminides 537 4.16.3.3.3 Nitrides 538 4.16.3.3.4 Carbides 539 4.16.3.3.5 Low permeation metals 541 4.16.4 Application of Barriers 542 4.16.4.1 Expected In-Reactor Performance 542 4.16.4.2 How Barriers Work and Why Radiation Affects Them 543 4.16.4.3 Why Barriers Are Needed for Fusion Reactors 545 4.16.5 Summary 545 References 546 Abbreviations bcc Body-centered cubic CLAM China low activation martensitic steel CVD Chemical vapor deposition fcc Face-centered cubic HFR High flux reactor HIP Hot isostatically pressed ITER International Thermonuclear Experimental Reactor PCA Prime candidate alloy PRF Permeation reduction factor RAFM Reduced activation ferritic/martensitic steel 4.16.1 Introduction As fusion energy research progresses over the next several decades, and ignition and energy production 511 512 Tritium Barriers and Tritium Diffusion in Fusion Reactors are attempted, the fuel for fusion reactors will be a combination of deuterium and tritium. From a safety point of view, these are not the ideal materials. The reaction of deuteriumwith tritium produces a-particles and 14.1MeV neutrons. These neutrons are used not only to breed the tritium fuel, but also interact with other materials, making some of them radioactive. Although the decay of tritium produces only a low- energy b-radiation, it is difficult to contain tritium. Additionally, being an isotope of hydrogen, tritium can become part of the hydrocarbons that compose our bodies. From the tritium point of view, the fusion facility can be divided into three components: the inner vessel area where the plasma is formed, the blanket where tritium production occurs, and the tritium exhaust and reprocessing system. There is the potential for tritium release in all the three sections of the facility. The tritium cycle for a fusion reactor begins in the blanket region. It is here that the tritium is produced by the interaction of neutrons with lithium. Specifically, the reaction is given symbolically as 6Li(n,a)3H. A neutron that has been thermalized, or lowered in energy by interaction with surrounding materials, is absorbed by 6Li to produce both an a-particle (helium nucleus) and a triton. Elemental lithium contains �7.5% 6Li. As a breeder material in a fusion plant, lithium is enriched in the 6Li isotope to various degrees, depend- ing on the particular blanket design. The 7Li isotope also produces a small amount of tritium via the 7Li (n,a)3H þ n reaction. The cross-section for this endo- thermic reaction is much smaller than that for the 6Li reaction. Upon release from the lithium breeder, the tritium is separated from other elements and other hydrogen isotopes. It is then injected as a gas or frozen pellet into the torus, where it becomes part of the plasma. A fraction of the tritium fuses with deuterium as part of the fusion process, or it is swept out of the chamber by the pumping system. If tritium is removed from the torus by the pumping system and sent to the reprocessing system, it is again filtered to separate other elements and other isotopes of hydrogen. All through the different steps, there is the potential for permeation of the tritium through the materials containing it and for its release to the environment. The probability of this occurring depends on the location in the tritium cycle. This chapter describes hydrogen permeability through two categories of materials that will be used in fusion reactors: candidate plasma-facing and structural materials. The plasma-facing materials in future fusion devices will be heated by high-energy neutrons, by direct interaction of the plasma particles, and by electromagnetic energy released from the plasma. These plasma-facing materials must be cooled. It is primarily through the cooling tubes passing through the plasma-facing materials that tritium losses can occur in the primary vacuum vessel. The three mate- rials typically used for plasma-facing applications are carbon, tungsten, and beryllium. In this report, we describe the behavior of these materials as plasma- facing materials and how tritium can be lost to the cooling system. The term ‘structural material’ is used here to describe materials that serve as the vacuum boundary in the main chamber, as the containment boundary for the blanket region, and as the piping for cooling and vacuum lines. These materials can be ferritic and austenitic steels, vanadium alloys, and zirconium alloys, as well as aluminum alloys in some locations, or potentially ceramics. We give a complete list of the different types of structural materials and review their tritium permeation characteristics. Materials with a low permeability for tritium are being considered as barriers to prevent the loss of tritium from fusion plants. There are a few metals with relatively small values of permeability, but as a whole, metals themselves are not good barriers to the transport of tritium. Ceramics, on the other hand, are typically very good barriers if they are not porous. In most cases, the low permeation is due to extremely low solubility of hydrogen isotopes in ceramic mate- rials. Bulk ceramics, such as silicon carbide, may one day be used as tritium permeation barriers, but most of the current barrier development is for coatings of oxides, nitrides, or carbides of the metals themselves. We show in this review that many such oxides and nitrides may exhibit extremely good permeation behavior in the laboratory, but their performance as a barrier is significantly compromised when used in a radiation environment. We review the permeation parameters of materials being considered for barriers. This report begins with a review of the processes that control the uptake and transport of hydrogen isotopes through materials. The parameters used to define these processes include diffusivity, solu- bility, permeability, trapping characteristics, and recombination-rate coefficients. We examine the transport of hydrogen isotopes in plasma-facingmate- rials, discuss the conditions that exist in the main torus, and look at the ways in which tritium can be lost there. Next, we consider the tritium transport properties of structural materials, followed by the transport properties in barrier materials, including Tritium Barriers and Tritium Diffusion in Fusion Reactors 513 oxides, nitrides, and carbides of structural metals, as well as low-permeation metals. The application of tritium barriers is discussed in some detail: both the theoretical performance of barriers and their observed performance in radiation environments, as well as an example of tritium permeation in the blanket of a fusion reactor. We conclude by summar- izing the tritium permeation properties of all the materials, providing the necessary parameters to help designers of fusion reactors to predict tritium losses during operation. 4.16.2 Background Hydrogen and its isotopes behave similarly in many regards. Gaseous protium, deuterium, and tritium are all diatomic gases that dissociate, especially on metal surfaces, and dissolve into the metal lattice in their atomic form (in some materials, such as polymers and some ceramics, the molecules may retain their diatomic character during penetration of the material). The isotope atoms readily recombine on the free surfaces, resulting in permeation of the gaseous hydro- gen isotopes through metals that support a gradient in hydrogen concentration from one side to the other. In order to understand this process, it is necessary to characterize the source of the hydrogen isotope as well as its transport within the materials. For the purposes of the presentation in this sec- tion, we focus on tritium and its transport through materials. Much of the discussion is equally valid for the deuterium and protium as well (and subsequent sections normalize data to protium). In this section, we provide background on the diffusivity and solu- bility of tritium in metals and relate these thermody- namic parameters to the permeability. In addition, we discuss the role of trapping of hydrogen isotopes on transport of these isotopes, as well as kinetically limited transport phenomena such as recombination. 4.16.2.1 Equation of State of Gases In the case of gaseous exposure, the ideal gas equation of state characterizes the thermodynamic state of the gas: V 0m ¼ RT=p ½1� where V 0m is the molar volume of the ideal gas, T is the temperature of the system in Kelvin, p is the partial pressure of the gaseous species of interest, and R is the universal gas constant equal to 8.31447 J mol�1 K�1. The ideal gas equation of state provides a good estimate for most gases, particularly at low pressures (near ambient) and elevated temperatures (greater than room temperature). In the context of materials exposed to hydrogen isotopes in fusion technologies, the assumption of ideal gas behavior is a reasonable estimate for gaseous hydrogen and its isotopes. More details about the equation of state for real gaseous hydrogen and its isotopes can be found in San Marchi et al.1 4.16.2.2 Diffusivity Tritium diffusion in metals is simply the process of atomic tritium moving or hopping through a crystal lattice. Tritium tends to diffuse relatively rapidly through most materials and its diffusion can be measured at relatively low temperatures. Diffusivity, D, is a thermodynamic parameter, and therefore, fol- lows the conventional Arrhenius-type dependence on temperature: D ¼ D0 expð�ED=RT Þ ½2� where D0 is a constant and ED is the activation energy of diffusion. Measuring tritium diffusion is nontrivial because of the availability of tritium. Therefore, hydrogen and deuterium are often used as surrogates. From the classic rate theory, it is commonly inferred that the ratio of diffusivities of hydrogen isotopes is equivalent to the inverse ratio of the square root of the masses of the isotopes: DT DH ¼ ffiffiffiffiffiffi mH mT r ½3� where m is the mass of the respective isotope, and the subscripts tritium and hydrogen refer to tritium and hydrogen, respectively. When this approximation is invoked, the activation energy for diffusion is gener- ally assumed to be independent of the mass of the isotope. Diffusion data at subambient temperatures do not support eqn [3] for a number of metals;2 however, at elevated temperatures, the inverse square root dependence on mass generally provides a rea- sonable approximation (especially for face-centered cubic (fcc) structural metals).3–9 Although eqn [3] provides a good engineering estimate of the relative diffusivity of hydrogen and its isotopes, more advanced theories have been applied to explain experimental data; for example, quantum corrections and anharmo- nic effects can account for experimentally observed differences of diffusivity of isotopes compared to the predictions of eqn [3].3,10 For the purposes of this report, we assume that eqn [3] is a good approximation 514 Tritium Barriers and Tritium Diffusion in Fusion Reactors for the diffusion of hydrogen isotopes (as well as for permeation) unless otherwise noted, and we normal- ize reported values and relationships of diffusivity (and permeability) to protium. 4.16.2.3 Solubility The solubility (K) represents equilibrium between the diatomic tritium molecule and tritium atoms in a metal according to the following reaction: 1=2T2 $ T ½4� The solubility, like diffusivity, generally follows the classic exponential dependence of thermodynamic parameters: K ¼ K0 expð�DHs=RT Þ ½5� where K0 is a constant and DHs is the standard enthalpy of dissolution of tritium (also called the heat of solution), which is the enthalpy associated with the reaction expressed in eqn [4]. A word of caution: the enthalpy of dissolution is sometimes reported per mole of gas (i.e., with regard to the reaction T2 $ 2T as in Caskey11), which is twice the value of DHs as defined here. Assuming a dilute solution of dissolved tritium and ideal gas behavior, the chemical equilibrium between the diatomic gas and atomic tritium dissolved in a metal (eqn [4]) is expressed as 1=2 m0TT þ RT ln pTT p0TT � � ¼ m0t þ RT ln c0 ½6� where c0 is the equilibrium concentration of tritium dissolved in the metal lattice in the absence of stress, m0TT is the chemical potential of the diatomic gas at a reference partial pressure of p0TT, and m 0 T is the chemical potential of tritium in the metal at infinite dilution. This relationship is the theoretical origin of Sievert’s law: c0 ¼ K ðpTTÞ1=2 ½7� where to a first approximation, the solubility is equiv- alent for all isotopes of hydrogen. It is important to distinguish between solubility and concentration: solubility is a thermodynamic property of the material, while the concentration is a dependent variable that depends on system conditions (including whether equilibrium has been attained). For example, once dissolved in a metal lattice, atomic tritium can interact with elastic stress fields: hydrostatic tension dilates the lattice and increases the concentration of tritium that can dissolve in the metal, while hydrostatic compression decreases the concentration. The relationship that describes this effect in the absence of a tritium flux12–14 is written as cL ¼ c0 exp VT�� RT � � ½8� where cL is the concentration of tritium in the lattice subjected to a hydrostatic stress (�� ¼ �ii=3), and VT is the partial molar volume of tritium in the lattice. For steels, the partial molar volume of hydrogen is�2 cm3 mol�1,15 which can be assumed to first order to be the same for tritium. For most systems, the increase of tritium concentration will be relatively small for ordi- nary applied stresses, particularly at elevated tempera- tures; for example, hydrostatic tension near 400 MPa at 673 K results in a �15% increase in concentration. On the other hand, internal stresses near defects or other stress concentrators can substantially increase the local concentration near the defect. It is unlikely that local concentrations will significantly contribute to elevated tritium inventory in the material, but locally elevated concentrations of hydrogen isotopes become sites for initiating and propagating hydrogen- assisted fracture in structural metals. 4.16.2.4 Trapping Tritium can bond to microstructural features within metals, including vacancies, interfaces, grain bound- aries, and dislocations. This phenomenon is generally referred to as ‘trapping.’15–18 The trapping of hydro- gen and its isotopes is a thermally governed process with a characteristic energy generally referred to as the trap binding energy Et. This characteristic energy represents the reduction in the energy of the hydro- gen relative to dissolution in the lattice16,19 and can be thought of as the strength of the bond between the hydrogen isotope and the trap site to which it is bound. Oriani16 assumed dynamic equilibrium between the lattice hydrogen and trapped hydrogen yT 1� yT ¼ yL 1� yL exp Et RT � � ½9� where yT is the fraction of trapping sites filled with tritium and yL is the fraction of the available lattice sites filled with tritium. According to eqn [9], the frac- tion of trap sites that are filled depends sensitively on the binding energy of the trap (Et) and the lattice con- centration of tritium (yL). For example, traps in ferritic steels, which are typically characterized by low lattice concentrations and trap energy 1000 K). Tritium Barriers and Tritium Diffusion in Fusion Reactors 515 For materials with strong traps and high lattice con- centration of tritium, trapping can remain active to very high temperatures, particularly if the trap energy is large (>50 kJ mol�1). The coverage of trapping sites for low and high energy traps is shown in Figure 1 for two values of K: one material with relatively low solubility of hydrogen and the other with high solubility. The absolute amount of trapped tritium, cT, depends on yT and the concentration of trap sites, nT 15: cT ¼ anTyT ½10� where a is the number of hydrogen atoms that can occupy the trap site, which we assume is one. If multiple trapping sites exist in the metal, cT is the sum of trapped tritium from each type of trap. A similar expression can be written for the tritium in lattice sites, cL: cL ¼ bnLyL ½11� where nL is the concentration of metal atoms and b is the number of lattice sites that hydrogen can occupy per metal atom (which we again assume is one). Substituting eqns [10] and [11] into eqn [9] and recognizing that yL � 1, the ratio of trapped tritium to lattice tritium can be expressed as cT cL ¼ nT½cL þ nLexpð�Et=RT Þ� ½12� 200 0 0.2 Fr ac tio na l c ov er ag e of t ra p s q T 0.4 0.6 0.8 1 300 400 5 Tempe Et= 10 kJ mol –1 Et= 5 Figure 1 Fraction of filled traps as a function of temperature f (modeled as reduced activation ferritic/martensitic steel and aus Table 1). The pressure is 0.1 MPa, the molar volume of the stee to be one lattice site for hydrogen per metal atom. Therefore, the ratio of trapped tritium to dissolved (lattice) tritium will be large if cL is small and Et is large. Conversely, the amount of trapped tritium will be relatively low in materials that dissolve large amounts of tritium. The transport and distribution of tritium in metals can be significantly affected by trapping of tritium. Oriani16 postulated that diffusion follows the same phenomenological form when hydrogen is trapped; however, the lattice diffusivity (D) is reduced and can be replaced by an effective diffusivity, Deff, in Fick’s first law. Oriani went on to show that the effective diffusivity is proportional to D and is a function of the relative amounts of trapped and lattice hydrogen: Deff ¼ D 1þ cT cL ð1� yTÞ ½13� If the amount of trapped tritium (cT) is large relative to the amount of lattice tritium (cL), the effective diffusivity can be several orders of magnitude less than the lattice diffusivity.20 Moreover, the effective diffusivity is a function of the composition of the hydrogen isotopes, depending on the conditions of the test as well as sensitive to the geometry and microstructure of the test specimen. Thus, the intrin- sic diffusivity of the material (D) cannot be measured directly when tritium is being trapped. Equation [13] is the general form of a simplified expression that is commonly used in the literature: 00 rature (K) 0 kJ mol–1 ‘Low’ solubility ‘High’ solubility 600 700 800 or ‘low-solubility’ and ‘high-solubility’ materials tenitic stainless steel, respectively, using relationships from ls is approximated as 7 cm3 mol�1 and there is assumed 516 Tritium Barriers and Tritium Diffusion in Fusion Reactors Deff ¼ D 1þ nT nL exp Et RT � � ½14� Equation [14] does not account for the effect of lattice concentration, and is therefore inadequate when the concentration of tritium is relatively large. For materials with high solubilities of tritium (such as austenitic stainless steels), trapping may not affect transport significantly and Deff � D as shown in Figure 2. For materials with a low solubility and relatively large Et, the effective diffusivity can be substantially reduced compared to the lattice diffu- sivity (Figure 2). The wide variation of reported diffusivity of hydrogen in a-iron at low temperatures is a classic example of the effect of trapping on hydrogen transport2,20: while the diffusivity of hydro- gen at high temperatures is consistent between studies, the effective diffusivity measured at low tem- peratures is significantly lower (in some cases by orders of magnitude) compared to the Arrhenius relationship established from measurements at ele- vated temperatures. Moreover, the range of reported values of effective diffusivity demonstrates the sensi- tivity of the measurements to experimental technique and test conditions. For these reasons, it is important to be critical of diffusion data that may be affected by trapping and be cautious of extrapolating diffusion data to experimental conditions and temperatures 200 10–3 10–2 10–1 100 300 nT= 10 –7 D ef f/ D n 400 Tempe Figure 2 Ratio of effective diffusivity to lattice diffusivity (Deff/D (squares with varying nT, modeled as reduced activation ferritic/ material (triangles, modeled as austenitic stainless steel with Et pressure is 0.1 MPa, the molar volume of the steels is approxim site for hydrogen per metal atom. different from those measured, especially if trapping is not well characterized or the role of trapping is not known. 4.16.2.5 Permeability Permeability of hydrogen and its isotopes is generally defined as the steady-state diffusional transport of atoms through a material that supports a differential pressure of the hydrogen isotope. Assuming steady state, semi-infinite plate, and Fick’s first law for dif- fusion J ¼ �Dðdc=dxÞ, we can express the steady- state diffusional flux of tritium as J1 ¼ �D cx2 � cx1ð Þ x2 � x1 ½15� where cx is the concentration at position x within the thickness of the plate. Using chemical equilibrium (eqn [7]) and assuming that the tritium partial pres- sure is negligible on one side of the plate of thickness t, the steady-state diffusional flux can be expressed as J1 ¼ DK t p 1=2 TT ½16� and the permeability, F, is defined as: F � DK ½17� Substituting eqns [2] and [5] into eqn [17], the per- meability can be expressed as a function of tempera- ture in the usual manner: T= 10 –5 nT= 10 –3 500 rature (K ) 600 700 800 ) as a function of temperature for ‘low-solubility’ material martensitic steel with Et ¼ 50 kJ mol�1) and ‘high-solubility’ ¼ 10 kJ mol�1 and nT ¼ 10�3 traps per metal atom). The ated as 7 cm3 mol�1 and there is assumed to be one lattice Tritium Barriers and Tritium Diffusion in Fusion Reactors 517 F ¼ K0D0 exp½�ðDHs þ EDÞ=RT � ½18� Permeability is a material property that charac- terizes diffusional transport through a bulk material, that is, it is a relative measure of the transport of tritium when diffusion-limited transport domi- nates; see LeClaire21 for an extensive discussion of permeation. By definition, the permeability (as well as diffusivity and solubility) of hydrogen isotopes through metals is independent of surface condition, since it is related to diffusion of hydrogen through the material lattice (diffusivity) and the thermody- namic equilibrium between the gas and the metal (solubility). In practice, experimental measurements are strongly influenced by surface condition, such that the measured transport properties may not reflect diffusion-limited transport. Under some conditions (such as low pressure or due to the presence of resid- ual oxygen/moisture in the measurement system), the theoretical proportionality between the square root of pressure and hydrogen isotope flux does not describe the transport;21,22 thus, studies that do not verify diffusion-limited transport should be viewed critically. In particular, determination of the diffusivity of hydrogen and its isotopes is partic- ularly influenced by the surface condition of the specimen, since diffusivity is determined from tran- sient measurements. While permeation measurements (being steady-state measurements) are relatively less sensitive to experimental details, the quality of reported solubility relationships depends directly on the quality of diffusion, since solubility is typically determined from the measured permeability and dif- fusivity.1 In addition, trapping affects diffusivity and must, therefore, be mitigated in order to produce solubility relationships that reflect the lattice disso- lution of hydrogen and its isotopes in the metal. These characteristics of the actual measurements explain the fidelity of permeation measurements between studies in comparison with the much larger variation in the reported diffusivity and solubility. 4.16.2.6 Recombination As shown earlier, steady-state permeation of hydro- gen through materials is normally governed only by solubility and diffusivity. It has been shown23 that at low pressures, permeation can also be limited by dissociation at the surface. Due to limited data in the literature on this effect (and questions about whether this condition ever really exists), we do not consider this effect in this chapter. It is also possible for permeation to be limited by the rate at which atoms can recombine back into molecules. With the exception of extremely high temperatures, this recom- bination is necessary for hydrogen to be released from a material. Wherever the release rate from a surface is limited by recombination, the boundary condition at that boundary is given by: Jr ¼ krc2 ½19� where kr is the recombination-rate constant and c is the concentration of hydrogen near the surface (for this discussion, we assume that there is no surface rough- ness). The units for k and c are m4 s�1 per mol of H2 and mol H2 m �3, respectively. There are two specific types of conditions that can lead to the hydrogen release being rate limited by recombination. One of them occurs for plasma- facing materials in which the recombination-rate coefficient is relatively low, and the implantation rate is high. With this condition, the concentration of hydrogen in the very-near plasma-exposed surface will increase to the point at which Jr is effectively equal to the implantation rate. It is not exactly equal to the implantation rate because there is permeation away from that surface to the downstream surface. The other condition that can lead the hydrogen release being controlled by recombination is when the rate of ingress at the upstream boundary is very low. This condition can occur either when the upstream pressure is extremely low or a barrier is placed on the upstream surface, and the downstream surface has a relatively low recombination-rate con- stant. In the extreme case, the release rate at the downstream side is so slow that the hydrogen con- centration becomes uniform throughout the material. The release rate from the downstream surface will be krc 2, where c now represents the uniform concentra- tion. From c ¼ Kp1=2TT and eqn [19], it can easily be shown that the recombination-limited permeation is linearly dependent on pressure, rather than having the square root of pressure dependence of diffusion- limited permeation. There are various derivations and definitions of the recombination-rate constant. In the case of intense plasma exposure in which extreme near-surface concentrations are generated, Baskes24 derived the recombination-rate constant with the assumption that the rate was controlled by the process of bulk atoms jumping to the surface, combining with surface atoms, and then desorbing. His expression for the recombination-rate constant is 518 Tritium Barriers and Tritium Diffusion in Fusion Reactors kr ¼ C 8 mT � �1=2 s K 20 exp 2DHs � EX RT � � ½20� where C is a constant, s is the sticking constant, which depends on the cleanliness of the material surface, and EX ¼ DHs þ ED > 0, otherwise EX ¼ 0. The sticking constant can be anywhere from 1 for clean surfaces to 10�4 or smaller for oxidized surfaces. Pick and Sonnenberg25 solved the recombination- rate constant for the case where the near-surface con- centration of hydrogen is small. In the limit of low surface concentration, the rate of atom jump to the surface does not play an important role in the recom- bination rate, thus eliminating EX from the exponen- tial. The sticking constant in the Pick and Sonnenberg model is thermally activated: s ¼ s0 expð�2EC=RTÞ, where s0 is the sticking coefficient and EC is the activation energy for hydrogen adsorption. Wampler26 also studied the case of low near- surface concentration to arrive at an expression for the recombination-rate constant. He assumed equi- librium between hydrogen atoms in surface chemi- sorption sites and atoms in solution, deriving the recombination-rate constant as kr ¼ nsn bnLð Þ2 exp 2DHs RT � � ½21� where ns is the area density of surface chemisorption sites, and n is the jump frequency. These expressions differ, but also display many similarities. Unfortunately, the surface cleanliness dominates the rate of recombination and these theo- retical relationships are relevant only for sputter- cleaned surfaces and very low pressures. For example, Causey and Baskes27 showed that the Baskes24 model predicts fairly accurate results for plasma-driven per- meation of deuterium in nickel. Comparison with values in the literature for nickel showed other results to differ by as much as four orders of magnitude and to have significantly different activation energies. 4.16.2.7 Irradiation and Implantation Irradiation and implantation can affect the transport of hydrogen isotopes in materials. Since these effects can be complex and depend on the conditions of the mate- rials and the environment, it can be difficult to draw broad conclusions from the literature. Nevertheless, changes in apparent transport properties are generally attributed to damage and the creation of hydrogen traps28–31 (see also Chapter 1.03, Radiation-Induced Effects on Microstructure). Therefore, the effects of irradiation and implantation will depend sensitively on the characteristics of the traps that are created by these processes. The density of damage is an important con- sideration: for example, it has been shown that helium bubbles are not effective trapping sites for steels,32 likely because in these experiments, the density of helium bubbles was relatively low. The energy of the trap will determine the coverage as a function of tem- perature (eqn [9]): generally, the effect of trapping will be stronger at low temperatures, especially in materials with a low solubility (Figure 2), which can result in substantial increases in hydrogen isotope inventory compared to hydrogen content predicted from lattice solubility. Additionally, irradiation may increase ioni- zation of hydrogen isotopes, thus enhancing apparent permeation.29 Reactor environments can defeat per- meation barriers, for example, by damaging the integ- rity of oxide layers; this is discussed at the end of this chapter. 4.16.3 Fusion Reactor Materials 4.16.3.1 Plasma-Facing Materials Tritium generated in the fusion-reactor blanket will be fed directly into the plasma in the main vacuum cham- ber. There, the tritium will be partially consumed, but it will also interact with the materials composing the first wall. Materials used to line the first wall will be exposed to energetic tritium and deuterium escaping from the plasma. Particle fluxes in the range of 1021 (D þ T) m�2 s�1 will continuously bombard the plasma-facing materials. Materials used for the diver- tor at the top and/or bottom of the torus will be exposed to lower energy particles with a flux of 1023 (D þ T) m�2 s�1 or higher. While the neutral gas pressure of tritium will be relatively low at the outer vacuum wall boundary, some minor permeation losses will occur. In reality, the primary concerns in the plasma-facing areas are tritium inventory and perme- ation into the coolant through coolant tubes. While future power reactors are likely to have primarily refractory metals such as tungsten, present-day devices are still using carbon and beryllium. In this section on plasma-facing materials, we examine the interaction of tritium with carbon, tungsten, and beryllium. 4.16.3.1.1 Carbon In many ways, carbon is ideal for fusion applications. It is a low-Z material with a low vapor pressure and excellent thermal properties. The carbon used in fusion applications comes in two forms, graphite and carbon Tritium Barriers and Tritium Diffusion in Fusion Reactors 519 composites. Graphite is described in Chapter 2.10, Graphite: Properties and Characteristics; Chapter 4.10, Radiation Effects in Graphite, and Chapter 4.18, Carbon as a Fusion Plasma-Facing Material and is typically made using the Acheson process.33 Calcined coke is crushed, milled, and then sized. The properties of the graphite are determined by the size and shape of these particles. Coal tar is added to the particles and the batch is heated to �1200 K. This process is repeated several times to increase the density of the compact. The final bake is at temperatures between 2900 and 3300 K and takes �15 days. The final product is quite porous, with a density of around 1.8–1.9 g cm�3 (compared to a theoretical density of 2.3 g cm�3). Graphite is composed of grains (from the original coke particles) with a size of 5–50 mm, which are in turn composed of graphite subgrains with a typical size of 5 nm. Carbon composites are made by pyrolyzing a composite of carbon fibers in an organic matrix. These fibers have a high strength-to-weight ratio and are composed of almost pure carbon. As with graphite, carbon composites are quite porous with a density of 1000 K). As a plasma-facing material, graphite will be exposed to atomic tritium and deuterium, and these hydrogen isotopes will migrate inward along the open porosity. Several research groups have measured the diffusion coefficient for hydrogen on carbon surfaces. Robell et al.39 inferred an activation energy of 164 kJ mol�1 for the diffusion during measurements on the uptake of hydrogen on platinized carbon between 573 and 665 K. Olander and Balooch40 used similar experi- ments to determine the diffusion coefficient for hydro- gen on both the basal and prism plane: 6� 10�9 exp (�7790/T) m2 s�1 for the basal plane and 6� 10�11 exp(�4420/T) m2 s�1 for the prism plane. Causey et al.41 used tritium profiles in POCOAXF-5Q graphite exposed to a tritium plasma to extract a diffusivity for tritium on carbon pores of 1.2� 10�4 exp(�11 670/T) m2 s�1. An example of the deep penetration of hydro- gen isotopes into the porosity of graphite was reported by Penzhorn et al.42 Graphite and carbon composite tiles removed from the Joint European Torus (JET) fusion reactor were mechanically sectioned. The sec- tions were then oxidized, and tritiated water was collected for liquid scintillation counting. Relatively high concentrations of tritium were detected tens of millimeters deep into the tiles. The diffusion or migration of hydrogen isotopes on carbon surfaces occurs at lower temperatures. The solubility, diffusion, and trapping of hydrogen in the carbon grains are higher temperature processes than adsorption and surface diffusion. At higher tempera- tures (>1000 K), hydrogen molecules can dissociate and be absorbed at chemically active sites on carbon surfaces. Some of these sites are located on the out- side of the grains, but many exist along the edges of the subgrains that make up the larger grains. Hydrogen isotopes dissociating on the outer grain boundary are able to migrate along the subgrain boundaries, entering into the interior of the grain. It is the jumping from one moderate energy site (�240 kJ mol�1) to another that determines the effec- tive diffusion coefficient. Traps on the grain bound- aries pose a binding energy barrier (�175 kJ mol�1) that must be overcome in addition to this normal lattice activation. Atsumi et al.43 used the pressure change in a constant volume to determine the solu- bility of deuterium in ISO 88 graphite. They deter- mined the solubility to be given by K ¼ 18.9 exp (þ2320/T) mol H2 m�3 MPa�1/2 over the tempera- ture range of 1123–1323 K. This solubility is shown in Figure 3 along with two data points by Causey44 at 1273 and 1473 K. A negative heat of solution is seen in both sets of data, suggesting the formation of a bond between hydrogen and carbon. 0.65 10 100 S ol ub ili ty (m ol m –3 M P a– 1/ 2 ) 1000 0.7 0.75 Causey et al. Atsumi et al. Temperature, 1000/T (K–1) 0.8 0.85 0.9 Figure 3 Solubility of hydrogen in carbon. Adapted from Atsumi, H.; Tokura, S.; Miyake, M. J. Nucl. Mater. 1988, 155, 241–245; Causey, R. A. J. Nucl. Mater. 1989, 162, 151–161. 0.4 10–21 10–19 10–17 D iff us iv ity (m 2 s– 1 ) 10–15 10–13 10–11 0.6 Causey et al. Rohrig et al. Atsumi et al. Malka et al. Causey (best estimate) 0.8 Temperature, 1000/T (K–1) 1 1.2 1.4 Figure 4 Diffusivity of hydrogen in graphite. The bold line is the best estimate given in Causey44 based on uptake data for tritium in POCO AFX-5Q graphite. Adapted from Atsumi, H.; Tokura, S.; Miyake, M. J. Nucl. Mater. 1988, 155, 241–245; Causey, R. A. J. Nucl. Mater. 1989, 162, 151–161; Röhrig, H. D.; Fischer, P. G.; Hecker, R. J. Am. Ceram. Soc. 1976, 59, 316–320; Causey, R. A.; Elleman, T. S.; Verghese, K. Carbon 1979, 17, 323–328; Malka, V.; Röhrig, H. D.; Hecker, R. Int. J. Appl. Radiat. Isot. 1980, 31, 469. 520 Tritium Barriers and Tritium Diffusion in Fusion Reactors The variation in the diffusivities of hydrogen in graphite determined by various researchers is extreme. This variation results primarily from differ- ences in interpretation of the mechanism of diffusion (e.g., bulk diffusion or grain boundary diffusion). Representative values for the diffusion is shown in Figure 4. Röhrig et al.45 determined their diffusion coefficient using the release rate of tritium from nuclear grade graphite during isothermal anneals. They correctly used the grain size as the real diffusion distance. Causey et al.46 measured the release rate of tritium recoil injected into pyrolytic carbon to deter- mine the diffusivity. Malka et al.47 used the release rate of lithium-bred tritium in a nuclear graphite to deter- mine the diffusion coefficient. Atsumi et al.43 used the desorption rate of deuterium gas from graphite samples that had been exposed to gas at elevated tem- peratures to determine a diffusivity. Building on the workof others, Causey44 proposed an alternative expres- sion for the diffusivity that he labeled ‘best estimate.’ Tritium Barriers and Tritium Diffusion in Fusion Reactors 521 The result was based on uptake experiments for tri- tium into POCOAFX-5Q graphite, and the assump- tion that the total uptake is determined by the product of the diffusivity and the solubility. The uptake data were analyzed assuming that the expression given by Atsumi et al.43 for the solubility was correct. The expression ‘best estimate’ was used because it prop- erly took into consideration that the grain size was the effective diffusion distance, and that diffusivity was more properly determined by uptake than release. Release rates are strongly affected by trapping. The trapping of hydrogen isotopes at natural and radiation-induced traps has been examined by sev- eral research groups.41,48–50 Causey et al.41 exposed POCO-AFX-5Q graphite to a deuterium/tritium mixture at elevated temperatures in an examination of the kinetics of hydrogen uptake. For temperatures above 1500 K, it was discovered that increasing the pressure did not increase the retention. It appeared that the solubility hit an upper limit at 17 appm. Anal- ysis of the data revealed that the 17 appm did not represent solubility, but a trap density. The trap was determined to have a binding energy of 175 kJ mol�1. Atsumi et al.48 found that radiation damage increased the apparent solubility of deuterium in graphite by a factor of 20–50 with saturation at a radiation damage level of 0.3 dpa. A significant decrease in the apparent diffusivity was also noted. In a later study, Atsumi et al.49 reported graphites and composites to vary sig- nificantly in their natural retention values. They saw that the saturation retentionwas inversely proportional to the lattice constant (which relates to the degree of graphitization, and so grain size). Radiation damage was seen to decrease the apparent lattice constant and increase the saturation retention. 6MeVCþ ions were used byWampler et al.50 to simulate neutron damage to different graphites. The trap density increased with damage levels up to 0.04 dpa at which the saturation retention was 650 appm. In a defining set of experi- ments, Causey et al.28 examined tritium uptake in unir- radiated and radiated pitch-based carbon composites. Pitch-based carbon composites have a large lattice parameter due to the sheet-like configuration of the fibers. These composites retained significantly less tri- tium before and after irradiation than other carbon materials. From these results, it is apparent that high- energy trapping occurs at the edges of the hexagonal crystals on the prism plane. Hydrogen isotope permeation in the normal sense does not apply to graphite and carbon composites. There is inward migration of atomic hydrogen isotopes along porosity at lower temperatures. To calculate the potential tritium inventory for this process, one can obtain an upper bound by assuming monolayer coverage of the pore surfaces. Typical nuclear grade graphite has a specific surface area of 1 m2 g�1. Complete loading of that amount of surface area yields 2� 1025 Tm�3, or about 2 g for a 20 m2 carbon wall that is 10-mm thick. At higher temperatures, molecular hydrogen isotopes, which are moving through the graphite pore system, are able to disso- ciate and enter the multimicron-sized graphite grains. This migration into the grains occurs along the edges of the nanometer scale subgrains. As the hydrogen migrates inward, it decorates high-energy trap sites. The density of these trap sites is higher than the effective solubility derived from the migra- tion rate. Permeation into graphite will not lead to tritium release from a fusion device, but will affect tritium inventory. If one assumes that radiation dam- age from neutrons has increased the concentration of traps with a binding energy of 175 kJ mol�1 up to 1000 appm, and that tritium is occupying 100% of those traps, the same 20 m2 wall 10-mm thick listed above would now contain an additional 100 g of tritium. Occupation of all of the traps is difficult to achieve: at low temperatures, kinetics makes it impos- sible to achieve saturation, while at substantially higher temperatures, the traps do not remain filled. There is another process called carbon codeposi- tion that can strongly affect tritium inventory in a fusion device. In the codeposition process, carbon eroded from the walls of the tokamak is redeposited in cooler areas along with deuterium and tritium from the plasma. Because carbon codeposition is not a diffusion or permeation process, it will not be covered in this review. The interested reader is referred to a review of this process by Jacob.51 4.16.3.1.2 Tungsten Tungsten is another of the plasma-facing materials, described in Chapter 4.17, Tungsten as a Plasma- Facing Material. Like carbon, it will not be a vacuum barrier. Thus, permeation through the tungsten will not lead to tritium release directly into the environ- ment. It can lead to tritium permeation into the coolant through the coolant tubes inside the tungsten facing materials. Permeation will also affect the tri- tium inventory of the fusion device. Tungsten has excellent thermal properties with a very high melting point of 3683 K. The problem that tungsten presents to the tokamak designer is the deleterious radiation losses if tungsten is present in the plasma. Fortu- nately, the energy threshold for sputtering by 522 Tritium Barriers and Tritium Diffusion in Fusion Reactors hydrogen ions is quite high, 700 eV for tritium.52 For that reason, tungsten will be used primarily in the divertor region where the energy of the impacting particles can be limited. There are a limited number of reports on the diffusivity of hydrogen isotopes in tungsten. Frauen- felder53 measured the rate of hydrogen outgassing from saturated rolled sheet samples at temperatures over the wide range 1200–2400 K. His material was 99.95% pure tungsten. Zakahrov and Sharapov54 used 99.99% pure tungsten samples in their permeation techniques to determine the hydrogen diffusivity over a limited temperature range of 900–1060 K. In the Benamati et al.55 experiments using tungsten con- taining 5% rhenium, a gaseous permeation technique was also used. These experiments were performed over a very limited temperature range of 850–885 K. Reported diffusivities are shown in Figure 5. There are a couple of reasons why the diffusion coefficient reported by Frauenfelder53 is widely accepted as most correct. The first of these reasons is the wide temperature range over which experiments were per- formed. The second reason is that the experiments were performed at a temperature above that where trapping typically occurs. It can be seen in Figure 5 that Zakahrov’s54 diffusivity agrees quite well with Frauenfelder’s at the highest temperatures, but falls below his values at lower temperatures, where trapping would occur. The database on hydrogen solubility in tungsten is also limited. The results of the two experimental 0.4 D iff us iv ity (m 2 s– 1 ) 10–11 10–10 10–9 10–8 10–7 0.6 0.8 Temperatu Frauenfelder B Figure 5 Diffusivity of hydrogen in W. Adapted from Frauenfe A. P.; Sharapov, V. M. Fiziko-Khimicheskaya Mekhanika Materialo Mater. 2000, 283–287, 1033–1037. studies are shown in Figure 6. In the same experi- ments used to determine the diffusivity, Frauenfelder53 also measured solubility. Over the temperature range 1100–2400 K, samples were saturated at fixed pressures and then heated to drive out all of the hydrogen. Over a more limited temperature range of 1900–2400 K, Mazayev et al.56 also examined hydrogen solubility in tungsten. The agreement with the Frauenfelder’s53 data is quite good in magnitude, but not good in apparent activation energy. As with his diffusivity, the solubility reported by Frauenfelder is typically the value used in predicting the migration of hydrogen in tungsten. Hydrogen trapping in tungsten has been studied by several research groups. van Veen et al.57 used bom- bardment by 2 keV protons in their study of the bond- ing of hydrogen to voids in single-crystal tungsten. Thermal desorption from the samples with appms of voids revealed a broad release peak at 600–700K. It was stated that the release could be modeled as gas going back into solution from the voids with a trap binding energy of 96.5–135 kJ mol�1 controlling the process. Eleveld and van Veen,58 in a similar study, used a lower fluence of 30 keV Dþ ions in desorption experiments. In these samples containing vacancies but no voids, the release occurred at 500–550 K. The authors reported a value of 100 kJ mol�1 for the trap binding energy of vacancies. Pisarev et al.59 used lower fluences of 7.5 keV deuterons into 99.94% pure tungsten samples. During thermal desorption ramps, peaks in the release rates were seen at 350, 480, 600, and 750K.The release at the highest temperature was seen only in the highest re, 1000/T (K–1) 1 1.2 1.4 Zakharov et al. enamati et al. lder, R. J. Vac. Sci. Technol. 1969, 6, 388–397; Zakahrov, v 1973, 9, 29–33; Benamati, G.; Serra, E.; Wu, C. H. J. Nucl. 0.6 0.7 Temperature, 1000/T (K–1) 0.8 0.9 10.4 0.01 0.1 S ol ub ili ty (m ol m −3 M P a− 1/ 2 ) 1 10 0.5 Mazayev et al. Frauenfelder Figure 6 Solubility of hydrogen in W. Adapted from Frauenfelder, R. J. Vac. Sci. Technol. 1969, 6, 388–397; Mazayev, A. A.; Avarbe, R. G.; Vilk, Y. N. Russian Metallurgy-Metally-USSR 1968, 6, 153–158. Tritium Barriers and Tritium Diffusion in Fusion Reactors 523 fluences. Garcı́a-Rosales et al.60 used 100 eV deuterium implantation to study the trapping and release rate of hydrogen isotopes from wrought and plasma-sprayed tungsten. Two broad desorption peaks at 475–612 K and 670–850 Kwere seen in the thermal desorption spectra. Modeling of the release data suggested the lower temperature peak to be controlled by both diffu- sion and trapping at a binding energy of 44 kJ mol�1. The second release peak was reported to correspond to trapping at defects with a binding energy of 97 kJ mol�1. In experiments with 99.99% pure tungsten and tungsten with 1% lanthanum oxide, Causey et al.61 examined tritium retention in plasma-exposed samples. Modeling of the results suggested two traps, one with a binding energy of 97 kJ mol�1 and another with 204 kJ mol�1. The density of the trapped tritium averaged 400–500 appm. Anderl et al.62 used deuterium im- plantation into polycrystalline tungsten to determine the correlation between dislocation density on cell walls and deuterium trapping. Annealing tungsten at 1673K reduced the dislocation density by a factor of 7, subsequently reducing the deuterium trapping by a similar factor. The binding energy of these traps was estimated to be 88–107 kJ mol�1. As-received 99.95% pure tungsten was used by Sze et al.63 in experiments with intense deuterium plasma exposure. Exposure at 400K resulted in blisters with diameters of tens of microns. Elevating the temperature to 1250K elimi- nated the blisters. Venhaus et al.64 used high-purity foils in experiments to examine the effect of annealing temperature on blistering by deuterium plasma expo- sure. An unannealed sample and one annealed at 1473 K both exhibited blisters after the plasma exposure. The sample annealed at 1273K did not blister. There have been a multitude of other reports on blister for- mation on tungsten samples exposed to various forms of hydrogen implantation.65–68 Anderl et al.62 used 99.95% tungsten in 3 keV D3 þ ion implantation to determine the recombination- rate coefficient. Over a temperature range of 690– 825 K, the recombination rate coefficient was given as kr ¼ 3.85� 109 exp(�13 500/T )m4 s�1 per mol of H2. This expression is shown in Figure 7, where it is plotted along with the expression given by the Baskes24 model. It can be seen that there is very little correlation between the measured Anderl value and the calculated Baskes value. This is not entirely unusual. Impurities on the surface, especially oxide layers, can have a very strong effect on this coefficient. While tungsten has excellent low permeability for gaseous tritium, it will be used only in fusion devices as a plasma-facing material. As a plasma- facing material, tungsten will be exposed to intense fluxes of energetic tritium and deuterium. With traps for hydrogen at binding energies of 97 and 203 kJ mol�1(57–62) at natural and radiation-induced defects, it would appear that a substantial tritium inventory could be generated in divertor tungsten. There are several reasons why this high inventory is not likely to occur. The first reason is the high recombination- rate coefficient given earlier. For a recombination- rate constant of 10�1 m4 s�1 per mol of H2 or higher (see Figure 7), the recombination rate on the surface is so rapid as to generate the equivalent of c ¼ 0 at Baskes model Anderl et al. 1 10 100 1000 R ec om b in at io n- ra te c oe ffi ci en t (m 4 s– 1 p er m ol H 2) 104 105 106 107 10.8 1.2 1.4 Temperature, 1000/T (K–1) 1.6 1.8 2 2.2 Figure 7 Recombination-rate coefficient of hydrogen in W. Adapted from Baskes, M. I. J. Nucl. Mater. 1980, 92, 318–324; Anderl, R.; Holland, D. F.; Longhurst, G. R.; et al. Fusion Technol. 1991, 21, 745–752. 524 Tritium Barriers and Tritium Diffusion in Fusion Reactors the boundary. With the very limited penetration dis- tance of energetic hydrogen in the dense tungsten, most of the implanted material is immediately released back out of the surface. There are also recent reports suggesting that ruptured blisters and very fine cracks near the surface69–71 will even further reduce the inward migration of deuterium and tritium into the tungsten. 4.16.3.1.3 Beryllium Beryllium is a low-Zmaterial with good thermal char- acteristics, described in Chapter 2.11, Neutron Reflector Materials (Be, Hydrides) and Chapter 4.19, Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices. Additionally, it is a good getter for oxygen impurities in the plasma. The low-Zminimizes the radiation losses from the plasma, and the oxygen removal keeps the plasma clean. For these reasons, beryllium has been used in the JET fusion reactor and will be the first wall material for the International Thermonuclear Experimental Reactor (ITER). Beryllium has interesting hydrogen retention behavior. Beryllium may also be used as a neutron multiplier in the blanket area of future fusion devices to increase the tritium breeding ratio. Abramov et al.72 used two grades of beryllium in their permeation–diffusion experiments. These were high-purity (99%) and extra grade (99.8%). Adding to the validity of their experimental result was the fact that the authors used multilayer permeation theory analysis to take permeation through the outer oxide layer into consideration. For a lower purity material (98%), Tazhibaeva et al.73 also used the multilayer permeation analysis to determine dif- fusivity. Jones and Gibson74 studied tritium diffusiv- ity and solubility for arc-cast beryllium in the temperature range of 673–1173K. Beryllium was exposed to tritium gas for various temperatures, durations, and pressures during isothermal anneals. After removing the samples to another experimental system, the samples were heated to various tempera- tures. For the initial heating, the tritium release would rise, but soon fall to zero. Elevating the tem- perature would reestablish the tritium release, but again the release would fall. While this behavior is not typical of diffusion controlled release, the data were analyzed to extract an effective diffusivity. The different reported diffusivities are shown in Figure 8. It can be seen that the diffusivity reported by Abramov et al.72 is considerably larger than those of Tazhibaeva et al.73 and Jones and Gibson.74 It is apparent that the purity of the beryllium played a strong role in deter- mining the effective diffusivity. Oxygen, the primary impurity in beryllium provides a strong trap for hydrogen. Thompson and Macaulay-Newcombe75,76 examined the diffusion of deuterium in single-crystal and polycrystalline beryllium. The effective diffusivity in the single-crystal material was lower than that for the polycrystalline material. The polycrystalline results agreed quite well with the results reported by Abramov et al.72 They suggested that the lower diffusiv- ity seen for the single-crystal samples was the true diffusivity for beryllium, and that the polycrystalline results represented diffusion along the grain boundaries. 0.8 10–13 10–12D iff us iv ity (m 2 s– 1 ) 10–11 10–10 1 1.2 Tazhibaeva et al. Abramov et al. Extra grade High grade Jones and Gibson Temperature, 1000/T (K–1) 1.4 1.6 1.8 Figure 8 Diffusivity of hydrogen in beryllium. ‘Extra’ grade is 99.8% pure and ‘high’ is 99.0% pure. Other authors did not specify purities. Adapted from Abramov, E. I. L.; Riehm, M. V. P.; Thompson, D. A.; et al. J. Nucl. Mater. 1990, 175, 95427– 95430; Tazhibaeva, I. L.; Shestakov, V. P.; Chikhray, Y. V. In Proceedings of the 18th Symposium of Fusion Technology; Elsevier: Karlsruhe, Germany, 1990; pp 427–430; Jones, P. M. S.; Gibson, R. J. Nucl. Mater. 1967, 21, 353–354. 0.8 Shapovalov and Dukel’skii Swansiger Jones and Gibson 0.6 0.1 S ol ub ili ty (m ol m –3 M P a– 1/ 2 ) 1 10 1 1.2 Temperature, 1000/ T (K–1) 1.4 1.6 1.8 2 Figure 9 Solubility of hydrogen in beryllium. Adapted from Jones, P. M. S.; Gibson, R. J. Nucl. Mater. 1967, 21, 353–354; Shapovalov, V. I.; Dukel’skii, Y. M. Izvestiva Akademii Nauk SSR Metally 5, 201–202; Swansiger, W. A. J. Vac. Sci. Technol. A 1986, 4, 1216–1217. Tritium Barriers and Tritium Diffusion in Fusion Reactors 525 If hydrogen isotopes migrate along the grain bound- aries, it is logical that the rate of migration would be affected by oxygen segregated to those boundaries. The very limited results for hydrogen isotope solubility in beryllium are shown in Figure 9. In the earlier described experiments by Jones and Gibson,74 the solubility was seen to be effectively independent of temperature in the temperature range 550–1250 K. For sintered, distilled a-beryllium, Shapovalov and Dukel’skii77 reported similar values of solubility for the temperature range 673–1473 K. In experiments using 98.5% and 99.8% pure beryl- lium samples, Swansiger78 used gaseous uptake of tritium to determine the solubility. The amount of tritium uptake did not increase with increasing sample size. The solubility for the two purity materi- als was also seen to be the same. For temperatures below 650 K, the apparent solubility increased; this strange effect was attributed to trapping. It is interesting to note the fact that the apparent 526 Tritium Barriers and Tritium Diffusion in Fusion Reactors solubility over temperatures at which the three research groups74,77,78 performed their experiments varied by less than one order of magnitude even though the activation energies varied by 96 kJ mol�1. It should be questioned whether the reported values really represent the solubility of hydrogen isotopes in bulk beryllium. For all plasma-facing materials, there is concern that implantation of energetic deuterium and tritium could lead to excessive retention and permeation. Implantation of hydrogen isotopes into a material with a low recombination-rate constant can lead to a majority of the hydrogen being pushed into the bulk of the material. In the limiting case of slow recombi- nation, 50% of the hydrogen exits the front face and 50% exits the rear face. Langley79 implanted 25 keV deuterium into 99.1% pure hot isostatic pressed beryllium. The retention was seen to be 100% until the particle fluence reached 2� 1022 D m�2. The retention flattened to a limit of �2.8� 1022 D m�2. Wampler80 recorded similar results for his implanta- tion of 0.5 and 1.5 keV deuterium into 99.6% pure beryllium samples. Saturation occurred at 0.31 D/Be in the implant zone. Yoshida et al.81 used 99% pure beryllium in his implantation experiments with 8 keV deuterons. Transmission electron microscopy revealed bubble formation at all temperatures between room temperature and 873K. The bubbles were not removed even by annealing at temperatures up to 973K. Plasma exposure was used by Causey et al.82 and by Doerner et al.83 in low-energy, high-fluence deuterium exposures to beryllium. In both studies, the fractional retention was extremely low and decreased with increasing temperature. Open poros- ity in the implant zone was listed as the likely cause of the low retention. Chernikov et al.84 and Alimov et al.85 showed bubbles and microchannels to be responsible for the behavior of implanted hydrogen in beryllium. At 300K, very small bubbles with a high volume density are formed even at low fluences. As the fluence is increased, the bubbles agglomerate into larger bubbles and then form microchannels that eventually intersect with the surface. For irradiation at 500–700K, small facetted bubbles and large oblate, gas-filled cavities are formed. This microstructure was seen to extend well beyond the implant zone. Alimov et al.85 postulated that the hydrogen retention in the porous region was due to binding to the beryl- lium oxide that forms on the pore surfaces. Beryllium is known as a neutron multiplier because of the reaction 9Be þ n! 8Beþ 2n. Another neutron reaction for beryllium is 9Be þ n ! 4He þ 6He, followed by 6He decaying to 6Li. 6Li has a very large cross-section to absorb a thermal neutron and produce a helium atom and a tritium atom. Baldwin and Billone86 calculated the amount of tritium that could be produced in a large fusion device of the future. In an experiment, they exposed beryllium to a neutron fluence of 5� 1026 n m�2 with 6% of the neutrons having energy >1MeV. The resulting tri- tium level was determined to be 2530 appm. Scaling up to a fusion reactor with 50 Mg of beryllium exposed to 3MWy m�2 results in the production of 5.5 kg of tritium. This is a sizeable quantity of tritium. The relevant question is whether this tritium would be released during normal operation of the fusion plant. Baldwin and Billone86 examined exactly that question in their experiment. The samples containing the 2530 appm of tritium were heated in stepped anneals to determine the release rate of tritium from beryllium materials with different densities. The annealing began with a very long anneal at 773 K, and the temperature was increased in incre- ments of 100 K. For each temperature, there was a nondiffusional burst of release followed by a rapid decrease in the release rate. The release behavior for the different materials was similar, but the fractional release was greater for the less dense materials. Andreev et al.87 irradiated hot-pressed beryllium at 373K. After neutron irradiation, thermal desorption spectroscopy was used with a heating rate of 10 K s�1. Release began to occur at �773K. The temperature at which maximum release occurred depended on the neutron fluence. The sample irradiated to a fluence of 3� 1025 nm�2 had a peak release at 1080K, while the sample irradiated to the higher fluence of 1� 1026 nm�2 exhibited a peak release at a lower temperature of 1030 K. The authors examined the microstructure of the samples after the release anneals. If the anneal was stopped at 973 K, pores with a diameter of 2–16 mm were formed. If the anneal was taken to 1373 K, the pore diameters increased to 25–30 mm. Due to the toxicity of beryllium, there have been relatively small numbers of experiments performed on the behavior of hydrogen isotopes in beryllium. The apparent diffusion coefficient of hydrogen in beryllium is strongly affected by purity levels. The values determined for the solubility of hydrogen in beryllium all fall within one order of magnitude even though the apparent activation energy differs by 96 kJ mol�1. Implantation of hydrogen into beryllium results in the formation of bubbles and eventually open channels or porosity. Connection of Tritium Barriers and Tritium Diffusion in Fusion Reactors 527 the porosity to the surface facilitates the release of hydrogen from the beryllium as the particle fluence is increased. The tendency to form bubbles would suggest that the solubility of hydrogen in beryllium is extremely small. It is possible that the values determined for the solubility of hydrogen in beryl- lium actually represent the amount of hydrogen absorbed on the external surface and on the grain boundaries. The measured diffusivity may represent migration along the grain boundaries. More experi- ments, and experiments with single crystals, are needed to answer these questions. For beryllium used for long times in future fusion devices, tritium produced by neutron reactions on the beryllium is likely to dominate tritium retention in beryllium. Tritium inventory from eroded beryllium codepos- ited with tritium may play a strong role in tritium inventory, but that effect is not covered in this review. 4.16.3.2 Structural Materials Structural materials for a fusion reactor are simply those that comprise a majority of the plant. They are not directly exposed to the plasma, but most are exposed to high doses of neutrons and electromag- netic radiation. Many of these materials are used in the reactor blanket where the tritium is bred by the nuclear reaction 6Li(n,a)3H. It is in the blanket and in the fuel reprocessing area that the structural materials are most likely to be exposed to tritium. The following sections review the structural materi- als that have been considered for fusion reactors. 4.16.3.2.1 Austenitic stainless steels Austenitic stainless steels, particularly type 316, have been used extensively as a construction mate- rial for nuclear reactors (see Chapter 2.09, Proper- ties of Austenitic Steels for Nuclear Reactor Applications). The type 300-series austenitic stain- less steels (Fe–Cr–Ni) have relatively high nickel content (8–12 wt% for the 304 family of austenitic stainless steels and 10–14 wt% for 316 alloys), which is a detriment for fusion applications for several reasons including the susceptibility of nickel to acti- vation (induced radioactivity).88–90 The solubility and the diffusivity of gaseous hydrogen and its iso- topes through type 300-series austenitic stainless steels have been extensively studied and reviewed in San Marchi et al.1 Higher strength austenitic stainless steels (such as the Fe–Cr–Ni–Mn alloys, which have not been widely considered for fusion applications) feature solubility and diffusivity that differ by a factor of about 2 compared to the type 300-series alloys.1 The so-called prime candidate alloy (PCA) is a variant of type 316 austenitic stain- less steel modified for fusion applications (although interestingly enough with higher nickel content); from a permeation perspective PCA is anticipated to behave in a manner essentially similar to conven- tional type 316 alloys.91 The Fe–Cr–Mn austenitic stainless steels have been considered as a substitute for the more com- mon grades of austenitic stainless steels since they have only a nominal nickel content,88,89,90 although low-activation ferritic/martensitic steels have received more attention (see subsequent sec- tion). Alloys that have been considered typically contain both chromium and Mn in the range 10–20 wt%, often with small amounts of other alloying elements (Sahin and Uebeyli90 provides a list of a number of alloys that have been explored for fusion applications). Unlike the Fe–Cr–Ni austenitic stainless steels, there are few reports of transport properties for the Fe–Cr–Mn austenitic alloys; data for oxidized Fe–16Cr–16Mn are reported in Gro- mov and Kovneristyi.92 Austenitic stainless steels can contain ferritic phases in the form of residual ferrite from alloy production, ferrite in welds formed during solidifica- tion, and in some cases, strain-induced martensite from deformation processing. The ferritic phases can result in a fast pathway for the transport of hydrogen and its isotopes at a relatively low temper- ature because the ferritic phases have a much higher diffusivity for hydrogen and its isotopes than austen- ite.93,94 In the absence of ferritic second phases, how- ever, hydrogen transport in austenitic stainless steels is independent of whether the material is annealed or heavily cold-worked95–97 and relatively insensitive to composition for the type 300-series alloys.1 Reported values of hydrogen diffusivity in aus- tenitic stainless steels are less consistent than per- meability as a consequence of surface effects and trapping, as mentioned earlier and elsewhere.1 Figure 10 shows the reported diffusivity of hydro- gen from a number of studies in which special pre- cautions were taken to control surface conditions. The activation energy for diffusion is relatively large for austenitic stainless steels (ED ¼ 49.3 kJ mol�1), and thus the diffusivity is sensitive to temperature, approaching the values of the ferritic steels at very high temperatures (>1000K), while being many orders of magnitude lower at room temperature. 21 10–12 10–11D iff us iv ity (m 2 s– 1 ) 10–10 10–9 1.2 Temperature, 1000/T (K–1) 1.4 1.6 1.8 Figure 10 Diffusivity of hydrogen in austenitic stainless steels from gas permeation studies that confirmed diffusion-limited transport. The bold line represents the average relationship determined in Perng and Altstetter93 for several austenitic stainless steels. Adapted from Quick, N. R.; Johnson, H. H. Metall. Trans. 1979, 10A, 67–70; Gromov, A. I.; Kovneristyi, Y. K.Met. Sci. Heat Treat. 1980, 22, 321–324; Perng, T. P.; Altstetter, C. J. Acta Metall. 1986, 34, 1771–1781; Louthan, M. R.; Derrick, R. G. Corrosion Sci. 1975, 15, 565–577; Sun, X. K.; Xu, J.; Li, Y. Y. Mater. Sci. Eng. A 1989, 114, 179–187; Grant, D. M.; Cummings, D. L.; Blackburn, D. A. J. Nucl. Mater. 1987, 149, 180–191; Grant, D. M.; Cummings, D. L.; Blackburn, D. A. J. Nucl. Mater. 1988, 152, 139–145; Mitchell, D. J.; Edge, E. M. J. Appl. Phys. 1985, 57, 5226–5235; Kishimoto, N.; Tanabe, T.; Suzuki, T.; et al. J. Nucl. Mater. 1985, 127, 1–9. 528 Tritium Barriers and Tritium Diffusion in Fusion Reactors The exceptionally low diffusivity of hydrogen near room temperature results in austenitic stainless steels having significantly lower permeability of hydrogen than other structural steels. The solubility of hydrogen and its isotopes in the type 300-series austenitic stainless steels is high rela- tive to most structural materials. Compilation of data from gas permeation studies shows that most studies are consistent with one another,93,95,96 while studies that considered a variety of alloys within this class show that the solubility of hydrogen is essentially the same for a wide range of type 300-series austenitic stainless.93,95,96 The heat of solution of hydrogen in austenitic stainless steels is relatively low (DHs ¼ 6.9 kJ mol�1), and thus the equilibrium content of hydro- gen in the metal remains high even at room tempera- ture. The solubility of hydrogen and its isotopes is plotted in Figure 11, while Table 1 lists the recom- mended transport properties for austenitic stainless steels (and a number of other metals and alloys). The primary traps in type 300-series austenitic stainless steels are dislocations with relatively low binding energy �10 kJ mol�1.112 Therefore, the amount of trapped hydrogen (in the absence of irra- diation and implantation damage) is relatively low at elevated temperatures. Moreover, due to the high sol- ubility of hydrogen and its isotopes in austenitic stainless steels, the density of trapping sites would need to be impractically high to measurably increase the inventory of hydrogen and its isotopes in the metal.20 For these reasons, trapping from a microstruc- tural origin is anticipated to have little, if any, impact on the transport and inventory of hydrogen and its isotopes in austenitic stainless steels at temperatures greater than ambient. The recombination-rate constant (kr) for austenitic stainless steels near ambient temperature is typically less than about 10�9 m4 s�1 per mol of H2. 113 At higher temperatures (�700K), the value varies between �10�5 and 10�7 m4 s�1 per mol of H2, depending on the surface condition.80,113–116 4.16.3.2.2 Ferritic/martensitic steels There is significant interest in reduced activation ferritic/martensitic (RAFM) steels to replace nickel- bearing austenitic stainless steels in reactor applica- tions117. There are many RAFM steels that have been proposed and investigated in the literature specifi- cally for fusion applications; these typically contain between 7 and 12 wt% chromium, relatively low carbon ( 21 1.2 10 S ol ub ili ty (m ol m – 3 M P a– 1/ 2 ) 100 Temperature, 1000/T (K–1) 1.4 1.6 1.8 Figure 11 Solubility of hydrogen in austenitic stainless steels from gas permeation studies that confirmed diffusion-limited transport. The bold line represents the average relationship determined in Perng and Altstetter93 for several austenitic stainless steels. Adapted from Quick, N. R.; Johnson, H. H. Metall. Trans. 1979, 10A, 67–70; Gromov, A. I.; Kovneristyi, Y. K. Met. Sci. Heat Treat. 1980, 22, 321–324; Perng, T. P.; Altstetter, C. J. Acta Metall. 1986, 34, 1771–1781; Louthan, M. R.; Derrick, R. G. Corrosion Sci. 1975, 15, 565–577; Sun, X. K.; Xu, J.; Li, Y. Y. Mater. Sci. Eng. A 1989, 114, 179–187; Grant, D. M.; Cummings, D. L.; Blackburn, D. A. J. Nucl. Mater. 1987, 149, 180–191; Grant, D. M.; Cummings, D. L.; Blackburn, D. A. J. Nucl. Mater. 1988, 152, 139–145; Mitchell, D. J.; Edge, E. M. J. Appl. Phys. 1985, 57, 5226–5235; Kishimoto, N.; Tanabe, T.; Suzuki, T.; et al. J. Nucl. Mater. 1985, 127, 1–9. Tritium Barriers and Tritium Diffusion in Fusion Reactors 529 and niobium content). The transport of hydrogen and its isotopes has been extensively studied in MANET (MArtensitic for NET, including the so-called MANET II) and modified F82H (generally referred to as F82H-mod). Some of the other desig- nations of RAFM steels that can be found in literature include EUROFER 97, Batman, OPTIFER-IVb, HT- 9, JLF-1, and CLAM steel. In general, studies of RAFM steels report rela- tively consistent transport properties of hydrogen and its isotopes; some of these studies are reviewed in Serra et al.118 Despite the consistency of the data available in literature from several research groups, few studies verify the expected pressure dependence of the transport properties that is expected for diffusion- controlled transport. Pisarev and coworkers119,120 have suggested that the literature data may underestimate diffusivity and solubility due to surface limited trans- port. Similar suggestions have been presented to explain some of the data for the austenitic stain- less steels1; however, the work on austenitic stainless steels has been cognizant of the issues with surface effects; generally surface effects are mitigated by coating specimens with palladium or other surface catalyst. Such precautions have not been systemati- cally employed for permeation studies of the RAFM steels, although the need to control the surface condition (and confirm the square root dependence on pressure) has been widely acknowl- edged.29,30,118,121 While the apparent transport prop- erties in the absence of trapping are relatively consistent for all the RAFM steels, the issue of sur- face effects and the suggestions of Pisarev et al. need further validation in the literature because the trans- port of tritium is less likely to be affected by surface conditions compared to deuterium and protium. The diffusivity of hydrogen is shown in Figure 12 along with an average relationship (Table 1). The literature data are generally within a factor of 2 of the average relationship. The MANET alloys tend to have lower diffusivity of hydrogen and its isotopes than F82H-mod. Differences in permeability between these two alloys has been attributed to Chromium content;29,30 however, a clear correlation of transport properties with Chromium content cannot be estab- lished on the basis of existing data.122 At temperatures less than about 573K, the apparent diffusivity is signif- icantly less than the exponential relationship extrapo- lated from higher temperatures. This is attributed to the effect of trapping on the transport of hydrogen and its isotopes. The reported values of apparent solubility of hydrogen and its isotopes in RAFM varies very little in the temperature range from 573 to 873K. Pisarev Table 1 Recommended diffusivity and solubility relationships for protium in various metals and classes of alloys in the absence of trapping Alloy Diffusivity Solubility, F/D References D ¼ D0 exp (�ED/RT) K ¼ K0 exp (�DHs/RT) D0 (m 2 s�1) ED (kJ mol �1) K0 (mol H2 m �3 MPa�1/2) DHs (kJ mol �1) Beryllium 3�10�11 18.3 18.9a 16.8a 74, 43 5.9�106a 96.6a 78 Graphite 9�10�5 270 19 �19.2 43 Aluminum 2�10�8 16 46 39.7 98, 99 Vanadium 3�10�8b 4.3b 138 �29 100, 101 RAFM steelsc 1�10�7 13.2 436 28.6 Austenitic stainless steel 2�10�7 49.3 266 6.9 93 Nickel 7�10�7 39.5 564 15.8 102 Copper 1�10�6 38.5 792 38.9 103 Zirconium 8�10�7 45.3 3.4�107 35.8 104, 105 Molybdenum 4�10�8 22.3 3300 37.4 106 Silver 9�10�7 30.1 258 56.7 107, 108 Tungsten 6�10�4 103.1 1490 100.8 53 Platinum 6�10�7 24.7 207 46.0 109 Gold 5.6�10�8 23.6 77 900d 99.4d 111 aPer text, the solubility of hydrogen in beryllium is very low and there is not good agreement between the few studies of the material. bData for isotopes other than protium does not scale as the square root of mass. cValues are averaged over the data presented in Figures 12 and 13. dEstimated using the permeability from Caskey and Derrick110 and the quoted diffusivity. 530 Tritium Barriers and Tritium Diffusion in Fusion Reactors and coworkers report values that are three to four times higher on the basis of their assessment of sur- face effects. Here we recommend a relationship for the apparent solubility (Table 1) that is consistent with the majority of the literature data with DHs ¼ 28.6 kJ mol�1, which is based on a simple curve fitting of the data shown in Figure 13. The values of the solubility are about an order of magnitude less than the austenitic stainless steels in the temperature range between 500 and 1000 K, although the solubil- ity of hydrogen is more sensitive to temperature for the RAFM steels since DHs is four times the value for the austenitic stainless steels. The trapping characteristics of the RAFM steels have been estimated for several alloys.19,118,121,123–126 Although binding energies and densities of hydrogen traps vary substantially, the majority of reported values for RAFM steels are in the range 40–60 kJ mol�1 and 10�3–10�5 traps per metal atom, respectively. The traps are attributed primarily to boundaries118 and result in a significant reduction in the apparent diffu- sivity at temperatures less than about 573K. At higher temperatures, the traps are essentially unoccupied and do not affect diffusion.20 The measured recombination coefficient is many orders of magnitude lower than theoretical predic- tions; moreover, the measured values can also vary substantially from one study to another.118,127,128 Measured values for the recombination coefficient for deuterium on MANET alloys are approximately in the range 10�2–10�4m4 s�1 per mol of H2 for the temperature range 573–773K.127,128 Oxidation of MANETwas shown to induce surface-limited trans- port of deuterium and reduce the recombination coefficient kr � 10�6 m4 s�1 per mol of H2.128 Furthermore, it is suggested that structure and composition of the oxide may also affect the recom- bination coefficient and that oxidation can increase the energy barrier associated with dissociation of the gaseous diatomic hydrogen isotopes.128 In summary, the diffusivity and the solubility of hydrogen and its isotopes are consistently similar for all the RAFM steels that have been tested for fusion applications. RAFM steels show a relatively rapid diffusion and low solubility of hydrogen and its isotopes at ambient temperature. The diffusivity is six orders of magnitude greater than that of the austenitic stainless steels at 300 K, while the solubility is more than three orders of magnitude lower than that of the austenitic stainless steels. The diffusivity of hydrogen and its isotopes is not strongly sensitive to temperature compared to most other metals. On the other hand, the heat of solution (DHs) for the RAFM steels is quite large, 21 D iff us iv ity (m 2 s– 1 ) 10–9 10–8 10–7 1.2 Temperature, 1000/T (K–1) 1.4 1.6 1.8 Figure 12 Diffusivity of hydrogen in reduced activation ferritic/martensitic steels from gas permeation studies that confirmed diffusion-limited transport. The bold line represents an approximate relationship estimated from the plotted data. Adapted from Forcey, K. S.; Ross, D. K.; Simpson, J. C. B.; et al. J. Nucl. Mater. 1988, 160, 117–124; Serra, E.; Perujo, A.; Benamati, G. J. Nucl. Mater. 1997, 245, 108–114; Serra, E.; Benamati, G.; Ogorodnikova, O. V. J. Nucl. Mater. 1998, 255, 105–115; Pisarev, A.; Shestakov, V.; Kulsartov, S.; et al. Phys. Scripta 2001, T94, 121–127; Esteban, G. A.; Perujo, A.; Douglas, K.; et al. J. Nucl. Mater. 2000, 281, 34–41; Dolinski, Y.; Lyasota, I.; Shestakov, A.; et al. J. Nucl. Mater. 2000, 283–287, 854–857; Kulsartov, T. V.; Hayashi, K.; Nakamichi, M.; et al. Fusion Eng. Des. 2006, 81, 701–705. 21 1 10 S ol ub ili ty (m ol m –3 M P a– 1/ 2 ) 1.2 Temperature, 1000/T (K–1) 1.4 1.6 1.8 Figure 13 Solubility of hydrogen in reduced activation ferritic/martensitic steels from gas permeation studies that confirmed diffusion-limited transport. The bold line represents an approximate relationship estimated from the plotted data. Adapted from Forcey, K. S.; Ross, D. K.; Simpson, J. C. B.; et al. J. Nucl. Mater. 1988, 160, 117–124; Serra, E.; Perujo, A.; Benamati, G. J. Nucl. Mater. 1997, 245, 108–114; Serra, E.; Benamati, G.; Ogorodnikova, O. V. J. Nucl. Mater. 1998, 255, 105–115; Pisarev, A.; Shestakov, V.; Kulsartov, S.; et al. Phys. Scripta 2001, T94, 121–127; Esteban, G. A.; Perujo, A.; Douglas, K.; et al. J. Nucl. Mater. 2000, 281, 34–41; Dolinski, Y.; Lyasota, I.; Shestakov, A.; et al. J. Nucl. Mater. 2000, 283–287, 854–857; Kulsartov, T. V.; Hayashi, K.; Nakamichi, M.; et al. Fusion Eng. Des. 2006, 81, 701–705. Tritium Barriers and Tritium Diffusion in Fusion Reactors 531 and thus the solubility of hydrogen approaches that of austenitic stainless steels at temperatures >1000 K. Consequently, at elevated temperatures (e.g., >700K), the permeability is less than an order of magnitude greater than that of the austen- itic stainless steels and within a factor of 5 at tem- perature >1000K. Trapping is significant in the RAFM steels at temperatures less than about 532 Tritium Barriers and Tritium Diffusion in Fusion Reactors 573K, and thus the apparent diffusivity is much lower than expected from tests that are performed at higher temperatures. 4.16.3.2.3 V–Cr–Ti alloys Body-centered cubic (bcc)-structured V–Cr–Ti alloys (particularly composition ranges of around V–4Cr–4Ti and V–15Cr–5Ti) have low neutron cross-sections and the isotopes that do form with neutron capture have short half lives (51V has a half-life of 0 10–12 10–11 10–10 D iff us iv ity (m 2 s– 1 ) 10–9 10–8 10–7 2 4 Temperature, 1000/T (K–1) 6 8 10 Figure 14 Diffusivity of hydrogen in vanadium and its alloys. The bold line represents the relationship for pure vanadium, reported in Freudenberg et al.100 Adapted from Schaumann, G.; Völki, J.; Alefeld, G. Phys. Status Solidi B 1970, 42, 401–413; Cantelli, R.; Mazzolai, F. M.; Nuovo, M. J. Phys. Chem. Solid. 1970, 31, 1811–1817; Tanaka, S.; Kimura, H. Trans. Jpn. Inst. Met. 1979, 20, 647–658; Eguchi, T.; Morozumi, S. J. Jpn. Inst. Met. 1977, 41, 795–802; Hashizume, K.; Masuda, J.; Otsuka, K. T.; et al. Fusion Sci. Technol. 2008, 54, 553–556; Klepikov, A. K.; Romanenko, O. G.; Chikhray, Y. V.; et al. Fusion Eng. Des. 2000, 51–52, 127–133; Lottner, V.; Heim, A.; Springer, T. Zeitschrift für Physik B 1979, 32, 157–165; Masuda, J.; Hashizume, K.; Otsuka, T.; et al. J. Nucl. Mater. 2007, 363–365, 1256–1260; Pine, D. J.; Cotts, R. M. Phys. Rev. B 1983, 28, 641; Freudenberg, U.; Völkl, J.; Bressers, J.; et al. Scripta Metall. 1978, 12, 165–167; Qi, Z.; Volkl, J.; Lasser, R.; et al. J. Phys. F 1983, 13, 2053–2062; Boes, N.; Züchner, H. Phys. Status Solidi A 1973, 17, K111–K114; Anderl, R. A.; Longhurst, G. R.; Struttmann, D. A. J. Nucl. Mater. 1987, 145–147, 344–347; Romanenko, O. G.; Tazhibaeva, I. L.; Shestakov, V. P.; et al. J. Nucl. Mater. 1996, 233–237, 376–380; Fujii, K.; Hashizume, K.; Hatano, Y.; et al. J. Alloys Compd. 1998, 270, 42–46; Hashizume, K.; Masuda, J.; Otsuka, T.; et al. J. Nucl. Mater. 2007, 367–370, 876–881; Heller, R.; Wipf, H. Phys. Status Solidi (a) 1976, 33, 525–529. Tritium Barriers and Tritium Diffusion in Fusion Reactors 533 increase the lattice parameter of vanadium.133 How- ever, titanium is a much stronger trap than other elements that increase the lattice parameter as much or more (including niobium, molybdenum, and zirconium).133 Pine and Cotts139 assert that tita- nium solute atoms trap not only hydrogen isotopes at nearest-neighbor interstitial sites, but also hydrogen substitutionally. They demonstrated that the binding energy varied from 3 kJ mol�1 in V–3Ti to 9.84 kJ mol�1 in V–8Ti. The trapping energy for D is larger than that for hydrogen for both alloys. However, it should also be noted that there is considerable short- range ordering in V–Ti alloys with more than �4 at. % Ti.133 This ordering means that trapping will not obey an Oriani-type behavior, in which trapping would be linearly dependent on the number of solute atoms, because the solid solution is not random. The elements chromium, iron, and copper all reduce the lattice parameter of vanadium and the diffusivity change in alloys containing these elements is also much lower than that in alloys with Ti.133 In fits to the apparent diffusivity in tritium diffusion experi- ments in ternary V–Cr–Ti alloys, Hashizume et al.135 show that, in addition to single titanium atoms, the most likely secondary trap is not chromium. Instead, the secondary trap has much higher energy and a lower concentration when compared to the monomer titanium trap. They also speculated that this was due to solute dimers and larger clusters. Interstitial oxygen, carbon, and nitrogen are also com- mon in vanadium alloys. One or more hydrogen atoms bind with single carbon or nitrogen atoms readily, and oxygen atoms tend to trap at least two hydrogen atoms each.144 Other defects, such as dislocations, may still be effective traps at 773 K.145 As with other materials, vanadium can be damaged by radiation, and this will likely be the dominant trap in fusion reactors.146,147 The recombination coefficient for hydrogen is over five orders of magnitude slower in vanadium than in nitrogen in the range of operating tempera- tures, and is relatively insensitive to the surface concentration of sulfur.129 Because of this and the high diffusivity of tritium, release is recombination limited in vanadium alloys. Deuterium ion-driven permeation experiments148 of V–15Cr–5Ti have Temperature, 1000/T (K–1) 0.8 1000 104 S ol ub ili ty (m ol m –3 M P a– 1/ 2 ) 105 106 1 1.2 1.4 1.6 1.8 2 Figure 15 Solubility of hydrogen in vanadium and its alloys. The bold line represents the relationship for pure vanadium, reported in Steward.101 Adapted from Klepikov, A. K.; Romanenko, O. G.; Chikhray, Y. V.; et al. Fusion Eng. Des. 2000, 51–52, 127–133; Heller, R.; Wipf, H. Phys. Status Solidi (a) 1976, 33, 525–529; Steward, S. A. Review of Hydrogen Isotope Permeability Through Materials; Lawrence Livermore National Laboratory: Livermore, CA, 1983; Buxbaum, R. E.; Subramanian, R.; Park, J. H.; et al. J. Nucl. Mater. 1996, 233–237, 510–512; Maroni, V. A.; Van Deventer, E. H. J. Nucl. Mater. 1979, 85–86, 257–269; Zaluzhnyi, A. G.; Tebus, V. N.; Riazantseva, N. N.; et al. Fusion Eng. Des. 1998, 41, 181–185. 534 Tritium Barriers and Tritium Diffusion in Fusion Reactors estimated the recombination-rate coefficient to be 2.4� 10�29 m4 s�1 (although this measurement is three orders of magnitude lower than measurements on more dilute alloys149 and two orders of magnitude higher than measured in pure vanadium150). It should be noted that most measured recombination rates are lower bounds due to surface oxides. In environ- ments in which this native oxide layer may be dam- aged (such as by radiation in a fusion reactor), the actual recombination rate may be higher.147 V–Cr–Ti alloys have hydrogen permeabilities that are at least two orders of magnitude more than nearly any other blanket material and form detrimental hydrides.129,130,151–156 The ongoing stud- ies of permeation barriers may allow mitigation of this significant disadvantage so that V’s positive traits in a high-energy neutron environment can still be utilized. 4.16.3.2.4 Zirconium alloys Zirconium alloys are described more fully inChapter 2.07, Zirconium Alloys: Properties and Character- istics. They are used in fusion reactors partly because of their corrosion resistance in aqueous environments and low neutron cross-sections.157 However, zirco- nium readily forms embrittling hydride precipitates. Zirconium alloys oxidize and the surface ZrO2may be an effective permeation barrier, preventing both hydrogen release and formation of detrimental hydrides. Andrieu et al.158 demonstrated that the rate of tritium release of zircaloy-4 (Zry4) decreased sub- stantially upon oxide formation in tritiated water. Zirconium has multiple phases at temperatures of interest: for example, a-, b-, and g-Zr coexist in equilibrium at 833K. Most solubility and diffusivity studies have been conducted on the single-phase a-Zr generally at 773K and below (Figure 16). Above this temperature, zirconium alloys dissolve up to 50 at.% hydrogen and this solubility decreases rapidly with decreasing temperature, causing hydride precipitates within the alloys. The solubility has been found to vary slightly with the alloying content. Yamanaka et al.159 note that the solubility in the b-phase decreases with alloying additions, while the solubility in the a-phase increases with alloying additions. The solubility of hydrogen in ZrO2, regardless of the crystal structure (10�4 to 10�5 mol hydrogen per mol oxide), is much lower than in the base metal and is even lower than that in Al2O3. a-ZrO2 exhibits a solubility almost an order of magnitude lower than b-ZrO2. 160 Greger et al.161 have reviewed hydrogen diffusion in zirconium. The diffusivities reported in studies they cite and in others is plotted in Figure 17. At 623 K, the diffusivity of hydrogen in zirconium is 10�10 m2 s�1,104,158,161–164 while the diffusivity in ZrO2 is only 10 �19 to 10�20 m2 s�1.158,163,165 Austin 1.2 1000 S ol ub ili ty (m ol m –3 M P a– 1/ 2 ) 104 105 106 1.3 1.4 1.5 1.6 1.7 1.8 1.9 2 Temperature, 1000/T (K–1) Figure 16 Solubility of hydrogen in zirconium and its alloys. The bold line represents the average for 13 studies on pure zirconium and Zr-based alloys, reported in Kearns.105 Adapted from Mallett, M. W.; Albrecht, W. M. J. Electrochem. Soc. 1957, 104, 142–146; Kearns, J. J. J. Nucl. Mater. 1967, 22, 292–303; Giroldi, J. P.; Vizcaı́no, P.; Flores, A. V.; et al. J. Alloys Compd. 2009, 474, 140–146; Khatamian, D. J. Alloys Compd. 1999, 293–295, 893–899; Khatamian, D. J. Alloys Compd. 2003, 356–357, 22–26; Khatamian, D.; Pan, Z. L.; Puls, M. P.; et al. J. Alloys Compd. 1995, 231, 488–493; Sawatzky, A.; Wilkins, B. J. S. J. Nucl. Mater. 1967, 22, 304–310; Une, K.; Ishimoto, S.; Etoh, Y.; et al. J. Nucl. Mater. 2009, 389, 127–136; Vizcaı́no, P.; Rı́os, R. O.; Banchik, A. D. Thermochim. Acta 2005, 429, 7–11. 0 10–17 10–16 10–15 10–14D iff us iv ity (m 2 s– 1 ) 10–13 10–12 10–11 10–10 10–9 1 2 3 4 5 6 Temperature, 1000/T (K–1) Figure 17 Diffusivity of hydrogen in zirconium and its alloys. The bold line represents the relationship for pure zirconium, reported in Kearns.104 Adapted from Mallett, M. W.; Albrecht, W. M. J. Electrochem. Soc. 1957, 104, 142–146; Greger, G. U.; Münzel, H.; Kunz, W.; et al. J. Nucl. Mater. 1980, 88, 15–22; Austin, J. H.; Elleman, T. S.; Verghese, K. J. Nucl. Mater. 1974, 51, 321–329; Cupp, C. R.; Flubacher, P. J. Nucl. Mater. 1962, 6, 213–228; Kearns, J. J. J. Nucl. Mater. 1972, 43, 330–338; Gulbransen, E. A.; Andrew, K. F. J. Electrochem. Soc. 1954, 101, 560–566; Kunz, W.; Münzel, H.; Helfrich, U. J. Nucl. Mater. 1982, 105, 178–183; Khatamian, D.; Manchester, F. D. J. Nucl. Mater. 1989, 166, 300–306; Sawatzky, A. J. Nucl. Mater. 1960, 2, 62–68. Tritium Barriers and Tritium Diffusion in Fusion Reactors 535 et al.163 were able to measure the diffusivity in both a- and b-phases by measuring the activity, due to tritium, in tomographic slices of samples. The diffu- sivity values do not have a very strong dependence on crystallographic orientation or on alloy composition. On the basis of observations of tritium segregation to some precipitates,158,164 many authors158,166,167 argue that intermetallic precipitates in zircaloy could be paths for short-circuit diffusion due to large reported values of solubility and diffusivity in 536 Tritium Barriers and Tritium Diffusion in Fusion Reactors some of these phases. However, these quantities appear to be relatively large for the zirconium-matrix material. Further, autoradiography shows depletion in some iron-rich precipitates and at 623K, the diffusivity in ZrFe2 is 2.5� 10�11 m2 s�1, slower than in bare zirconium.168 The permeability values through hydrides might be larger because of the high solubil- ity of hydrogen isotopes in the hydride phase. How- ever, the volume fraction of hydrides tends to be small and the activation energy has been shown to be independent of the presence of the hydride.169 Zirconium alloys that lack an oxide layer are not useful in hydrogen environments that exceed the solubility of hydrogen in zirconium, because of hydride formation. At relatively low use temperatures ( Tritium Barriers and Tritium Diffusion in Fusion Reactors 537 otherwise coated with a second metal, which forms a low-permeability oxide. Such coatings may reduce hydrogen permeability by five orders of magnitude or more.176 Chromia, alumina, and rare-earth oxides have been studied extensively. The low dissociation pressure of Cr2O3 makes it a common native oxide on steels when allowed to form at elevated temperatures and relatively low oxygen partial pressures.177 Chromia is a better barrier (offering a permeation reduction of about an order of magnitude)175,178 than various Cr2MO4 spinels that may also form (M ¼ Ni, Fe, Co).177 Chromia is also present in mixed oxides in chemical densified coatings, which help to give reduction factors of four orders of magnitude.179 Aluminum forms a self-passivating native oxide that has been shown to be resistant to hydrogen iso- tope permeability, because of a very low solubility for hydrogen. Because this layer is very thin (�4 nm) and the hydrogen permeability of the base metal is very low, it remains debatable whether this amorphous native oxide, which can be grown by anodization, has a lower permeability than aluminum or not.180,181 Cleaned stainless steel may be hot-dipped in alumi- num (forming both a relatively pure surface layer and mixed aluminides between the surface and substrate) and then oxidized. This hot-dip aluminizing proces- sing is simple and generally forms coatings that have excellent adhesion properties (although substantially different thermal expansion coefficients),182 which reduce permeation rates by at least one order of Temperatu 0.5 10–22 10–20 10–18 10–16 D iff us iv ity (m 2 s– 1 ) 10–14 10–12 10–10 1 Figure 18 Diffusivity of hydrogen in alumina. The bold line rep single crystal aluminas, reported in Fowler et al.184 Adapted from J. Am. Ceram. Soc. 1977, 60, 155–161; Serra, E.; Bini, A. C.; Co Roberts, R. M.; Elleman, T. S.; Iii, H. P.; et al. J. Am. Ceram. Soc magnitude, and sometimes more than five orders of magnitude.175,183 The basic properties of hydrogen transport in alu- mina have been characterized and are presented in Figures 18 and 19. Roy and Coble185 hot-isostatically pressed high-purity (>99.99%) alumina powders and charged the dense alumina with hydrogen at elevated temperatures to determine solubility. Fowler et al.184 obtained diffusion coefficients for single-crystal, poly- crystalline, and powdered alumina, and for alumina that was doped with MgO. They observed faster diffu- sion in powdered specimens, suggesting that the grain boundaries may provide short-circuit diffusion paths. They also noted that the diffusivity of MgO-doped alumina was four to five orders of magnitude greater than that of pure alumina. This suggests that the purity of barrier coatings matters a great deal and transmuta- tion of barriers in a fusion environment may increase the permeability from the ideal case measured in the laboratory. Yttria and erbia have been deposited on specimens through a number of physical deposition techniques, including plasma spray, arc deposition, and sol–gel deposition.186–188 The advantage of these oxides is not the magnitude of permeation reduction (one to three orders of magnitude), but their high thermal and mechanical stability in a reducing atmosphere. 4.16.3.3.2 Aluminides In addition to forming Al2O3, which is known to decrease hydrogen permeation, aluminization of steels re, 1000/T (K–1) 1.5 2 2.5 resents the average for many sintered, powdered, and Fowler, J. D.; Chandra, D.; Elleman, T. S.; et al. soli, G.; et al. J. Am. Ceram. Soc. 2005, 88, 15–18; . 1979, 62, 495–499. Temperature, 1000/T (K–1) 0.4 0.1 1 10 0.5 Serra et al. Roy and coble 0.6 0.7 0.8 0.9 1 S ol ub ili ty (m ol m –3 M P a– 1/ 2 ) Figure 19 Solubility of hydrogen in alumina. Adapted from Serra, E.; Bini, A. C.; Cosoli, G.; et al. J. Am. Ceram. Soc. 2005, 88, 15–18; Roy, S. K.; Coble, R. L. J. Am. Ceram. Soc. 1967, 50, 435–436. 538 Tritium Barriers and Tritium Diffusion in Fusion Reactors forms aluminide intermetallics that are believed to also lower permeability. Most studies of aluminized samples either have intentionally grown an oxidized layer in order to achieve greater PRFs or at least have not attempted to suppress the formation of surface Al2O3 prior to permeation testing. To our knowledge, no permeation measurements on oxide-free alumi- nides have been performed. However, different pro- cesses lead to oxide scales of differing composition, thickness, and defect density, and the PRF may not be attributable to oxides alone. Steels have been alumi- nized by the hot dip process (described earlier), as well as by various chemical vapor deposition (CVD), spray, packed cementation bed, and hot isostatic press- ing (HIP) techniques. For those techniques that lay down a substantial amount of materials that does not react with the matrix (such as HIP), an aluminum- containing iron alloy can be used in preference to aluminum to offer a higher temperature barrier per- formance.189 Oftentimes, the aluminized layer will be made up of mixed FeAl, FeAl2, Fe2Al3, Fe2Al5, FeAl3, Fe3Al, and even Fe4Al13 intermetallics. 190 Nickel, chro- mium, and mixed-aluminides are also formed.191,192 Due to the aluminum-rich intermetallics on the sur- face, aluminizedmaterial will often have a mixed oxide scale that is rich in Al2O3. 193,194 PRFs are generally larger than for a pure aluminum layer, varying between 10 and 10 000,175,195 while barriers containing a clean Al2O3 surface often have the greatest PRF. 196 It should be noted that aluminum additions can also stabilize ferrite in some austenitic stainless steels, producing a duplexmicrostructure and an increased permeation.197 4.16.3.3.3 Nitrides As with oxides, either the native nitrides of the base metal or depositions of other nitrides can be made to serve as barriers. One of the most common native nitrides is Fe2N, which forms when the upstream face of steel samples are nitrided and reduces permeabil- ity by one to three orders of magnitude.198–200 After oxides and aluminide, TiN coatings are one of the most researched barriers because of their good adhesion and the ease of deposition.175 Reported PRFs for nitride barriers vary widely, from less than an order of magnitude to six orders of magnitude. TiN barriers reduce permeability the most when they are placed on the high-pressure side of samples, and much less change in permeation is observed when they are placed downstream.201,202 Boron nitride has been shown to reduce the permeability of hydrogen in 304SS by one to two orders of magnitude.175,203 It forms both cubic and hexagonal structures. Checchetto et al.204 noted that the hexagonal structures absorb a greater amount of hydrogen isotopes (because of a larger number of trapping sites; particularly dangling B and nitrogen bonds), but also display a greater diffusivity of hydro- gen. The latter effect may be due to preferential diffusion along the a direction of the hexagonal lattice. Bazzanella et al.205 found that a 1.7 mm (Al,Ti)N coating reduces the deuterium permeability of 0.1-mm thick 316L by two to three orders of magnitude. From permeation transients, they speculated that this reduction was primarily due to the very low diffusiv- ity of D in (Al,Ti)N. Tritium Barriers and Tritium Diffusion in Fusion Reactors 539 4.16.3.3.4 Carbides The only report on the solubility of hydrogen in boron carbide is that by Shirasu et al.206 They exposed crystals of boron carbide to hydrogen gas at various temperatures and pressures for 20 h, and subse- quently outgassed them during anneals in which the temperature was linearly increased at the rate of 20 K min�1 up to 1273 K. The uptake was seen to increase with the square root of pressure, and to decrease with increasing temperature (exothermic). Schnarr and Munzel207,208 measured the diffusivity of tritium in both irradiated and unirradiated boron car- bide. While the actual expression for the diffusivity for each case was not given, it can be extracted from the figures. It was noted that the apparent diffusivity decreased with increasing radiation damage until the percentage of 10B exceeded 10%. Elleman et al.209 used the 6Li-neutron reaction to generate tritium profiles in samples of boron carbide. Diffusivity was determined by examining the rate of release of the tritium during isothermal anneals at elevated temperatures. The diffusivity and the solubility of hydrogen in silicon carbide (a material described inChapter 2.12, Properties and Characteristics of SiC and SiC/ SiC Composites and Chapter 4.07, Radiation Effects in SiC and SiC-SiC) have been measured twice by Causey et al.210,211 In the first set of experi- ments,210 various grades of silicon carbide were implanted with tritium, using the neutron reaction with 6Li on the sample surfaces. The diffusivity for each material was then determined by fitting the release curves determined during isothermal anneal to those predicted by the analytical solution to the diffusion equation. The results were seen to differ strongly depending on the type and purity of the silicon carbide. As an example, the measured diffusiv- ity in hot pressed and aluminum-doped a-silicon car- bide was approximately five orders of magnitude greater than that in vapor-deposited b-silicon carbide at 1273K. The lowest diffusivities were reported for vapor-deposited b-silicon carbide and single-crystal a-silicon carbide. In all cases, the activation energy of the diffusivity was >200 kJ mol�1 (suggesting that chemical bonding plays a strong role in the diffusion). For the diffusion of tritium in vapor-deposited silicon carbide, the diffusivity was given as D¼ 1.58� 10�4 exp(�37 000/T ) m2 s�1. Deuterium solubility was also determined for the vapor-deposited silicon car- bide. The valueswere determined by exposing samples at elevated temperatures to deuterium gas followed by outgassing to determine the amount of uptake. Because equilibrium retention was not obtained in the experiments, inherent in the calculations was the assumption that the diffusivity values determined in the implantation experiments were valid in the gas- eous uptake experiments. The amount of uptake was assumed to be the product of diffusivity, the solubil- ity, the sample area, and the square root of pressure. The solubility was given as K ¼ 1.1� 10�3 exp (þ18 500/T ) mol H2 m�2 MPa�1/2. Again, the nega- tive value of the activation energy would suggest chemical bonding of the hydrogen to the host mate- rial. In the later work by Causey et al.,211 vapor- deposited silicon carbide was again tested. In these experiments, the implantation of energetic particles into the silicon carbide was avoided. Samples were exposed to gas containing 99% deuterium and 1% tritium at a temperature of 1573 K for 1 h. The sam- ples were subsequently outgassed at temperatures from 1373 to 1773 K. The outgassing rates were then fitted to release curves predicted by the solution to the diffusion equation to determine the diffusivity. In this case, the diffusivity was given by the expres- sion D ¼ 9.8� 10�8 exp(�21 870/T ) m2 s�1, one to two orders of magnitude faster than the values deter- mined earlier with energetic particles.210 The solu- bility was also determined in this study. Samples were exposed to the deuterium/tritium gas at tem- peratures from 1273 to 1873K for sufficient duration to achieve equilibrium loading. The samples were then outgassed to determine this equilibrium amount. The expression for the solubility in this case was K ¼ 2.2� 10�2 exp(þ7060/T ) mol H2 m�3 MPa�1/2. This solubility is one to two orders of magnitude lower than the one determined in the earlier experi- ments.210 If one assumes the migration of hydrogen in silicon carbide to occur along active sites on the edges of the grains, it is not unexpected that radiation damage produced by the implantation of energetic particles would increase the apparent solubility and proportionately decrease the apparent diffusivity. If hydrogen can exist only on the grain bound- aries by being attached to trap sites, higher trapping means higher apparent solubility. Conversely, higher trapping means slower diffusion. It was the apparent higher solubility on small-grained samples that led Causey et al.211 to propose the trap-controlled grain boundary diffusion model. The permeation of hydrogen isotopes through sili- con carbide has beenmeasured by several groups.212–214 Verghese et al.213 measured the permeation of a hydro- gen/tritium mixture through aKT silicon carbide tube that was manufactured by wet extrusion and sin- tering. The permeability reported for the experiments 540 Tritium Barriers and Tritium Diffusion in Fusion Reactors is given byF¼ 3.8� 108 exp(�66 000/T ) molH2m�1 s�1 MPa�1/2. Sinharoy and Lange212 measured the permeation of hydrogen through a tungsten tube with a CVD coating of silicon carbide. The retarding effect of the tungsten was taken into consideration in the calculation. The recorded permeation for these experiments was F ¼ 2� 10�4 exp(�6830/T ) mol H2 m �1 s�1 MPa�1/2. Yao et al.214 performed perme- ation experiments on a steel sample that had been RF sputter-coated with silicon carbide. The thickness of the coating was estimated to be 1.3 mm and contained several percent oxygen and traces of iron. The coating was seen to decrease the permeation rate of steel by about two orders of magnitude, but did not change the activation energy. In this case, the coating was clearly porous, and the reduction in permeation was simply due to a reduction in the effective permeation surface area. The plot of the permeation values for the vapor- deposited silicon carbide by Causey et al.211 (calcu- lated as the product of diffusivity times solubility), KT silicon carbide by Verghese et al.,213 and CVD silicon carbide by Sinharoy and Lange212 is shown in Figure 20. The differences in the absolute values of the permeability as well as the differences in the activation energy of the process are extreme. It is difficult to even imagine that the values are for the same material. In fact, the materials are not the same. Asmentioned for the original study byCausey et al.,210 differences in impurities play a significant role in determining the behavior of hydrogen in silicon car- bide. If hydrogen does migrate along the grain bound- aries, impurity metals along those grain boundaries Temperatur 0.55 10–14 10–12 P er m ea b ili ty (m ol H 2 m –1 s –1 M P a– 1/ 2 ) 10–10 Sinharoy a Verghese et Causey et 10–8 10–6 0.6 0.65 0.7 Figure 20 Permeability of hydrogen in SiC. Adapted from Caus 1993, 203, 196–205; Sinharoy, S.; Lange, W. J. J. Vac. Sci. Tech Zumwalt, L. R.; Feng, C. P.; et al. J. Nucl. Mater. 1979, 85–86, 1 reduce the fraction of migrating hydrogen chemically bound to the silicon carbide. Likewise, the apparent diffusivity would be much more rapid if hydrogen trapping at the grain boundaries is reduced. In the case of the permeability measured by Sinharoy and Lange,212 it is difficult to believe that the measured permeation is not really controlled by permeation through the underlying tungsten with the specific surface area limited by the porous silicon carbide coating. The activation energy for the permeation in the report by Verghese et al.213 is difficult to under- stand. The value of 555 kJ mol�1 is even greater than the chemical bond of hydrogen to carbon.215 The permeation was seen to vary by as much as an order of magnitude at the same temperature. There is no apparent explanation for the rapid change in perme- ation with temperature. Titanium carbide has also been tested as a perme- ation barrier. Due to adhesion problems with direct deposition on steel, titanium nitride was used as an intermediate layer between the steel and titanium carbide. Forcey et al.202 measured deuterium perme- ation through 3-mm thick layers of TiC and TiN on steel, observing a PRF of ten. For the experiments performed over the temperature range of 550–740K, extended defects were listed as the reason for the relatively small improvement over bare steel. Checchetto et al.201 used ion-beam assisted deposition of TiN–TiC films on steel in their permeation experiments. When the film was deposited on the downstream side, little reduction in permeation was seen. Using the deposited film on the upstream side e, 1000/T (K–1) nd Lange al. al. 0.75 0.8 0.85 0.9 0.95 ey, R. A.; Wampler, W. R.; Retelle, J. R.; et al. J. Nucl. Mater. nol. A Vac. Surf. Films 1984, 2, 636–637; Verghese, K.; 161–1164. Tritium Barriers and Tritium Diffusion in Fusion Reactors 541 did yield a PRF of �50. Shan et al.216 used a CVD process to deposit their 2.5-mm thick film on steel and noted a permeation reduction of five to six orders of magnitude. It is obvious from these three studies that the deposition of theoretically dense thin films is very difficult. There is also the question of cracking of such thin films during thermal cycling. This is dis- cussed later in this chapter. 4.16.3.3.5 Low permeation metals The permeation of hydrogen and its isotopes through many of the transition metals is lower than that dis- played by iron and the ferritic steels; the notable excep- tions include groups 4 and 5 as well as palladium. Figure 21 shows the permeability of several metals; the diffusivity and the solubility are listed in Table 1 for these metals. In general, the activation energy asso- ciated with permeability ðDHs þ EDÞ is larger for the materialswith lower permeability and the permeability tends to converge at elevated temperatures. We do not attempt to comprehensively review the data for nonferrous metals. However, gas permeation studies are considered the standard for transport prop- erties, particularly studies that report permeability, diffusivity, and solubility. Permeation of tritium through metals and alloys was reviewed by Steward.101 4.16.3.3.5.1 Molybdenum Several reviews of the literature on hydrogen transport in molybdenum have noted variability of the transport 10–3 10–6 10–9 10–12 1 1.2 1.4 Temperature, 1000/T (K–1) P er m ea b ili ty (m ol H 2 m –1 s– 1 M P a– 1/ 2 ) 1.6 1.8 2 V RAFM Mo Cu AlAg 1 Figure 21 Permeability of hydrogen in variousmetals using dat for clarity. Adapted from Frauenfelder, R. J. Vac. Sci. Technol. 1 1986, 34, 1771–1781; Freudenberg, U.; Völkl, J.; Bressers, J.; et of Hydrogen Isotope Permeability Through Materials; Lawrence L J. J. J. Nucl. Mater. 1967, 22, 292–303; Kearns, J. J. J. Nucl. Mat 1998, 46, 6337–6349; Louthan, M. R.; Donovan, J. A.; Caskey, G Technol. 1978, 15, 1146–1154; Tanabe, T.; Yamanishi, Y.; Imoto McLellan, R. B. Scripta Metall. 1979, 13, 65–66; McLellan, R. B. Kass, W. J.; O’Keeffe, M. J. Chem. Phys. 1968, 49, 3329–3332; 1962, 17A, 355; Ransley, C. E.; Neufeld, H. J. Inst. Met. 1948, 7 parameters.101,106,118 The reported permeability values are relatively consistent between the majority of studies, while the diffusivity and solubility values range over several orders of magnitude. The results of Tanabe et al.106 are proposed here as they appear to represent nearly upper bounds of both diffusivity and solubility, without overestimating permeability. The study of Tanabe and coworkers also has the advantage that permeability and diffusivity were measured over a wide range of temperature and pressure, confirming the appropriate pressure depen- dencies of permeability and diffusivity for diffusion- limited transport. 4.16.3.3.5.2 Silver The available data for hydrogen permeation through silver are limited. The diffusivity of hydrogen is reported by Katsuta and McLellan.107 McLellan also reports the solubility of Group IB metals from satura- tion experiments.108 Although these saturation experi- ments do not appear to provide reasonable values for other Group IB metals and are not consistent with other reported solubility measurements,217 Stew- ard, nevertheless, suggests estimating the permeability of hydrogen using these reported relationships.108 4.16.3.3.5.3 Platinum There are relatively few gas permeation studies of platinum. Ebisuzaki et al.109 report the permeability, diffusivity, and solubility of both hydrogen and 10–3 10–6 10–9 10–12 P er m ea b ili ty (m ol H 2 m –1 s– 1 M P a– 1/ 2 ) Temperature, 1000/T (K–1) W Au Pt Aus. SS Ni Zr 1.2 1.4 1.6 1.8 2 a from Table 1. Data is distributed across two separate plots 969, 6, 388–397; Perng, T. P.; Altstetter, C. J. Acta Metall. al. Scripta Metall. 1978, 12, 165–167; Steward, S. A. Review ivermore National Laboratory: Livermore, CA, 1983; Kearns, er. 1972, 43, 330–338; Young, G. A.; Scully, J. R. Acta Mater. . R. Acta Metall. 1975, 23, 745–749; Begeal, D. R. J. Vac. Sci. , S. J. Nucl. Mater. 1992, 191–194, 439–443; Katsuta, H.; J. Phys. Chem. Solids 1973, 34, 1137–1141; Ebisuzaki, Y.; Eichenauer, W.; Liebscher, D. Zeitschrift fur Naturforschung 4, 599–620. 542 Tritium Barriers and Tritium Diffusion in Fusion Reactors deuterium through single crystals of high-purity plat- inum. The permeability of hydrogen in platinum is similar to that in copper. 4.16.3.3.5.4 Gold Caskey and Derrick110 report the permeability of deuterium through gold; Begeal103 reports a similar relationship. Diffusivity measurements, however, differ depending on the conditions of the measure- ment and the microstructural state of the gold.110,218 Cold-worked gold tends to give a higher activation for diffusion, suggesting that trapping is active to relatively high temperatures. Caskey and Derrick110 speculate that trapping is related to vacancies. 10–12 10–15 10–18 10–21 10–241 1.2 1.4 Temperatu P er m ea b ili ty (m ol H 2 m –1 s– 1 M P a– 1/ 2 ) Figure 22 Permeability of hydrogen in various ceramics usin De Van, J. H.; Röhrig, D. H.; et al. J. Nucl. Mater. 1999, 273, 102 1992, 27, 2848–2856; Tanabe, T.; Tamanishi, Y.; Sawada, K.; e K. S.; Ross, D. K.; Simpson, J. C. B.; et al. J. Nucl. Mater. 1989 2006, 54, 1525–1532; Perujo, A.; Kolbe, H. J. Nucl. Mater. 199 Mater. 1997, 246, 139–143. Table 2 Recommended diffusivity and solubility relationship of trapping Material Diffusivity Solubili D = D0 exp (�ED/RT) K = K0 e D0 (m 2 s�1) ED (kJ mol �1) K0 (mol Al2O3 1.1�10�8 132 5.5 a-ZrO2 4� 10�18 30.1 2.5�10 UO2 3.7�10�6 59.8 9.6�10 B4C 1.2�10�11 80.8 3.8 SiC 9.8�10�8 182 2.2�10 The diffusivity shown in Table 1 is from Eichenauer and Liebscher,111 while the solubility is estimated from this diffusivity and the permeability reported by Caskey and Derrick.110 4.16.4 Application of Barriers 4.16.4.1 Expected In-Reactor Performance As implied by Figures 21 and 22,Tables 1 and 2 and in the earlier sections, permeation barriers can be used to reduce the effective permeation in laboratory testing.175–179,183,195,196,198–203,205 PRFs from labora- tory experiments have been reported to be many 1.6 1.8 2 Al2O3 SiC UO2 B4C α–ZrO2 re, 1000/T (K–1) g data from Table 2. Adapted from DiStefano, J. R.; –110; Spitzig, W. A.; Owen, C. V.; Reed, L. K. J. Mater. Sci. t al. J. Nucl. Mater. 1984, 122&123, 1568–1572; Forcey, , 161, 108–116; Wolarek, Z.; Zakroczymski, T. Acta Mater. 8, 263, 582–586; Song, W.; Du, J.; Xu, Y.; et al. J. Nucl. s for protium in various nonmetallic materials in the absence ty, F/D References xp (�DHs/RT) H2 m �3 MPa�1/2) DHs (kJ mol �1) 22.5 184, 233 �2 �28.2 160, 163 4 100 234 �29.8 206 �2 �58.7 211 Tritium Barriers and Tritium Diffusion in Fusion Reactors 543 thousands in certain barrier systems.175,183,195,196,201,202 However, while the data available in the open litera- ture are quite limited, there is significant evidence that the effectiveness of the permeation barriers decreases in radiation environments. There were three sets of experiments219–222 performed in the high flux reactor (HFR) Petten reactor in the Neth- erlands. In the first of these experiments219,222 reported in 1991 and 1992, tritium was produced by the liquid breeder material Pb–17Li. Permeation of tritium through a bare 316 stainless steel layer was compared with that through an identical layer cov- ered with a 146-mm thick aluminide coating. Over the temperature range 540–760K, the barrier was reported to decrease the permeation by a factor of 80 compared to the bare metal, that is, PRF ¼ 80. In the LIBRETTO-3 experiments,220 three different permeation barrier concepts were tested with the tritium again produced by the liquid breeder mate- rial. One irradiation capsule for tritium breeding was coated on the outside with a 6–8-mm thick CVD layer of TiC. A second capsule was coated on the inside with a 0.5–1.5-mm thick layer of TiC followed by a 2–3-mm thick layer of Al2O3. The third barrier was an aluminide coating produced by the cementation process. The aluminum-rich layer was �120-mm thick with about 5 mm of Al2O3 on the outside. The single TiC layer reduced the tritium permeation by a factor of only 3.2, the TiC and Al2O3 layer reduced the permeation by 3.4, and the pack cementation aluminide coating reduced the permeation by a factor of 14.7. These are surprisingly small reductions PRFs compared to laboratory experiments. In a third set of experi- ments,221 the tritium production was achieved with the solid ceramic breeder materials. Both double-wall tubes and single-wall tubes with a permeation bar- rier were tested. The double wall configuration had an inner layer of copper. The permeation barrier on the other system was an aluminide coating with a thickness of 7mm. The aluminide coating was reported to be 70 times more effective than the double-wall configuration in suppressing permeation. Unfortunately, different breeder materials were used for the two different experiments, and the results could have been strongly affected by the amount of tritium released from the ceramic as well as the form of release (T2 vs. T2O). The bottom line on the irradiation testing of barriers is that barriers do not perform as well in a reactor environment as expected from laboratory experiments: a PRF > 1000 has not been achieved in reactor environments. 4.16.4.2 How Barriers Work and Why Radiation Affects Them To understand the effects of radiation on the perfor- mance of permeation barriers, we need to first exam- ine how barriers work. For tritium to permeate through a material with or without a coating, the tritium must absorb on the surface, dissociate into atoms, dissolve into the material, diffuse through the material, and then recombine into molecules on the downstream side. In the simple case in which diffu- sion through the structural material is rate limiting, the permeation rate is controlled by the ratio of the permeability and the thickness of the pressure boundary (eqn [16]); as described earlier, the perme- ability is the product of the diffusivity and the solu- bility, which can be thought of as the velocity times capacity. These parameters are dependent on tem- perature, and should not be affected by radiation effects or nominal surface cracking. In the simple case of diffusion-limited permeation through the structural boundary, the experimental result deter- mined in the laboratory cannot be extrapolated to the radiation environment. Permeation barriers, by their basic nature, consist of a thin layer adhered to the structural material. The performance of barriers depends on the integrity of the barrier as well as the physical interaction of the barrier material with tritium. What is it about many barriers and how they operate that causes labo- ratory and reactor data to disagree? In their review, Hollenberg et al.175 considered these aspects of bar- riers and their performance in radiation, proposing three models that describe distinct physics of the interactions between tritium and the barrier material. The most basic model is the Composite Diffusion Model, in which hydrogen transport is diffusion- controlled in both the barrier and the base metal. The steady-state permeation rate (Q1) through a pressure boundary in this case is Q1 ¼ A ffiffiffiffiffiffiffiffi pTT p tB FB þ tM FM ½22� where A is the surface area of the boundary, and the subscripts B and M refer to the barrier and structural metal, respectively. Considering the intent of the barrier, the ratio tB=FB should be much larger than tM=FM, thus the permeation is controlled simply by the permeation through the barrier. The second model proposed by Hollenberg et al. considers the barrier to be effectively impermeable to 544 Tritium Barriers and Tritium Diffusion in Fusion Reactors tritium and is called the Area Defect Model. In this case, hydrogen is transported through the metal, reaching the metal surface through a limited number of cracks or other defects in the barrier layer. The permeation rate for this case is Q1 ¼ AdFM teff ffiffiffiffiffiffiffiffi pTT p ½23� where Ad is the area of the defects and teff is the effective distance the hydrogen isotope must traverse to reach the other side of the metal. The third model proposed by Hollenberg et al. is the Surface Desorption Model, in which case, perme- ation is controlled by the recombination rate of hydrogen isotope atoms into molecules on the back surface and the recombination-limited flux of tritium is described by eqn [19]. Surface desorption does not make sense by itself; as show, it is actually part of the Area Defect Model. As reported by Hollenberg et al.175 and as revealed by a review of the literature on barriers and oxides,196,223–225 the activation energy of permeation is generally not altered by the addition of the barrier layer onto the substrate. This means that, in practice, the permeation process itself is being controlled by the substrate, not the barrier, strongly supporting the Area Defect Model described earlier. In short, the barrier works simply by limiting the area of the metal exposed to the driving pressure. Pisarev et al.226 provide particularly intriguing insights into the effects of cracks on permeation barriers. Their report showed that permeation reduc- tion for the Area Defect Model is difficult to achieve when the distance between defects is not larger than the combined thickness of the barrier and substrate. Inherent in this conclusion is the assumption that the dissociation rate at the defect is sufficiently fast to maintain the equilibrium concentration dictated by Sievert’s law. If this condition is not met, then the activation energy for the process would be that asso- ciated with the dissociation, and not that of perme- ation through the substrate. Thus, barriers that can provide a significant permeation reduction in the laboratory must be essentially defect free. The physics of hydrogen transport in metals with permeation barriers can be further understood by examining the pressure dependence of perme- ation. As discussed earlier, diffusion-controlled per- meation through metals is proportional to the square root of the hydrogen partial pressure. Perujo et al.227 reported that the pressure dependence of permeation through MANET plasma sprayed with aluminum changed from the classic square root dependence to linear as the pressure was decreased below 20 000 Pa. Mcguire228 also noted the transition to near-linear pressure dependence in the pressure range from 200 to 1000 Pa. Linear pressure dependence is symp- tomatic of permeation limited by absorption or recombination. For example, if recombination limits permeation, the concentration of hydrogen in the metal will be almost constant and uniform, and it will be established by equilibrium at the upstream side of the pressure boundary. Thus, Sievert’s law (eqn [7]) can be substituted into eqn [24], leading to linear pressure dependence: Jr ¼ krK 2pTT ½24� While the same relationship will be found if the permeation is limited by absorption on the upstream surface, known values for the recombination-rate constant for MANET can explain the linear pressure dependence seen in permeation measurements.128 The conclusion is that a combination of the Area Defect Model and the Surface Desorption Model is needed to properly model permeation though barrier materials. If barriers work by limiting the area available for the gas to contact the underlying metal surface, and possibly by creating low enough permeation to have recombination even further reduce the permeation, how does radiation affect this process? One possible answer is by increasing the porosity or cracking of the barrier. According to Arshak and Korostynska229 properties of metal oxide materials are directly or indirectly connected to the presence of defects, oxy- gen vacancies in particular. Oxygen vacancies are also known as color centers, and these color centers are stabilized by hydrogen trapped at the defects. The hydrogen can come from preexisting OH� groups or from hydrogen isotopes migrating through the oxide, possibly increased by the enhanced electrical con- ductivity generated by the radiation damage and the oxygen vacancies. While cracking was not considered by Arshak and Korostynska, one can speculate that the radiation damage with increased oxygen vacan- cies and trapped hydrogen would lead to a more brittle oxide layer. In metals, lateral stress from hydrogen or helium trapping can lead to blisters.230 Without the required ductility to allow blistering, the oxide layer could experience significantly increased cracking. The cracking would then increase the area available for hydrogen to reach the metal surfaces. Tritium Barriers and Tritium Diffusion in Fusion Reactors 545 4.16.4.3 Why Barriers Are Needed for Fusion Reactors In this chapter, the materials for the blanket region have been reviewed, and their permeation para- meters described. In this section, the need for barriers is evaluated. For example, consider the tritium migra- tion processes that might be associated with the liq- uid Pb–17Li systems. In a set of experiments, Maeda et al.231 found the solubility of hydrogen in Pb–17Li to be on the order of 10�7 Pa�1/2 atom fraction (�10�6 mol H2 m�3 MPa�1/2). As tritium is pro- duced in the blanket, some of the tritium will be in solution and some will be in the vapor phase. It is the tritium in the vapor phase that will drive the perme- ation through the metal used to contain the liquid. For an 800MW fusion reactor, Maeda et al. state that 1.5 MCi (150 g) of tritium will have to be bred each day. That means that 1.5 MCi of tritium will be flowing around in stainless steel or similar metal tubes at a temperature >600K. To estimate tritium permeation in a generic 800MW plant, we will scale the design parameters proposed by Farabolini et al.232 for a much larger plant. Approximately 10 000 m2 of surface area will be needed for the tubes passing through the liquid Pb–17Li to extract the heat. We will assume that a sufficient number of detritia- tion cycles per day are performed to keep the amount of tritium in the liquid breeder at 10% of the 150 g listed above. Scaling to 800MW, the amount of Pb–17Li will be�750 000 kg. This leads to a molar fraction of tritium equal to 6.7� 10�7. Using the solubility of Maeda et al.231 for Pb–17Li yields a tritium pressure of �45 Pa. Assuming the contain- ment metal to be 1mm of MANET with an alumi- nized coating, a temperature of 700 K and an effective PRF of 1000, a permeation rate of 2.7� 10�10 T2 mols m�2 s�1 will occur.29,118,122,127,128 With the 10 000 m2 surface area, the daily permeation rate is 0.23mol or 1.4 g of tritium per day. To prevent subsequent permeation through the steam generator tube walls, a tritium clean up unit will have to be applied to this helium loop. Because the steam gen- erator tube wall must be thin to permit effective heat transfer, the tritium cleanup loop will have to be extremely effective to limit release of tritium to the environment. This calculation was performed simply to show the extreme need for barriers in the blanket region of fusion reactors. Even with an active detritiation unit and a barrier providing a PRF of 1000, 1.4 g or 14 000 Ci of tritium end up in the cooling system each day. The situation is not much better for the solid breeders. The same amount of tritium will obviously be required for that system. To minimize the tritium inventory in the ceramic breeder materials, temperatures equal to or greater than that of the liquid breeder will be maintained. The tritium will be released into the helium coolant as elemental tritium (T2) and tritiated water (T2O); the relative concentrations of these forms depend on the type of ceramic breeder. The steel or similar containment metal will be exposed to nontrivial pres- sures of tritium gas. We can conclude that effective barriers are needed for the blanket. It is difficult to imagine that, even with double-walled designs, fusion reactor facilities can meet radioactive release requirements for tritium without an effective barrier. 4.16.5 Summary In this chapter, we have presented tritium permeation characteristics and parameters for materials used in fusion reactors. These materials have included those used to face the plasma in the main chamber as well as materials used as structural materials for the main chamber and blanket. A description of the conditions that exist in those locations has also been provided. Reasons were given why direct contact of the plasma with the plasma facing materials would not lead to sizeable quantities of tritium being lost to the envi- ronment or to the cooling system. The same was not concluded for the blanket region. The need for per- meation barriers there was stressed. A number of materials were listed as possible tritium barriers. These materials included a few metals with some- what reduced permeation and a larger number of ceramics with very low tritium permeability. Due to the difficulty of lining large chambers with bulk ceramics, much of the tritium permeation barrier development around the world has been dedicated to thin ceramic layers on metal surfaces. Unfortu- nately, radiation testing219–222 of these materials has shown that these thin layers lose their ability to limit tritium permeation during exposure to radiation damage. It was suggested, but not proved, that this increase in permeation was due to cracking of the ceramics or the increase in defects. Tomake this chapter more useful to the reader with a need for permeation data, tables and plots of the permeation coefficients are provided. 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All rights reserved. 4.17.1 Introduction 552 4.17.2 Functional Requirements 553 4.17.3 Material Selection 554 4.17.3.1 Fabrication and Microstructure 554 4.17.3.2 Advantages and Limitations for Fusion Application 555 4.17.3.2.1 High atomic number: material erosion/melting 556 4.17.3.2.2 Recrystallization 556 4.17.3.2.3 Machinability, mechanical properties, and DBTT 556 4.17.3.2.4 Component fabrication: CTE mismatch with heat sink 557 4.17.3.2.5 Neutron embrittlement 557 4.17.3.2.6 Neutron activation and radiological hazards 558 4.17.3.2.7 Material availability 558 4.17.3.3 Tungsten Grades 559 4.17.4 Influence of In-Service Conditions 561 4.17.4.1 Thermal Shock Resistance 561 4.17.4.1.1 Microstructure, composition, and mechanical properties 561 4.17.4.1.2 Power density and pulse duration 562 4.17.4.1.3 Base temperature 562 4.17.4.1.4 Repetition rate 563 4.17.4.1.5 Thermal shock during off-normal events: disruptions 563 4.17.4.1.6 Thermal shock during normal operation: ELMs 564 4.17.4.2 Thermal Fatigue Resistance 566 4.17.4.2.1 ITER 566 4.17.4.2.2 Prototype and commercial reactors 567 4.17.4.3 Neutron Irradiation 568 4.17.4.3.1 Thermophysical properties and swelling 568 4.17.4.3.2 Mechanical properties 569 4.17.4.3.3 Thermal shock on irradiated W 569 4.17.4.3.4 Thermal fatigue on irradiated W components 570 4.17.4.4 Ion Irradiation and Retention 570 4.17.4.4.1 He-irradiation 570 4.17.4.4.2 Hydrogen-irradiation and retention 571 4.17.4.4.3 Combined loading conditions 575 4.17.5 Conclusion 576 References 576 Abbreviations APS Atmospheric plasma spraying AUG ASDEX-upgrade CFC Carbon fiber composite CTE Coefficient of thermal expansion CVD Chemical vapor deposition DBTT Ductile to brittle transition temperature DEMO Demonstration fusion reactor ECAE Equal-channel angular extrusion ECAP Equal-channel angular pressure ELMs Edge localized modes fpy Full power years FTU Frascati tokamak upgrade (Frascati, Italy) ICRH Ion cyclotron resonance heating IFE Inertial Fusion Experiment 551 552 Tungsten as a Plasma-Facing Material IFMIF International Fusion Materials Irradiation Facility ITER Tokamak, Latin for ‘the way’ JET Joint European Torus (Culham, UK) LPPS Low-pressure plasma spraying MIM Metal injection molding NIF National Ignition Facility (Livermore, CA, USA) PFC Plasma-facing component PFM Plasma-facing material PS Plasma spraying PVD Physical vapor deposition SC Single crystal SPS Spark plasma sintering TEXTOR Tokamak EXperiment for Technology Oriented Research (Jülich, Germany) TZM Ti–Zr–Mo VPS Vacuum plasma spraying Symbols cp The specific heat Tm Melting temperature l Thermal conductivity r Density 4.17.1 Introduction Until the mid-1990s, only few fusion devices used high-Z elements in plasma-facing materials (PFMs).1 These devices either operated at high plasma cur- rents and high plasma densities such as Alcator C-Mod2 and Frascati tokamak upgrade (FTU)3,4 or used high-Z materials only as test limiters such as Tokamak EXperiment for Technology Oriented Research (TEXTOR).5–9 Since then, high Z refractory metals have been attracting growing interest as candidates for PFMs because of their resistance against erosion and the need for low erosion and stability against neutron irradiation.10 Considerable effort has been made to study the behavior of high Z impurities in the core and edge plasmas, erosion/redeposition processes at the limiter/divertor surfaces, hydrogen isotope retention, and on material development and testing. In particular, the modification of ASDEX-upgrade (AUG) into a fully tungsten machine,11–17 which was achieved in 2007, provided positive answers to critical questions on the reliability of tokamak operation with high-Z plasma-facing components (PFCs) and the compatibility with standard and advanced H-mode scenarios and with the available heating methods.10 Among the challenges, for tokamak devices, that still remain are the strong increase of the W source and W concentration resulting from ion cyclotron reso- nance heating (ICRH) and the need for rigorous mod- eling to support the extrapolation of current results to ITER conditions. Clearly, not all questions posed by ITER can be answered by AUG only. For example, the effects of material mixing with Be, the melt behav- ior under transients, or the change of the hydrogen retention due to damage by high-energy neutron irra- diation18 cannot be addressed in AUG. Answers to some of these issues may be provided by the ITER- like wall project in Joint European Torus (JET), which is installing a bulk tungsten component for the strike point and physical vapor deposition (PVD)-W-coated carbon fiber composite (CFC) tiles for the remaining parts of the divertor.19–21 The remaining questions have to be answered by dedicated experiments in other plasma devices or can only be assessed by mod- eling. However, the results obtained so far do not exclude the use of W in ITER as a standard PFM.10 Further investigations related to future fusion power plants such as demonstration fusion reactor (DEMO) have to focus on the minimization of plasma heat loads to the PFCs to increase their lifetime. In particular, transient heat loads caused by instabilities significantly decrease the operation domain of PFCs, due to thermal stresses and consequent enhanced erosion.22 There- fore, it is also important to mitigate all instabilities, such as edge localized modes (ELMs), that cause sig- nificant plasma transient heat losses.23 Plasma scenar- ios need to be developed, such that the conditions for achieving the required fusion yield are maintained in steady state, while at the same time sustaining tolerable heat loads on the PFCs. The above-mentioned upgrades to the JET24 and AUG15 will allow further optimization of the plasma scenarios under these con- ditions, in particular with DEMO relevant tungsten PFCs.25 These investigations will show how the iden- tified deficiencies of W can be overcome or how they have to be dealt with. In addition to the application of tungsten in ITER and in potential future tokamak devices such as DEMO,26–29 tungsten also became an interesting alter- native for the divertor of stellarators, for example, Advanced Reactor Innovation Evaluation Studies – Compact Stellerator (ARIES-CS),30 and as a first Tungsten as a Plasma-Facing Material 553 wall material for inertial fusion devices.31 Due to similar demands on the PFMs during the operation of all these devices, similar problems have to be solved for each application. 4.17.2 Functional Requirements In the current design of the ITER divertor32–34 for the start-up phase, tungsten has been selected as armor for the divertor dome and the upper part of the divertor vertical targets. In addition, due to exces- sive co-deposition of tritium in CFC raising regu- latory concerns related to tritium inventory limits, a full tungsten divertor will be installed before the D–T phase of operation.32 The PFC design for ITER consists of bulk W bonded to an actively pressurized water-cooled Cu alloy heat sink. Here W has no primary structural function. However, due to the operating conditions listed in Table 1, the PFMs face large mechanical loads particularly at the interface to the heat sink material during cyclic steady state heat loads (see Section 4.17.4.2) and at the plasma-loaded surface during transient thermal events (see Section 4.17.4.1). Furthermore, the material response to these loads is influenced by the material damage or degradation due to neutron irradiation (see Table 1, Sections 4.17.4.3.3 and 4.17.4.3.4). Table 1 Operating conditions assumed for the design of the Material Number of replacements Baking temperature (�C) Normal operation Lifetime (number of cycles) Peak surface heat flux (MWm�2) Peak particle flux (1023m�2 s�1) ELM energy density (MJm�2) controlled/uncontrolled ELM duration (ms) ELM frequency (Hz) controlled/uncontrolled Maximum radiation damage (dpa) Operation temperature design window during normal operation Off normal operation: disruptions Peak surface heat load (MJm�2) Duration (ms) rise time/decay time Frequency (%) Source: Federici, G.; Wuerz, H.; Janeschitz, G.; Tivey, R. Fusion Eng. D In Proceedings of the 22nd IAEA Fusion Energy Conference, Geneva, S Whyte, D. G.; et al. Fusion Eng. Des. 2010, 85, 93–108. aSlow transients lasting 10 s up to 20MWm�2 (10%). bWithout replacement. Along with thermally induced loads, the interac- tion of the PFM with the plasma, that is, the hydro- gen isotopes D and T as well as the fusion product He, is of importance (see Section 4.17.4.4) because they have an influence on material erosion and near- surface material degradation. The further development of the ITER design led to four conceptual designs for the DEMO divertor.25,35 These designs include either water (inlet 140 �C/outlet 170 �C) or, due to the higher achievable efficiency, more probably He-cooling (inlet 540 �C/outlet 700 �C). In all cases bulk W is foreseen as the armor material that will have to face peak steady state heat loads of 15MWm�2 in case of the water-cooled design and 10MWm�2 for the He-cooled designs. In contrast to ITER, off-normal events such as disruptions have to be avoided completely and transient thermal events during nor- mal operation, for example, ELMs, have to be miti- gated below the damage threshold of the material (see Section 4.17.4.1). This may be particularly important considering the expected neutron damage that will amount up to 40–60 dpa during the planned operation of the fusion reactor35 leading to a signifi- cant amount of transmutation products.36 However, the main limiting factor is expected to be the materi- al’s erosion leading to a maximum lifetime of 2 years for the divertor armor.35 ITER PFCs during D–T operation Divertor target Divertor baffle/dome CFC/W W �3 �3 240 240 3000–10000 3000–10000 �10a 3 �10 554 Tungsten as a Plasma-Facing Material In comparison to tokamaks, calculations for a device such as ARIES-CS predict steady state heat loads between 5 and 18MWm�2.30,37 Similar to DEMO, a He-cooled W divertor is anticipated with a maximum heat removal capability of 10MWm�2. The design limits for neutron irradiation at the shield of ARIES-CS are up to 200 dpa at 40 fpy (full power years).38 The component lifetime limits are similarly dictated by the material’s expected erosion. Finally, tungsten or more specifically tungsten coatings find their application also in the dry wall concept for inertial fusion devices, for example, the National Ignition Facility (NIF). In future, inertial confinement devices, thermal loads will occur only in the form of transient thermal loads (P¼ 0.1MJm�2, t¼ 1–3 ms, f¼ 5–15Hz, Tbase� 500 �C).31 These are similar to those expected during ELMs and almost identical to those occurring in an X-ray anode39 and, therefore, affect a thin surface layer only. (a) (b) 200 mm 200 mm Figure 1 Light microscopy images of etched cross-sections of (a) a deformed rod and (b) a rolled plate. 4.17.3 Material Selection 4.17.3.1 Fabrication and Microstructure Tungsten and tungsten alloys are commercially available in many forms, for example, as bulk rods, plates and discs, or thin coatings on various kinds of substrates. For each of these tungsten products, opti- mized production routes exist involving mainly pow- der metallurgical techniques for bulk materials and PVD and chemical vapor deposition (CVD) as well as plasma spraying (PS) for coatings. Each of these processes has its own advantages and disadvantages as well as an individual influence on the material’s microstructure and subsequently the material prop- erties. In addition to the fabrication method, the raw materials, the alloying elements and dopants/impu- rities, pre- and postthermomechanical treatments, and the final shape/geometry have a strong impact on the achieved microstructure. Focusing on the powder metallurgy fabrication route, tungsten powder is obtained from ammonium paratungstate ((NH4)2WO4), tungsten oxide (WO3), and tungsten blue oxide (WO3�x) by hydrogen reduc- tion at temperatures in the range of 700–1100 �C. Vari- ous grain sizes can be produced depending on the reduction temperature and the hydrogen dew-point. The purity of the metal powder obtained is >99.97%. In the manufacture of doped or alloyed tungsten pro- ducts, the dopants or alloying elements are either introduced into the raw materials before reduction or they can be added to the metal powder after reduction. Following the reduction stage, the powder is sieved and homogenized. The initial densification of the powder in various plate and rod geometries takes place predominantly through die pressing and cold isostatic pressing. The pressed compacts are subsequently sintered at temperatures between 2000 and 2500 �C (2273–2773K), mostly using furnaces with hydrogen flow. This increases the density and the strength of the pressed blanks.40 After sintering, the products have a rather low density of about 80% of the theoretical value and poor mechanical properties. To increase density and improve mechanical properties, the sintered products are subject to a mechanical treatment such as rolling, forging, or swaging at temperatures up to 1600 �C. Intermediate annealing, leading to recovery and recrystallization, is necessary to maintain sufficient workability. The working temperature can be reduced as the degree of deformation increases. In this way, forged parts such as rods and discs as well as sheets and foils are produced.40 The final step, that is, the mechanical treatment, changes the microstructure from isotropic with grain sizes determined by the initially used powder size into anisotropic. Depending on the deformation method, the grains may show either: � an elongated, needle-like structure along the deformation direction for radially forged rods and uniaxially rolled plates (see Figure 1(a)), or Tungsten as a Plasma-Facing Material 555 � a flat disc-shaped structure for axially forged discs or blanks and cross-rolled plates (see Figure 1(b)). In addition to bulk materials, research and develop- ment is also directed on tungsten coatings. One possi- bility would be the plasma-spraying process, in which powders are injected into a plasma flame, melted, and accelerated toward the (heated) substrate. The depos- ited layers are splat-cooled, leading to a flat disc- shaped microstructure. Depending on the atmospheric conditions, the result may be layers with high porosity and oxygen content (water stabilized and atmospheric plasma spraying, APS, see Figure 2(a))41,42 or low porosity and good thermal contact (low-pressure or vacuum plasma spraying, LPPS/VPS).26,43–47 In contrast, PVD and CVD coatings show a columnar structure perpendicular to the coated sub- strate with grain sizes in the range of the coating thickness (see Figure 2(b)). PVD coatings, which are also used as thin intermediate layers below a plasma-sprayed tungsten top layer,48 are deposits of tungsten vapor on the substrate surface, which is in the source’s line of sight.43,49 CVD coatings are reactions of a W-containing gaseous phase and have the ability to coat complex geometries.6,50–52 In both cases, a high density (�100%) of the coatings is achieved. The coated substrate can be graphite as used for AUG (PS),12,13 CFC as used for the ITER-like wall (a) (b) 50 mm 300 mm Figure 2 Light microscopy images of etched cross- sections of (a) atmospheric plasma spraying W and (b) chemical vapor deposition W on a graphite substrate. project in JET (PVD),21,53–55 or copper and steel as it might be used for first wall applications in future fusion devices (PS, PVD, CVD).44,49,50,56–60 4.17.3.2 Advantages and Limitations for Fusion Application For fusion plasma-facing applications, the most essential properties are thermal conductivity, strength and ductility, thermal shock and thermal fatigue resistance, structural stability at elevated tem- perature, and stability of the properties under neu- tron irradiation. The advantages and disadvantages of tungsten for these conditions are manifold and opposed to each other as shown in Table 2. While the advantages of the material are mainly related to its high temperature-handling capability, the limitations are associated with manufacturing and handling at low temperatures (below ductile to brittle transition temperature, DBTT61–63), plasma com- patibility including neutron irradiation, and radio- logical issues. However, with regard to other potential PFMs, for example, Be (see Chapter 4.19, Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices), CFC (see Chapter 4.18, Carbon as a Fusion Plasma-Facing Material), and Mo, tungsten is still the most promising, offering an advantageous combination of physical properties and, therefore, has become the material of choice for ITER and DEMO. Since this decision was made, R&D efforts for investigating newly developed tungsten grades Table 2 Features of W armor materials Advantages Disadvantages � High melting point � Low erosion (high energy threshold for sputtering) � High thermal stress resistance � High thermal conductivity � Low swelling � Low tritium retention � High Z (low allowed concentration in plasma) � Potential loss of melt layer during transient events � Recrystallization � Poor machinability � High DBTT � High CTE mismatch with Cu or stainless steel heat sink � Neutron embrittlement � Irradiation-induced transmutation � High radioactivity (short-term waste, decay afterheat) � Explosion dust potential � Limited resistance to grain growth 556 Tungsten as a Plasma-Facing Material and alloys that are able to overcome or at least miti- gate some of the above-mentioned disadvantages have significantly increased. 4.17.3.2.1 High atomic number: material erosion/melting As the high atomic number is an intrinsic material property that cannot be changed, the only possibility to avoid plasma contamination by tungsten is to adapt to the loading realities, that is, thermal loads and plasma wall interaction conditions, and the energy of the incident plasma particles. In particular, surface crack formation, loosening of particles, and particle ejection or melting are addressed (see Section 4.17.4). Concerning the latter, the addition of suitable alloying elements or dispersoids (see Section 4.17.3.3) reduces the material’s thermal conductivity causing a reduction of allowed applied heat fluxes. From this point of view only low-alloyed grades should be considered and the best grade is tungsten of high purity. 4.17.3.2.2 Recrystallization Recrystallization is a thermally activated process. Therefore, it is expected that the activation energy of nucleation is dominated by small angle grain boundaries. The activation energy of grain growth is dominated by large angle grain boundaries.64 The temperature of recrystallization depends mainly on the deformation history, that is, the higher the degree of deformation, the lower the recrystallization tem- perature,65,66 and the chemical purity. When heated above the recrystallization temperature, the structure of tungsten is altered due to grain growth causing an increase in DBTT and reducing other mechanical properties, that is, strength and hardness.67 There are several possibilities for increasing the recrystallization temperature. Particle reinforcement and controlled formation of porosity are the best and most investigated options.68 For example, the higher recrystallization temperature of dispersion strength- ened alloys results from the interaction between dispersoids and dislocations during hot-working; the higher the amount of hot-work, the finer are the dispersoid particles and the higher is the recrystalli- zation temperature. During recrystallization, these particles prevent secondary grain growth and conse- quently, the recrystallization temperature of disper- sion strengthened alloys may increase compared to pure W.67 Another example is highly creep-resistant doped/nonsag materials with aligned porosity acting as obstacles for dislocation movement as they are used in the lighting industry.69 Experience shows that incomplete recrystalliza- tion often helps to achieve the desired balance in material properties. If the operating temperature is well known, controlled recrystallization during application might be feasible as well.67 However, for operational conditions in nuclear fusion devices, it is expected that the very high thermal strain rates experienced in the thin layer heated by plasma dis- ruption or any other transient thermal event will significantly affect the material’s microstructure and properties. 4.17.3.2.3 Machinability, mechanical properties, and DBTT Mechanical properties of W strongly depend on vari- ables such as production history, alloying elements, impurity level, thermomechanical treatment, and form of material. Depending on the production his- tory and heat treatment, W and W-alloys could have anisotropic mechanical properties. This is expressed by showing significantly better properties in the direction of elongated grains (by rolling, forging, or due to deposition processes for coatings) but poorer properties in other directions.70 While reported data on single crystals (SCs) (e.g., Gumbsch62) and for isotropic materials (e.g., Kurishita et al.71) give a clear indication of the material’s performance, typi- cally the reported data refer to the best orientation of the material as shown for fusion relevant tungsten grades in numerous publications.44,51,57,72–81 The prop- erties in other directions, particularly the DBTT, could significantly differ.76 This will affect the operational performance, which is reflected by the orientation- dependent thermal shock response.82 Tungsten is a body-centered cubic (bcc) refrac- tory metal, with a comparatively low fracture toughness,61,83 high DBTT, and poor machinability, which is directly correlated to the material’s low ductility and low grain boundary strength.67 How- ever, DBTT is an ill-defined property and depends strongly on purity, alloying elements, thermo- mechanical treatment, and, most essentially, the testing/loading conditions due to its deformation rate dependence.62,63 The obtained values vary over a broad temperature range from room temper- ature (RT) to several hundreds of degrees Celsius. The exact value depends on the stress state, for example, a three-dimensional state of stress in the sample leads to a lower DBTT. Although many other parameters influence the fracture of bcc metals, the DBTT is usually asso- ciated with the thermal activation of dislocation Tungsten as a Plasma-Facing Material 557 kink pairs. Below this characteristic temperature the separation of a screw dislocation into three partial dislocations (which cannot easily recombine and are therefore more or less immobile) is responsible for the brittle behavior. Increasing temperature leads to thermal activation of the kink mechanism and increased ductility due to shielding of the crack tip.84 There is an empirical correlation between tem- perature and activation energy for brittle-to-ductile transitions in single-phase materials suggesting that the ratio between the activation energy and the DBTT gives approximately a value of 25.63 Another factor is the occurrence of interstitial solute elements, such as oxygen, carbon, and nitro- gen, which even in very small amounts tend to segregate at grain boundaries, promoting intergran- ular brittleness and increasing the DBTT. Two ways can be used to get rid of or mitigate the negative effects of interstitial impurities: either a reduction of the grain size,84 to dilute their effect on a larger grain boundary surface, or the complete elimination of grain boundaries, as in SCs. The development of W-alloys essentially follows the first route, as the SC technique, although effective, is too costly. The conventional method to decrease the grain size of tungsten or tungsten alloys is to deform the material at an intermediate temperature, above the DBTT and below the recrystallization tempera- ture.81,84–86 The formation of oxides and carbides of the alloy constituents helps to stabilize the grain boundaries and to dispersion strengthen the matrix at high temperature. Recently, mechanical alloying followed by powder densification has been applied to refractory alloys. Materials with a stabilized fine-grained structure and with the grain boundary strengthened by even finer dispersoids of TiC improve the low-temperature impact toughness of refractory alloys, leading to increased ductility even down to RT and create superplasticity at high temperatures.71,87–89 Another reliable method to increase the ductility at low temperatures and therefore reduce the DBTT is to alloy tungsten with the rather expensive element rhenium, which is a substitutional solute in the W lattice.67,83 As mentioned before, both material deformation and heat treatment influence the DBTT. A heat treatment slightly below the recrystallization tem- perature is able to significantly reduce the DBTT. In contrast, annealing above the recrystallization temperature reduces strength and hardness and increases the DBTT.67 4.17.3.2.4 Component fabrication: CTE mismatch with heat sink A mismatch between the coefficients of thermal expansion (CTEs) can lead to thermal stresses at the interface, which are detrimental to the compo- nent lifetime. This can occur with either Cu-based alloys or steels (steel is more likely to be used in case of coatings) such as that used for water-cooled designs, or to W and W-alloys in the He-cooled design. In particular W and W alloys, in the cold- worked and stress-relieved condition, tend to delam- inate in the direction parallel to the direction of deformation. Such delamination can occur during machining or during operation. To avoid failure due to delamination, the orientation of the texture has to be perpendicular to the surface of the joints,90 raising the question of the suitability of plasma-sprayed W coatings. Two possible options are recommended to mitigate the thermal stresses, that is, reducing the joint interface by introducing castellations or using smaller tiles,91–93 or introducing soft and chemically stable interlayers94,95 or graded layers.96–101 Despite the fact that surface finish has no direct effect on the performance of ITER-related compo- nents,94 it is recommended to avoid possible crack initiators in the armor design, such as castellations ending in the tile and to ensure accurate surface finishing.102–104 Designs that have been proven to reduce the tile and interface thermal stresses and to extend the component lifetime beyond the design limits are the macrobrush or the monoblock. The latter is the reference design for ITER105 because it provides the most reliable attachment and therefore a reduced risk of catastrophic cascade failure.106 Finally, the thermal treatment of Wduring joining manufacturing cycles might have an influence on the material’s properties. While the process tempera- tures during joining of W and Cu do not lead to any significant change of the W properties, in the case of high-temperature brazing of W to W alloys for the He-cooled divertor design,102 the recrystallization temperature of W has to be taken into account. 4.17.3.2.5 Neutron embrittlement There are little data available for irradiated tungsten (see Chapter 1.04, Effect of Radiation on Strength and Ductility of Metals and Alloys). Based on results for other refractory alloys and limited data on tungsten, one would expect neutron irradiation to increase the strength and decrease the ductility of the tungsten armor largely through increases in the DBTT. To minimize the neutron-induced material 558 Tungsten as a Plasma-Facing Material degradation, it is reasonable to limit the operational conditions for components in a neutron environment to temperatures above �900 �C where recovery of tungsten takes place,107 as the ductility loss is more pronounced belowabout 0.3Tm. This is possible in the region close to the plasma-facing surface, but it is impossible in the heat sink region as tungsten will be in contact with materials that cannot operate at this temperature and stress concentrations in these ‘cold’ areas have to be avoided.108 In the case of ITER, Cu will be employed in the heat sink while steel is more likely to be used in DEMO, which has a higher operating temperature.29 Hence, a greater under- standing of the irradiation response of tungsten at temperatures between 700 and 1000 �C is needed.109,110 The effect of embrittlement is alle- viated when operating above 250 �C, although in the presence of He (produced by transmutation reactions) somewhat higher temperatures may be required.83 Although at intermediate temperatures (0.3–0.6Tm), void swelling and irradiation creep are the dominant effects of irradiation, the amount of volumetric swelling associated with void formation in refractory alloys is generally within engineering design limits ( Tungsten as a Plasma-Facing Material 559 4.17.3.3 Tungsten Grades Within current R&D programs for the selection and characterization of candidate grades of W and W alloys for fusion applications, many materials pro- duced according to the schemes outlined above were investigated. These are discussed in the following sec- tion, which introduces some of their characteristics. The manifold production processes described below for pure W are also applicable to W alloys. Pure tungsten (undoped) � Sintered W is the most readily available and cheapest grade with a grain size that depends on the initially used W powder. However, it is characterized by high porosity, low recrys- tallization temperature (1000–1200 �C), and low strength at elevated temperature.96 The option of improving the sinterability by add- ing small amounts of activators (Ni, Fe)121 increases the radiological hazard due to addi- tional activation products that have to be taken into account.119 � Forged or swaged W offers an increased density and a refined microstructure compared to sin- tered material, resulting in higher ductility and mechanical strength. Forging and swaging are therefore the industrial production processes that are typically applied not only for pure tungsten but also for most kinds of tungsten alloys (see below). This grade of W is manu- factured in block shape or more commonly in the form of rods with different diameters (�90mm)40 showing an anisotropic micro- structure122 with elongated grains along the axial direction and an increasing grain size and porosity with increasing rod diameter. Thus, increasing rod diameter leads to a decrease in mechanical strength and ductility. For the production of monoblock tiles, such as those planned for ITER, rods with a minimum diameter of 30–35mm are necessary. � Rolled W is applied in the form of plates or foils with thicknesses from 0.02 to 20mm.40,123,124 It offers a densified but layered microstructure that is strongly anisotropic, with flat disc-shaped grains parallel to the rolled surface affecting the mechanical properties (see Section 4.17.3.2.3) and resulting in the risk of delamination. � Double-forged W is in the form of blanks with a diameter of 140mm and a height of 45mm. The double-forging process, first in the radial and then in the axial direction, provides a more isotropic microstructure than it is generated by single forging. This material should act as a reference grade for establishing a reliable mate- rials database for finite element calculations.82 � SC W provides higher ductility than polycrys- talline W, higher thermal conductivity, lower neutron embrittlement, higher thermal fatigue resistance, and a more stable structure at ele- vated temperatures. The disadvantages are high cost and low industrial availability.96,125,126 � Metal injection molded (MIM)-W127–129 provides a dense and isotropic microstructure with grain sizes on the order of the powder particle sizes used. A final densification by hot isostatic pressing (HIP) at temperatures>2000 �C leads to an improvement of the mechanical proper- ties; recrystallization and grain growth do not play a role. Furthermore, the production pro- cess offers the possibility of net shaping. � Spark plasma sintered (SPS)-W and resistance sin- tering under ultra-high pressure.130–132 The mate- rial is characterized by a short fabrication time of only a few minutes keeping the initial fine microstructure determined by the powders used. The finer the grain size, the higher the microhardness and the bending strength but also the lower the achievable density. The applica- tion of alternatively uni-, two-, or three- directional orthogonally applied forces for the material’s densification during the process leads to internal stresses, which have an influence on the recrystallization behavior. Recrystallization and grain growth occur at �1500 �C. Depend- ing on the amount of porosity, the finer the initial grain size of tungsten, the smaller is the grain growth. � Severe plastically deformed W (and W alloys, see below) with ultra-fine grains in the nm range are produced by either high-pressure torsion at 400 �C84,133 or by the equal-channel angular extrusion or pressure (ECAE or ECAP) pro- cess at high temperatures (1000–1200 �C).134 The material shows stable, that is, deformation- independent, recrystallization temperatures and exhibits considerably enhanced ductility and fracture toughness.61,85,86,135,136 � Plasma-sprayed W involves, in general, application of VPS, more precisely also called low-pressure plasma spraying (LPPS), which provides a significantly reduced oxygen content and improved thermophysical properties com- pared to atmospheric (APS) or water-stabilized 560 Tungsten as a Plasma-Facing Material plasma spraying.42 However, LPPS-W is typi- cally characterized by a lower thermal con- ductivity (up to 60% of bulk tungsten is reported67) and a lower strength than bulk W particularly when deposited on large sur- faces. The recrystallization temperature is similar to pure W.48,137 Although the thick- ness of the plasma-sprayed coatings required for fusion applications are flexible, coat- ings with 200 mm or thicker are commonly produced.26,43,67,138,139 Furthermore, PS is the only production method that offers the possibil- ity to produce and repair W components.57,60,96 � CVD W provides a microstructure with a columnar grain structure parallel to the sur- face, high thermal conductivity similar to bulk W, and a very high density and purity.6,140,141 Thicknesses up to 10mm were produced,67 but its high cost is a significant drawback for prac- tical applications.52,96 � PVD W provides a featureless structure that is extremely dense and pore free. In contrast to plasma sprayed and similar to CVD-coatings, the deposition rates are low. Economic and process-related restrictions generally limit the depositedW thickness to 10–50 mm.13,54,55,67,142 � W foam for Inertial Fusion Experiment (IFE) applications provides structural flexibility dur- ing quasivolumetric loading. The material is microengineered with a relative density of �21% and can be simultaneously optimized for stiffness, strength, thermal conductivity, and active surface area.143 Tungsten alloys � Oxide dispersion strengthened W alloys such as W–La2O3, W–Y2O3, and W–CeO2 with oxide additions �2% are processed by powder met- allurgy methods similar to pureW.40 The insol- uble dispersoids, which are influenced in shape and distribution by the thermomechanical treatments during the production process,73,144 improve the grain boundary strength and machinability and play an important role in controlling recrystallization and the morphol- ogy of the recrystallized grains.68 This results in a higher recrystallization temperature by 100–350 K by suppression of secondary grain growth (i.e., grain boundary migration), lower grain size, higher strength after recrystalliza- tion, and better machinability than sintered W even at RT. This permits fabrication at lower costs.67 The size of the dispersoids in commercially available alloys is �10 mm; how- ever, research on mechanically alloyed materi- als using submicron dispersoids is currently being performed.145 However, the presence of oxide particles with a melting temperature below those of tungsten has a negative effect on the erosion resistance.146,147 � W–3–5% Re is, compared to sintered pure W, characterized by a higher recrystallization tem- perature and strength even after recrystalliza- tion,148 better machinability, and improved ductility at low temperatures.67 The addition of Re, which has a high solubility in W, how- ever, reduces thermal conductivity, increases embrittlement after neutron irradiation, and significantly increases the cost and safety con- cerns because of the high Re activation under neutron irradiation.96 � W–1–2% Mo (þY and Ti) cast alloy. The addi- tion of Mo and the reactive elements Y and Ti, which reduce the amount of free oxygen and carbon and form obstacles to grain growth, improves the mechanical properties compared to large grained pure cast W.67,73 � W–TiC produced by mechanical alloying and slow deformation techniques provides, similar to all other Walloys, higher strength and recrys- tallization temperature, better machinability, and improved ductility compared to pure W with superplastic behavior at temperatures of 1400–1700 �C.89 The addition of Ti-carbide par- ticles stabilizes the grains during the material’s production process. This generates an isotropic grain structure and has the additional effect of keeping a fine grain structure even in the recrystallized condition, but the alloy is more expensive. After recrystallization, the finer dispersoids of TiC particles improve the low- temperature impact toughness of refractory alloys following low-dose neutron irradia- tion.71,87–89,149–154 Other carbides, for example, ZrC155,156 or HfC (in combination with Re and Mo),96 can be used instead of TiC. � K-doped W is a nonsag material that contains a maximum of 40 ppm of potassium.40 Originally known from the lighting industry, it provides high creep strength due to its aligned pore structure, high recrystallization temperature >1600 �C, and good machinability.68,77,78,157 � W–Si–Cr as a ternary or even by the addition of another element as a quarternary alloy is a Tungsten as a Plasma-Facing Material 561 newly developed and not yet optimized mate- rial that is being investigated as a wall protec- tion material due to its favorable oxidation resistance, preventing excessive material ero- sion in case of accidental air ingress.158,159 � Severe plastically deformed W alloys offer, similar to pure tungsten (see above), significantly improved fracture toughness and ductility.61,84 The addition of alloying elements to the starting material (any developmental or commercial produced W alloy), such as Re or dispersoids, leads to an increasing stability of the grains and therefore a higher recrystallization tempera- ture and less grain growth.133 Any of the bulk materials mentioned above could be used and are being investigated in its cold-worked, stress-relieved, or recrystallized state. The latter is of particular interest due to in situ recrystallization of surface near regions during operation.108 In spite of the fact that a large variety of tungsten grades and alloys already exist, the attempts to fur- ther optimize these materials are ongoing. The fabri- cation and successful testing of He-cooled divertor mock-ups for DEMO and ARIES-CS102,160 under a heat flux of 10MWm�2 are important driving forces for the present development ofWalloys with improved performance in the fusion environment.25 However, R&D has to address many different issues related to the performance of the material when exposed to thermal loads, neutron irradiation, and the plasma; these will be discussed in the following section. 4.17.4 Influence of In-Service Conditions 4.17.4.1 Thermal Shock Resistance Tungsten-armored PFCs will be subjected to differ- ent types of heat fluxes dependent on their field of application (see Section 4.17.2). Among others, this includes thermal transient loads (e.g., ELMs and disruptions). The behavior of the material under these conditions, that is, the combination of cyclic steady state and transient heat loads, is a key factor that has to be considered for the selection of a suitable grade of W. The machines simulating these operational conditions are electron and ion beam facilities, quasi- stationary plasma accelerators, plasma guns, and high-energy lasers. A most critical issue is the compa- rability of such simulations. Therefore, a round robin test involving some representative facilities was made for investigating the influence of the different time regimes and different power density levels. The results showed that when compared on the basis of a heat flux parameter P (MWm�2 s1/2), which is directly propor- tional to the temperature increase, the cracking and melting thresholds are almost identical. This permits a direct transfer of the qualitative results obtained in any of these facilities.161 In contrast, quantitative results representative of the operational conditions in large fusion devices can only be obtained when the loads are applied in the desired time range. The reason for this is the heat penetration depth and the related stress field that is produced, which influences crack and melting depth. There are several parameters influencing the thermal shock behavior of tungsten that will be dis- cussed in the following sections for the different materials under disruption and ELM-like loads. 4.17.4.1.1 Microstructure, composition, and mechanical properties During thermal shock loads, steep temperature gradients of hundreds to several thousand degrees Celsius on a length scale of millimeter or even micrometer (depending on the pulse length) are formed, influencing only a limited volume near the loaded surface. While the heat load is applied, due to thermal expansion and the decreasing strength of the material at the surface compared to the bulk material, compressive stresses are formed in the surface plane. These stresses can lead to permanent plastic defor- mation that might, during cool down, generate tensile stresses sufficiently high to initiate crack formation perpendicular to the surface and thereby cause stress relaxation at the surface. Depending on the mechanical properties in the surface plane, the amount and starting point of crack formation can be influenced. Based on this and on the fact that the mechanical properties are strongly dependent on the material’s microstructure (see Section 4.17.3.2.3), a grain orientation parallel to the surface and therefore high strength in the surface plane might be preferred.162 However, grains oriented parallel to the surface, such as in rolled materials or plasma-sprayed coatings, might result in delamination (see Figure 3(a)), which causes over- heating and subsequently surface melting if they have a lower strength in the depth direction and exhibit preferential cracking along the weak grain boundaries. Therefore, a grain orientation perpendicular to the surface and parallel to the direction of the heat (a) (b) 200 mm 500 mm Figure 3 Light microscopy images of the etched cross-sections of thermal shock–loaded specimens with grains oriented (a) parallel and (b) perpendicular to the loaded surface; cracks follow the grain orientation/ deformation direction. 100 mm Figure 4 Light microscopy images of the etched cross-sections of thermal shock–loaded metal injection molding tungsten with isotropic grain structure. 562 Tungsten as a Plasma-Facing Material flow is recommended.90 This will cause cracks to form along the grain boundaries toward the cooling structure (see Figure 3(b)) causing no degradation or only a negligible degradation of the material’s ther- mal transfer capabilities. Due to the lower mechani- cal properties in the surface plane, more or larger cracks will form during thermal shocks, running per- pendicular to the surface and following the grain orientation. In contrast to deformed materials, crack formation and crack orientation in materials with isotropic or almost isotropic grain structures, for example, MIM-W or recrystallized W, is rather unstable and is strongly enhanced for the weakened recrystallized material. Depending on the applied power densities, the formed temperature gradient, and the resultant stress fields within the material, cracks initially running perpen- dicular to the surface might deflect at zones with compressive stresses and keep running parallel to the surface (see Figure 4). 4.17.4.1.2 Power density and pulse duration The material’s response is strongly related to the applied temperature fields and by this to the absorbed power density and the pulse duration. This results in a material-related surface temperature increase and heat penetration depth.163 A classification of the impact of the temperature field is made by establishing three parameters: the damage, the cracking, and the melting threshold. While the latter depends on the ther- mal conductivity and the melting temperature (for alloys or mixed materials formed during tokamak operation) of the material, the damage and cracking threshold are determined mainly by the material’s mechanical properties. Damage here means that the material’s surface has undergone a visible and measur- able modification, for example, by surface roughening, recrystallization, or pore/void formation. 4.17.4.1.3 Base temperature The base temperature influences the thermal shock behavior in two ways. First, a higher base tempera- ture influences the damage, cracking, and melting threshold. All of them are essential because they limit the operational conditions and when exceeded cause enhanced material degradation. Therefore, life- time estimates based on RT data will yield unrealistic conclusions. Second, crack formation strongly depends on the plastic deformation at high temperatures and even more on the stress developed during cool down. To understand the influence of a higher base temperature, one has to be aware of the typical shape of the yield and tensile strength curve for W or a Walloy.105,157 While the decrease in strength is rather high at low temperatures, the curve flattens at high temperatures despite a drop in strength when exceed- ing the recrystallization temperature. As a result, the high temperature plastic deformation induced by the combination of a heated surface and ‘cool’ basematerial can be significantly reduced by a small increase in base temperature. Combining this effect with the increased ductility of Wat the given base temperature, brittle crack formation can be avoided when heating Cracking roughening Homogeneous melting Incident beam Incident beam Melt ejection Boiling and droplet formation Increasing energy density Figure 5 Performance of tungsten and metals in general under transient thermal loads. Tungsten as a Plasma-Facing Material 563 the material above a certain threshold.82,157,164,165 This temperature threshold is related to the DBTT but is not necessarily identical to it. 4.17.4.1.4 Repetition rate In addition to the parameters mentioned above, the damage, cracking, and melting thresholds are deter- mined by the number of load repetitions, because of continuing material degradation such as hardening and recrystallization. This is of particular interest for short transient events with a high repetition rate in magnetic (ELMs) and inertial fusion devices. Up to now the simulation of submillisecond events (ELMs, IFE) has been performed only up to a rela- tively low number of cycles; large numbers of pulses (e.g., >106 ELM pulses during the life-time of the ITER divertor) are not feasible in the majority of the above-mentioned test facilities. 4.17.4.1.5 Thermal shock during off-normal events: disruptions Disruptions still occur frequently in operating tokamaks, and therefore they are also expected for ITER with an anticipated occurrence in �10% of the ITER pulses (3000 pulses per expected compo- nent lifetime). During a disruption in which the plasma undergoes a partial or full thermal quench, most of the plasma thermal energy will be dumped on the divertor plates.166 Taking into account the resultant loading conditions (see Section 4.17.2), significant material loss from the tungsten plasma- facing surface should occur by melting and evapora- tion particularly in the dome area.167,168 In simulating these events, the amount of melting, the melt motion and subsequent roughening of the surface, the mate- rial erosion by droplet emission, the resolidification behavior, and finally, the crack formation occurring in the loaded area or at the boundary between melted and unmelted zone are the most important para- meters to be determined. The underlying mechanisms for the above- mentioned material degradation are well described (see Figure 5).169 Thermal loading of tungsten and metals, in general, at ‘moderate’ energy densities (up to a few MJm�2) will result in a homogeneous, localized melting of the sample surface. When higher energy densities are applied, surface evaporation occurs; the momentum transfer due to evaporating atoms from the surface generates an effective pressure on the melt layer, which finally results in the formation of a melting ridge. Increasing the incident energy density even fur- ther, the material’s response is characterized by intense boiling and convection of the melt layer resulting in droplet formation and ejection.170–172 Open pores in the recrystallized material have a strong impact on the thermophysical properties. The melting threshold and subsequently the amount of melt formation depend on the material’s thermal conductivity, which is lower for porous materials such as plasma-sprayed tungsten, and for tungsten alloys. In particular, it has to be taken into account that dispersoids such as La2O3 (Tm¼ 2578K) have a lower melting temperature than tungsten. This may result in early melting and increased evaporation causing the formation of a porous and depleted sur- face layer, which becomes even more important when applying loads below the melting threshold (see below and Section 4.17.4.1.2). On the other hand, the melting threshold is correlated with the base temperature of the PFM. When the base temperature increases, the melting threshold energy decreases and the amount of melt formation, the obtained cra- ter depth, and the evaporation losses for the same applied loading conditions increase significantly.169 As it cools, the material resolidifies in a recrystallized state providing a columnar grain structure typical of PVD orCVD coatings.With further cooling, depending on the base temperature of the material/component (see ‘Base Temperature’ in Section 4.17.4.1.3), brittle crack formation will not take place above a certain threshold temperature. However, with fast cooling after loading below this temperature, the material will undergo severe cracking with crack lengths that can reach the order of millimeters.169 When the qualification of different W grades and alloys108,141,147 is done in combination with thermal fatigue loading,90 materials with high thermal con- ductivity in combination with superior mechanical properties, that is, with high ductility, performed best with regard to melt material loss and crack formation. This comprises low-alloyed W materials with increased ductility such as W–Re or W–Ta, or fine-grained pure Wor W alloys. 564 Tungsten as a Plasma-Facing Material Disruption simulation experiments on bulk tungsten and tungsten coatings have also been described in the literature. These were performed not only to investigate the melting behavior but also for the purpose of characterizing the cracking behavior.26,42,60,101,122,130,131,162,165,173–176 Although these experiments are more related to those on the characterization of ELM conditions (see Section 4.17.4.1.2) and were often performed only at RT, the results indicate that the use of highly ductile SC materials is preferred.90,177 Alternately, in case of cheaper polycrystalline materials, it is necessary for the material to have the proper microstructure orientation as described above, that is, the grain ori- entation perpendicular to the loaded surface. The reason for this is that crack formation occurs mainly along the grain boundaries and follows the orienta- tion of the deformed microstructure. The crack depth is, in general, related to the applied loading conditions and therefore the pulse length, which determines the heat penetration depth and the temperature and stress gradient induced during loading. The temper- ature gradient also determines the recrystallization zone, which is generated below the loaded area as a function of temperature (1million of events at a frequency of 1–25Hz183) that, although yet unexplored, will impose high demands on the PFMs. While it is the desire of plasma physicists to oper- ate in H-mode regimes with high-energy ELM depo- sition (�1MJm�2), the response of bulk tungsten, tungsten coatings, and tungsten alloys to such loading conditions, that is, surface melting, melt motion, material erosion, and vaporization,170,171,184–189 is detrimental. To obtain further insight into material behavior under these conditions, modeling of experi- mental conditions was carried out.9,167,168,190–195 It has been shown that with regard to melt motion/ erosion, the results of the different facilities cannot be directly compared196 and none of the testing facil- ities used provides identical conditions to those that will occur in a tokamak. However, mitigation techni- ques have been explored for reducing the applied ELM energy, which, in general, can only be done at the expense of a higher repetition rate.183 The extent to which the ELMs have to be mitigated depends on the melt formation at tile edges due to the shallow plasma impact, which was experimentally found to be between 0.4 and 0.6MJm�2 for pure forged tung- sten.189,197 On the other hand, the effect of crack formation during ELMs on the lifetime behavior of the PFCs has to be taken into account. As mentioned before, this behavior is yet unexplored at high repe- tition rates. Typical investigations on various grades of W,82,187,189,198 coatings21,54,186,199 and alloys146,157,187 were in the range of 10–100 repetitions. In a few cases up to 1000 repetitions and in single experiments even on the order of tens of thousands of repetitions have been obtained, depending on the testing facility used. As the repetition rate is still rather low compared to the expected millions of events, the main interest of 0 (a) (b) H ea t flu x p ar am et er (M W m –2 s– 1/ 2 ) H ea t flu x p ar am et er (M W m –2 s– 1/ 2 ) 0 5 10 15 20 25 30 35 40 45 200 Temperature (�C) Damage threshold Cracking th re sh ol d 400 600 Heat flux 0 0 5 10 15 20 25 30 35 40 45 200 Temperature (�C) 400 600 C ra ck in g th re sh ol d Cracks Surface modification No damage Heat flux Figure 6 Thermal shock testing results of double forged W as a function of temperature and the heat flux parameter; grain orientation (a) perpendicular and (b) ‘parallel’ (one direction still perpendicular, indicated by the orientation of the large cracks) to the heat flux. Tungsten as a Plasma-Facing Material 565 these investigations was the qualification of different W grades and alloys (see Section 4.17.3.3) with regard to their damage and cracking thresholds. The characterization was done as a function of the main parameters described in Section 4.17.4.1, that is, microstructure, power density, and base temperature. The results obtained so far showed that crack formation200 vanishes above a certain base tem- perature (see Figure 6).82,157,198 This temperature decreases with increasing material ductility, indicat- ing that the use of W alloys or fine-grained W is preferred. In the case of an anisotropic microstructure, this effect strongly depends on the material’s orienta- tion. Better results are obtained for grain orientations parallel to the loaded surface (see Section 4.17.4.1), yielding differences in the threshold temperature compared to the orthogonal direction of up to several hundred K (cf. Figure 6(a) and 6(b)). Recrystallization leads to a slight homogenization of the material’s microstructure and therefore the mechanical proper- ties; however, there is no full convergency of the orientation-dependent thresholds.82 Despite the fact that for the currently limited number of applied pulses no crack formation was observed above a material and orientation-dependent temperature, the material is still damaged by plastic deformation and surface roughening. The evolution of this plastic deformation and of the related material hardening as a function of the applied number of loads is still unclear and has to be investigated. However, there are also heat load levels (at least up to Tbase� 800 �C), at which no visual material degradation could be determined and the future goal will be to investigate if these damage thresholds are still valid for high repetition rates, at higher base temperatures, and particularly in combination with neutron irradiation (see Section 4.17.4.3) and plasma wall interaction (see Section 4.17.4.4). All the information given above on the effect of ELMs is also directly transferable to the short tran- sient events expected for inertial fusion applications and has been verified by IFE-related tests on dif- ferent W-based materials.201–204 There are coating parameters of high interest besides those mentioned above; these include the manufacturing-induced residual stresses at the surface, which are dependent on the used substrate, and the coating thickness. As mentioned in Section 4.17.4.1.1, the applied loading conditions and therefore the pulse length determine the heat penetration depth.163 As a result, the tem- perature and stress gradient induced under IFE applications should be similar to those in X-ray anodes (see Section 4.17.2). In case of thin coatings, residual and induced stresses might affect the coating to substrate interface and could lead to interfacial crack formation and delamination. This leads to minimum requirements for coating thicknesses that depend on the applied loading conditions.54 For example, in industrially produced X-ray anodes, W–Re coatings are typically used with a thickness of 200–700 mm205,206 to provide better mechanical and thermal-shock properties compared to pure W.204 However, the first experience on the influence of ELMs on coatings under real plasma operational con- ditions will be gained in the ITER-like wall project in JET, which involves testing relatively thin PVD- tungsten coatings (14–20mm) on a CFC substrate that provides a strong and anisotropic CTE difference.19,142 The behavior of this material under the above outlined transient heat loads is of course a key factor for the lifetime assessment of PFCs. However, the 566 Tungsten as a Plasma-Facing Material results obtained for pure thermal shock testing might underestimate the material damage and by this over- estimate its lifetime. Only a combination of thermal shock, thermal fatigue (see Section 4.17.4.2), neu- tron irradiation (see Section 4.17.4.3), and plasma wall interaction (see Section 4.17.4.4) will be able to give appropriate answers for the selection of suitable grades of W. 4.17.4.2 Thermal Fatigue Resistance The thermal fatigue resistance of tungsten is strongly related to its performance as part of current inertially, and future actively cooled components for applica- tion in magnetic fusion devices. The functional requirements these components have to fulfill are listed in Section 4.17.2. State of the art inertially cooled components include W coatings on graphite, CFC, and TZM.12,13,46,47,54,140,207,208 These concepts are used or are going to be used in the large operating toka- maks, for example, AUG and JET. During thermal loading, they mainly suffer from the problem of high interfacial stresses as a result of the CTE difference between the W coating and the substrate. Further- more, interfacial reaction products and their poten- tial reduced power handling capability have to be taken into account. Besides coatings, recent development of an iner- tially cooled bulk tungsten divertor for JET20,123 showed that under thermal fatigue loads theW quality is of minor importance for the integrity of the compo- nent. The major issue for the tungsten PFM was found to be the necessary shadowing of the plasma-loaded surface to avoid overheating and melting at tile edges as a result of the shallow angle of the incident plasma. This was realized by surface shaping.209 In the design of actively cooled components, tungsten is joined to a water-cooled Cu-based heat sink (ITER) or He-cooled steel or W-based heat sinks (e.g., DEMO, ARIES-CS). Direct cooling of the tungsten armor should be avoided, particularly without castellation, as the induced stresses might cause catastrophic material failure with subsequent water or He-leakage.124 Therefore, the only perfor- mance requirements are a sufficiently good surface quality to reduce possible crack initiation points and therefore suitable fabrication and surface finish- ing technologies,103,210,211 the chemical compatibility with the heat sink and, if present, the joining interface material, and the cyclic stability of the joint(s). The latter is influenced by the temperature gradient applied during steady state heat loads, the difference of the CTEs, the quality of the joining process and, perhaps most important for reducing induced stres- ses, the tile size, or the dimensions of the castellated segments (see Section 4.17.3.2.4). Smaller tile sizes significantly improve the stress situation at the interface, and also at the top surface of the PFM. This has to be taken into account when comparing the thermal fatigue results of various kinds of components and the response of different grades and alloys of W, as shown by Makhankov et al.,90 where smaller tile sizes resulted in little or no crack formation. Furthermore, variations in the size of the component investigated can often explain the contradictory results presented in the literature that show good behavior of a material in one test while it fails in another. However, there are limitations to the minimum size of tiles and a compromise between operational and economical needs has to be made. Despite the fact that design and manufacturing technique seems to be more important than the mechanical properties and the microstructure of the particular W grade or alloy, the latter should still be considered. Similar to thermal shock results, the risk of delamination parallel to the loaded surface at the interface90 or anywhere in the bulk material has to be minimized. Therefore, the grain orientation of the PFM microstructure should be perpendicular to the loaded surface, although this still bears the risk of crack formation toward the cooling structure.108 To avoid subsequent water or He-leakage in case of crack propagation into the heat sink material, particularly in the He-cooled divertor design (see Section 4.17.4.2.2), suitable material and design solutions still have to be found. Furthermore, sha- dowing of adjacent tiles similar to the JET bulk W divertor has not yet been included in the design of the actively cooled components. 4.17.4.2.1 ITER The performance of bulk W for ITER has been investigated using water-cooled divertor designs, that is, flat tile, macrobrush, and, most relevant, monoblock options. One important factor in these design solutions is the maximum allowable distance between the front surface and coolant to accommodate the heat without melting96 and, if possible, to avoid recrystallization during normal operating conditions. The ability to estimate this parameter requires not only the thermal conductivity of the materials but also the amount of allowable damage at the interface. This requires knowing not only the damage produced during Tungsten as a Plasma-Facing Material 567 operation but also understanding the manufacturing accuracy and reproducibility because tens of thousands of armor/heat sink joints will be produced. Studies on this issue have shown that the current W monoblock design with a defect extension up to 50� appears to be suitable for the upper part of the vertical target (P¼ 10MWm�2), but is not well adapted to a heat flux of 20MWm�2, which is neces- sary for application at the strike point of the vertical target, as systematic defect propagation was observed. A tungsten flat tile design with 6-mm long defects in the material interface was studied and proved to be compatible with fluxes of 5MWm�2 but was unable to sustain cyclic fluxes of 10MWm�2.212 These results confirm that the monoblock geome- try generally proves to have superior behavior under high heat flux testing when compared with flat tile geometry. However, it is worthwhile to continue the investigation of the flat W tile design for low-flux regions despite the hazard of cascade failure of the flat tiles106 for two reasons: cost and weight. Besides this characterization, a number of high heat flux tests have been carried out on mock-ups and prototypes without artificial defects representing the different design options to assess the ‘fitness for pur- pose’ of the developed technologies.33,90,161,213–218 The results obtained for small test mock-ups of the flat-tile and monoblock design can be transferred to large-scale prototypes for the divertor vertical target. Independent of the type of pureWorW–La2O3 armor material used in these prototypes, theWparts survived in the nonneutron-irradiated condition up to 1000 cycles at 20MWm�2 in the monoblock design213,219 and up to 1000 cycles at 18MWm�2 in the flat tile Figure 7 Thermal fatigue testing results of W macrobrush and W design (see Figure 7).213 This is far beyond the design requirements for use in the upper part of the vertical target (P¼ 5MWm�2) and, in case of the monoblock design, even meets the design requirements for the strike point area of the vertical target. Alternative concepts such as explosive bonding of tungsten to a heat sink material,220 PS on a Cu-alloy216 or on EUROFER steel44 could probably be of use in the divertor but even more for first wall applications for fusion machines beyond ITER. How- ever, these concepts often suffer from high interfacial stresses as a result of the CTE difference between the W coating and the substrate. 4.17.4.2.2 Prototype and commercial reactors There are many design proposals for a He-cooled first wall and divertor concept for DEMO and ARIES-CS.37,160,221 Among these, the He-cooled modular design with jet cooling (HEMJ)102 is the most developed and qualified in terms of microstruc- tural response,103 having survived at reduced coolant temperatures of 450–550 �C at least 100 cycles at 11MWm�2 without failure. In contrast to the results obtained for water-cooled components for ITER, no influence of grain orientation on the components performance was observed.102 This might be a result of the higher temperature, which was always above the DBTT. Nevertheless, some difficulties in the design still have to be resolved. First, there are problems related to temperature with a desired coolant temperature of �600 �C; these include material recrystallization at the top surface and the necessary high temperature joining to the W-based heat sink material. Second, W macrobrush 0 dpa: 1000 cycles at 18 MW m−2 0 dpa: 1000 cycles at 20 MW m−2 0.6 dpa: 1000 cycles at 10 MW m−2 (increasing of Tsurf) 0.6 dpa: 1000 cycles at 18 MW m−2 (no degradation) W monoblock monoblock mock-ups before and after neutron irradiation. 568 Tungsten as a Plasma-Facing Material issues related to the material’s mechanical properties must be solved, in particular for the ductility of W-based structural material and its neutron irradia- tion resistance (see Section 4.17.4.3). Finally, the manufacturing reproducibility has to be at a high level because of the large number of small units (1 unit 3� 10�4m2) necessary for cladding the DEMO divertor. Temperature (�C) Th er m al d iff us iv ity (m m 2 s– 1 ) 200 400 600 800 1000 1200 1400 16000 70 65 60 55 50 45 40 35 30 0.6 dpa at 200 �C Nonirradiated Figure 8 Thermal diffusivity of W–1% La2O3 in nonirradiated and irradiated condition. 4.17.4.3 Neutron Irradiation The irradiation of tungsten and tungsten alloys with energetic neutrons (14MeV) resulting from the D–T reaction causes radiological hazards that were already discussed in Section 4.17.3.2.6. In addition, the neu- tron irradiation affects the material composition by transmutation of tungsten to Re and subsequently osmium (transmutation of W isotopes to Ta and Hf are negligible222). The amount of transmutation strongly depends on the applied neutron wall load and neutron spectrum223 and for theW to Re transmu- tation reaction reaches values between 0.3 and 5 at.% per dpa.222 The subsequent transmutation of Re to Os is expected to occur faster than the production of Re from W resulting in a steadily proceeding burnup of Re. The neutron fluence on the first wall varies strongly with location. For the full lifetime of ITER a maximum of �0.3MWam�2 is achieved224 ( 1.35 dpa in tungsten225). As the divertor PFCs will be exchanged 3 times and only the last three will operate in a D–T environment, a neutron fluence of �0.1MWm�2 is expected during the lifetime of each PFC. For DEMO, an average neutron wall load of 2MWm�2 is assumed for the main chamber, which would result in �45 dpa after 5 full power operation years. These conditions yield a transmutation of 100% W into 75% W, 12% Re, and 13% Os.36 For geometrical reasons, that is, larger surface to angular extension ratio, it will be roughly a factor 2 less in the divertor region. Furthermore, neutron irradiation damages the material properties by the formation of vacancies and interstitials (see Chapter 1.03, Radiation-Induced Effects onMicrostructure). Their behavior including analysis of displacement cross-sections,226,227 diffu- sion, mutual recombination, and clustering are being assessed by atomistic modeling.228–231 Both transmutation and defect generation influ- ences the material properties and subsequently the material response to steady state and transient thermal loads. 4.17.4.3.1 Thermophysical properties and swelling The influence of neutron irradiation on the thermo- physical properties is related to the irradiation tem- perature and the number of defects generated in the crystal structure. At temperatures 1200 �C (see Figure 8). In addition to defect generation, material degrada- tion is also related to the formation of transmutation products such as Re and Os, which in general exhibit poorer thermophysical properties. Transmutation- induced degradation increases with increasing tem- perature and irradiation dose, which makes it the most relevant process for the degradation of material properties for future fusion reactors such as DEMO. Despite the potential for full recovery of the mate- rial defects mentioned above, void-induced swelling occurs. The results235,236 of tungsten and tungsten alloys show that the material’s volume increases with increasing irradiation temperature (�1050 �C).237 W–Re alloys exhibit significantly improved swelling behavior compared to pure W, with a local maxi- mum at �750 �C. However, the swelling only amounts to �1.7% at 9.5 dpa.237 Therefore, a negligi- ble effect of swelling can be expected for the opera- tion of ITER. Experimental values do not exist at temperatures >1050 �C as expected for the operation of DEMO. Tungsten as a Plasma-Facing Material 569 4.17.4.3.2 Mechanical properties Data in the literature on mechanical properties of neutron-irradiated tungsten are very limited.234,238,239 However, in combination with experimental results obtained for other refractory metals, it has been shown that in metals with a bcc lattice structure, irradiation hardening causes a steep increase in yield stress and a decrease in ductility.110 Conse- quently, the material fails by brittle cleavage frac- ture as soon as the yield stress exceeds the cleavage strength. Therefore, the increase of the DBTT depends on the neutron fluence, the neutron spec- trum (will be addressed by the International Fusion Materials Irradiation Facility, IFMIF), and the irradiation temperature. The radiation hardening in bcc alloys at low temperatures (1000 �C would be preferred as full or at least partial recovery of defect-induced material degradation is achieved by annealing at 1200 �C.234 This implies that the near- surface part of a W component will retain its ductility, which has a beneficial effect on the crack resistance at the plasma loaded surface. However, such temperatures are not feasible at the interface to the heat sink where tungsten will be in contact with copper (ITER) or steel (DEMO), which are limited to significantly lower operational tempera- tures. Hence, better understanding of the irradiation effects on tungsten at temperatures between 700 and 1000 �C is needed, particularly related to reactor application in DEMO.109,110,240 In addition to the influencing factors on the DBTT mentioned above, that is, neutron fluence, neutron spectrum, and irradiation temperature, the material’s composition also plays an important role. While the addition of Re has a beneficial effect on the material’s ductility in the nonirradiated state, under neutron irradiation it results in more rapid and severe embrit- tlement than it is observed for pure W.239 Similarly, less mechanical strength and an increased loss of ductility compared to pure W is found for particle- strengthened W alloys (e.g., W–1% La2O3) when tested up to 700 �C. The only exception among all explored tungsten alloys might be mechanically alloyed W–TiC (see Section 4.17.3.3) that showed no irradiation hardening as measured by Vickers hard- ness at 600 �C.87 Finally, the mechanical properties are influenced by neutron-induced He-generation and the transmu- tation of tungsten. While He generation inW is, com- pared to CFC and Be, very small (�0.7 appm He per dpa) and its influence on the mechanical properties of W negligible,73,83,224 the transmutation of W into Re and subsequently Os significantly alters the mate- rial structure and its properties. The generation of significant amounts of ternary a and subsequently s-phases results in extreme material embrittlement and will cause shrinkage. In combination with ther- mally induced strains, this might produce high tensile stresses causing the extremely brittle material to extensively crack and perhaps even crumble to powder.36 4.17.4.3.3 Thermal shock on irradiated W The simulation of disruptions ( (a) (b) 500 mm 500 mm Figure 9 Thermal shock response of W–0.2% Re (HF¼ 41MWm�2 s1/2, P¼ 1.31GWm�2, t¼1ms, n¼ 10); (a) before and (b) after neutron irradiation (0.6 dpa at 200 �C). 570 Tungsten as a Plasma-Facing Material 4.17.4.3.4 Thermal fatigue on irradiated W components Information on the thermal fatigue resistance of W components is limited to the experience obtained in two irradiation campaigns for ITERwhich reached neutron doses of 0.15 and 0.6 dpa at 200 �C. Refer- ence and irradiated actively cooled mock-ups with W–1% La2O3 as the PFM were exposed to 1000 cyclic steady state heat loads at power densities up to 18MWm�2.213,216,217 The results obtained indicate that at these neutron fluences the material changes occurring in tungsten do not have any significant influence on the component’s performance. However, mock-ups based on the macro- brush design experience a degradation of themaximum achievable power density from 18 to 10MWm�2, which is related to neutron embrittlement and subsequent cracking failure of the Cu-heat sink mate- rial. In contrast, monoblock mock-ups show identical high-level performance before and after irradiation, which makes it the favored design for ITER. Despite these positive results, based on the irradiation-induced mechanical property changes outlined above, the use of tungsten in any highly stressed component at low temperatures Tungsten as a Plasma-Facing Material 571 varies as a function of the implantation fluence.247 The higher the ion energy, the lower the fluence and the temperature required to create material damage beyond vacancies and vacancy clusters. For example, very small blisters were observed for 8 keV ions at RT and a fluence of 4� 1021 Heþ�m�2 and above.247 In contrast, for low energies (1300 K and fluences of about 1026 Heþ m�2 and more are necessary to form bubbles and surface holes.149,248 The reason for the lack of blistering at low temperature and low ion energy is assumed to be the trapping of He-ions at defects/vacancies in the very near-surface range. With increasing temperature, the defects and He-atoms debond and the He-atoms diffuse toward the bulk, agglomerate, and result in blistering.249 Similarly, with increasing energy, the penetration depth of the He-atoms increases from nm for eV-ions up to 1.7 mm for 1.3MeV He-ions and the probability of blister formation correspondingly rises. In both cases, whether the process is driven by He-diffusion or high penetration depth, blister- ing and exfoliation are expected to occur when the amount of He locally reaches 4 at.% and 20–40 at.%, respectively.250 4.17.4.4.1.2 Influence of temperature Vacancy mobility is dependent on temperature and starts at 523–573 K.251,252 As the mobility of vacancies and the formation of thermal vacancies are driving forces for the formation of bubbles, holes, and blis- ters, an increase in temperature increases the size and decreases the density of material damage.149,253 How- ever, it is not only the temperature during ion irradi- ation but also the annealing temperature during experiments such as thermal desorption measure- ments, which can influence the damage characteristics.247 The formation of holes and porous structures observed after thermal treat- ments,254 particularly for temperatures above the material-dependent recrystallization temperature, is related to the movement of vacancies, accelerating the expansion and coalescence of He bubbles, their migration to the surface,253 and subsequently the release of He. The latter is also a function of temper- ature, showing several release peaks between 400 and 1600K related to different trapping sites247 and determines the amount of He retained as a func- tion of the incident fluence.242,247,255 However, helium retention may be mitigated by cyclic He-implantation and high temperature heating, for example, flash heating to 2000 �C, because He flows away before critical amounts accumulate and form complex He-vacancy clusters with higher binding energy.250 4.17.4.4.1.3 Influence of material’s microstructure The impact of the material’s microstructure is related to the amount of intrinsic defects at which He can be trapped and therefore determines the amount of He-retention.250 SC tungsten contains fewer defects than powder metallurgically (PM) produced tungsten (grain boundaries) and plasma-sprayed tungsten (grain boundaries and porosity), which is directly related to the thickness of porous/spongy structures (porosity about 90%,256 see Figure 10) that form depending on energy and temperature.257,258 However, investiga- tions at 1650K have shown that at such high tempera- tures there exists no difference between SC-W and PM-W, even for ion energies as low as 25 eV. Themain trap sites at such high temperatures are thermal vacan- cies while intrinsic defects play a minor role.149,251 With regard to the material’s lifetime under He-exposure, the migration of He-bubbles toward the surface and the formation of pores and porous/ spongy structures seem to prevent the rupture and exfoliation that can accompany blistering. This is important as the exfoliation of blisters creates dust,199 which limits the plasma performance. For an antici- pated flux of 2� 1018 Heþ m�2 s�1 at 850 �C in iner- tial confinement devices, this may lead to a removal of 20mmyear�1 from the wall.109 Therefore, one way to increase the material’s lifetime might be to operate it at higher temperature.253 Another approach would be to develop advanced microengineered materials that have typical feature sizes less than the classical helium migration distance (20 nm).109 However, bubbles, holes, and porous/spongy structures signifi- cantly influence the material’s performance by reducing its thermal conductivity in the near-surface layer. This might play an important role when deter- mining the erosion and melt formation under com- bined He-irradiation and transient thermal loads, which will be shortly addressed in Section 4.17.4.4.3. 4.17.4.4.2 Hydrogen-irradiation and retention Besides He, hydrogen isotopes, particularly the fuel elements deuterium and tritium, are the main inci- dent ion species contacting PFMs and PFCs. The energy of these particles corresponds with the plasma temperature at the edge, which is in the range of some eV, but also includes highly energetic particles (�10 keV) escaping from the inner core of the SR W (a) (b) (c) (f)(e) (h) (i) (d) (g) SC W (100) ASTM B760 (ITER) W-TiC(1.5% wt)W- La2O3(1% wt) W-Re(5% wt) W-Re(10% wt) VPS-W (EAST) RC-W Figure 10 Cross-sectional scanning electron microscopic images for nine different grades of W relevant to fusion engineering practice. All target specimens were exposed to consistent pure He plasmas at 1120K for 1 h. The Heþ impact energy was �40eV; (a) PLANSEE stress-relieved W, (b) single crystal h100i W, (c) ITER ASTM B760 compliant W, (d) PLANSEE W–5% Re, (e) PLANSEE W–1% La2O3, (f) ultrafine-grained W–1.5% TiC, (g) ULTRAMET CVD W–10% Re, (h) VPSW (EAST), and (i) recrystallizedW. Reproduced fromBaldwin, M. J.; Doerner, R. P. J. Nucl. Mater. 2010, 404, 165–173, with permission from Elsevier. 572 Tungsten as a Plasma-Facing Material plasma. The impact of the energetic hydrogen ions is influenced by the incident ion energy, the ion flu- ence, the temperature, and the material’s composition and microstructure. The resulting damage, that is, vacancy formation, vacancy clustering, bubble formation, and blistering, determine not only the amount of material degradation and erosion but also the hydrogen (tritium) retention in the material. For active control of the hydrogen retention, short-term thermal treatments of the surface are being investi- gated. However, the short thermal load required to effectively remove the deuterium and tritium may also destroy the thin material layer (nm to low-mm range) that is responsible for the majority of the retention.259 4.17.4.4.2.1 Influence of ion energy, fluence, and temperature Independent of the ion energy, blistering (see Figures 11 and 12) due to H occurs only during irradiation at temperatures below 900–950 K and as a function of the ion fluence199,260 at 500 K.261–263 This temperature dependence of blistering is attrib- uted to the formation, movement, and agglomeration of vacancies containing trapped hydrogen,264 which is dominant at temperatures 500K.263 The fluence threshold for blister formation increases with decreasing ion energy and signifi- cantly increases to values >1025 Dþ m�2 at ion ener- gies (a) (b) (c) Large blisters Large blisters Small blisters Crack/void along grain boundary Lids of small blisters partially or fully removed by FIB fabrication Figure 11 Scanning electron micrographs of tungsten exposed to a hydrogen fluence of 1026Dm�2 at 480K (45� tilt). (a) Overall image; (b) cross-sectional image of a large blister; (c) internal image of small blisters. Reproduced from Shu, W. M.; Kawasuso, A.; Yamanishi, T. J. Nucl. Mater. 2009, 386–388, 356–359, with permission from Elsevier. (a) (b) (c) Figure 12 Scanning electron micrographs of small blisters appearing at tungsten exposed to a hydrogen fluence of 1026Dm�2 at 480K (45� tilt). (a) Initial stage; (b) growing; (c) bursting. Reproduced from Shu, W. M.; Kawasuso, A.; Yamanishi, T. J. Nucl. Mater. 2009, 386–388, 356–359, with permission from Elsevier. Tungsten as a Plasma-Facing Material 573 In contrast, from the point of view of H retention, blistering is favorable because tritium accumulates in blisters, which act as a diffusion barrier for hydrogen, and which can otherwise penetrate deep into the material even at RT269–274 until it finally ends up in the heat sink structure and the coolant. Accordingly, as tritium retention is, in general, strongly correlated with the generation of blisters,275 it shows a maxi- mum at an irradiation temperature of about 500 K.45,261–263,266,268,276 However, the retention of tritium and deuterium is also dependent on the trapping sites existing in the material. These are, in ascending order of their trapping potential, resid- ual impurities, from which slow desorption occurs even at RT,277 grain boundaries and dislocations, 574 Tungsten as a Plasma-Facing Material radiation-induced vacancies and vacancy clusters, and pores. Depending on the occurrence and domi- nance of particular sites, the temperature at which the maximum hydrogen retention is observed varies between 450 and 850K.138,262,264,268,277–281 With increasing temperature (>1000K), such as that occurring at the strike point of the divertor, and with the lack of blisters, continuously decreasing hydrogen retention is observed.2,255,278,282,283 The remaining amount of retained hydrogen might be attributed to the presence of hydrogen as a solute, which depends only slightly on the incident ion energy, but scales with the implantation fluence and which is assumed to be of the same order of mag- nitude as the trapped concentration. It decreases only slightly with increasing temperature and at 1600K still amounts to about 10% of the initial hydrogen content retained at 300 K.274 In addition to blisters, the high amount of hydrogen out-gassing at temperatures of 873K and above results in the formation of bubbles and pores.253,277,284,285 This effect depends on the ion energy and fluence, which determines the amount and penetration depth of trapped hydrogen. Even though a beneficial smoothing effect on the surface quality is observed in comparison to pure annealing without hydrogen impact, at high temperatures up to 2500 K the surface smoothening might be accompanied by detrimental crack formation.284 4.17.4.4.2.2 Influence of material’s composition and microstructure As mentioned above, hydrogen retention depends on the trapping sites available in the material and their relative energies. Their existence and concentrations are influenced not only by the impinging H-ions, but also by the manufacturing process, thermal pre- treatments, the material’s composition and micro- structure, and the surface quality. Accordingly, the retention increases with the amount of porosity in the material, as it allows a deep penetration of hydrogen and the voids and pores provide the highest trapping energies264,275 with thermal desorption occurring at temperatures >700K.45,262 Another material parameter that increases hydro- gen retention is the number of dislocations,266,286 particularly those introduced during deformation processes used for material densification. However, the recrystallization of the material removes not only dislocations but also vacancies and vacancy clus- ters, which have been introduced by the impinging H-ions286,287 and as grain boundaries. This effectively reduces the trapping sites for hydrogen retention and, consequently, the lowest retention is observed for high-purity SC materials, particularly due to the low diffusion rate compared to polycrystalline tungsten.275,288,289 This low diffusion rate results in a near-surface accumulation of hydrogen, which acts as a diffusion barrier and leads to a saturation of hydrogen retention with increasing fluence.290 Such saturation is not observed for pure polycrystalline tungsten due to the possible hydrogen migration along grain boundaries.291 Finally, the hydrogen retention is influenced by impurities277 and dopants. The addition of La2O3 and TiC particles as well as the formation of pores, for example, by potassium doping, not only introduce traps and increase hydrogen retention,293 but also decrease the diffusion rate.291 In contrast, alloying with up to 10% Re has no measurable effect on the H retention properties of the material,279 as it only creates a slightly deformed crystal lattice structure but no additional hydrogen traps. In addition to hydrogen retention, material dam- age and particularly blistering is influenced by the material’s microstructure. Blistering occurs preferen- tially when the crystal is oriented with the h111i direction perpendicular to the surface292 and the blisters develop in different shapes from low, large, and spherical to high, small, and dome or cone- shaped.45,262,267 The blisters in recrystallized materi- als are mainly plateau-shaped, often multilayer structures, which indicate a step-wise build-up, and in few cases also small blisters on top of large ones are formed.263,293 However, it is significant that the blister size is commonly limited by the grain size45,294 indicating that the grain boundaries play an important role in the formation of blisters. Accordingly, SCs and nanostructured materials such as W–TiC provide the strongest resistance against blistering, although the particular reason is different. For SCs, the hydrogen diffusion and accumulation is limited and there is a fast desorption from low- energy traps at elevated temperatures. In contrast, for nanostructured materials the size of individual grains is extremely small and so is the volume for blister formation. Furthermore, the migration of hydrogen is significantly increased by the large num- ber of grain boundaries.292 Further material parameters that reduce blister formation are open porosity and the surface finish, particularly the number of random or artificially introduced scratches that might act similar to grain boundaries.45 In contrast, the introduction of Tungsten as a Plasma-Facing Material 575 impurities and dopants in commercially available grades of tungsten increases the number of blisters and exfoliation in both their stress relieved and recrystallized states.277,293 4.17.4.4.3 Combined loading conditions As described above, the damage mechanisms of hydrogen and He-irradiation are rather similar, although they occur in different temperature ranges. Accordingly, their mutual interaction is also strongly influenced by the implantation tempera- ture. Therefore, the testing sequence plays a role in the behavior, as for He preirradiation followed by hydrogen implantation, the implantation tempera- ture of He determines the amount and kind of produced damage and the He-retention, which sub- sequently influences the hydrogen uptake occur- ring as described in Section 4.17.4.4.2. For example, He-implantation at RT either does not change the retention or may increase it due to the formation of additional trapping sites295–297 and the lower diffusion rate of He compared to H. With increasing He implantation temperature up to 800 K, hydrogen retention significantly decreases compared to pure hydrogen irradiation.261,292,296 This may be attributed to the occupation of trap sites by He as a result of its increasing mobility.298 Potential trap sites are the numerous He-induced nanosized bubbles acting as a diffusion barrier.292 A further increase in temperature to 1600 K does create significant material damage by He due to pore and bubble formation or even blistering. This tremendously increases the number of trap sites in the material and leads to He desorption during implantation and accordingly increases the hydro- gen retention.299 For simultaneous loading of He and hydrogen, the fraction of He should reach at least 5 at.% to observe significant changes in the material’s response.261,292 Furthermore, for implantation tem- peratures below 900 K, results similar to those described above are observed for sequential ion beam loading.299 However, due to desorption of hydrogen at high temperatures >1000 K, no hydro- gen retention takes place and the damage mechan- isms are dominated by the He-irradiation during such temperature excursions. Correlated with hydrogen retention, blister for- mation at temperatures 576 Tungsten as a Plasma-Facing Material 4.17.5 Conclusion Along with other favorable properties, tungsten is characterized by the highest melting temperature among all metals, a low energy threshold for sputter- ing, and a low tritium inventory compared to carbon- based materials. These characteristics make tungsten the most promising material for the plasma-facing inner wall of future nuclear fusion devices based on the magnetic confinement principle, and it is also under consideration for inertial fusion applications. Accordingly, it has been selected as the PFM for a large part of the ITER divertor during its start-up phase and will be used for the full divertor as soon as tritium operation starts; in addition, it is the reference material for DEMO. However, tungsten also offers less favorable prop- erties. Related to these, there are some material issues that have to be resolved before operating tungsten in a fusion environment in an economically reasonable way, which means in DEMO and beyond. These are � recrystallization, which influences the mechanical properties by reducing the ductility and increasing the DBTT � embrittlement as a result of neutron-induced damages and transmutation � resistance to crack formation, depending on the mechanical properties, which is particularly important during transient thermal loads � He-induced sputtering and modification of a thin surface layer, which is influenced by existing mate- rial damage as well as by temperature and temper- ature gradients, for example, those occurring during transient thermal loads � melting, which is related to crack formation and the degradation of thermophysical properties as a result of He-irradiation-induced surface modifica- tion; melt splashing and droplet ejection will influ- ence the stable operation of the fusion plasma. As all grades of tungsten investigated so far have their own individual drawbacks, R&D programs world- wide are aiming for a deeper understanding of the parameters that influence the degradation of tung- sten, and the development of new tungsten grades that are capable of dealing with the above-mentioned requirements. Therefore, the materials are character- ized and qualified with regard to their microstructure before and after recrystallization by � mechanical tests : evaluation of the material’s strength and DBTT � thermal shock loading : determination of tempera- ture- and power density–dependent damage, cracking and melting thresholds, which are related to the mechanical and physical properties � thermal fatigue loading : evaluation of the material’s performance as part of an actively cooled component � neutron irradiation : characterization of the degrada- tion of the material’s strength and the DBTTas well as the thermal shock and thermal fatigue response � He- and H-irradiation : determination of the damage mechanisms such as blister, void, and bubble forma- tion as a function of ion energy, fluence, and temper- ature as well as addressing hydrogen retention issues. 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All rights reserved. 4.18.1 Introduction 584 4.18.1.1 Background 584 4.18.1.2 Plasma-Facing Materials 584 4.18.1.3 Particle–Matter Interactions 586 4.18.2 The Advantages of Carbon as a PFC 586 4.18.2.1 Plasma Impurities and the Need for Graphite Materials 586 4.18.2.2 Thermomechanical Loading of PFMs 587 4.18.2.3 Transient Loading of PFMs 587 4.18.3 Irradiation Effects on Thermophysical Properties of Graphite and CFCs 590 4.18.3.1 Graphite Irradiation Damage 590 4.18.3.2 Surface Effects 591 4.18.3.3 Properties and Property Evolution of Graphite Fiber Composite 591 4.18.3.3.1 Irradiation-induced dimensional changes in CFCs 594 4.18.3.3.2 Irradiation-induced changes in strength and modulus 597 4.18.3.3.3 Thermal conductivity degradation 598 4.18.4 Plasma–Particle Interactions 602 4.18.4.1 Chemical Erosion 602 4.18.4.2 Doping of Graphite to Suppress Erosion 604 4.18.4.3 Physical Sputtering 606 4.18.4.4 Radiation-Enhanced Sublimation 607 4.18.4.5 Erosion of Graphite in Simulated Disruption Events 608 4.18.5 Tritium Retention in Graphitic Materials 608 4.18.6 HHF Component Technology 611 4.18.6.1 Joining of CFC to Heat Sink 611 4.18.6.2 Evaluation of HHF Joint 617 4.18.7 Summary and Conclusions 617 References 618 Abbreviations ASTM ASTM International CFC Carbon(graphite) fiber composite CTE Coefficient of thermal expansion CVD Chemical vapor deposition DPA Displacement per atom EU European Union FoMd Disruption figure of merit FoMth Thermal figure of merit GMP Galvanic Metallization Process HIP Hot isostatic press ITER International Thermonuclear Experimental Reactor JET Joint European Torus LAM Low activation materials PAN Polyacrylonitrile PFC Plasma facing component PFM Plasma facing material(s) PVD Physical vapor deposition RES Radiation enhanced sublimation RT Room temperature SATIR Transient infrared thermography SEM Scanning electron microscopy XRD X-ray Diffraction 583 584 Carbon as a Fusion Plasma-Facing Material 4.18.1 Introduction 4.18.1.1 Background Graphite-moderated, gas-cooled reactors led the way into the nuclear age starting with the Chicago Pile-1 reactor, where the first controlled and sus- tained critical nuclear reaction was initiated in December 1942. The first commercial nuclear power plant, Calder Hall in the United Kingdom, went critical in 1956. As the graphite moderator was literally at the core of these early reactors, graphite became one of the first and most exten- sively studied nuclear materials. As discussed in Chapter 4.10, Radiation Effects in Graphite, the fission-born neutron results in significant thermo- physical property changes in graphite. Moreover, depending on the type of fission reactor, other environmental factors such as graphite oxidation become extremely important. In addition to being the moderator of gas-cooled reactors, graphite has found a number of new nuclear power applications. As examples, pyrolytic graphite is a key functional element in TRi ISOtropic (TRISO) fuels, which con- tinue to be developed and utilized for gas-cooled reactors; carbon fiber composites (CFCs) are now under development for core application in high- temperature gas-cooled reactors1 and have been widely used as plasma-facing components (PFCs) in fusion reactors.2 The latter application began in 1978, when the Princeton Large Torus made a transition from tungsten to graphite ‘limiters.’ This enabled the first thermonuclear temperatures, beginning the widespread application of graphite materials in fusion systems, the subject of this chapter. As will be discussed, the primary motivation for the use of graphite in fusion systems is not (as in fission reac- tors) for neutron moderation, but for reasons related to its exceptional high temperature performance and its relatively innocuous interaction with the plasma. However, the fusion reactor environmental effects on graphite, including irradiation-induced property evolution, are very similar to those of their fission reactor analogs. In contrast to the fission of heavy elements such as uranium or plutonium, which releases a large amount of energy in their fission fragments and a moderate amount of energy in the form of neutron kinetic energy (mean about 1MeV), fusion can occur for a number of light elements, some of which have reac- tions that release very high-kinetic-energy neutrons. Several possible routes to fusion are shown below in eqn [1]: 1H 1 þ1 H1 !1 D2 þ positron ¼ 1:4MeV 1H 1 þ1 D2 !2 He3 ¼ 5:5MeV 1H 1 þ1 T3 !2 He4 ¼ 19:9MeV 1D 2 þ1 D2 !2 He3 þ neutron ¼ 3:3MeV 1D 2 þ1 D2 !1 T3 þ1 H1 ¼ 4:0MeV 1D 2 þ1 T3 !2 He4 þ neutron ¼ 17:6MeV 1D 2 þ2 He3 !2 He4 þ1 H1 ¼ 18:2MeV ½1� For any of these reactions to take place, the ionized atoms must be brought together with sufficient force to overcome the coulombic barrier. In thermonuclear fusion, this is accomplished by heating the ‘plasma’ of these atoms to the point where the kinetic energy is sufficient to overcome that barrier. Currently, it is thought that DþT fusion is the most accessible route to fusion, though the gaseous temperature required for DþT reaction is more than 50million Kelvin. Control and containment of high-temperature, high-density fusion plasmas is the primary challenge and obstacle to fusion power. Many reactor concepts have been studied in the past and attention is now focused on the ‘tokamak’ system. This toroidal con- finement machine system was developed in the mid-1960s in Russia. In this design, a high-strength twisted helix of magnetic lines forms a magnetic bottle. Ions, which are trapped within a certain gyro- radius, travel along these lines circulating around the helix in opposition to the plasma electrons. For non- collisional plasmas, the ions can be heated by mag- netic induction or through various external means to the extreme temperature necessary for the fusion reac- tion to take place. This concept is the basis for the four largest present-day fusion machines (Table 1), and is the premise for the ITER machine currently under construction. To give an idea of scale, in all of the present-day machines listed in Table 1, the helical cavity is big enough in size for an adult to walk within, and the radius from the center of the machine to the middle of the helix is typically several meters. A depiction of the inside of the JET torus, complete with beryllium-coated CFC wall, is given in Figure 1. 4.18.1.2 Plasma-Facing Materials Perfect containment of the high-density plasma needed for power production, where perfection means no interaction of the high-energy plasma and its sur- roundings, is not a practical reality. Whether through normal operation, or in off-normal incidents such as plasma ‘disruptions,’ plasma–material interaction (PMI) will occur in fusion devices. Components Table 1 Materials and heat loads for the major fusion machines worldwide Fusion device Location Fuel system First wall heat load (MWm�2) First wall material Divertor or limiter heat load (MWm�2) Divertor or limiter material ITER Cadarache, France D/D 0.88 normal Beryllium 5–10 Sepcarb NB41, or CX 2002U CFCsD/T 1.75 off-normal 10 (ELM) DIII-D San Diego, USA D/D 0.6 Poco ATJ graphite 5.3 Poco ATJ graphite JT-60U Naka, Japan Dunlap DMS704 CFC Hitachi HCB-18S Ibiden EPT-10 Showa-Denko CC312 JET Culham, England D/D Negligible Dunlop DMS-704 18 Dunlop DMS-704 Sepcarb N11 Sepcarb N11-S (3D CFC) TFTR Princeton, USA D/D Poco AXF-5Q graphite FMI 4D D/T Figure 1 Inside the JET torus. Beryllium-coated carbon fiber composite. Carbon as a Fusion Plasma-Facing Material 585 in line of sight with the plasma, and therefore impacted by the hot gasses and particles, are referred to as plasma-facing components (PFCs) or materials (PFMs). The reactions between the fusion plasma and the PFMs are quite severe and typically cause melting or sublimation, component mechanical fail- ure due to high thermal stress, and excessive surface erosion. The plasma ion flux and associated heat loading to the PFMs can be highly nonuniform and quite dependent on the tokamak design. The hot plasma gasses are made up of unburned hydrogen fuel, fusion byproducts such as helium, plasma electrons, and impurities, which include ele- ments previously removed from PFCs. As can be seen in eqn [1], the types of particles that may strike the PFMs are dependent on the fusion fuel. For the DþT fuel system, the plasma will contain not only the DþT fuel, but also high-energy alpha particles (3.5MeV He) and neutrons (14.1MeV). The parti- tioning of the reaction energy between helium and the neutron is both an advantage and a disadvantage for the DþT fuel system. Because the energetic helium nucleus quickly collides with the surrounding gasses, most of its energy remains in the plasma and helps to sustain the high plasma temperature. Con- versely, the neutron has very little chance of collision in the low-density plasma and loses its energy outside of the plasma, usually over meters of path length inside the structure of the reactor. Because less than 30% of the DþT reaction energy remains in the plasma, only this fraction is eventually dumped on the PFCs, thus reducing the heat load handling requirement and material erosion. However, as discussed in Section 4.18.3, the material damage associated with the 14.1MeV neutron collisions is significant and perhaps offsets the advantages of reduced DþT heat loading. A characteristic classifying the fusion device type is the manner in which the plasma edge is defined and the plasma power handled. The classic approach is to define the plasma edge by placing a sacrificial com- ponent in contact with the plasma. This component, which intercepts the plasma edge particle flux, is known as a bumper or bumper limiter, and extends circumferentially around the torus. A second approach to defining the plasma edge is magnetically capturing and diverting the edge plasma onto a divertor plate well removed from the central plasma. Once the plasma gasses strike the surfaces and are thus cooled, they are pumped away. Unless mitigated, the energy 586 Carbon as a Fusion Plasma-Facing Material deposited locally on the ‘divertor’ can be excessive. Many techniques, such as magnetic sweeping to spread the load and puffing of gas to ‘soften’ the ion impact, have been used to reduce the particle flux and energy. Regardless of whether the limiter or divertor design is employed, the majority of the par- ticle and heat flux is intercepted by these components (Table 1). However, a significant flux also impacts the balance of the torus lining, generally referred to as the first wall. A convenient comparison for the heat load- ings given in Table 1 is that the maximum output from a conventional propane torch is approximately 10MWm�2, or about the maximum seen in current fusion devices. 4.18.1.3 Particle–Matter Interactions Of the flux of particles that will impact the PFMs, the highest particle flux will be the ionized fuel itself. The energy of the impacting fuel ions on the various plasma-facing areas depends on many variables. For the divertor, where most of the interactions will occur, the majority of particles will have energies in the eV range. On the first wall where the interaction intensity is less, the charge exchange particles will mostly be in the keV range. For larger fusion devices of the future where dense plasmas and higher mag- netic fields result in the thermalization of the ener- getic helium (eqn [1]), those helium ions will have energies similar to the fuel ions. Electrons, which are in number density equilibrium with the plasma ions, also travel along the plasma field lines, albeit in the opposite direction. The high-energy neutrons pres- ent in the DþT reaction (14.1MeV), or those for the DþD reaction (2.4MeV) have mean free paths of several centimeters in graphite and so will not interact strongly with the first wall. However, these neutrons will be scattered and slowed down within and behind the first wall, resulting in a nearly isotro- pic flux of high-energy neutrons throughout the fusion device. The reaction of the plasma neutrons, ions, and electrons with graphite PFMs, which is discussed in some detail in the following sections, can have a wide range of effects. These effects include physical and chemical erosion of the first wall and thermomechanical property degradation of the bulk and surface material. The discussion thus far has been limited to the operation of tokamaks in the quasi-steady state (long pulse). All present-day large tokamaks are pulsed machines with pulse lengths of seconds, where the plasma discharge consists of a rapid heating phase, a steady state, and a cool down phase. In this case, the heat flux is approximately uniform around the cir- cumference of the machine and scales with the machine power. However, a significant number of these plasma shots end in an abrupt and somewhat violent fashion referred to as disruption. When this occurs, the plasma rapidly becomes unstable and instantaneously ‘dumps’ its energy onto the PFC. This causes significantly larger heat loads than dur- ing normal operation, and in many cases, defines the design limits for these components. 4.18.2 The Advantages of Carbon as a PFC 4.18.2.1 Plasma Impurities and the Need for Graphite Materials The fusion plasma is maintained through a combina- tion of internal heating, (i.e., the 3.5MeV helium nucleus from the DþT reaction) and externally, by means of induction, radio frequency waves, or neutral particle injection. Plasma heating is balanced by plasma-cooling mechanisms among which electro- magnetic radiation dominates. In fully ionized plasma, the radiative cooling comes from the Bremsstrahlung that occurs when the energetic ions interact with the plasma electrons. A fraction of the electromagnetic radiation released from this interaction is lost from the plasma. The energy lost in this manner is signifi- cantly increased by low concentrations of impurities. The plasma power loss in the Bremsstrahlung channel, Pbrem, is determined through: PbremðMWm�3Þ� 4:8�10�43Z2i NiNeT 1=2e /Z2i Ni ½2� where Zi, Ni, Ne, and T are the atomic number of the radiating species, their density, the electron density, and the plasma temperature, respectively. Clearly, from the linear dependence on the plasma impurity concentration, and the square dependence on the atomic mass of the impurity, the ideal PFMs com- prise light elements that have a low tendency to erode and migrate into the plasma. Carbon and beryl- lium are two low atomic number elements commonly used in tokamaks. The next suitable element is alu- minum, which would have almost a factor of five higher radiative loss on an atom-per-atom basis com- pared to carbon. On the same basis, molybdenum, which has been used in many tokamak experiments, has a radiative loss 49 times that of carbon, and tung- sten 150 times the radiative loss of carbon. However, Carbon as a Fusion Plasma-Facing Material 587 this is based on the assumption that the same number of impurity atoms find their way into the plasma (i.e., Ni), which, as discussed later, is not the case. 4.18.2.2 Thermomechanical Loading of PFMs As seen in Table 1, the first wall must handle high plasma surface heat fluxes under normal operation and volumetric heat loadings due to the penetrating neutron and electromagnetic radiation. Surface heat loading is dependent on line-of-sight distance from the plasma and can be as high as several MWm�2. These surface and volumetric heat loadings will induce temperature gradients on the PFMs and corresponding thermal stress, and stresses at the interface between the PFM and the heat sink. For example, if one assumes the ideal case of a 2.5 cm thick, infinitely wide graphite plate that is perfectly bonded to a 50 �C copper heat sink, the thermal stress at the graphite–copper interface for a heat flux of 5MWm�2 has been shown to be 200MPa.3 The ability of the PFC to withstand this heat flux and thermal stress will depend both on the material prop- erties and the component design. The two most obvi- ous design parameters are the thickness of the PFM and how it is attached to the heat sink. The critical material property of thermal conductivity, which to a great extent can be engineered to optimize con- duction to the heat sink, is a strong function of temperature. As discussed later in Section 4.18.3, this property and other performance properties such as elastic modulus and strength are also highly depen- dent on radiation-induced displacement damage. A typical design for a fusion reactor divertor is shown in Figure 2. In this design, the heat flux strikes the surface of CFC composite blocks and the heat flows into a water-cooled copper tube that has been brazed inside the block. The PFC is bolted to a stain- less steel support structure. This configuration of PFC is called the monoblock structure, as compared to the flat plate and saddle types inset into Figure 2. To provide a quantitative comparison of candidate PFMs, a number of figures of merit (FoMs) have been derived, one of which may be written as follows: FoMth ¼ KsyaEð1� vÞ ½3� where K is the thermal conductivity, sy the yield strength, a the thermal expansion coefficient, E the Young’s modulus, and n the Poisson’s ratio. High values of FoMth provide guidance to superior performing candidate materials. Figure 3 shows a comparison of the three primary candidate PFMs: graphite, beryllium, and tungsten. Graphite has been further broken down into fine and coarse-grained (Poco and H451 respectively) graphites, and a high-quality one-dimensional (1D) fiber architecture (MKC-1PH) and a balanced weave 3D fiber architecture (FMI-222) CFC. In Figure 3, it has been assumed that the high thermal conductivity direction for the 1D CFC is oriented at a normal angle to the surface of the PFC. From Figure 3, it is apparent that the graphites and graphite fiber composites, which possess higher strength and thermal conductivity, exhibit thermal FoMs considerably higher than either beryllium or tungsten. Thus, strictly from a thermal stress point of view, high-conductivity and high-strength graphite materials would be considered superior under normal operating conditions for fusion PFCs. 4.18.2.3 Transient Loading of PFMs The disruption, or collapse, of the fusion plasma causes a potentially intense thermal load to the PFC of all large fusion devices. As discussed later in Section 4.18.4, such events will cause very high thermal stresses and significant material erosion. As these events are transient in nature, the ability of the PFC to withstand the disruption depends on the material’s ability to conduct, and its ability to absorb the deposited heat, before reaching a temperature or stress limit. For comparative purposes, a disruption figure of merit takes this into account: FoMd ¼ suðCprK Þ1=2 aE ½4� where su is the ultimate tensile strength, Cp the specific heat, and r the density. Figure 4 reports this disruption figure of merit for the materials in Figure 3. Consistent with the results of the thermal FoMth, high-quality, high-thermal conductivity composites and fine-grained graphites perform better than standard and larger grained gra- phites, and exhibit an order of magnitude better FoMd than beryllium and tungsten. As discussed later in Section 4.18.4, the erosion of graphite and beryllium is very high and dictates the use of thick tiles in high flux areas. This is in contrast to tungsten, which has a relatively low erosion yield, potentially allowing an armor thickness of only a few millimeters. Because the FoMs are essentially calculated on a per unit tile thickness, comparing tungsten with graphite can be somewhat misleading. However, because graphite Graphite tiles Coolant tubes Ion flux Support structure Monoblock Saddle type Flat plate GraphiteGraphiteGraphite Coolant tube Figure 2 Schematic diagram of the proposed monoblock first wall structure for the ITER reactor. Redrawn from Kuroda, T. et al. ‘‘ITER Plasma Facing Components,’’ ITER Documentation Series, No. 30, International Atomic Energy Agency (1991). 588 Carbon as a Fusion Plasma-Facing Material and beryllium are erosion-limited, the FoMs and the melting temperatures are useful evaluation tools. While the sublimation temperature of graphite (�3350 �C) is comparable to the melting point of tungsten (�3400 �C), it is clear that beryllium, which has a melting point of�1300 �C, is at a distinct disadvantage. Removal of beryllium, as well as other metallic PFCs, by melting has been seen in several large experiments. Performance calculations for graphite and CFCs have been conducted in both laboratory test stands and operating tokamaks. Some experimental data generated using electron beam simulation are given in Figure 5. Here, the power is deposited by a ras- tered electron beam for approximately one second up to surface heat loads of 11MWm�2. The samples were 2.5� 2.5 cm tiles, 1 cm in thickness, facing the beam. Each sample had a large notch machined into one edge (the highest stressed area) to serve as a stress intensifier. It was noted that, without the notch, the graphites did not crack. Figure 5 gives the maximum heat flux of which each material was tested and whether cracking of the tile occurred. The data indicate that CFC materials and graphites with a higher thermal conductivity and high density are superior. No cracking occurred in either the three composites studied, or the two FMI graphites, at the maximum power density applied. The superior per- formance of the composite materials agrees with the performance of CFCs in the large tokamaks such as TFTR and JT-60U. The reason for the superior performance of the CFCs and the graphites is most likely their low thermal expansion coefficient, high thermal conductivity, and high strength. In addition, the presence of the fibers in the CFCs may serve to 200 1000 Th er m al s tr es s fig ur e of m er it, F oM th Brush Wellman S65-C: Wrought be 316L stainless steel Pure tungsten Sigri Great Lakes H451: graphite Mitsubishi Kasei MFC-1: 1D-C/C Fiber Materials Inc. FMI-222: 3D-C/C Unocal, poco AXF-5Q: Graphite 104 105 106 300 400 500 Application temperature (�C) 600 700 800 900 1000 Figure 3 Thermal stress figure of merit for selected plasma-facing materials. 200 Th er m al s ho ck fi gu re o f m er it, F oM d 105 106 107 108 300 400 500 Application temperature (�C) 600 Brush Wellman S65-C: Wrought be Pure tungsten Sigri Great Lakes H451: Graphite 316L stainless steel Mitsubishi Kasei MKC-1PH: 1D-C/C Fiber Materials FMI-222: 3D-C/C Unocal, poco AXF-5Q: graphite 700 800 900 1000 Figure 4 Thermal shock figure of merit for selected plasma-facing materials. Carbon as a Fusion Plasma-Facing Material 589 blunt and arrest cracks, thus increasing toughness. All monolithic graphites shown in Figure 5, with the exception of the two FMI-HDFG materials, cracked. It is interesting to note that this graphite possessed the highest FoMd, even higher than that of the composites. However, strict correlation of improved performance with increased FoMd was not seen, although a loose correlation was noted. As pointed Power density (kW cm−2) 0 2 Isograph 880 Stackpole 1336 Stackpole 1225 Poco AXF5Q SGL H-478 Toyo Tanso IG-110 SCHUNK&EBE FE219 Carbone Iorainne 589 Poco ZXF5Q IBIDEN ETP-10 RINGS DORFF EK-98 Toyo carbon AX 280-K SGL H-489 SGL H-490 Stackpole 2204 Stackpole 2191 Union carbide TS1909 Union carbide TS1792 Union carbide ATJS Union carbide CGW Union carbide CGW II FMI HDFG FMI HDG BFG. 2-D staple KNIT FMI 4-D C/C fine FMI 4-D C/C coarse 4 6 8 10 Cracking range G ra p hi te o r C FC t yp e Failure range No cracking 12 Figure 5 The performance of several grades of graphite and graphite composites subject to thermal shock loading. Redrawn from Croessmann, C. D.; Gilbertson, N. B.; Watson, R. D.; Whitley, J. B. Fusion Technol. 1989, 127–135. 590 Carbon as a Fusion Plasma-Facing Material out by Watson,4 the CTE may be the most dominant property, with the lowest CTE graphites showing the best resistance to thermal shock. Finally, it should be noted that there are many issues regarding the selection of carbon materials as PFCs other than simply their thermal shock behavior. The issues of radiation damage, erosion, and hydro- gen retention are the three leading issues/drawbacks to the use of graphite as a PFC and they are discussed in the following sections. 4.18.3 Irradiation Effects on Thermophysical Properties of Graphite and CFCs 4.18.3.1 Graphite Irradiation Damage Gross physical property changes can occur in the graphite PFMs through two generic routes: (1) near- surface damage caused by interaction with plasma ions and, to a lesser extent, electrons, and, (2) bulk dis- placements caused by neutrons emanating from the plasma or back scattered by the surrounding struc- ture. Of the tokamaks, only TFTR had significant DþT fusion reactions and, therefore, experienced a significant flux of fusion neutrons (see eqn [1]). Even so, the dose from that TFTRwas not high enough for the structural materials to experience appreciable neutron effects. However, machines such as the ITER will see a significant neutron dose from both DþD and DþT reactions. As energetic particles travel through matter, they can interact with their surroundings, losing energy (per unit path length) in three ways: elastic collisions, electron excitations, and nuclear inter- actions. The interaction of primary interest from the materials property evolution point of view results from the particle elastic collisions with the graphite crystal. This has also been discussed in Chapter 4.10, Radiation Effects in Graphite (Section 4.10.4). If an ion or a neutron can provide sufficient energy to overcome an atom’s binding energy (Ed carbon�20–30 eV), the carbon can be displaced from its original lattice position. If the energy transferred to the displaced atom is sufficient to displace further atoms, a series of displacement events or a ‘cascade’ occurs. In the simplest interpre- tation, the Kinchin–Pease5 model is used to calculate the total number of atoms displaced. For example, if a carbon atom were ejected by the plasma and Carbon as a Fusion Plasma-Facing Material 591 reimpacted onto the carbon tile with a kinetic energy (Ecarbon) of 1 keV, the estimated number of atoms displaced (n) would be estimated as follows: n ¼ Ecarbon 2 Ed � � ¼�20 atoms ½5� The interaction of high-energy neutrons with matter is very similar to that of high-energy ions. The pri- mary difference between the two is the amount of energy transferred in a single collision and the dis- tance over which the interactions take place. An ion, which has a relatively large coulombic interaction radius, loses its energy over a short path length (typi- cally less than a micron). In contrast, the compara- tively small uncharged 14.1MeV fusion neutrons undergo only simple elastic or ‘billiard ball’ collisions with a mean free path between collisions of �10 cm. So, on average, a fusion neutron will have an elastic collision with a carbon atom once in 10 cm of graph- ite. The amount of energy transferred to the carbon in this first collision (Ec) is calculated by simple elastic theory as: Ec ¼ 4mcmnðmc þ mnÞ2 " # Eocos 2a ¼ 4� 12� 1ð12þ 1Þ2 " # 14:1MeVcos2a ½6� where mc and mn are the carbon and neutron mass (in amu), respectively, Eo is the neutron energy, and a is the angle between the neutron path before and after the collision. For a totally back scattered neu- tron (the maximum imparted energy), the energy transferred to the displaced carbon is �4.7MeV. Again, from eqn [5], the number of displaced carbon atoms in this 14.1MeV neutron collision event is nearly 100 000. The vast majority of these atoms do not stay ‘displaced,’ but condense back into the gra- phitic structure within a few picoseconds. To assess the effects such collision events will have on a mate- rial, a convention has been adopted to compare irra- diation doses. The displacement per atom (dpa) gives the average number of times an atom has been knocked from its original lattice position. The dpa is an integrated average quantity, and takes into account the atomic density, the interaction cross-section, and the neutron energy spectrum. For the next-genera- tion fusion reactors such as ITER, peak end of life values for PFMs due to neutrons will be on the order of tenths of a dpa, while power fusion reactors could potentially be subjected to greater than 10 dpa year�1. 4.18.3.2 Surface Effects While fast neutrons will produce relatively uniform atomic displacements, ions will produce very high near-surface damage. This damage can be on the level of hundreds of dpa, even for the experimental machines in use today and certainly for machines such as ITER and beyond. However, the damage is typically limited to much less than a micron in depth. The effect of this high damage level will be the reduction of a well-graphitized structure into a struc- ture that appears amorphous. However, these near- surface regions are subjected to erosion either by physical sputtering (caused by elastic collisions), or by chemical interactions. Both these effects are addressed in Section 4.18.4. A second surface radia- tion damage issue, that is, the ability of the thin damaged surface layer to retain and transport hydro- gen, is discussed in Section 4.18.5. 4.18.3.3 Properties and Property Evolution of Graphite Fiber Composite As mentioned earlier, the first wall materials in next- generation machines will receive many tens of dpa. At low doses ( 250 2.3 � 6 � 30 mm bend bars 4-point bending, 6.45/19.05 mm load/support spans 200 150 FMI-222 3-D CFC Poco AXF-5Q graphite100 50 B en d s tr es s (M P a) 0 0 0.2 0.4 Displacement (mm) 0.6 0.8 1 Figure 6 Comparison of the loading behavior of a typical graphite and carbon fiber composite. 592 Carbon as a Fusion Plasma-Facing Material exists on CFCs, mainly from the same source, but with some additional data from fusion research. These data suggest that CFCs have very similar irradiation behav- ior compared to graphite. In Chapter 4.10, Radiation Effects in Graphite, Burchell discusses radiation dam- age mechanisms in graphite, and some of the specific property changes that occur in fission reactor appli- cations. Because they are of special significance to fusion energy, the radiation effects in CFCs in gen- eral and the radiation-induced degradation in thermal conductivity in graphite and CFCs in particular will be focused on in the remainder of this section. How- ever, it is first important to contrast nuclear graphite (essentially a form of purified structural graphite) with that of graphite composites. For the purposes of dis- cussing graphite materials for fusion applications, the term composites is applied specifically to continuous fiber composites, typically woven, and infiltrated with pitch or some other resin that is graphitized to form a highly crystalline graphite matrix. The fibers compris- ing these composites are, as comparedwith most forms of graphite, highly crystalline and of comparatively high strength, elastic modulus, and thermal conduc- tivity. The fibers themselves are typically either poly- acrylonitrile (PAN) or Pitch derived. In general, one would select the PAN-based fiber, which is some- what less expensive, if the application required higher strength while the Pitch-based fibers would produce a product with superior elastic modulus and thermal conductivity. As observed in Sections 4.18.2.2 and 4.18.2.3, the composite materials, due to their typically higher strength and elastic modulus, have a superior performance in terms of thermal stress and thermal shock. Another key advantage of these materials stems from the fact that they tend to fail in a less abrupt manner than seen for graphite or ceramics in general due to the presence of the reinforcing fibers, which bridge evolving crack fronts. This can be seen by casual inspection of Figure 6, which compares the nuclear graphite Poco AXF-5Q (historically used in TFTR and for other nuclear applications) and the FMI-222 balanced weave, 3D CFC. From Figure 6, and by a comparison of the graphite and composite data of Table 2, it is clear that the FMI-222 CFC material has both higher bend strength and higher elastic modulus (greater slope) as compared to this Poco graphite. Moreover, it is clear fromTable 2 that other engineering properties of importance, such as strength and thermal conductivity, are superior for the CFC. These superior properties are primarily attributable to the exceptional quality of graphite fiber. Unlike nuclear graphite, which is on the order of 20% porosity with a relatively imperfect, heavily faulted, inhomogeneous amalgam of filler particles (such as coke) and graphitized binder (such as pitch; see Section 4.10.2 in Chapter 4.10, Radiation Effects in Graphite for a discussion of graphite man- ufacture), graphite fibers, while somewhat different depending on the starting material (PAN, Pitch, Rayon, etc.), are extremely uniform, and highly crys- talline with density that can approach theoretical den- sity. This leads to exceptional properties. For example, the PAN-based T-300 fiber has a tensile strength of 3.66GPa, slightly higher than the 2.41GPa strength of the P120 fiber of the FMI-222 composite of Table 2, or more than 40 times that of the Poco AXF-5Q graphite. Similarly, the elastic moduli of T-300 and P-120 fibers are 21 and 75 times the elastic modulus of the Poco AXF-5Q graphite. In the case of the P-120 fiber, which has been graphitized at a very high temperature, very long, defect-free basal planes oriented along the axis of the fiber result in excep- tional 1D thermal conductivity (640Wm�1K�1, twice that of copper). This property is the primary reason for the twofold increase in ambient thermal conduc- tivity of the FMI-222 composite as compared to the Poco AXF-5Q graphite. Clearly from this example of thermal conductivity, the architecture (fiber weave or loading) will determine the composite properties. Examples of practical fusion CFCs are the mate- rials chosen for consideration and application by the ITER project. Table 3 provides the nonirradiated thermophysical property data for selected CFCs Table 2 Comparison of thermophysical properties of a typical graphite and carbon fiber composite Poco AXF-5Q nuclear graphite FMI-222 3D carbon fiber composite Manufacturer Poco specialty Fiber Materials Inc. Architecture Near isotropic Balanced 3D weave Precursor Pitch Amoco P-120 fibers pitch matrix Grain size/unit cell size (mm) 9 900 Ambient thermal conductivity (Wm�1 K�1) 95 200 Apparent density (g cm�3) 1.78 1.96 Flexural strength (MPa) 86 175 Elastic modulus (GPa) 11 52 http://www.poco.com/MaterialsandServices/Graphite/IndustrialGrades/GradeChart/tabid/95/Default.aspx. Table 3 Thermophysical properties of CFCs of interest to fusion CX- 2002U MFC-1 INOX Sepcarb NS31 Sepcarb NB31 Dunlop concept 1 Constituents Pitch fiber X: 18%, Y, Z: 6% HIP pitch matrix K139 pitch fiber pitch matrix Amoco P55 pitch fiber CVI pyrocarbon matrix SiC by liquid Si infiltrate 2800 �C final heat treatment X: Amoco P55 pitch fiber Y, Z PAN fiber X: 27%, Y, Z: 4% CVI and then pitch matrix 2800 �C graphitization temperature X: Amoco P120 pitch fiber Y, Z PAN fiber volume 30%CVI pyrocarbon matrix 2450 �C graphitization temperature Density 20 �C 1.65–1.7 1.96 2.116 1.96 1.88 Specific heat (J Kg�1 K�1) 20 �C 0.71 0.76 0.73 0.7 CTE� 106 RT-400 �C X: 1.6, Y, Z: 5.2 X: �0.9, Y, Z: 12 X: �1.036, Y: 0.64, Z: 1.199 X: �0.339, Y: �1.376, Z: �0.018 X: �1.32, Y: 0.07, Z: 3.08 Thermal conductivity (Wm�1 K�1, 20 �C) 20 �C X: 368 X: 640 X: 265, Y: 124, Z: 109 X: 319, Y: 115, Z: 113 X: 413, Y: 102, Z: 78 500 �C X: 196, Y: 76, Z: 64 X: 196, Y: 72, Z: 68 X: 245, Y: 65, Z: 53 800 �C X: 146, Y: 58, Z: 49 X: 151, Y: 55, Z: 53 X: 78, Y: 52, Z: 38 Elastic modulus (GPa) 20 �C X: 100, Y, Z: 0.8 X: 120, Y: 55, Z: 40 X: 107, Y: 15, Z: 12 Ultimate strength (MPa) 25 �C X: 400, Y, Z: 3 X: 160, Y: 46, Z: 25 X: 130, Y: 30, Z: 19 1000 �C X: 200, Y: 56, Z: 36 X: 165, Y: 42, Z: 27 1500 �C X: 230, Y: 67, Z: 40 X: 185, Y: 50, Z: 30 Bend strength (MPa) 20 �C X: 39 X: 480, Y, Z: >5 Compressive strength (MPa) 20 �C X: 48 X: 216, Y, Z: >16 X: 102, Y: 31 X: 102, Y: 31 Shear strength (MPa) 20 �C XZ: 25, YZ: 15 Poisson’s ratio 20 �C XZ: 0.15, XY: 0.09, YZ: 0.15 XZ: 0.2, XY: 0.1, YZ: 0.1 Carbon as a Fusion Plasma-Facing Material 593 of the ITER project. A review of these properties emphasizes the anisotropic nature of the composite system, which is engineered through selection of the fiber type and route to matrix infiltration, fiber architecture, and final heat treatment of the system. All the materials for this application have been engineered with a preferred thermal conductivity direction (the x direction in the table), and in order to maximize thermal conductivity, the composites will tend to have a higher volume fraction of fibers in that direction and the fibers will be of the higher conduc- tivity pitch-based type. In the directions normal to http://www.poco.com/MaterialsandServices/Graphite/IndustrialGrades/GradeChart/tabid/95/Default.aspx Ir ra d ia tio n- in d uc ed d im en si on al c ha ng e (% ) Perpendicular Parallel 0 ITER range −0.1 −0.2 −0.3 −0.1 −0.2 −0.3 −0.4 −0.5 −0.6 0 0 0.2 0.4 0.6 Neutron fluence (1025n m–2), E > 0.1 MeV 0.8 1 A05 CX 2002U DMS 678 Figure 7 Dimensional change at low (ITER-relevant) fluence for select two-dimensional (2D) (AO5 and DMS 678 polyacrylonitrile (PAN)-based composites) and CX-2002U 3-C composite. Parallel and perpendicular refer to parallel and perpendicular to fabric lay-up for the 2D composite and y–z plane of PAN fibers for the CX-2002U. Reproduced from Bonal, J. P.; Wu, C. H. J. Nucl. Mater. 1996, 228, 155–161. 594 Carbon as a Fusion Plasma-Facing Material this preferred thermal conductivity direction, for strength, cost, and fabricability reasons PAN-based fibers are typically chosen. The composite INOX Sepcarb NS31 underwent a final processing step of 10� 2% liquid silicon infiltration. This silicon reacted with carbon-producing SiC, which is thought to mitigate chemical erosion and tritium retention while enhancing oxidation resistance. Also observed from Figure 6 is the clear differ- ence in the shape of the load–displacement curves for the two materials. Clearly, the composite material has significant nonelastic behavior, which is attributed to the progressive load transfer from the composite to the high-strength fiber as the matrix becomes exten- sively microcracked. This contrasts with the graphite material, which undergoes abrupt failure when the load exceeds some critical stress adequate to propa- gate a crack through the test article. This added toughness of the composite is another key attribute to the systems that make it particularly attractive for fusion applications where disruption (shock) loading tends to produce interconnected cracking in materi- als leading to loss of material mass. 4.18.3.3.1 Irradiation-induced dimensional changes in CFCs As discussed in Chapter 4.10, Radiation Effects in Graphite, irradiation-induced dimensional changes in graphite are highly anisotropic, and a strong function of irradiation temperature and neutron dose (dpa). The temperature range of interest for fusion applications varies from 100 �C in areas well removed from the plasma of experimental devices, to over 1000 �C for the surface of PFCs, which experience appreciable plasma flux, and for future power-producing machines. As described in detail in Chapter 4.10, Radiation Effects in Graphite, the mechanism of graphite irradiation-induced dimensional change is a combi- nation of intra- and intercrystallite effects. Within the crystallites, displacement damage causes an hai-axis shrinkage (within the basal plane) and a hci-axis growth (perpendicular to the basal plane.) The upcoming ITER reactor will be the first fusion reactor to provide a flux of neutrons to pro- duce measurable thermophysical effects to fusion structural materials. Even so, this will be a relatively modest fluence machine, with the maximum fast dose accumulating less than 1� 1025 nm�2 (E> 0.1MeV), or less than a displacement per atom, over its lifetime. The work of Bonal provides data on the dimensional changes in CFCs, which are expected in this dose range. Specifically, his work11 irradiated 2D and 3D composites to doses approaching 1 dpa in the tempera- ture range of 610–1030 �C. Figure 7 shows the dimen- sional instability that occurs in these materials in the sub-dpa region, specifically indicating a shrinkage. The work of Burchell12 in Figure 8 shows the dimensional change behavior of 1, 2, and 3 direc- tional composites for doses somewhat in excess of the ITER lifetime. In this example, solid cylinders were irradiated at 600 �C to doses ranging to 5 dpa and the resulting diameter and length measured. The behavior of each material can be explained by the accepted theory for dimensional change in graphite (Chapter 4.10, Radiation Effects in Graphite) after taking into account the individual fiber architec- tures, and by the observation that a model for fibers describes them as graphite fiber, filaments of circum- ferential or radial basal planes running parallel to the fiber axis. The irradiation-induced dimensional change of such a fiber is therefore a shrinkage in length and a growth in diameter. However, at doses Neutron dose (dpa) D im en si on al c ha ng e (% ) Unidirectional fiber composite (UFC) Random fiber (RFC) composite Three-dimensional balance weave 1 0 0.5 0 0 0 1 2 3 4 5 –1 –2 –3 –4 –0.5 –1 –1.5 –2 –2.5 –0.5 –1 –1.5 –2 –2.5 –3 –3.5 –4 Fiber axis Fiber axis Fiber axis Composite diameter-PAN composite II Fiber axis - PAN composite II Fiber axis - pitch composite Fiber axis Axis parallel to fiber axes Axis parallel to a set of fiber axes Axis perpendicular to fibers axes Figure 8 Dimensional change in carbon fiber composites at a moderately high neutron dose. Reproduced from Burchell, T. D. In Physical Processes of the Interaction of Fusion Plasmas with Solids, Plasma-Materials Interactions; Hofer, W. O., Roth, J., Eds.; Academic Press: New York, 1996; pp 341–382. Carbon as a Fusion Plasma-Facing Material 595 less than 1 dpa the dimensional change is relatively minor (Figure 7). As the dose is increased, the direc- tion perpendicular to the fiber axis is more or less unchanged while a significant shrinkage along the direction parallel to the fiber axis occurs. At about 2–3 dpa, swelling in the composite occurs in the perpendicular direction. The random fiber composite of Figure 8 has a random orientation of chopped PAN fibers in the plane of the composite. The speci- men diameter shows practically no change perpen- dicular to the fiber axis to about 4.5 dpa, though it exhibits �2% shrinkage parallel to the fiber axis. The 3D balanced PAN weave fiber has essentially isotropic shrinkage to a dose of �2 dpa, at which point the diameter of the fibers, and hence the sam- ple, begins to swell. Also given in the 3D composite plot in Figure 8 is the radiation-induced dimensional change parallel to the fiber axis of an Amoco P55 pitch fiber com- posite. This material was processed in an identical manner to the PAN fiber composite. From the plot, it appears that the pitch fibers, and thus the compos- ite, undergo slightly less shrinkage, possibly due to the higher fiber crystallinity. This hypothesis is also supported by the observation that fibers with higher final heat treatment temperatures tend to exhibit less dimension change13 and it is also consistent with the observation that elevating the heat treatment temper- ature of graphite reduces the irradiation-induced shrinkage.14 The irradiation-induced dimensional changes are of fundamental importance to the design and perfor- mance of the fusion structure, and even more so of the PFCs. This is due to the need to precisely define the plasma edge. For this reason, it is instructive to look at the irradiation effects at the higher dose and temperature conditions representative of the next- generation fusion power devices. The data shown in Figures 9 and 10 provide higher temperature dimen- sional swelling data for the FMI-222 3D CFC and MKC-1PH 1D CFC, which were model, high ther- mal conductivity CFCs studied in the early phases of the ITER composite development program.15 In Figure 10, the dimensional change of the 1D composite yields substantial swelling perpendicular to the fiber axis and equally impressive shrinkage parallel to the fiber. The FMI-222 of Figure 10, a nearly isotropic orthogonal weave pitch-fiber com- posite with equivalent fiber volume fraction in the x, y, and z directions, undergoes a positive dimensional change (swelling) parallel to the cylindrical axis of the sample, which increased with increasing temper- ature. The magnitude of swelling was in excess of 10% at the highest temperatures studied at the 2 dpa dose level. This is in contrast to the FMI-222 swelling data reported by Burchell12 and Snead,16 also for HFIR irradiation, though at a lower irradiation tempera- ture. Specifically, a contraction of 0.6% is interpolated from the data of Burchell for FMI-222 irradiated 600 0D im en si on al c ha ng e (% ) 5 10 15 Three-dimensional CFC (FMI-222) 2 dpa irradiation Length change (parallel to axis) −5 800 1000 Irradiation temperature (�C) 1200 1400 1600 Figure 10 Dimensional change at high irradiation dose and temperature for a one-dimensional carbon fiber composite. Reproduced from Snead, L. L.; Burchell, T. D.; Katoh, Y. J. Nucl. Mater. 2008, 381, 55–61. 800 −4 −2 0 D im en si on al c ha ng e (% ) 2 4 6 900 1000 1100 Irradiation temperature (�C) 1200 1300 1400 Length change (parallel to axis) One-dimensional CFC MKC 1-PH 2 dpa irradiation Diameter change (perpendicular axis) 1500 Figure 9 Dimensional change at high irradiation dose and temperature for a balanced three-dimensional carbon fiber composite. Reproduced from Snead, L. L.; Burchell, T. D.; Katoh, Y. J. Nucl. Mater. 2008, 381, 55–61. 100μm33-N1 Figure 11 SEM image of the top surface of an FMI-222 composite following irradiation to 980 �C, 2dpa. 596 Carbon as a Fusion Plasma-Facing Material at 600 �C to an equivalent fluence as the data of Figure 9. Snead16 reports on an 800 �C irradia- tion to a substantially higher dose (7.7� 1025 nm�2) than the Figure 8 dose (�2.4� 1025 nm�2). In this case, the material underwent a contraction of 3.6% along the length of a bend-bar (2.3� 6� 30mm). It was also noted in this work that the width and thickness direction exhibited swelling. Specifically, swelling parallel to the width direction (6mm) was 1.4% and swelling parallel to the thickness direc- tion (2.3 mm) was 5.9%. The overall dimensional effects were related to the effect of (measured) gross changes in the dimension of fiber bundles noting that gaps were evident on the surface of the bend bars. Figure 11 shows an example of the top surface of an FMI-222 composite irradiated in the present work to 980 �C, 2.4 dpa. This composite underwent very low swelling. By inspection of the figure the contraction of the fiber tows below the free surface of the sample is evident. However, there is evidence from this micro- graph that some of the fibers (particularly at the tow edge) have not withdrawn into the sample. This is evidence of shear within the fiber bundle as opposed to the tow–matrix interface. This observation is evi- dence of the large stresses that must be building in the composite under irradiation. The fact that the bundles are not failing at the tow–matrix interface also sup- ports the previous finding that, at least in the initial period of gross dimensional change, the load-carrying capacity of the composite has not been degraded. In fact, previous measurement of FMI-222 irradiated to a dose of �7.7� 1025 nm�2 (E> 0.1MeV) at 800 �C described a 54% increase in strength.16 Figure 12 shows an scanning electron microscopy (SEM) image comparing the 2 dpa surface of the FMI-222 composite of Figure 11 with cylindrical samples of the same size, also irradiated near 1000 �C, though at progressively higher doses. Clearly the dimensional instability continues with dose leading to gross changes in the composite. 2 dpa, ~980 �C 22-A 1 mm 1 mm 29-J1 1 mm11N 6 dpa, ~1025 �C 10 dpa, ~1065 �C Figure 12 Evolution of macrostructure for a three-dimensional composite under relatively high-dose, high-temperature irradiation. Table 4 Properties before and after irradiation of IG-110 graphite and ITER-relevant composites IG-110 graphite CC-312 MFC-1 CX-2002U X Y Z Young’s modulus (GPa) Unirradiated 8.83 34 74 87.6 14.9 Irradiated 11.5 31.3 98 87.2 18.4 Bending strength (MPa) Unirradiated 35.2 90.5 103.9 5.8 99.2 36.3 Irradiated 38.4 110.8 98.4 88.9 46.7 Bending fracture strain (%) Unirradiated 0.532 0.324 0.332 0.312 0.294 0.44 Irradiated 0.364 0.394 0.174 0.325 0.435 Compressive strength (%) Unirradiated 85 65.1 59.8 76.7 59.6 33.3 Irradiated 82 93.7 55.9 51.0 41.2 Compressive fracture strain (%) Unirradiated 2.67 0.79 0.165 3.72 0.71 1.94 Irradiated 1.49 0.393 0.156 0.11 1.82 Source: Eto, M.; Ishiyama, S.; Ugachi, H.; Fukaya, K.; Baba, S. J. Nucl. Mater. 1994, 212–215, 1223–1227. Carbon as a Fusion Plasma-Facing Material 597 4.18.3.3.2 Irradiation-induced changes in strength and modulus Significant increases in both strength and elastic modulus occur in graphite at dose levels as low as 0.01 dpa.8 This increase continues to high displace- ment levels until volumetric expansion and extensive micro cracking occurs. For graphite, the reversal to property degradation typically occurs at tens of dpa depending on the graphite type and irradiation tem- perature. The increase in modulus is a result of dislo- cation pinning by lattice defects produced by neutron irradiation. The magnitude of the increase is depen- dent on the perfection of the graphite. For most graph- ite types, a maximum modulus increase of 2–2.5 times the nonirradiated value is typical for irradiation tem- peratures less than 300 �C, with the change becoming less pronounced at higher irradiation temperatures. Irradiation-induced increase in strength occurs in a similar fashion as in the elastic modulus. Several authors17–23 report the effect of neutron irradiation on the elastic modulus of CFC. For example, Sato18 reports an increase of 42% and 30% in modulus following neutron irradiation to 1.2� 1025 nm�2, E> 0.18MeV in the temperature range of 750–810 �C for a 2D pitch fiber and PAN fiber composite, respec- tively. Similar to the irradiation-induced increase in strength, the absolute increase and percent increase in elastic modulus is highly dependent on starting material and irradiation condition. The irradiated and nonirradiated mechanical properties of some can- didate ITER PFC materials are shown in Table 4 for ITER relevant temperatures and doses somewhat higher than ITER neutron doses. Specifically, these materials were irradiated at approximately 1000 �C to a dose of about 2 dpa.20 The change in properties is relatively small because of the high irradiation tem- perature and the relatively low dose. As with elastic modulus, reported data on the effect of irradiation on the strength of CFCs are somewhat sparse.16–22 Snead23 has reported the strength and elastic modulus of the 3D pitch fiber composite FMI-222 for doses higher than expected for ITER, or more consistent with a fusion power reactor. Figure 13 gives the modulus as a function of dose to 32 dpa at 800 �C, exhibiting a marked increase to at least 10 dpa followed by a degradation by the 32 dpa 0 30 35 40 E la st ic m od ul us (G P a) 45 50 Thermal control Sonic elastic modulus Strain gage 55 60 10 20 Dose (dpa) 30 40 Figure 13 Effect of neutron irradiation on the elastic modulus of a balanced three-dimensional carbon fiber composite at high neutron dose. Reproduced from Snead, L. L.; Katoh, Y.; Ozawa, K. J. Nucl. Mater. 2010. 0 0 50 100 S tr en gt h (M P a) 150 Ultimate flexural strength Proportional limit strength 200 5 10 15 20 Dose (dpa) 25 30 35 Figure 14 Effect of neutron irradiation on the strength of a balanced three-dimensional carbon fiber composite at intermediate and high neutron dose. Reproduced from Snead, L. L.; Katoh, Y.; Ozawa, K. J. Nucl. Mater. 2010. 598 Carbon as a Fusion Plasma-Facing Material value. The same samples, as seen in Figure 14, exhibit more than a 50% increase in strength, which is retained even at the 32 dpa value. This is particularly remarkable given that the composite had undergone significant dimensional change in this dose range. 4.18.3.3.3 Thermal conductivity degradation CFCs for fusion applications are specifically designed to maximize thermal conductivity and for this reason irradiation-induced thermal conductivity degradation is of primary importance. As with ceramics, graphite thermal conductivity is dominated by phonon trans- port and is therefore greatly affected by lattice defects, such as those caused by neutron irradiation. The extent of the thermal conductivity reduction is therefore directly related to the efficiency of creating and annealing lattice defects, and is therefore related to the irradiation temperature. The effect of neutron irradiation on the thermal conductivity of graphite has been widely studied. The majority of the literature10,13,23–31 in this area has been in support of the gas-cooled, graphite-moderated, fission reactor program in the United States and United Kingdom and has focused on ‘nuclear’ gra- phites as well as more fundamental work on pyrolytic graphite.8,27,32,33 In recent years, the emphasis of graphite radiation effects research has switched to its use in PFCs of graphite fusion reactors.10,11,34 Because of the significant advances in carbon– carbon composite (CFC) processing and fiber develop- ment, very high thermal conductivity materials have been recently demonstrated and they have become attractive for high heat flux applications. The highest thermal conductivity CFCs are made from highly crys- talline graphite fibers having intrinsic conductivities approaching those of pyrolytic graphite. For example, vapor-grown carbon fibers35 have a thermal conductiv- ity of 1950Wm�1K�1. Along with advances in fiber properties, improvements have occurred in bothmono- lithic graphite and the CFC matrix-processing areas, which also have enhanced thermal conductivities. The physical processes governing the thermal conductivity of graphites, as well as the mechanisms responsible for the radiation-induced degradation in conductivity, are well established.8 For all but the poorest grades of carbon, thermal conductivity is dominated by phonon transport along the graphite basal planes and is reduced by scattering ‘obstacles’ such as grain boundaries and lattice defects. For graphites with the largest crystallites, that is, pyro- lytic graphite or natural flake, the in-plane room temperature thermal conductivity is approximately 2000Wm�1 K�1.36 The thermal conductivity of graphite-based mate- rials can be written as a summation of the thermal resistance due to scattering obstacles: K ðxÞ ¼ bðxÞ 1 Ku þ 1 Kgb þ 1 Ki � ��1 ½7� where b(x) is a coefficient that includes terms due to orientation (with respect to the basal plane), porosity, and some other minor contributors. This coefficient 0.1 0 150 �C 200 �C 250 �C 300 �C 450 �C 600 �C 920 �C Tirr= 1150 �C 0.05 N or m al iz ed t he rm al c on d uc tiv ity K irr /K un irr 0.1 0.15 0.2 0.25 0.35 Pile grade A graphite Measured at ambient temperature 0.3 0.4 1 10 dpa Figure 15 Degradation in thermal conductivity as a function of irradiation dose and temperature. Reproduced from Kelly, B. T. Plot Constructed from Personally Communicated Data. Carbon as a Fusion Plasma-Facing Material 599 is assumed in most cases to be constant with tem- perature, with a value of around 0.6. The first two terms inside the parentheses are the contributions to the thermal conductivity due to umklapp scattering (Ku) and grain boundary scattering (Kgb). The grain boundary phonon scattering dominates the thermal resistance (1/Kgb) at low temperatures and is insig- nificant above a few hundred degrees Celsius, depending on the perfection of the graphite. The umklapp scattering, which defines the phonon– phonon scattering effect on the thermal conductivity, dominates at higher temperatures and scales nearly as T2.8 The umklapp scattering therefore defines the upper limit to the thermal conductivity for a ‘perfect’ graphite. Following Taylor’s analysis,37 the umklapp- limited thermal conductivity of the graphite crystal would be �2200Wm�1 K�1 at room temperature, in close agreement with the best pyrolytic graphites or the vapor grown carbon fibers mentioned earlier. The third term in eqn [7], Ki, is the contribution to the basal plane thermal resistance due to defect scattering. Neutron irradiation causes various types of defects to be produced depending on the irradia- tion temperature. These defects are very effective in scattering phonons, even at flux levels that would be considered modest for most nuclear applications, and would quickly dominate the other terms in eqn [7]. Several types of irradiation-induced defects have been identified in graphite. For irradiation tem- peratures lower than 650 �C, simple point defects in the form of vacancies or interstitials, along with small interstitial clusters, are the predominant defects. Moreover, at an irradiation temperature near 150 �C,27 the defect that dominates thermal resistance is the lattice vacancy. Due to its sensitivity to the presence of defects, the temperature at which graphite is irradiated has a profound influence on the thermal conductivity degradation. As an example, Figure 15 shows one of the most complete sets of irradiation data on Pile Grade A nuclear graphite.38 This graphite is a medium- grained, extruded, anisotropic material with a room temperature thermal conductivity of 172Wm�1 K�1 in the extrusion direction. Figure 15 presents the normalized room temperature thermal conductivity of this graphite of various irradiation temperatures. It is seen that as the irradiation temperature is decreased, the degradation in thermal conductivity becomes more pronounced. For example, following irradiation at 150 �C, the thermal conductivity of this graphite appears to approach an asymptotic thermal conductivity of �1% of the original. The reason for this is that as the irradiation temperature is decreased, the fraction of vacancies surviving a cascade event increases, and thus the number of vacancies available to scatter phonons is much higher for the lower temperature irradiation. Data have been published for CFCs whose thermal conductivities are similar to those of nuclear graphites, showing degradation similar to that expected from the graphite literature. For example, Burchell34 has shown that the saturation thermal conductivity for a 3-directional composite (FMI-222, Kunirr¼ 200W m�1 K�1 at RT) is reduced to �40% of the origi- nal room temperature conductivity following fast neutron irradiation at 600 �C. Published data for the degradation of thermal conductivity in highly con- ductive CFCs have led to the conclusion that a higher initial conductivity composite results in higher abso- lute conductivity after irradiation.39,40 Figure 16 demonstrates this point. At the extremely damaging irradiation temperature of �150 �C, it is observed that the absolute reduction (Kunirr� Kirr) is substan- tially greater for the high thermal conductivity mate- rials compared to the lower conductivity CFCs and graphite, as seen in Figure 16, although the normal- ized fraction (Kirr/Kunirr) is approximately the same for all the carbon materials in the figure. Moreover, saturation in thermal conductivity degradation occurs at a neutron dose of �1 dpa. Data for higher irradia- tion temperatures11,31 show that the higher thermal conductivity materials have a slightly larger fractional 600 Carbon as a Fusion Plasma-Facing Material change in thermal conductivity (Kirr/Kunirr) compared to lower conductivity materials, although the absolute value of the irradiated thermal conductivity is still greater for the higher conductivity materials. A com- parison of thermal conductivity degradation for a nuclear graphite (CH-45) with the composites FMI-222 and MFC-1 is given in Figure 17.31 0.0 0 100 200 300 Th er m al c on d uc tiv ity (W m –1 K –1 ) 400 500 600 700 FMI 1D Tirr= 150 �C, HFIR core Measurements at ambient temperature MFC-1 RGTi graphite Copper Hercules 3D FMI 222 H451 graphite 0.01 Damage level (dpa) 0.1 Figure 16 Comparison of absolute degradation in thermal conductivity for various graphite and carbon fiber composite materials irradiated at low temperature. Reproduced from Snead, L. L.; Burchell, T. D. J. Nucl. Mater. 1995, 224, 222–229. 0 50 100 Th er m al c on d uc tiv ity (W m –1 K –1 ) 150 200 250 300 350 400 Time in reactor (h) 0.1 1 10 100 1 H451 graphite; Tirr~430 �C CFC : FMI-222; Tirr~310 �C CFC : MFC-1; Tirr~430 �C Figure 17 Comparison of the effect of neutron dose on the th high conductivity carbon fiber composites. Reproduced from Sn For the low-dose regime relevant to machines such as ITER (less than 1 dpa or about 500 h in this figure), the conductivity is seen to decrease by a factor of two for the highest conductivity material (MFC-1) and by about 30% for the nuclear graphite. An algorithm has been developed to predict the thermal conductivity degradation in a high thermal conductivity composite (�555Wm�1 K�1 at room temperature) as a function of radiation dose and temperature.41 The absence of irradiation data on CFCs of this type required the use of data from intermediate thermal conductivity materials and pyrolytic graphite to derive an empirical radiation damage term.24,27,28,39,42 An analysis of the effects of temperature and neu- tron dose on the thermal conductivity is shown in Figure 18. Specifically, the algorithm assumed the nonirradiated properties of the unidirectional fiber composite MFC-1 material compiled with an empir- ical radiation damage term. As with the experimental data of Figures 15 and 16, it is seen in Figure 18 that an enormous loss in thermal conductivity occurs at low irradiation temperatures. Presently, only a few data points exist that are relevant to the validation of this algorithm, and these are also plotted on the figure.39 The data do agree within the errors of irradi- ation temperature and thermal conductivity measure- ment, with the algorithm predictions. However, they 000 H451 graphite; Tirr~430 �C CFC : FMI-222; Tirr~310 �C CFC : MFC-1; Tirr~430 �C ~1 dpa ~2 dpa Time in reactor (h) 0 200 400 600 800 1000 ermal conductivity degradation of a nuclear graphite and ead, L. L. J. Nucl. Mater. 2008, 381, 76–82. 300 0 Data from Bonal, et al. dpa = displacement per atom 100 200 Th er m al c on d uc tiv ity (W m –1 K –1 ) 300 400 Conductivity at measurement temperature Unirradiated Irradiated 0.0 dpa 0.001 dpa 0.005 dpa 0.01 dpa 0.05 dpa 0.1 dpa 0.5 dpa 1 dpa 400 500 600 Composite temperature (�C) 700 800 900 1000 1100 1200 Figure 18 Model for thermal conductivity degradation as a function of irradiation temperature and dose for high conductivity one-dimensional carbon fiber composite. Monoblock design Flat plate design Tmax= 740 �C MFC-1PH 1D-C/C GlidCop DC copper Water coolant Unirradiated 1–3 dpa irradiation 1–3 dpa irradiation Tmax= 1240 �C Tmax= 920 �C – 700 – 1200 – 1100 – 940 – 810 – 690 – 570 – 440 – 320 – 260 – 200 – 130 – 170 – 220 – 300 – 390 – 520 – 700 – 870 – 640 – 570 – 500 – 390 – 360 – 320 150 220 290 140 33 .6 m m 8 m m Figure 19 Effect of neutron irradiation on thermal conductivity-driven temperature evolution in a monoblock and flat-plate divertor design. Carbon as a Fusion Plasma-Facing Material 601 are insufficient to validate the algorithm and the need clearly exists for additional data for this purpose. To illustrate the usefulness of such an algorithm, and the significance of the issue of thermal conductivity degradation to the design and operation of PFCs, this algorithm has been used to construct Figure 19, which shows the isotherms for a mono- block divertor element in the nonirradiated and 0 0 0.2 0.4 ~200 �C irradiation 0.01 dpa 0.6 0.8 1 200 400 600 Annealing temperature (�C) N or m al iz ed t he rm al c on d uc tiv ity K irr /K un irr 800 1000 1200 1400 1600 FMI-1D FMI-222 H451 MFC-1 RGTi Figure 20 The effect of annealing on recovery in irradiation-degraded thermal conductivity of graphite materials. Reproduced from Snead, L. L.; Burchell, T. D. In Reduction in Thermal Conductivity due to Neutron Irradiation, 22nd, San Diego, CA, July 16–21, 1995; AmericanCarbon Society: San Diego, CA, 1995; pp 774–775. 602 Carbon as a Fusion Plasma-Facing Material irradiated state and the ‘flat plate’ divertor element in the irradiated state. In constructing Figure 19, the thermal conductivity saturation level of the 1 dpa given in Figure 18 is assumed, and the flat plate and monoblock divertor shown are receiving a steady state flux of 15MWm�2. Both composite materials have been assumed to be in perfect contact with a copper coolant tube or plate. Figure 19 clearly shows two points. First, a very high conductivity composite is required to handle the extreme heat fluxes expected if the temperature is to be limited to 0 200 0.001 0.01 S p ut te rin g yi el d (Y ; a to m s p er io n) 0.1 1 1 keV H 1 keV D 3 keV He Physical sputtering Chemical erosion Radiation-enhanced sublimation 400 600 800 Temperature (�C) 1000 1200 1400 1600 Figure 21 Mechanisms of carbon removal from a graphite plasma-facing material as a function of temperature. Reproduced from Roth, J.; Bodhansky, J.; Wilson, K. L. J. Nucl. Mater. 1982, 111–112, 775–780. 0 0.0001 0.001 C he m ic al e ro si on y ie ld p er in ci d en t io n 0.01 0.1 H� onto soft amorphous C:H film H� onto pyrolytic preirradiated by D+ -and- hard amorphous C:H films H� onto redeposited carbon H� onto pyrolytic (0.4 keV) H+ onto pyrolytic 200 400 Temperature (�C) 600 800 Figure 22 Chemical erosion of graphitic materials as a function of temperature. Carbon as a Fusion Plasma-Facing Material 603 The combination of energetic damage plus chemical reaction, which is sometimes referred to as ‘chemi- cal sputtering,’ is discussed by Jacob and Roth48 and others. For low(RT) and intermediate temperatures, from 400 to 1000 �C (Figure 21), the volatilization of car- bon atoms by energetic plasma ions becomes impor- tant. As seen in the upper curve of Figure 21, helium does not have a chemical erosion component of its sputter yield. In the currently operating machines, the two major contributors to chemical erosion are the ions of hydrogen and oxygen. The typical chemi- cal species that evolve from the surface as measured by residual gas analysis49 and optical emission50 are hydrocarbons, carbon monoxide, and carbon dioxide. The interaction of hydrogenwith graphite appears to be highly dependent on the ion species, on mate- rial temperature, and on the perfection and type of the graphite. This is illustrated in Figure 22, which shows typical bell-shaped thermally acti- vated erosion yield curves for hydrogen and deute- rium ions on graphite. The shape of the yield curve is influenced by the competition for hydrogenation from the sp2 and sp3 hybridization states.51,52 Hydrogen ions incident to the surface are slowed down and pref- erentially attach to sp2 carbon atoms (such as graph- ite edge plane atoms) forming sp3 CH3 complexes. Above approximately 400K, these CH3 complexes can be released, thus returning the structure to the sp2 state. It is important to note that this phenomenon will only happen in the presence of simultaneous ion damage. It will not occur simply due to a thermal process. This step leads to chemical erosion products (a host of erosion species are possible). The ability of hydrogen to continue to be bonded to carbon drops as the temperature goes up. If there are no CH or CH2 precursors on the surface, then no volatile CH3 or CH4 complexes can be formed, and thus there is no chemical erosion. This balance yields a maximum erosion rate, which for undamaged pyrolytic graphite resides at�280–600 �C.53 It is noted that more recent work by Balden54 has determined the maximum to be in the range of 872–1222 �C. This mechanism was first elucidated by Horn55 and Wittmann.56 The rate of formation of CH2, CH3, and complex hydrocar- bons from atomic hydrogen in well-graphitized mate- rial is fairly low unless the material is altered (damaged) in the near-surface layer. For preirra- diated pyrolytic graphite (i.e., damaged graphite, meaning that a carbon atom has been removed from its lattice position, thus increasing the available sp2 sites) preirradiated by high-energyDþ of Hþ ions, the total erosion yield following exposure to low-energy hydrogen increases dramatically. This is illustrated in the upper curves of Figure 22 that show more than an order of magnitude increase in erosion yield over D 4/ D -i on ) 0.1 Pyrolytic SEP-CFC (9% B) BE/C (JET) S2508 (3%B) USB 15(15% B) GB 120 (20% B) B4C 604 Carbon as a Fusion Plasma-Facing Material the undamaged case. This increased carbon loss has been attributed to the creation of active sites for Ho attachment.57,58 This structurally dependent mech- anism is supported by the data of Phillips et al.,59 which shows a factor of two difference in erosion yield between high- and low-quality pyrolytic graphite. 100 0.001 0.01 C he m ic al y ie ld (C 200 300 400 Temperature (�C) 500 600 700 800 Figure 23 Effect of including the dopant atom boron on the suppression of chemical erosion of graphite. Reproduced from Roth, J.; et al. J. Nucl. Mater. 1992, 191–194, 45–49. 4.18.4.2 Doping of Graphite to Suppress Erosion Surface treatment of PFMs, while extremely effec- tive for the current day short-pulse tokamaks (pulses typically less than a few seconds), are of limited value for the next-generation (quasi-steady state) machines because of the significant surface erosion expected. However, forming graphite or CFC homogenously with various erosion-mitigating elements is possible. The mechanisms for this mitigation are twofold: (1) geometric shielding by low erosion yield par- ticulate, and (2) changes in local chemical reactivity due to the presence of doping atoms. Promising results have been obtained by doping of graphite with boron, which resides substitutionally in the graphite lattice, trapping migrating interstitials and altering the electronic structure of the material. Boron doping has been shown60 to both reduce the erosion due to oxygen, and to significantly reduce the sputtering yield due to methane formation. How- ever, other factors, such as the drastic reduction in thermal conductivity that is unavoidable in boro- nized graphite, need to be factored into the overall picture. Boron is discussed in some detail here, mostly because it had received great early attention. However, newer dopant combinations61 have served to suppress erosion, and they have also not had such a negative impact on thermal conductivity, and are therefore considered superior for PFC application. As discussed by Garcia-Rosales,62 until the mid-1990s, boron was the primary ‘dopant’ of interest in fusion PFM graphites for chemical erosion mitiga- tion. As seen in Figure 23,63 the inclusion of up to 15% boron in graphite can result in significant (an order of magnitude in the peak regions) reduction in the erosion yield. Studies to date indicate that for effective suppression a minimum of 3% boron in graphite is required. The mechanism behind this suppression may include the reduced chemical activ- ity of the boronized material, as demonstrated by the increased oxidation resistance64 or the suppressed diffusion caused by the interstitial trapping at the boron sites. From the mid-1990s onward, many other metallic element additions to graphite were studied for their possible beneficial effect on erosion mitigation. Specifically, elements such as silicon, titanium, tung- sten, and vanadium have been studied with varying levels of success.62 These elements are somewhat less effective in erosion mitigation than boron, though a factor of two reduction is to be expected.65–67 In a recent review by Balden,61 considerably higher reduc- tions in erosion are noted. More recently, emphasis has been on the use of multielement doping strategies.61 Because the removal of graphite is significant both in terms of gross material loss (possible consumption of the entire wall for power devices) and enhanced tritium retention for the resulting carbon dust, the effect of any additive to graphite in terms of physical properties or impact on plasma performance when eroded (Section 4.18.2.1) needs to be considered. With the exception of boron, additive elements dis- cussed in the previous paragraph will all have a negative impact on plasma performance in compari- son with carbon atoms, and therefore the balance of reduced mass loss compared to enhanced parasitic radiative plasma loss (eqn [2]) must be considered. As for physical properties, at levels well below the threshold at which they are effective for erosion suppression (�3%), they are direct substitutional elements in the graphite lattice, effecting significant reduction in thermal conductivity (due to their mass-defect phonon scattering.) In contrast, titanium doping, as evidenced in materials such as the Russian RGTi material, serves to enhance the graphitization Carbon as a Fusion Plasma-Facing Material 605 process, resulting in very well-crystallized materials of high (though somewhat anisotropic) thermal con- ductivity. A comparison of several element additions to graphite and their effects on the properties of graphite has been carried out by Paz68 and discussed by others.47,69 It is seen (Figure 2468) that all the metallic inclusions studied, with the exception of silicon,70 had, at these graphitization temperatures, the effect of enhancing the effective length (perfec- tion) of the basal plane of the graphite crystals, which is directly linked to enhanced thermal conduc- tivity. In comparison the basal crystal lengths of the Poco nuclear graphite see Table 2 and the Russian RGTi,71 which is processed at a similar graphitiza- tion temperature but with an applied electric current are shown in Table 2. The right hand side of Figure 24 shows the effect of varying the amount of titanium on the crystallite size, indicating that there is an increase in crystallite size with an increase of up to a few atomic percent of titanium. In addition to the thermal process for chemical erosion described in the previous section, a second route to erosion, which is limited to the regions very near the surface, is also of importance.62 Specifically, for ions of Temperature (K) 2000 1500 1000 800 700 600 Codeposits and implants TFTA N3-15 DIII-D ASDEX a-C:D Textor (estimate) DIARC film Upper limit [2] (HPG99, EK98, AXF-5Q, ATJ, JET 2D C/C) Pure graphites HPG99 E ro si on y ie ld /( C /O 2) EK98 Okada and Ikegawa Duval Olander et al. basal plane, CO Olander et al. basal plane, CO 2 Olander et al. prism plane, CO 2 Lang and Magnier Rodriguez-Reinoso et al. graphite M Rodriguez-Reinoso et al. graphite P Vietzke et al. CO Penzhorn et al. EK98 Penshorn et al. JET 2D C/C McKee C/C Bacos 0.4 10−11 10−10 10−9 10−8 10−7 10−6 10−5 10−4 10−3 10−2 10−1 0.6 0.8 1.0 1.2 (1000/T) / (1 K–1) 1.4 1.6 1.8 2.0 Olander et al. prism plane, CO Ion implant (lower limit) 500 Figure 25 Temperature effect on erosion yield in various graphite grades. Reproduced from Davis, J. W.; Haasz, A. A. J. Nucl. Mater. 1999, 266–269, 478–484. 606 Carbon as a Fusion Plasma-Facing Material with titanium,75 boron,76,77 beryllium,78 and silicon.79 Comprehensive reviews can be found for the chemi- cal erosion of graphite46,47,55,56,80 doped graphite by hydrogen,62 and also an article on the surface treat- ment of a graphite wall by Winter.75 4.18.4.3 Physical Sputtering When a relatively high-energy impacting particle transfers energy to a near-surface carbon atom in an amount sufficient to overcome the lattice bond energy or surface binding energy, some carbon atoms may be displaced and these may move in a direction defined by the angle between their path and the initial path of the impacting atom. Analogous to striking a billiard ball, this angle must be between 0� and 90�. The energy imparted to the displaced atom follows that given in eqn [6]. For a relatively high-energy atom striking a surface normally, the recoiling atom cannot be sputtered from the surface. However, in off-normal impact or displacement cas- cade events from fusion neutrons that occur near the surface, some fraction of atoms will be emitted (physically sputtered) from the graphite surface. The amount of material lost from the surface is defined by the sputtering yield (Y ), which is the number of target PFM atoms emitted per plasma ion impacting the surface. From eqn [6], the energy transferred to a target PFM atom, which is directly related to the erosion yield, is a strong function of the impacting particle mass and the mass of the material being sputtered. The impact angle also has a large effect on the number of atoms that receive adequate kinetic energy normal to the PFM surface to be physically sputtered. The plasma ions travel along the magnetic field lines that are at a shallow (grazing) angle with the PFM, typically 1–5�, though the ion impact angle will be modified by surface potentials and collisional processes.81 The quantitative effect of the mass, energy, and angle of impact on the sputter yield of impacting deuterium ions is shown in Figure 26(a) and 26(b). As the kinetic energy of deuterium increases, the total amount of energy transferred to the target atoms increases, and the average amount of energy per colli- sion results in greater erosion. From Figure 26(a), it may be seen that the physical sputtering yield of light target atoms is considerably greater than that for the heavy atoms, primarily due to the reduced impact energy required to overcome the displacement energy of the higher target atoms. For example, in purely kinetic terms, approximately 20 eV is required to dis- place an atom of carbon from the surface, while 220 eV is required for an atom of tungsten. In the sub-keV energy range of plasma fuels, the high yield materials are therefore carbon and beryllium. As the impacting ion energy increases, the sputtering yield for all materials decreases as the depth of interaction of 0.1 0.01 Be C Fe Mo W 0.001 S p ut te rin g yi el d (Y ) Be Fe Mo WC 0.0001 10 100 1000 (a) Impacting deuterium energy (eV) 104 105 10 1 0.1 S p ut te rin g yi el d (Y ) 0.01 0 20 40 60 Angle (a) 80 Poco graphite Pyro : polished Pyro : basal Pyro : edge Graphite (a) (b) 3 keV C+ 2 keV D+ Trim.SP code Figure 26 Sputtering yields for plasma-facing material atoms of various candidate materials (a); and as a function of angle of incidence for various graphite materials. Carbon as a Fusion Plasma-Facing Material 607 the impacting ion becomes too great for displaced atoms to back scatter to the surface. In the case of graphite, the majority of the sputtered material comes from the top few atomic layers.82 With the correct combination of incident energy and target mass, it is possible for the sputtering yield to exceed unity, that is, more than one atom leaves the surface for every impacting particle. This quickly leads to what is called the catastrophic ‘carbon bloom,’ that is, the self-accelerating sputtering of carbon. As can be seen in Figure 26(b), this problem is the worst for carbon self-impacts at grazing angles to the surface. 4.18.4.4 Radiation-Enhanced Sublimation The limiting temperature for graphite use in fusion systems is defined by thermal sublimation (�1500– 2000 �C). However, a process that is very similar to thermal sublimation (in cause and in effect) appears to define the current temperature limit. This phe- nomenon, which is known as radiation-enhanced sublimation (RES), is not clearly understood yet, but dominates above a temperature of about 1000 �C, and increases exponentially with increasing temperature. One theory says that the process responsible for initiating RES follows from the earlier discussion of radiation damage in graphite. Specifically, in a displacement event, a Frenkel pair is created. The interstitial has a low (�0.5 eV) migration energy, is quite mobile between the basal planes, and thus diffuses readily. Some fraction of these interstitials are condensed at vacancy sites, which are essentially immo- bile below about 700 �C (migration energy �4 eV). Other migrating interstitials can be trapped by microstructural defects or can coalesce into simple clusters, which limits their mobility. However, some fraction of the interstitials diffuse to the surface of the graphite and thermally sublime. The thermal sublimation of radiation-induced interstitials is RES, and must be distinguished from both physical and chemical sputtering. Time of flight measurements have shown that the thermal energy of RES ions has a Maxwellian energy distribution, which is directly coupled to the mean surface temperature.83 This clearly distinguishes RES atoms from physically sput- tered atoms, which exhibit highly anisotropic energy distributions. RES atoms are also distinguished from thermally sublimed species in that only single carbon atoms are detected, whereas single atoms and atom complexes (C2, C3, . . .) are found during thermal sublimation. Another theory for the explanation of RES is simply that the bombarding hydrogen ions turn the very near-surface region into a low-density amorphous zone. A very large fraction of the carbon atoms in this zone are now edge atoms with weak bonding to the connecting atoms. These edge atoms are much more easily thermally volatized into the plasma. The effect of RES in the next generation of high surface particle flux fusion systems is presently unclear. Evidence suggests that the erosion yield does not scale linearly with flux, as physical sput- tering does, but may in fact decrease significantly with increasing flux.84 Moreover, as with chemical erosion, the inclusion of interstitial boron into the crystal lattice can decrease RES and shift the thresh- old to higher temperatures. Boron will volatilize above �1500 �C, thus limiting the PFM temperature to 608 Carbon as a Fusion Plasma-Facing Material 4.18.4.5 Erosion of Graphite in Simulated Disruption Events Finally, the effect of plasma disruptions needs to be considered. Section 4.18.2 discussed the thermo- mechanical response of the PFCs to the excessive plasma energy in a disruption. This large thermal energy dump can additionally cause enhanced ero- sion due to the increased particle flux, elevated sur- face temperature, or simply by exfoliation of the surface due to thermal shock. The latter two material losses are reduced for materials with high thermal conductivity. This has been demonstrated experi- mentally, and is shown in Figure 27,3 which gives weight loss as a function of thermal conductivity for a number of graphites and composites of varying thermal conductivities subjected to one electron beam pulse at 4.1MWm�2. As discussed in Section 4.18.2, and as seen in the data of Figure 27, high thermal conductivity materials reduce the surface tempera- ture, and hence the overall erosion yield, during a disruption. 4.18.5 Tritium Retention in Graphitic Materials Tritium retention and transport is a critical phenom- enon for graphite in fusion systems in general, and it is the subject of a chapter by Causey, San Marchi, and Karnesky in this series. In the previous section of this 50 0.5 1 1.5 2 W ei gh t lo ss (m g) 2.5 3 3.5 4 4.5 100 Mean ambient thermal conductivity (W m–1K–1) 150 200 250 300 350 Figure 27 Weight loss as a function of graphite thermal conductivity. Reproduced from Akiba, M.; Madarame, H. J. Nucl. Mater. 1994, 212–215, 90–96. chapter, the interaction of the plasma particle flux with the surface of graphite was discussed. However, the fate of the implanted particles, most impor- tantly deuterium and tritium, following their impact with the graphite surface is also an important issue and is seen by some as the major impediment to the use of graphite as a PFM.85 Quantification of the problem and determination of possible mitigat- ing steps is complicated by experimental data, which can vary by orders of magnitude,86–92 as reviewed by Wilson.93 The primary concern over retention of fuel in the PFC is the inventory of hydrogen adsorbed into the graphite and the subsequent release of near-surface hydrogen (due to physical or chemical sputtering, etc.) as plasma discharge begins. The hydrogen sput- tered from the wall oversupplies the plasma edge with fuel, causing instabilities and making plasma control problematic. Tritium inventory concerns are generally safety-related but can have significant economic consequences because of the high cost of tritium. Tritium release to the environment in an accident situation had limited the allowed inventory in TFTR, and was a significant consideration for the sighting of the ITER. It has been estimated84 that as much as�1.5 kg of tritium would reside in the graph- ite PFM of ITER, corresponding to an additional fuel cost of 1.5–3million dollars. A source of trapped hydrogen, not discussed in detail here, which may dominate the tritium inventory in ITER-like machines, is the ‘codeposited layer.’94 This layer is formed by the simultaneous deposition of carbon, which is eroded from the first wall, and hydrogen. Thick layers of carbon redepos- ited to low erosion areas are common, and have been seen in all large tokamaks utilizing graphite PFMs. As this layer grows, the hydrogen contained therein cannot be liberated by surface sputtering and becomes permanently trapped. This problem is unique to graphite and will require continual sur- face conditioning to minimize the total inventory of trapped species. It represents a vast sink for tritium and therefore must be managed in some way. Below is a discussion of the retention of tritium in bulk graphite. The physical process involved in the retention of hydrogen, as it corresponds to graphite PFMs, is fairly well understood. The energetic hydrogen iso- topes are implanted to depths of less than a micron in the PFM surface. Once implanted, the hydrogen ions are either trapped, reemitted, or diffuse through the bulk. At temperatures less than 100 �C,95,96 Carbon as a Fusion Plasma-Facing Material 609 the majority of ions are trapped near the end of their range. These trapped ions are not in solution in the graphite, but are held97 in the highly defected struc- ture. The amount of hydrogen isotope that can be accommodated is largely dependent on the implan- tation temperature,96,98 the trap types and densities (defects), and, to a lesser extent, by the implantation depth.99,100 One model for bulk hydrogen trapping presented by Atsumi101 is shown in Figure 28. In his work, two distinct traps have been identified. The lower energy trap (2.6 eV) is associated with edge planes of the graphite crystals, with total trapping therefore depending on the effective size and accessi- bility of the crystals to diffusing hydrogen species. The second, higher energy trap (4.4 eV), is associated with dangling bonds, albeit at the edge of an interstitial loop. As the formation of small interstitial loops is one of the primary effects of neutron irradiation, the formation of these deeper traps is directly affected by neutron fluence. The total retained isotopic H can reach as much as 0.4–0.5 H/C in the implanted layer at room temperature.95,100,102 As the amount of implanted hydrogen increases toward its saturation value, a larger fraction of ions are released from the graphite surface. At intermedi- ate and high temperatures (>250 �C), diffusion of hydrogen in the graphite lattice occurs. This diffu- sion is most likely along internal surfaces such as micropores and microcracks, while transgranular dif- fusion has been seen above 750 �C.103,104 This bulk diffusion, along with the associated trapping of hydrogen at defect sites, has been studied widely Molecular diffusion (with a sequence of dissociation and recombination) Gas permeation through open pores Absorption (Rate-determining step) (r Desorption ∼mm 1–50mmH2 ED= 1.3 eV Figure 28 Schematic of the processing of hydrogen trapping J. Nucl. Mater. 2003, 313–316, 543–547. with quite variable results. This variation can be seen in Figure 29, where the temperature depen- dence of the hydrogen diffusion coefficient for sev- eral carbon and graphite materials is shown. It is expected that the diffusion of hydrogen through graphite would be highly dependent on the graphite microstructure, which may explain the wide range of the data of Figure 29. In any event, the transport of hydrogen through the bulk graphite and associated solubility limits can significantly increase the hydrogen inventory for fusion devices. The effect of the perfection of graphitic structure on the solu- bility of hydrogen is shown by Atsumi’s data105 in Figure 30. The data in Figure 30 indicate that the more defect-free, highly graphitized materials have a lower solubility limit. Further evidence for the role of structural perfection comes from the observation that materials that have been disordered by neutron irradiation have significantly higher solu- bility for hydrogen.105,106 The effect of atomic displacements on the hydro- gen retention of graphite was first shown by Wampler using 6MeV ion beams.107 Wampler used four types of intermediate and high quality graphites, which were irradiated with a high-energy carbon beam at room temperature, and then exposed to deuterium gas. Wampler’s results indicated that the residual deuterium concentration increased by more than a factor of 30–600 appm for displacement doses appropriate to ITER. However, for reasons that are not entirely clear yet, neutron-irradiated high-quality CFCs retain significantly less tritium than expected ate-determining step) (rate-determining step) (>90%) ( 10−9 10−10 10−11 10−12 10−13 10−14 10−15 1 - Atsumi et al. 8 6 4 2 1 3 5 7 D iff us io n co ef fic ie nt (c m 2 s− 1 ) 2 - Elleman 3 - Saeki 4 - Maika et al. 5 - Rohrig et al. 6 - Causey 7 - Morita et al. 8 - Tanabe and Watanabe 0.4 0.6 0.8 103/T (K) 1 1.2 1.4 10−16 10−17 10−18 Figure 29 Variation in hydrogen diffusion coefficient as reported in the literature. Adapted from Causey, R. A. J. Nucl. Mater. 1989, 162–164, 151; Atsumi, H.; Tokura, S.; Miyake, M. J. Nucl. Mater. 1988, 155–157, 241; Saeki, M. J. Appl. Radiat. Isot. 1983, 43, 739; Malka, V.; Rorhig, H. D.; Hecker, R. In Tritium Technology in Fission, Fusion and Isotope Application; Dayton, OH, 1980; Rohrig, H. D.; Rischer, O. G.; Hecker, R. J. Am. Ceram. Soc. 1976, 59, 316; Morita, K.; Ohtsuka, K.; Hasebe, Y. J. Nucl. Mater. 1989, 162–164, 990; Tanabe, T.; Watanabe, Y. J. Nucl. Mater. 1990, 179–181, 231–234. Graphitic perfection (%) 0 20 10 100 1000 H yd ro ge n so lu b ili ty (a p p m ) 104 40 60 80 100 Unirradiated Neutron irradiated Figure 30 Effect of perfection of graphite crystal on hydrogen solubility. Reproduced from Atsumi, H.; Iseki, M.; Shikama, T. J. Nucl. Mater. 1994, 212–215, 1478–1482. Radiation dose (dpa) 0.001 1 10 100 Tr iti um r et en tio n (a p p m ) 1000 N3M graphite FMI-222 CFC MFC-1 CFC 104 0.01 0.1 1 10 Figure 31 Effect of graphite perfection on tritium retention as a function of neutron irradiation. Reproduced from Causey, R. A. Phys. Scripta 1996, T64. 610 Carbon as a Fusion Plasma-Facing Material from the earlier work. This was reported by Atsumi105 and clearly shown in the work of Causey108 (Figure 31). Causey irradiated high thermal conduc- tivity MFC-1 unidirectional composite and FMI-222 3D composite at �150 �C (a particularly damaging irradiation temperature regime), to a range of dis- placement doses up to 1 dpa. As seen in Figure 31, tritium retention is more than one order of 20.04.0 IG–430U Trap 2 (crystallite edge surface) Trap 1 (interstitial cluster loop) 3.0 2.0 H yd ro ge n re te nt io n (� 1 0− 3 H /C ) 1.0 0.006 dpa 0.011 dpa 0.017 dpa 0.047 dpa 0.23 dpa 0.65 dpa 0 Unirradiated 4.7 � 1022 9.4 � 1022 1.4 � 1023 1.9 � 1024 5.4 � 10243.9 � 1023 15.0 10.0 5.0 0 Figure 32 Loading of defects in irradiated graphite with hydrogen. Reproduced from Atsumi, H.; Muhaimin, A.; Tanabe, T.; Shikama, T. J. Nucl. Mater. 2009, 386–388, 379–382. Carbon as a Fusion Plasma-Facing Material 611 magnitude less than that expected from the earlier work on GraohNOL-N3M.109 In some more recent work by Atsumi,110 different neutron-irradiated graphites were irradiated and the amount of hydrogen that could be entrained in the crystal was measured. Specifically, Atsumi110 irra- diated three isomolded grades of graphite (IG-110, IG-430U, and IG-880U) in the JMTR reactor below 200 �C to various fluences and then baked the sam- ples in an atmosphere of hydrogen. Figure 32 shows the relative abundance of hydrogen that can be loaded into the crystallite edge and interstitial cluster loop- type defects (Type 1 and Type 2 defects of Figure 28) of the IG-430U graphite. Clearly, both defects are produced during irradiation and are accessible to postirradiation loading of hydrogen. In the same work, Atsumi carried out a series of preloading anneal- ing of samples, which suggested that the edge-type defects would be preferentially annihilated. In this context, it is important to note that all work on hydrogen or tritium retention in irradiated graphite has followed the approach of irradiating the material at a relatively low temperature and then loading and unloading and measuring the released hydrogen from the sample at a comparatively higher tem- perature. This may be of significance in that, as inferred in the work of Atsumi and others,105,106,108 the relative crystalline perfection (amount of intrin- sic defects) is strongly related to hydrogen retention. As discussed in Section 4.18.3 and Chapter 4.10, Radiation Effects in Graphite, irradiation at low temperature may result in a significantly different microstructure (an abundance of simple interstitial and vacancy clusters) as compared to the more fusion relevant irradiation temperatures (formation of more perfect interstitial discs and collapsing of vacancy complexes). Moreover, the postirradiation annealing of a low-temperature-irradiated microstructure will likely not produce representative microstructures of irradiation at more relevant higher temperatures. For this reason, data generated to date should be consid- ered as a guide for the trends likely to occur rather than as quantitative information on the actual tritium retention that will occur in fusion devices. Moreover, they are likely overly conservative. 4.18.6 HHF Component Technology 4.18.6.1 Joining of CFC to Heat Sink CFCs, bonded either mechanically or otherwise to a metallic structure, are being used in most of the major fusion devices, including ITER,105,106 Tore Supra,111 Wendelstein 7-X,112 TFTR in United States,113 JT 60U, and JT60SU in Japan.110–113 Silicon-doped carbon is used as the first wall material for the Chinese reactor HT-7.114 See Table 3 for a description of ITER candidate composites (INOX Sepcarb NS-31, Sepcarb NB-31, and Dunlop Concept C1, for example). As mentioned, CFCs will be used in ITER (see Table 5 for performance specifications), and specifically a 3D CFC for the divertor in the initial phases of the ITER project. The diverter, which is among the most technically challenging ITER com- ponents, is located at the bottom of the plasma cham- ber where the CFCs (and tungsten) are bonded to a copper alloy (CuCrZr). CFCs have been selected Table 5 Operational conditions for CFC joined to copper alloy in the lower vertical target, component subjected to glancing incidence of heat flux Heat flux (MWm�2) Number of cycles Damage (dpa) Tmax of the C/C – CuCrZr joint (�C) Steady state Transient 10/20a 3� 103 0.2 300 400 a20MWm�2 transient events, duration 10 s, number 10% of normal shots. Reproduced from Ferraris, M., et al. In Ceramic Integration and Joining Technologies: From Macro to Nanoscale, 1st ed.; Singh, M., Ohji, T., Asthana, R., Mathur, S., Eds.; Wiley, 2011; Chapter 3, © 2010 The American Ceramic Society. (a) (b) Figure 33 (a) Flat tile design. Reproduced from Ferraris, M., et al. In Ceramic Integration and Joining Technologies: From Macro to Nanoscale, 1st ed.; Singh, M., Ohji, T., Asthana, R., Mathur, S., Eds.; Wiley, 2011; Chapter 3, © 2010 The American Ceramic Society. (b) Monoblock design. Courtesy of J. Linke, FZJ, Germany. 612 Carbon as a Fusion Plasma-Facing Material for the lower part of the vertical targets for the initial phase of ITER operation (without tritium), and they must be able to resist ‘steady state’ heat flux up to 10MWm�2 for at least three thousand 400 s pulses and up to 20MWm�2 in transient events. As discussed earlier, the use of graphite-based materials (particularly CFCs) in a divertor is believed to be an advantage for the first phase of ITER operations because CFCs have good demonstrated performance in the currently operational plants (e.g., Tore Supra). Their primary competitor, tung- sten, also suffers from macroscopic cracking, melting, and possible melt layer loss, thus making the poten- tial damage to the divertor components more serious than that of CFC. As conditions of additional heating and off-normal events will be very likely during initial operation of ITER, CFC is the present refer- ence design solution for the lower part of the vertical targets for the ITER initial phase, when tritium retention-related issues are not relevant. In the ITER design, CFC must be joined to the copper alloy CuCrZr-IG (ITER Grade)116–119 in order to transfer the heat loads. Two different designs being considered for this component are the CFC-bonded flat-tile (Figure 33(a)) and monoblock (Figure 33(b)). In order to join the mating surface of the CFC to the copper alloy heat sink, a pure copper interlayer (1–2mm thick, oxygen free high conduc- tivity copper, 99.95%, CTE: 15.4 at RT, up 20.6 at 700 �C and up to 21.6 at 800 �C) is used to relieve, by plastic deformation, the thermal expansion derived between the CFC and copper alloy heat sink. The CTE of CFC can be found in Table 3, with that of the copper alloy stresses being (16–19)� 10�6 K�1 at 700 �C).116 This joining design is shown schematically in Figure 34. Some alternatives to a pure copper interlayer have been proposed within the EU project ExtreMat. For example, an Mo interlayer (1–2mm), a Cu/W fiber interlayer, and CFC monoblocks with a Cu/W fiber interlayer (Figure 35) have been prepared and tested up to 10MWm�2. Results (unpublished) indicate that this method is less promising compared to the use of a pure Cu thin layer. Ti-doped CFC have also been prepared and tested for flat tiles.117,118 Among the several possible options, the flat tile and monoblock configurations (see Figure 33(a) and 33(b)) with a pure copper interlayer have yielded the most promising results. In particular, the monoblock gives a more robust solution in com- parison with the flat tile for the vertical target and it is now considered as ITER reference geometry.119,120 The monoblock design requires drilled blocks of CFC into which a CuCrZr tube is inserted and joined; also necessary in the monoblock is a pure copper interlayer between the CFC and the copper alloy to relax interface stresses, which have been (modeled and) measured at �45�, �90�, �135�. If 0� is considered to be the flux direction,121,122 the monoblock is preferred to the flat tile design. This design is also much easier to manufacture because of its better heat flux performances and because of its intrinsic ability to attach even in the presence of cracks at the interface of CFC and CuCrZr, caused by its preparation process. Methods to join CFC to CuCrZr derive from techniques developed to join carbon-based compo- nents to metals. Brazing is a well-known joining process recommended for joining dissimilar Flat tile CFC Pure copper interlayer CuCrZr (a) (b) CuCrZr Monoblock CFC Pure copper interlayer Figure 34 Carbon fiber composite joined to CuCrZr in a (a) flat tile and (b) monoblock configuration (drawings show test components and not full-scale components). Reproduced from Ferraris, M., et al. In Ceramic Integration and Joining Technologies: From Macro to Nanoscale, 1st ed.; Singh, M., Ohji, T., Asthana, R., Mathur, S., Eds.; Wiley, 2011; Chapter 3, © 2010 The American Ceramic Society. 1500mm Figure 35 Cross-section of a carbon fiber composite monoblock with a Cu/W fiber interlayer. Courtesy of EU-Extremat Project and Pintsuk, G. JFZ, Germany. Carbon as a Fusion Plasma-Facing Material 613 materials. If properly done, it results in good mechan- ical strength, high fatigue performance, and minimal thermal resistance at the interface.123–125 In the case of joining of CFC for nuclear applications, some additional restrictions must be considered. A primary consideration is that the joining materials should be low activation materials (LAMs), even if the volume occupied by these materials is negligible in respect of the total volume of the structural materials in the reactor.126 Also important is that the materials be irradiation-stable at the anticipated conditions. Furthermore, the use of pressureless joining techniques is preferred as the parts to be joined are relatively large. Some joining materials are not allowed, for example, elements with high vapor pres- sure (e.g., zinc or cadmium), or those giving dangerous transmutation reactions. The ITER project mandates thermodynamic and mechanical stability up to at least 800 �C under vacuum for the joint, in order to satisfy requirements of Table 5; the joint must survive the thermal, mechanical, and neutron loads faced by the component, and it is expected to operate in a cyclic mode with an acceptable reliability and lifetime. The joining technology must also be compatible with the overall component manufacture process and in particular with the preservation of the thermomechanical properties of the precipitation- hardened CuCrZr alloy.127 Wettability is a key factor in the joining of CFC to the heat sink. It is not within the scope of this chapter to review this subject, and an overview can be found elsewhere.124,128,129 However, to restate the most salient point, the Cu interlayer cannot be obtained by directly casting copper on the CFC surface, because Cudoes not wetCFC at all, the contact angle ofmolten copper on carbon substrate being about 140�.128,129 The poor wettability of CFC is related to the non- metallic character of its bonding, whereas the bonding electrons in copper are delocalized.123,124,128 Copper can be directly cast on CFC when the CFC surface is modified to form carbides by a solid state chemical reaction between the composite surface and elements such as Si, Al, Ti, Zr, Cr, Mo, Tungsten (upper part) Austenitic steel CFC (lower part) Copper/steel tube joint Blocks 1 Blocks 32 1000 m m Blocks 33 Blocks 64 Plansee A B C D Figure 36 Vertical target full-scale prototype manufactured by Plansee (high heat flux units) and Ansaldo Ricerche (support structure and integration). Reproduced from Missirlian, M., et al. J. Nucl. Mater. 2007, 367–370(2), 1330–1336. 614 Carbon as a Fusion Plasma-Facing Material orW; some metals can form carbides with ‘metal-like’ behavior128 and are usually well wetted by molten metals. Several patents refer to joints between carbon-based materials and copper130–135: for exam- ple, pure Cr and Ti react with C to form carbides. Cr and Ti wet carbon-modified surfaces very well with a contact angle of 35–40� at 1775 �C and of 50–60� at 1740 �C, in Ar, respectively.128 A joining technique based on CFC surface modi- fication is the Active Metal Casting (AMC®) tech- nique originally developed in the eighties by the Austrian company Plansee for nonnuclear purposes. ‘Active’ indicates the activation of the CFC surface to allow it to be wetted by Cu. Physical vapor deposition (PVD) or chemical vapor deposition (CVD) of Ti coating on the CFC surface is followed by a high temperature treatment to form TiC, which improves the wettability of CFC by molten copper. Active Metal Casting consists of casting a pure Cu layer onto a laser-textured and TiC-modified CFC surface.136,137 The laser texturing enhances Cu infiltration into the CFC, and the TiC-modified CFC surface improves the wetting. The special laser treatment of the CFC surface produces a large number of closed conical holes (diameter �50–500 mm, depth 100–750 mm), thus increasing the joined area and providing addi- tional crack growth resistance. Due to the open poros- ity of the TiC-modified CFC and laser machining, the cast Cu penetrates into the CFC up to 2mm. An example of a full-scale component produced by Plansee, Austria, is shown in Figure 36. AMC® was successfully applied for both flat tile and monoblock geometry. However, AMC® technol- ogy requires laser machining of CFC surfaces, which might not be economically attractive for large-scale production. Laser-induced stresses in the joined area and cracks induced during the joining process have recently been measured and modeled.121 However, Plansee has recently improved AMC® by using silicon and titanium to modify the CFC surface (TiSi-AMC) (Figure 37(c)).130 Ansaldo Ricerche- Genova, Italy, has proposed a joining technique based on a Cu–Ti-based (Cu ABA) commercial alloy, reinforced by 2D randomly oriented carbon fibers uniformly distributed in the brazing alloy. The joining is carried out at approximately 1000 �C. The Ti reacts with carbon to form a thin TiC layer that promotes wetting.131 Carbon fibers are expected to mitigate thermal expansion mismatch between CFC and the braze (Figure 38) and to react with titanium in the brazing alloy, resulting in beneficial thermal fatigue strength of the joint. This technique was successfully tested on CFC NB31-Cu flat-tile and monoblock joints. Several other solutions have been investigated to modify CFC surface, for example, by TiN or TiC, within the EU project ExtreMat.118 A method has been proposed based on the CFC surface modification by reaction of Cr, Mo, and W. Both Cr and Mo have been extensively used as active elements in brazing alloys for copper active brazing and in patents referring to nonnu- clear applications.130–135,138,139 As example W, Mo, and Cr powders were deposited on CFC (CFC NB31, Snecma Propulsion Solide, France) by the slurry technique: details related to the process can be found elsewhere.132–134 Cr-carbide-modified CFC appears to have yielded the best results (Figure 39). A 15-mm-thick carbide (Cr23C6, Cr7C3) layer has been identified by XRD on CFC; the CTE of the carbides lies between that of CFC and copper (reported above) (CTE of Cr7C3 is 10� 10�6K�1).135 A commercial brazing alloy Gemco® (87.75 wt% Cu, 12 wt% Ge and 0.25 wt% Ni; Wesgo Metals) has been used to braze Cr-modified CFC to Cu and CuCrZr (a) (b) Cu interlayer AMC CFC TiCTiC (c) 0.2 mm 0.25 mm 0.5 mm CFC100μm TiC + SiC Cu OFHC Figure 37 (a and b) Laser structuring of carbon fiber composite (CFC) in Active Metal Casting (AMCW) process and cross-section of the AMCW CFC–Cu joint. Courtesy of Chevet,G. Ph.D. Thesis 2010, University of Bordeaux, France. 150mm CFC Braze Cu Figure 38 Carbon fiber composite/Cu active brazing with Cu–Ti alloy and dispersed carbon fibers. Reproduced from Bisio, M., et al., Fusion Eng. Des. 2005, 75–79, 277–283. 38 mm 74 mm (a) (b) 31 mm Cr carbide Cu 10μm 20μm C/C Figure 39 Carbon fiber composite–Cu active brazing with Cu–Ti alloy and dispersed carbon fibers in the braze (b). Reproduced from Schedler, B.; Huber, T.; Eidenberger, E.; Scheu, C.; Pippan, R.; Clemens, H. Fusion Eng. Des. 2007, 82(15–24), 1786–1792. Reproduced from Ferraris, M., et al. In Ceramic Integration and Joining Technologies: From Macro to Nanoscale, 1st ed.; Singh, M., Ohji, T., Asthana, R., Mathur, S., Eds.; Wiley, 2011; Chapter 3, © 2010 The American Ceramic Society. Carbon as a Fusion Plasma-Facing Material 615 to CuCrZr in a single step process,140 which is an advantage in comparison with other joining technol- ogies that require two steps: first joining CFC to Cu, and then CFC–Cu to CuCrZr. Flat-tile (a) and monoblock (b) mock-ups have been obtained by this technique and tested (Figure 40). ENEA, Italy, has manufactured several actively cooled mock-ups of flat tile and monoblock type, by using different technologies; a new process (pat- ented) for the production of monoblocks is based on prebrazed casting and hot radial pressing (PBCþHRP) (Figure 41). The CFC surface modification is obtained by a titanium–copper–nickel commercial brazing alloy, which is followed by a Cu casting, then a radial diffusion bonding between the cooling tube and the CFC by pressurizing only the internal tube and keeping the joining zone in vacuum at the required bonding temperature.120,141,142 Complete manufacture and testing of this vertical target medium-scale mock-up (Figure 41) can be considered as a success for both PBC and HRP pro- cesses, which can be an alternative to current techniques. Several joining techniques are based on active braz- ing, which do not require any manner of CFC surface (a) (b) CuCrZr CuCrZr Cu Cu C/C Cu CFC 200mm Interface Chromium carbide Brazing alloy Brazing alloy Figure 40 Optical micrograph of the cross-section of a Cr-carbide modified carbon fiber composite (CFC)–Cu joint, (b) Cr-carbide modified CFC–Cu–CuCrZr mock-up. Reproduced from Ferraris, M., et al. In Ceramic Integration and Joining Technologies: From Macro to Nanoscale, 1st ed.; Singh, M., Ohji, T., Asthana, R., Mathur, S., Eds.; Wiley, 2011; Chapter 3, © 2010 The American Ceramic Society. Figure 41 ENEA-manufactured actively cooled monoblocks. 616 Carbon as a Fusion Plasma-Facing Material modification. In this case, the active brazing alloys for CFC include elements such as Ti, Zr, Cr, and Si, which allow CFC wettability by the molten brazing alloy. A drawback of active brazing can be that active elements may form brittle intermetallics or com- pounds of low melting point. In one study,132 a TiCuNi brazing alloy produced by Wesgo Metals in the form of sheets (70Ti–15Cu–15Ni) was used to join CFC and silicon-doped CFC to pure copper, but the presence of Ni and Ti brittle intermetallics at the joint interface had a detrimental effect on thermal fatigue resistance tests of the joined component. In Japan, a brazing technology was developed for ITER143 to join CFC to Cu by a NiCrP brazing alloy, followed by the joining of the CFC–Cu to the CuCrZr by low temperature HIP (500 �C). Recently, Carbon as a Fusion Plasma-Facing Material 617 a new brazing process was developed, based on the NiCuMn alloy, after metallization of the CFC surface. Several other brazing alloys have been devel- oped for CFC–Cu joints: Ag-based (63Ag–35Cu–2Ti, 59Ag–27Cu–13In–1Ti), or Cu-based brazing alloys (Cu–3Si–2Al–2Ti; Cu–Mn; Cu–Ti). Ag was discarded in view of nuclear transmutation-related issues.144 4.18.6.2 Evaluation of HHF Joint Reliable mechanical tests on CFCs joined to heat sinks are still an issue. As ASTM tests to measure the shear strength of CFCs joined to metals are not available, several laboratories have independently developed tests for CFC to metal joints, making interlab comparison of results almost impossible. Joints obtained by AMC® have been extensively tested,146 in particular for shear strength, with data ranging from 20 to 60MPa for prepared samples. At 600 �C, the shear strength dramatically decreases to �20MPa. The shear strength and tensile strength of the improved AMC® (TiSi-AMC) joint are in the range of 54–73MPa and of 39–64MPa, respectively.130,146 Monoblocks obtained by AMC® have been measured after HHF tests147: apparent shear strength has been measured in the range of 30–60MPa. Some cracks have been found at �45�, �90�, and �135�, considering 0� as the flux direction, leading to the de- tachment of CFC from the Cu layer before testing.148 The Cr-modified CFC–Cu joined samples148–152 have beenmeasured by single-lap test (adapted fromASTM C 1292, C 1425) and off-set single lap test (adapted fromASTMD905) at room temperature. Independent of the CFC surface machining and different casting process, results obtained133,138 for Cr-modified CFC on more than 50 samples yielded average values of apparent shear strength ranging from 26 to 32MPa. The average shear strength is in any case higher than the interlaminar shear strength of the CFC (15MPa). The shear strength of the CFC–Cu joints (flat-tile geometry) obtained by using a commercial Gemco® brazing alloy to braze CFC to pure copper was 34� 4MPa, measured by single lap tests at room temperature. This is comparable to the values obtained by other joining processes and higher than the intrinsic CFC shear strength.140 Mechanical tests on the monoblock braze require specific designs: some of them are adapted by ASTM D 4562-01 ‘‘Standard Test Method for Shear Strength of Adhesives Using Pin-and-Collar Speci- men’’, as in the compression test used by Plansee AG, but the joint is not stressed in uniform pure shear state. Reliable nondestructive tests (NDTs) are also extremely important for nuclear fusion components, especially for high heat flux PFCs. NDTs on CFC–Cu joints are complex because of the different response of CFC and copper to the physical excitations used to test the component. The aim of NDT is to identify and localize defects in the joined components before submitting them to high heat flux tests or actual appli- cation. It is also important to identify the maximum acceptable defect size, as a function of its position, defined as the largest defect that is stable under spe- cific loads in the fusion device.139,149 Several techni- ques are used for NDT150: X-ray microradiography, X-ray microtomography, ultrasonic inspection,151,153 lock-in thermography,152 and transient infrared ther- mography (SATIR). SATIR (Figure 42) is the French acronym for infra red acquisition and data proces- sing device: it is a dedicated facility developed in Cadarache-France at CEA. SATIR consists of record- ing the surface temperature evolution of the compo- nent with an infrared device during the circulation of hot (�95 �C) and cold (�5 �C) water through the cooling channel of the component. The transient ther- mal response is compared to a ‘defect-free’ component; defects such as debonding of CFC tiles from heat sinks are detected by a slower temperature surface response (Figure 42). Ultrasonic inspection has been applied to the flat tile and the monoblock design. Defects on joints between materials having very different acoustic impedance (e.g., copper and CFC) result in the gen- eration of a high reflected echo, making defect detectability more difficult.141 Lock-in thermogra- phy consists of applying a series of heat flashes on the CFC. The main advantage of this technique is that there is no need for an active cooling of the component and it can also be used as an inspection method during the manufacturing process.142,152 Several nondestructive tests have been performed on the Cr-modified CFC–Cu joined samples150 not only to test performance, but also to verify and compare the reliability of these tests on a CFC–Cu interface. However, the present conclusion is that nondestructive tests of joints should be validated by destructive tests such as morphological evidence of the detected defect and mechanical testing. 4.18.7 Summary and Conclusions Carbon and graphite materials have enjoyed consid- erable success as PFMs in current tokamaks because 5D 6D 7A Block number D Tr ef (� C ) Unit A Unit D 8A VTFS: CFC (straight) part (E = 0.9) Reference block per analyzed zone SATIR-2 (final examination) FE200 (final screening) 14D 15D3A 1A 2A14.00 (a) (b) 12.00 10.00 8.00 6.00 4.00 2.00 2 3 4 5 6 7 8 9 10 11 12 13 14 15 161 0.00 −2.00 −4.00 −6.00 B1–B6 D A 537.3 759.5 967.2 B7–B11 B12–B16 Figure 42 Scanning electron microscopy of the cross-section of CFC–Cu–CuCrZr sample brazed by Gemco alloy in a single-step process as in Salvo et al.155 (a) Flat-tile configuration and (b) monoblock configuration. 618 Carbon as a Fusion Plasma-Facing Material of their low atomic number, high thermal shock resis- tance, and favorable properties. However, their use is not without significant issues, and their application in next-generation fusion energy devices is by no means certain. Significant among the issues related to carbon and graphite PFMs are neutron irradiation damage, which causes significant dimensional change and degrades the thermal conductivity resulting in increased PFC surface temperatures; physical sputtering, chemical erosion, and radiation enhanced sublimation, which cause surface material loss to the plasma and redeposition of carbon with tritium; and tritium inventory, which constitutes both a safety problem and an economic impediment to the use of graphite. Joining of CFC to heat sinks has witnessed significant development in the past few decades, which has resulted in good performing designs for near-term test machines such as ITER. The high heat loads and surface temperatures that result after plasma disruptions are also problematic for carbons. 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All rights reserved. 4.19.1 Introduction 623 4.19.2 Background 624 4.19.2.1 Synopsis of PWIs in Tokamaks 624 4.19.2.2 Brief History of Plasma-Facing Materials in Fusion Devices 626 4.19.2.3 Experience with Beryllium in Tokamaks 628 4.19.3 Beryllium PWI Relevant Properties 629 4.19.3.1 Beryllium Erosion Properties 629 4.19.3.1.1 Physical sputtering of beryllium 629 4.19.3.1.2 Mixed-material erosion 631 4.19.3.1.3 Chemically assisted sputtering of beryllium 632 4.19.3.1.4 Enhanced erosion at elevated temperatures 632 4.19.3.2 Hydrogen Retention and Release Characteristics 633 4.19.3.2.1 Implantation 633 4.19.3.2.2 Beryllium codeposition 635 4.19.3.3 Mixed-Material Effects 637 4.19.3.3.1 Be–C phenomena 637 4.19.3.3.2 Be–W alloying 637 4.19.4 Main Physical and Mechanical Properties 638 4.19.4.1 General Considerations 638 4.19.4.1.1 Physical properties 639 4.19.4.1.2 Mechanical properties 639 4.19.4.2 Selection of Beryllium Grades for ITER Applications 640 4.19.4.3 Considerations on Plasma-Sprayed Beryllium 642 4.19.4.4 Neutron-Irradiation Effects 643 4.19.4.4.1 Thermal conductivity 643 4.19.4.4.2 Swelling 643 4.19.4.4.3 Mechanical properties 644 4.19.4.4.4 Thermal shock effects 644 4.19.4.4.5 Bulk tritium retention 644 4.19.5 Fabrication Issues 644 4.19.5.1 Joining Technologies and High Heat Flux Durability of the Be/Cu Joints 644 4.19.5.1.1 Be/Cu alloy joining technology 645 4.19.5.1.2 High heat flux durability of unirradiated Be/Cu joints 646 4.19.5.2 Thermal Tests on Neutron-Irradiated Joints 648 4.19.6 Tokamak PFC Design Issues and Predictions of Effects in ITER During Operation 650 4.19.6.1 PFC Design Considerations 650 4.19.6.1.1 Design of the beryllium ITER-like wall at JET 650 4.19.6.1.2 Design of the beryllium ITER wall 651 4.19.6.2 Predictions of Effects on the ITER Beryllium Wall During Operation 653 621 622 Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices 4.19.6.2.1 Safety issues in ITER 653 4.19.6.2.2 Erosion/damage of the ITER Be wall 654 4.19.6.3 Prospect of Using Beryllium in Beyond-ITER Fusion Reactors 659 4.19.7 Concluding Remarks 659 References 662 Abbreviations Alcator C-Mod The name Alcator was given to a class of tokamaks designed and built at the Massachusetts Institute of Technology; these machines are distinguished by high magnetic fields with relatively small diameters. The high magnetic field helps create plasmas with relatively high current and particle densities. The present incarnation is Alcator C-Mod ANFIBE Computer code for ANalysis of Fusion Irradiated BEryllium ASDEX- Upgrade Axially Symmetric Divertor Experiment. The original ASDEX, located in Garching, Germany, and decommissioned in about 1990, would qualify today as a medium sized tokamak. It was designed for the study of impurities and their control by a magnetic divertor. Its successor, ASDEX-Upgrade (a completely new machine, not really an ‘upgrade’), is larger and more flexible. ATC Adiabatic Toroidal Compressor CFC Carbon-fiber composite CIP Cold isostatic pressing CP Cold pressing DIII-D A medium-sized tokamak, but the largest tokamak still operational in the United States. Operated by General Atomics in San Diego DIMES Divertor Material Evaluation Studies, a retractable probe that allows the insertion and retraction of test material samples to the DIII- D divertor floor, for example, for erosion/deposition studies. DS-Cu Dispersion-strengthened copper EAST Experimental advanced superconducting tokamak – an experimental superconducting tokamak magnetic fusion energy reactor in Hefei, the capital city of Anhui Province, in eastern China ELMs Edge localized modes FISPACT Inventory code included in the European Activation System FZJ Forschungszentrum Juelich, Germany HIP Hot isostatic pressing INEEL Idaho National Engineering and Environmental Laboratory. Now Idaho National Laboratory (INL) ISX Impurity study experiment (ISX-A and ISX-B where two tokamaks operated at Oak Ridge National Laboratory) ITER ITER, the world’s largest tokamak experimental facility being constructed in the South of France to demonstrate the scientific and technical feasibility of fusion power. The project is being built on the basis of an international collaboration between the European Union, China, India, Japan, Russia, South Korea, and the United States. The international treaty was signed in November 2006 and the central organization established in Cadarache. Most of the components will be provided in kind by agencies set up for this purpose in the seven partners JET Joint European Torus – a large tokamak located at the Culham Laboratory in Oxfordshire, England, jointly owned by the European Community. First device to achieve >1W of fusion power, in 1991, and the machine that has most closely approached Q¼ 1 for DT operation (Q¼ 0:95 in 1997) Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices 623 JUDITH Juelich Divertor Test Facility in Hot Cells KSTAR Korea Superconducting Tokamak Advanced Reactor – a long-pulse, superconducting tokamak built in South Korea to explore advanced tokamak regimes under steady state conditions LANL Los Alamos National Laboratory LCFS Last closed flux surface MAR ITERMaterials Assessment Report MIT Massachusetts Institute of Technology MPH ITER Materials Properties Handbook NBI Neutral beam injection NRA Nuclear reaction analysis NRI Nuclear Research Institute in the Czech Republic PDX Poloidal divertor experiment PFCs Plasma-facing components PISCES Plasma Interaction with Surface Components Experimental Station. It is a plasma simulator located at the University of California San Diego in the United States (originally at University of California, Los Angeles) that is used to test materials and measure sputtering, retention, etc. expected in tokamaks PLT Princeton Large Torus PWIs Plasma–wall interactions RES Radiation enhanced sublimation RMP Resonance magnetic perturbation SNL Sandia National Laboratory SOL Scrape-off-layer ST Symmetric tokamak (in this chapter) STEMET 1108 Brazing alloy: Cu–Sn–In–Ni TFR Torus Fontenay-aux-Roses TPE Tritium plasma experiment TRIM Transport of ion in matter code UCSD University of California, San Diego UNITOR One of the first small tokamaks where beryllium was used UTIAS University of Toronto Institute for Aerospace Studies VDE Vertical displacement event VHP Vacuum hot pressing 4.19.1 Introduction Beryllium, once called ‘the wonder metal of the future,’ 1 is a low-density metal that gained early prominence as a neutron reflector in weapons and fission research reactors. It then found a wide range of applications in the automotive, aerospace, defense, medical, and electronic industries. Also, because of its unique physical properties, and especially favorable plasma compatibility, it was considered and used in the past for protection of internal components in various magnetic fusion devices (e.g., UNITOR, ISX-B, JET). Most important future (near-term) applications in this field include (1) the installation of a completely new beryllium wall in the JET tokamak, which has been completed by mid of 2011 and consists of �1700 solid Be tiles machined from 4 t of beryllium; and (2) ITER, the world’s largest experimental facility to demon- strate the scientific and technical feasibility of fusion power, which is being built in Cadarache in the South of France. ITER will use beryllium to clad the first wall (�700m2 for a total weight of about 12 t of Be). Although beryllium has been considered for other applications in fusion (e.g., as neutron multiplier in the design of some types of thermonuclear breeding blankets of future fusion reactors and for hohlraums in inertial confinement fusion), this chapter will be limited to discussing the use of beryllium as a plasma-facing material in magnetic confinement devices, and in particular in the design, research, and development work currently underway for the JET and the ITER tokamaks. Considerations related to health and safety procedures for the use of beryllium relevant for construction and operation in tokamaks are not discussed here. Designing the interface between a thermonuclear plasma and the surrounding solid material environ- ment has been arguably one of the greatest technical challenges of ITER and will continue to be a chal- lenge for the development of future fusion power reactors. The interaction between the edge plasma and the surrounding surfaces profoundly influences conditions in the core plasma and can damage the surrounding material structures and lead to long machine downtimes for repair. Robust solutions for issues of plasma power handling and plasma–wall interactions (PWIs) are required for the realization of a commercially attractive fusion reactor. A mix of several plasma-facing materials is currently pro- posed in ITER to optimize the requirements of areas with different power and particle flux characteristics 624 Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices (i.e., Be for the first wall, carbon-fiber composite (CFC) for the divertor strike point tiles, and W elsewhere in the divertor). Inevitably, this is expected to lead to cross-material contamination and the for- mation of material mixtures, whose behavior is still uncertain and requires further investigation. The use of beryllium for plasma-facing- component (PFC) applications has been the subject of many reviews during the last two decades (see, e.g., Wilson et al.2 and Raffray et al.3 and references therein). Much of this fusion-related work has been summarized in a series of topical workshops on beryl- lium that were held in the past, bringing together leading researchers in the field of beryllium tech- nology and disseminating information on recent progress in the field.4 Comprehensive reviews have also appeared recently in specialized journals5,6 con- taining state-of-the-art information on a number of topics such as manufacturing and development of coat- ing techniques, component design, erosion/deposition, tritium retention, material mixing and compatibility problems, safety of beryllium handling, etc. This chapter reviews the properties of beryllium that are of primary relevance for plasma protection applications in near-term magnetic fusion devices (i.e., PWIs, thermal and mechanical properties, fab- ricability and ease of joining, chemical reactivity, etc.) together with the available knowledge on perfor- mance and operation in existing fusion machines. Special attention is given to beryllium’s erosion and deposition, the formation of mixed materials, and the hydrogen retention and release characteristics that play an important role in plasma performance, com- ponent lifetime, and operational safety. The status of the available techniques presently considered for joining the beryllium armor to the heat sink material of Cu alloys for the fabrication of beryllium-clad actively cooled components for the ITER first wall is briefly discussed together with the results of the performance and durability heat flux tests conducted in the framework of the ITER first-wall qualification programme. The effects of neutron irradiation on the degradation of the properties of beryllium itself and of the joints are also briefly analyzed. This chapter is organized as follows. Section 4.19.2 provides some background information for the reader and briefly reviews (1) the problem of PWIs in tokamaks; (2) the history of plasma-facing materials in fusion devices and the rationale for choosing beryllium as the material for the first wall of JETand ITER; and (3) the experience with the use of beryllium in tokamaks to date. Section 4.19.3 describes in detail the beryllium PWI-relevant prop- erties such as erosion/deposition, hydrogen retention and release, and chemical effects such as material mixing, all of which influence the selection of beryl- lium as armor material for PFCs. Section 4.19.4 briefly reviews a limited number of selected phy- sical and mechanical properties of relevance for the fabrication of heat exhaust components and the effects of neutron irradiation on material properties. Section 4.19.5 describes the fabrication issues and the progress of joining technology and high heat flux durability of beryllium-clad PFCs. Section 4.19.6 describes the main issues associated with the JET and ITER first-wall designs and discusses some con- straints foreseen during operation. The prospects of beryllium for applications in fusion reactors beyond ITER are briefly discussed. Finally, a summary is provided in Section 4.19.7. 4.19.2 Background 4.19.2.1 Synopsis of PWIs in Tokamaks A detailed discussion on this subject is beyond the scope of this review. The relevant PWIs are compre- hensively reviewed by Federici et al.7,8 More recent interpretations of the underlying phenomena and impact on the ITER device can be found in Roth et al.9 Here we summarize some of the main points. PWIs critically affect tokamak operation in many ways. Erosion by the plasma determines the lifetime of PFCs, and creates a source of impurities, which cool and dilute the plasma. Deposition of material onto PFCs alters their surface composition and, depending on the material used, can lead to long-term accumula- tion of large in-vessel tritium inventories. This latter phenomenon is especially exacerbated for carbon- based materials but there are still some concerns with beryllium. Retention and recycling of hydrogen from PFCs affects fuelling efficiency, plasma density control, and the density of neutral hydrogen in the plasma boundary, which impacts particle and energy transport. The primary driver for the interactions between the core plasma, edge plasma, and wall is the power generated in the plasma core that must be handled by the surrounding structures. Fusion power is obtained by the reaction of two hydrogen isotopes, deuterium (D) and tritium (T), producing an a-particle and a fast neutron. Although the kinetic energy carried by the 14.1MeV neutron escapes the plasma and could be converted in future reactors beyond ITER to thermal energy in a surrounding blanket system, the Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices 625 kinetic energy of the a-particle is deposited in the plasma. The fraction of this power that is not radiated from the plasma core as bremsstrahlung or line radi- ation (and that on average is distributed uniformly on the surrounding structures) is transported across field lines to the edge plasma and intersects the material surfaces in specific areas leading to intense power loads. The edge plasma has a strong influence on the core plasma transport processes and thereby on the energy confinement time. A schematic represen- tation of the regions of the plasma and boundary walls in a divertor tokamak is portrayed in Figure 1 taken from Federici et al.7 The outermost closed magnetic field surface forms an X-point of zero poloidal magnetic field within the vessel. This boundary is called the ‘last closed flux Magnetic flux surfaces First wall Separatrix (LCFS) Separatrix (LCFS) X-point Plasma core Baffle Vertical divertor target platePrivate flux region Separatrix strike point Pump Divertor region Edge region Scrape-off layer Figure 1 Poloidal cross-section of a tokamak plasma with a single magnetic null divertor configuration, illustrating the regions of the plasma and the boundary walls where important PWIs and atomic physics processes take place. The characteristic regions are (1) the plasma core, (2) the edge region just inside the separatrix, (3) the scrape-off-layer (SOL) plasma outside the separatrix, and (4) the divertor plasma region, which is an extension of the SOL plasma along field lines into the divertor chamber. The baffle structure is designed to prevent neutrals from leaving the divertor. In the private flux region below the X-point, the magnetic field surfaces are isolated from the rest of the plasma.Reproducedwith permission fromFederici, G.; Skinner, C. H.; Brooks, J. N.; et al. Plasma-material interactions in current tokamaks and their implications for next-step fusion reactors. Nucl. Fusion 2001, 41, 1967–2137 (review special issue), with permission from IAEA. surface’ (LCFS) or ‘separatrix.’ Magnetic field surfaces inside the LCFS are closed, confining the plasma ions. The edge region, just inside the LCFS, contains signif- icant levels of impurities not fully ionized, and perhaps neutral particles. Impurity line radiation and neutral particles transport some power from here to the wall. The remaining power enters the region outside the LCFS either by conduction or, depending on the degree to which neutrals penetrate the plasma, by convection. This region is known as the scrape-off- layer (SOL) as here power is rapidly ‘scraped off ’ by electron heat conduction along open field lines, which are diverted to intersect with target regions that are known as ‘divertors.’ Poloidal divertors have been very successful at localizing the interactions of plasma ions with the target plate material in a part of the machine geometrically distant from the main plasma where any impurities released are well screened from the main plasma and return to the target plate. The plasma density and temperature determine the flux density and energy of plasma ions striking the plasma-wetted surfaces. These, in turn, deter- mine the rate of physical sputtering, chemical sput- tering, ion implantation, and impurity generation. The interaction of the edge plasma with the sur- rounding solid material surfaces is most intense in the vicinity of the ‘strike point’ where the separatrix intersects the divertor target plate (see inset in Figure 1). In addition, the plasma conditions deter- mine where eroded material is redeposited, and to what degree codeposition of tritium occurs at the wall. The plasma power flow also determines the level of active structural cooling required. Typical plasma loads and the effects expected during normal operation and off-normal operation in ITER are summarized in Table 1. Because of the very demanding power handling requirements (predicted peak value of the heat flux in the divertor near the strike-points is >10MWm�2) and the predicted short lifetime due to sputtering erosion arising from very intense particle fluxes (�1023–1024 particlesm�2 s�1) and damage during transient events, beryllium has been excluded from use in the ITER divertor and is instead the material selected for the main chamber wall of ITER. Recent observations in present divertor tokamaks have shown that plasma fluxes to the main wall are dominated by intermittent events leading to fast plasma particle transport that reaches the PFCs along the magnetic field (see Loarte et al.10 and references therein). The quasistationary heat fluxes to the main wall are thought to be dominated by convective Table 1 Major issues associated with operation of ITER PFCs PFCs Plasma loads Candidate armor Effects Issue Divertor – strike-point regions �Radiation and particle heat CFCa Chemical erosion evaporation brittle destruction and tritium codeposition �Erosion lifetime and component replacement� Large particle fluxes �Disruptions �High tritium inventory and safety�ELM’s �Slow-high power transients Divertor – baffle region �Radiation heat W High sputtering evaporation/ melting �Plasma contamination �Disruptions Dome �Radiation heat (MARFE’sb) W High sputtering evaporation/ melting �Erosion lifetime ��100eV ions and CX neutrals �Moderate power transients First wall �Plasma contact during VDEsc Be Evaporation/melting �T retention in beryllium codeposited layers �Disruptions and runaway electrons �Chemical reactivity especially with Be dust �ELMs Start-up limiters �High start-up heat loads Be High sputtering evaporation/ melting �Erosion lifetime �Plasma contact during VDEs �Disruptions aW is also considered as an alternative. bMultifaceted asymmetric radiation from the edge (MARFE). cVertical displacement event (VDE). 626 Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices transport,11 but still remain to be clarified. Although the steady-state parallel power fluxes associated with these particle fluxes will only be of the order of several MWm�2 in the ITER QDT¼ 10 reference scenario, local overheating of exposed edges of main wall PFCs can occur because of limitations in the achievable alignment tolerances. Similarly, transient events are expected to cause significant power fluxes to reach first-wall panels in ITERalong the field lines. Edge localized modes (ELMs) deposit large amounts of energy in a short time, and in some cases in a toroidally localized fashion, which can lead to strong excursions in PFC surface temperatures. While the majority of ELM energy is deposited on divertor surfaces, a significant fraction is carried to surfaces outside the divertor. There are obvious concerns that ELMs will lead to damage of the divertor and the first wall.12 An additional concern is that even without erosion, thermal shock can lead to degradation of material thermomechanical properties, for example, loss of ductility leading to an enhanced probability of mechanical failure or spalling (erosion). Research efforts to characterize the ELMs in the SOL are described elsewhere.13–15 There are still large uncer- tainties in predicting the thermal loads of ELMs on the ITER beryllium first wall and the range of parallel energy fluxes varies from 1.0MJm�2 (controlled ELMs) to 20MJm�2 (uncontrolled ELMs).16,17 Even for controlled ELMs, such energy fluxes are likely to cause melting of up to several tens of micrometers of beryllium at the exposed edges,18 which could cause undesirable impurity influxes at every ELM.10,11 4.19.2.2 Brief History of Plasma-Facing Materials in Fusion Devices PWIs have been recognized to be a key issue in the realization of practical fusion power since the beginning of magnetic fusion research. By the time of the first tokamaks in the 1960s in the USSR and subsequently elsewhere, means of reducing the level of carbon and oxygen were being employed.19,20 These included the use of stainless steel vacuum vessels and all-metal seals, vessel baking, and discharge cleaning. Ultimately, these improvements, along with improved plasma confinement, led to the first production of relatively hot and dense plasmas in the T3 tokamak (�1 keVand�3� 1019m�3).21,22 These plasmas, while being cleaner andwith low-Z elements fully stripped in the core, still had unacceptable levels of carbon, oxy- gen, and metallic impurities. The metallic contamina- tion inevitably consisted of wall and limiter materials. Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices 627 Early in magnetic fusion research, it was recog- nized that localizing intense PWIs at some type of ‘sacrificial’ structure was desirable, if only to ensure that more fragile vacuum walls were not penetrated. This led to the birth of the ‘limiter,’ usually made to be very robust, from refractory material and posi- tioned to ensure at least several centimeters gap between the plasma edge and more delicate struc- tures like bellows, electrical breaks, vacuum walls, etc. Typical materials used for limiters in these early days included stainless steel in Adiabatic Toroidal Compressor (ATC)23 and ISX-A24 and many others, molybdenum in Alcator A25 and Torus Fontenay-aux-Roses (TFR),26 tungsten in symmetric tokamak (ST)27 and Princeton Large Torus (PLT),28 and titanium in poloidal divertor experiment (PDX).29 Poloidal divertors have been very successful at loca- lizing the interactions of plasma ions with the target plate material in a part of the machine geometrically distant from the main plasma where any impurities released are well screened from the main plasma and return to the target plate.30 By the early 1980s, it was also recognized that in addition to these functions, the divertor should make it easier to reduce the plasma temperature immediately adjacent to the ‘limiting’ sur- face, thus reducing the energies of incident ions and the physical sputtering rate. Complementary to this, high divertor plasma and neutral densities were found. The high plasma density has several beneficial effects in dispersing the incident power, while the high neutral density makes for efficient pumping. Pumping helps with plasma density control, divertor retention of impurities and, ultimately, in a reactor, helium exhaust. By the late 1970s, various tokamaks were starting to employ auxiliary heating systems, primarily neutral beam injection (NBI). Experiments with NBI on PLT resulted in the first thermonuclear class temperatures to be achieved.28,31,32 PLT at the time used tungsten limiters, and at high powers and relatively low plasma densities, very high edge plasma temperatures and power fluxes were achieved. This resulted in tungsten sputtering and subsequent core radiation from partially stripped tungsten ions. For this reason, PLT switched limiter material to nuclear grade graphite. Graphite has the advantage that eroded carbon atoms are fully stripped in the plasma core, thus reducing core radia- tion. In addition, the surface does not melt if over- heated – it simply sublimes. This move to carbon by PLT turned out to be very successful, alleviating the central radiation problem. For these reasons, carbon has tended to be the favored limiter/divertor material in magnetic fusion research ever since. By the mid-1980s, many tokamaks were operating with graphite limiters and/or divertor plates. In addition, extensive laboratory tests and simulations on graphite had begun, primarily aimed at under- standing the chemical reactivity of graphite with hydrogenic plasmas, that is, chemical erosion. Early laboratory results suggested that carbon would be eroded by hydrogenic ions with a chemical erosion yield of Y� 0.1 C/Dþ, a yield several times higher than the maximum physical sputtering yield. Another process, radiation-enhanced sublimation (RES), was discovered at elevated temperatures, which further suggested high erosion rates for carbon. Carbon’s abil- ity to trap hydrogenic species in codeposited layers was recognized. These problems, along with graphite’s poor mechanical properties in a neutron environment (which had previously been known for many years from fission research33), led to the consideration of beryllium as a plasma-facing material. This was pri- marily promoted at JET.34 A description of the opera- tion experience to date with Be in tokamak devices is provided in Section 4.19.2.3. At present, among divertor tokamaks, carbon is the dominant material only in DIII-D. Alcator C-Mod at Massachusetts Institute of Technology (MIT), USA35 uses molybdenum. ASDEX-Upgrade (Axially Symmetric Divertor Experiment) is fully clad with tungsten,36 and JET has completed in 2011 a large enhancement programme37 that includes the installa- tion of a beryllium wall and a tungsten divertor. New superconducting tokamaks, such as Korea Supercon- ducting Tokamak Advanced Reactor (KSTAR) in Korea38 and experimental advanced superconducting tokamak (EAST) inChina,39 employcarbon asmaterial for the in-vessel components, but with provisions to exchange the material later on in operation. The current selection of plasma-facing materials in ITER has been made by compromising among a series of physics and operational requirements, (1) minimum effect of impurity contamination on plasma performance and operation, (2) maximum operational flexibility at the start of operation, and (3) minimum fuel retention for operation in the DT phase. This compromise is met by a choice of three plasma-facing materials at the beginning of opera- tions (Be, C, and W). It is planned to reduce the choices to two (Be and W) before DT operations in order to avoid long-term tritium retention in car- bon codeposits during the burning plasma phase. Beryllium has been chosen for the first-wall PFCs to minimize fuel dilution caused by impurities released from these surfaces, which are expected to 628 Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices have the largest contamination efficiency.40–44 More- over, the consequences of beryllium contamination on fusion performance and plasma operations are relatively mild. This has been demonstrated by experiments in tokamaks (see Section 4.19.2.3). The main issues related to the use of beryllium in ITER are (1) the possible damage (melting) during transients such as ELMs, disruptions, and runaway electron impact, and its implications for operations and (2) the codeposition of tritium with beryllium which is eroded from the first wall and deposited at the divertor targets (and possibly also locally rede- posited into shadowed areas of the shaped ITER first wall). Both issues are part of ongoing research, the initial results of which are being taken into account in the ITER design so that the influence of these two factors on ITER operation and mission is minimized. This includes ELM control systems based on pellets and resonance magnetic perturbation (RMP) coils, disruption mitigation systems, and increased temper- ature baking of the divertor to release tritium from the beryllium codeposited layers. Carbon is selected for the high power flux area of the divertor strike points because of its compatibility with operation over a large range of plasma conditions and the absence of melting under transient loads. Both of these characteristics are considered to be essential during the initial phase of ITER exploitation in which plasma operational scenarios will require development and transient load control and mitiga- tion systems will need to be demonstrated. 4.19.2.3 Experience with Beryllium in Tokamaks Only three tokamaks have operated with beryllium as the limiter or first-wall material. The first experi- ments were performed by UNITOR,45 which were then followed by ISX-B.46 Both tokamaks investi- gated the effects of small beryllium limiters on plasma behavior (UNITOR had side limiters at two toroidal locations and ISX-B had one top limiter) in support of the more ambitious beryllium experiment in JET (see below). The motivation to use beryllium came from the problem of controlling the plasma density and impurities when graphite was used. Both UNITOR and ISX-B showed that once beryllium is evaporated from the limiter and coated over a large segment of the first wall, oxygen gettering leads to significant reduction of impurities. When the heat load on the beryllium limiter was increased to the point of evaporating beryllium, the oxygen concentration was decreased dramatically. Although the concentration of beryllium in the plasma was increased, its contribution to Zeff (the ion effective charge of the plasma Zeff provides a measure for impurity concentration) was more than compensated by the reduction of oxygen, carbon, and metal impu- rities.45 The plasma Zeff was observed to be reduced from 2.4 to near unity with beryllium. It must be noted that there was a negative aspect associated with beryl- lium operation during the ISX-B campaign. The reduction in plasma impurities was not observed until the limiter surface was partially melted causing beryllium to be evaporated and coated on the first wall. Once melting did occur, the droplets made subsequent evaporation more likely but hard to con- trol. The consequent strong reduction in plasma impurities associated with gettering then made dis- charge reproducibility hard to obtain. However, if a much larger plasma contact area is already covered with Be, one does not need to rely on limiter melting to obtain the beneficial effect of beryllium. This effect could be achieved by using large area beryllium lim- iter, or coating the inside wall with beryllium which was the approach taken by JET when it introduced beryllium in 1989. Large tokamak devices such as JET had found it very difficult to control the plasma density with graphite walls as the discharge pulse length got longer. Motivated by the frequent occurrence of a phenomenon that plagued the earlier campaigns – the so-called carbon blooms due to the overheating of poorly designed divertor tiles and the subsequent influx of carbon impurities in the plasma due to evaporation – JET decided to use beryllium as a plasma-facing material. Thin evaporated beryllium layers on the graphite walls were used initially (�100 Å average thickness per deposition) on the plasma-facing surface of the device. Subsequently, beryllium tiles were installed on the toroidal belt limiter. The main experimental results with beryllium can be summarized as follows: 1. The concentration of carbon and oxygen in the plasma were 4–7% and 0.5–2%, respectively, when graphite was used as belt limiter. With a beryllium belt limiter, the carbon content was reduced to 0.5% and oxygen became negligible, because of oxygen gettering by beryllium. During ohmically heated discharges, the concentration of beryllium remained negligible even though beryllium was the dominant impurity. Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices 629 2. While the value of Zeff was �3 using the graphite limiter and auxiliary heating power of 10MW, Zeff was �1.5 even with additional heating powers of up to 30MW with a beryllium limiter. 3. The fuel density control was greatly improved with the beryllium limiter and beryllium evapo- rated wall. Gas puffing to maintain a given plasma density was typically 10 times larger when using beryllium than graphite. Following the beryllium limiter experience, diver- tor beryllium targets were installed in JET for two configurations. An extensive set of experiments with toroidally continuous X-point divertor plates was car- ried out in JET in the period 1990–1996 to characterize beryllium from the point of view of its thermomecha- nical performance, as well as its compatibility with various plasma operation regimes.47–50 In the JET Mk I experiments, melting of the beryllium tiles was reached by increasing (in a pro- gressive way) the power flux to a restricted area of the divertor target in fuelled, medium density ELMy H-mode discharges (Pinp� 12MW). Large beryllium influxes were observed when the divertor target tem- perature reached �1300 �C. In these conditions, it became difficult to run low-density ELMy H-mode discharges (Pinp� 12MW) without fast strike point movement (to achieve lower average power load) and the discharges either had very poor performance or were disrupted. However, no substantial plasma performance degradation was observed for medium density H-modes with fixed strike point position, or if fast strike point movement was applied in low- density H-modes, despite the large scale distortion of the target surface caused by the melt layer displace- ment and splashing due to the previous �25 high power discharges48,51 (see Figure 252). This demon- strated that it was possible to use the damaged Be divertor target as the main power handling PFC if the Figure 2 Melting of the Joint European Torus Mk I beryllium target plate tiles after plasma operation. Image courtesy of EFDA-JET. average power load was decreased, either by increas- ing plasma density and radiative losses, or by strike point sweeping. The damage did not prohibit subsequent plasma operation in JET, but would seri- ously limit the lifetime of Be PFCs in long-pulse ITER-like devices. The latest results of the operation of JET with beryllium have been reviewed recently by Loarte et al.10 4.19.3 Beryllium PWI Relevant Properties This section describes the present understanding of PWIs for beryllium-containing surfaces. First, it focuses on the erosion properties of ‘clean’ beryl- lium surfaces at different temperatures. Retention of plasma fuel species in both bulk and codeposited layers of beryllium is then described. As beryllium will not be used as the exclusive plasma-facing mate- rial in future confinement devices, issues associated with mixed, beryllium-containing surfaces are also addressed. 4.19.3.1 Beryllium Erosion Properties The term erosion is used to describe a group of processes that remove material from a surface sub- jected to energetic particle bombardment. Included under the general classification of erosion are pro- cesses such as physical sputtering, chemically assisted physical sputtering, chemical sputtering, and thermally activated release from surfaces. Of these processes, only chemical sputtering, where volatile molecular species are formed on the surface, appears to be inactive in beryllium. 4.19.3.1.1 Physical sputtering of beryllium Physical sputtering results from the elastic transfer of energy from incoming projectiles to atoms on the surface of the target material. Target atoms can be sputtered when the energy they receive from the collisional cascade of the projectile exceeds the bind- ing energy of the atom to the surface. The physical sputtering rate is usually referred to as the sputtering yield, Y, which is defined as the ratio of the number of atoms lost from a surface to the number of incident energetic particles striking the surface. The lower the binding energy of surface atoms, the larger the physical sputtering yield. As physical sputtering can be approximated using a series of binary collisions within the surface, it is relatively easy to estimate Measured yield in D plasma Calculated yield (D+ ions only) Calculated yield (D+, D2 +, D3 + ions) 0.1 0.01 0.001 0.0001 20 40 60 80 Ion energy (eV) P hy si ca l s p ut te rin g yi el d 100 120 140 1600 10-5 Figure 3 Calculated sputtering yields from pure Dþ bombardment at normal incidence compared to that calculated for a (0.25, 0.47, 0.28) mix of Dþ, D2 þ, and D3 þ; also shown is the measured yield from such a plasma. 630 Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices the physical sputtering yield of given projectile-target scenarios. Monte-Carlo based simulation codes (such as transport of ions in matter (TRIM))53 have been used to generate extensive databases of sputtering yields based on incident particle angle, energy, and mass, for a variety of targets54 including beryllium. Measurement of the physical sputtering yield from a beryllium surface is complicated by the natural affinity of beryllium for oxygen. A beryllium surface will quickly form a thin, stable, passivating oxide surface layer when exposed to atmosphere. In ion beam devices, it is possible to clean any oxides from the beryllium surface before a measurement and with careful control of the residual gas pressure, make the measurements before the oxide reforms on the surface and alters the measurement.55 It has also been shown that it is possible to deplete the beryllium surface of oxide by heating the sample to tempera- tures exceeding 500 �C, where the beryllium below the oxide can diffuse through the oxide to the surface,56 thereby allowing measurements on a clean beryllium surface. The comparison between the calcu- lated sputtering yield and measurements made using mass-selected, monoenergetic ion-beams devices impinging on clean beryllium surfaces is fairly good.57 Measurements of sputtering yields in plasma devices, however, are complicated by several factors. In plasma devices, the incident ions usually have a temperature distribution and may contain different charge state ions. Each different charge state ion will be accelerated to a different energy by the elec- trostatic sheath in the vicinity of the surface. When hydrogenic plasma interacts with a surface, one must also account for a distribution of molecular ions striking the surface. In the case of a deuterium plasma, for example, the distribution of molecular ions (Dþ, D2 þ, D3 þ) must be taken into account as the incident molecule disassociates on impact with the surface and a D2 þ ion becomes equivalent to the bombardment of two deuterium particles with one-half the incident energy of the original D2 þ ion. Figure 3 shows the change to the calculated sputter- ing yield when one includes the effects of molecular ions in a plasma, compared to the calculated sputter- ing yield from pure Dþ ion bombardment. The trajectory of the incoming ions can also be altered by the presence of electrostatic and magnetic sheaths. Plasmas also contain varying amounts of impurity ions, originating either from PWIs in other locations of the device, or ionization of residual back- ground gas present in the device and these impurity ions, or simply neutral gas atoms, may interact with the surface. Finally, the incident flux from the plasma is usually so large that the surface being investigated, and its morphology, becomes altered by the incident flux and a new surface exhibiting unique character- istics may result. Nevertheless, the physical sputtering yield from beryllium surfaces exposed to plasma ion bombard- ment has been measured in several devices. Unfortu- nately, there is little consensus on the correct value of the physical sputtering yield. In JET, the largest confinement device to ever employ beryllium as a PFC sputtering yield measurements range from values far exceeding47 to values less58 than one would expect from the predictions of TRIM. In the Plasma Interaction with Surface Components Experimental Station B (PISCES-B) device, systematic experiments to measure the physical sputtering yield routinely show values less59–61 than those expected from TRIM. This difference is shown in Figure 3, where the energy dependence of the calculated yield is com- pared to experimental measurements. Another primary difference between the condi- tions in an ion beam device and those encountered in a plasma device has to do with the neutral density near the surface being investigated. In an ion beam experiment, the background pressure is kept very low Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices 631 so that the surface being probed maintains its clean properties. On the other hand, the incident flux in a plasma device is usually several orders of magnitude larger than in an ion beam device, ensuring that the surface remains clean because of the large incident flux. However, this plasma-facing surface undergoes not only energetic ion bombardment, but also bom- bardment by neutral atoms and molecules. The neutral density in plasma generators is typi- cally on the order of 1020m�3 (a few millitorr) which is necessary for breakdown of the plasma. The esti- mated neutral atom flux is approximately equal to the incident ion flux to the surface61 and it is often not possible to alter significantly this flux ratio. In the case of a beryllium surface which can form a hydride (see Section 4.19.3.1.3), the presence of adsorbed deuterium on the surface could affect the measured sputtering yield by decreasing the beryllium concen- tration at the surface and altering the binding energy of surface beryllium atoms. Some evidence of this effect may be discerned in data from JET measurements of the beryllium sputtering yield. Two sets of sputtering yield mea- surements have been reported from JET; one from beryllium divertor plate measurements and the other from beryllium limiter measurements. In the divertor region, one expects a neutral density similar to that encountered in plasma generators (1020m�3 or more) and the measured sputtering yield is lower than that predicted by TRIM calculations.58 When sputtering measurements are made on the limiter, where the neutral density is typically lower, the sputtering yield agrees with, or exceeds, the calculated value.47 Of course, other issues such as impurity layers on the divertor plate and angle of incidence questions tend to confuse the results. However, the data sets from JET are consistent with the impact of neutral absorption on the beryllium plasma-facing surface. Effects associated with plasma operation will need to be taken into account when predicting sputtering yields from different areas of confinement devices. In addition to the low-energy neutral atom flux and higher-energy charge exchange neutral flux, the impact of small impurity concentrations in the inci- dent plasma flux will also have a large impact on the expected sputtering yield. Some of the implications of the formation of a mixed-material surface are discussed in the next section and in Section 4.19.3.3. 4.19.3.1.2 Mixed-material erosion As was pointed out in the previous section, it is impor- tant to have accurate knowledge of a target’s surface composition to predict its erosion rate. A small impu- rity concentration contained within the incident plasma can drastically alter the surface composition of a target subjected to bombardment by the impure plasma. Oxygen impurities in the plasma, either from ionization of the residual gas, or due to erosion from some other surface, will readily lead to the formation of beryllium oxide on the surface of a beryllium target. Depending on the arrival rate of oxygen to the surface compared to the erosion rate of oxygen off the surface, one can end upmeasuring the sputtering rate of a clean beryllium surface or a beryllium oxide surface. Careful control of the residual gas pressure in ion beam sput- tering experiments55 has documented this effect. Unfortunately, it is not always so easy to control the impurity content of an incident plasma. In the case of a magnetic confinement device composed of groups of different plasma-facing mate- rial surfaces, erosion from a surface in one location of the device can result in the transport of impurities to other surfaces throughout the device. Mixed- material surfaces are the result. To first order, a mixed-material surface will affect the sputtering of the original surface material in two ways. The first is rather straightforward, and is true even for materials which do not form chemical bonds, in that the surface concentration of the original material is reduced thereby reducing its sputtering rate. The second effect changing the sputtering from the surface results from changes in the chemical bonding on the surface, which can either increase, or decrease the binding energy of the original material. If the chemical bonds increase the binding energy, the sputtering rate will decrease. If the bonding acts to reduce the surface binding energy, the sputtering rate will increase (assuming the change in surface concentration does not dominate this effect). A recent review of mixed materials62 provides some background information on the fundamental aspects of general mixed-material behavior. If a plasma incident on a beryllium target contains sufficient condensable, nonrecycling impurities (such as carbon), it will affect the sputtering rate of the beryllium. This effect was first referred to as ‘carbon poisoning.’ 5,9,63 A simple particle balance model has been used to adequately explain the results for formation of mixed carbon-containing layers on beryllium at low surface temperature.64 However, as the target temperature increases, additional chemical effects, such as carbide formation, have to be included in the model. An interesting change occurs when the bombard- ing species is a mixture of carbon and oxygen. 632 Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices Measurements of the chemical composition of a beryllium surface bombarded with a COþ ion beam showed almost exclusive bonding of the oxygen to the beryllium in the implantation zone.65,66 The formation of BeO on the surface left the carbon atoms easily chemically eroded. The amount of oxy- gen present in the incident particle flux plays a strong role in the final chemical state of the surface atoms and their erosion behavior. The inverse experiment, beryllium-containing plasma incident on a carbon surface, has also been investigated.67–69 In the case of beryllium impurities in the plasma, a much more accurate measurement of the impurity concentration was possible. Contrary to the carbon in beryllium experiments, a simple particle balance model could not account for the amount of beryllium remaining on the surface after the plasma exposure. Clearly, the inclusion of chemi- cal effects on the surface needs to be taken into account to interpret the results. Beryllium carbide (Be2C) was observed to form on the surface of carbon samples exposed to beryllium- containing deuterium plasma even during bombard- ment at low surface temperature. Carbide formation will also act to increase the binding of beryllium atoms to the surface and decrease the binding of carbon atoms. This effect will result in an increase in the concentration of beryllium on the surface compared to a simple particle balance equation and must be included to understand the evolution of the surface. In addition, the formation of the carbide was correlated with the decrease of carbon chemical ero- sion70 (see Section 4.19.3.3.1 for more discussion of the chemical erosion of the beryllium–carbon system). 4.19.3.1.3 Chemically assisted sputtering of beryllium The term chemically assisted physical sputtering refers to the transfer of energy from an incident particle to a molecule on the surface. The energy gained is sufficient to break any remaining bonds of the molecule to other atoms on the surface resulting in the release of the molecule, or a fragment of the molecule, from the surface. In the case of beryllium bombarded by deuterium plasma, the sputtering of beryllium deuteride was first recorded in JET71 dur- ing operation with a beryllium divertor plate. Since that time, a series of systematic investigations were performed in PISCES-B to quantify the magnitude of this erosion term.72,73 The results from PISCES-B show a surface tem- perature dependence of the sputtering rate72 of BeD molecules. The maximum in the BeD sputtering rate (at �175 �C) corresponds with the onset of thermal decomposition of BeD2 molecules 73 from a standar- dized sample of BeD2 powder. Even at the maximum loss rate, the chemical sputtering remains smaller than the physical sputtering rate of beryllium atoms from the surface over the incident energy range examined (50–100 eV). Molecular dynamics simula- tions have predicted,74 and subsequent measure- ments have validated the prediction, that chemical sputtering can dominate physical sputtering of beryl- lium as the incident deuterium ion energy decreases below 50 eV. A distinction should be made between chemical sputtering and chemically assisted physical sputtering. Chemical sputtering involves the formation and loss of volatile molecules from a surface. In the case of beryl- lium deuteride, the molecule decomposes into a deu- terium molecule and a beryllium atom before it becomes volatile, so at least to date there is no evidence for chemical sputtering of beryllium during deuterium particle bombardment. Documentation of the chemi- cally assisted physical sputtering of beryllium may be important for determining material migration patterns in confinement devices and the identification of beryl- lium deuteride molecular formation in plasma- exposed surfaces may also help explain the hydrogenic retention properties of beryllium. 4.19.3.1.4 Enhanced erosion at elevated temperatures In addition to the temperature-dependent chemical sputtering of beryllium when exposed to deuterium plasma, another temperature-dependent loss term is present in beryllium exposed to plasma bombard- ment at elevated temperature. The classical picture of the temperature dependence of erosion from chemically inert surfaces exposed to energetic parti- cle bombardment is composed of the superposition of a constant physical sputtering yield with an expo- nentially varying thermal sublimation curve. The classical picture is contradicted, however, by experi- ments that show an exponential increase in erosion at lower temperature that cannot be explained by classical thermodynamic sublimation. First observed by Nelson75 for a variety of metal surfaces, similar results have been measured for Be,76,77 W,78 and C79,80 surfaces. In the case of carbon, this mechanism has been called RES. In the case of beryllium, two explanations have been proposed and both rely on the large flux of ions incident during plasma bombardment to modify the S ol ub ili ty (H /m et al a to m P a- 1/ 2 ) 1000/T(K) 0.5 W 1 2 3 Be 1.0 10-9 10-8 10-7 1.5 Figure 4 Measured solubility of hydrogen in tungsten (dashed line87) and beryllium (solid lines 1,88 2,89 and 390). Reproduced with permission from Federici, G.; Skinner, C. H.; Brooks, J. N.; et al. Plasma-material interactions in current tokamaks and their implications for next-step fusion reactors. Nucl. Fusion 2001, 41, 1967–2137 (review special issue), with permission from IAEA. W s− 1 ) -9 1´10-8 1´10-7 Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices 633 plasma-facing material surface. In the first, the inci- dent plasma ions, in addition to creating sputtered atoms from the surface, also create a population of surface adatoms. An adatom is an atom from a lattice site on the surface that has gained sufficient energy to leave its lattice location, yet does not have sufficient energy to escape from the surface as a sputtered atom. The atom then occupies a site on top of the regular lattice sites. Because an adatom does not have the same number of adjacent atoms as those in the lattice, it is less strongly bound to the surface and can therefore sublime at a lower temperature than one associates with equilibrium thermodynamic sublima- tion. In the second explanation, incident plasma ions that have thermalized somewhere below the surface of the lattice exert a stress on the surface atoms of the target again resulting in a lower binding energy of the surface atoms to the bulk of the material. Measurements show atoms are being lost from the surface at thermal energies,77 rather than the energy associated with sputtered particles (i.e., on the order of electron volts). This seems to verify the loss mech- anism that occurs because of the thermal release of an ensemble of particles with a lower surface binding energy than that of bulk atoms of that element. Addi- tional measurements at elevated temperature have documented the variation in Be atom surface loss rate with changes of the incident flux of energetic particles.81 The larger the incident flux, the lower the onset temperature for the enhanced erosion. The implication of this enhanced loss rate at ele- vated temperature is a reduction of the permissible operating temperature of any plasma-facing material, or alternatively that the lifetime of a component operating at extreme temperature may be less than that expected based on the predictions from classical surface loss terms. 1a 1b2 3 Be 1000/T(K) D iff us iv ity (m 2 0.5 1.0 1.5 10-13 10-12 10-11 1´10-10 1´10 Figure 5 Measured diffusivity of hydrogen in tungsten (dashed line87) and beryllium (solid lines 1a&b,91 2,89 and 392). Reproduced with permission from Federici, G.; Skinner, C. H.; Brooks, J. N.; et al. Plasma-material interactions in current tokamaks and their implications for next-step fusion reactors.Nucl. Fusion 2001, 41, 1967–2137 (review special issue), with permission from IAEA. 4.19.3.2 Hydrogen Retention and Release Characteristics 4.19.3.2.1 Implantation The use of beryllium as a plasma-facing material in tokamaks has prompted many experimental studies of retention and emission of hydrogen implanted into beryllium-like metals from ion-beams or plasmas. References and discussions of these studies can be found in reviews.82–85 Here, we review those studies which are relevant to H retention in Be in a fusion plasma environment. This section is mainly excerpted from Federici et al.7 Two basic parameters for under- standing H retention are the hydrogen diffusivity and solubility. Studies of solubility and diffusivity are reviewed in Causey and Venhaus85 and Serra et al.86 Figures 487–90 and 587,91,92 show experimental values for hydrogen solubility and diffusivity in W and Be. For Be there are significant differences between results from various studies. These differences may be due to effects of traps and surface oxide layers. The presence of bulk traps tends to increase the measured values of solubility and to decrease the mea- sured values of diffusivity (see Federici et al.7), especially under conditions where the concentration 634 Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices of hydrogen in solution is smaller than the concen- tration of traps. For this reason, studies done on materials of higher purity and crystalline perfection, and at higher temperatures and with higher concen- trations of hydrogen in solution, tend to give more reliable results. The porosity and oxide inclusions present in beryllium produced by powder metallurgy are also likely to lead to inconsistent results in mea- surements of hydrogen solubility and diffusivity. In the Be experiments, the effects of traps were not characterized and may be dominant. One firm con- clusion is that the solubility of hydrogen is very low in both Be and W. Many studies have been done on the retention and emission of H implanted into materials to provide data needed to predict H retention in fusion reactor environments. Figure 6 shows the retention of 1 keV deuterium implanted into Be at 300 K versus inci- dent fluence, measured by thermal desorption.93 D retention in Be was close to 100% at low fluences but saturated at high fluences. Earlier nuclear reac- tion analysis (NRA) measurements of D retained in Be within �1 mm of the surface gave very similar results.94 This saturation behavior indicates that D implanted into Be at 300 K does not diffuse, but accumulates until it reaches a limiting concentration of �0.3–0.4 D/Be within the implantation zone. At high fluences, the implanted zone becomes porous allowing additional implanted D to escape. This Incident flu R et ai ne d fl ue nc e (D m −2 ) 100% Be W Mo, slope 3 keV D3 + (1020 D m–2s) Be, M 1019 1020 1020 1021 1022 1021 1022 2 3 2 3 2 2 2 2 3 Figure 6 Retention of 1 keV deuterium implanted into Be and desorption. Adapted and reproduced with permission from Haas 1076–1081, Federici, G.; Skinner, C. H.; Brooks, J. N.; et al. Plas implications for next-step fusion reactors.Nucl. Fusion 2001, 41, saturationmechanism is confirmed by electronmicros- copy, which shows bubbles and porosity in the implan- tation zone after high fluence H implantation.95 Saturation of retention by the same mechanism is observed for D implanted into stainless steel at 150K where the D is not mobile.96 H retention in Be increases with increasing ion energy and decreases with increasing sample temperature.84,97 The retention of 1 keV deuterium implanted into W and Mo at 300K98 is also shown for comparison in Figure 6. Figure 784 shows retention of deuterium and tri- tium as a function of incident particle fluence from a set of high fluence experiments in which Be specimens were exposed to laboratory ion-beams (Idaho National Engineering and Environmental Laboratory (INEEL), University of Toronto Institute for Aerospace Studies (UTIAS)), linear plasma devices (Sandia National Laboratory (SNL)/Los Alamos National Laboratory- Tritium plasma experiment (LANL-TPE), University of California, San Diego-Plasma Interaction with Surface Components Experimental Station B (UCSD- PISCES-B)), a tokamak divertor plasma (DIII- D-DIMES), and a neutral beam (NB-JET). In some of these studies carbon deposition or formation of carbide or oxide surface layers occurred, which is likely to affect D retention. The figure shows the D retention in Be observed under a wide range of exposure conditions. The high fluence saturated con- centration tends to be lower at higher temperatures. ence (D m−2) retention = 0.345 o, or W (300 K) 1023 1024 10252 2 W, at 300K versus incident fluence, measured by thermal z, A. A.; Davis, J. W. J. Nucl. Mater. 1997, 241–243, ma-material interactions in current tokamaks and their 1967–2137 (review special issue), with permission from IAEA. 1019 1021 1022 1023 1024 Particle fluence (D (T ) m–2) R et ai ne d q ua nt ity (D (T )m –2 ) 1025 1026 1027 1020 1021 1022 1023 SNL/LANL-TPE PISCES-B-1 PISCES-B-2 PISCES-B-3 INEEL(200–400) UTIAS(27) DIII-D(DiMES)(130) NB-JET(120) 250 200 200 250 250 40200 100 100 500 500 500 500 500500 500 300 300200 200 540 700 Figure 7 Retention of deuterium and tritium in Be as a function of incident particle fluence. For purposes of comparing results from different experiments using different ion energies, the data have been scaled to correspond to an equivalent 100eV deuterium ion energy. Numerical values next to the symbols and in the legend are specimen exposure temperatures, in degrees Celsius. Reproduced with permission from Anderl, R. A.; et al. J. Nucl. Mater. 1999, 273, 1–26. Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices 635 It must be noted that this phenomenon is very important because it implies that tritium inventories and permeation due to implantation in beryllium for ITER PFC applications should be significantly lower than was previously estimated using classical recombination-limited release at the plasma surface. A first attempt to model this saturation by allowing the recombination coefficient to become exponen- tially large as the mobile atom concentration near the plasma-facing surface approaches a critical value was made by Longhurst et al.99 For Be, calcula- tions suggest that the critical concentration is related to the yield strength using Sieverts’ law of solubility. On the basis of the results of these calculations, it can be concluded that the inventory of tritium in the beryllium first wall of a device such as ITER, because of implantation, diffusion, trapping, and neutron- induced transmutation, will be of the order of 100 g rather than the kilogram quantities estimated previ- ously,100,101 and most of that will result from neutron- induced transmutations in the Be itself and from trapping in neutron-induced traps. Current predic- tions of tritium inventory in ITER are briefly dis- cussed in Section 4.19.6.2.1. Fusion neutrons will create vacancies and intersti- tials in plasma-facing materials. For metals at reactor wall temperatures, these defects will be mobile and will annihilate at sinks (e.g., surfaces or grain bound- aries), recombine, or agglomerate into defect clusters. Vacancy agglomerationmay also lead to the formation of voids. In beryllium, neutron-induced nuclear reac- tions produce helium and tritium, which may be trapped at defects or precipitate as gas bubbles. These defects, resulting from neutron irradiation, will increase the retention of hydrogen, by increasing the concentration of sites where diffusing hydrogen can precipitate as gas or become trapped as atoms. The effect of neutron irradiation on hydrogen reten- tion in metals is complex, but, in principle, this can be modeled, provided the material parameters are known, such as hydrogen diffusivity, solubility, trap binding energy, and defect microstructure produced by the neutron irradiation. For many metals, most of these parameters are known well enough to attempt such modeling. For beryllium, however, uncertainties in solubility, diffusivity, and trapping of hydrogen make such modeling of hydrogen retention difficult. The problem of production of helium and tritium by nuclear transmutation in beryllium itself is dis- cussed in Section 4.19.4.4.5. 4.19.3.2.2 Beryllium codeposition As deuterium retention in plasma-exposed beryllium targets saturates after a given ion fluence (see Section 4.19.3.2.1), it is apparent that retention in codeposits will eventually be the dominant accumulation mech- anism with respect to beryllium PFCs. This is pri- marily due to the fact that the thickness of a codeposit will continue to grow linearly with time. It is, therefore, critical to understand both the Temperature (K) 400 Present data PISCES Present data PISCES Causey et al. TPE Mayer et al. Mayer et al. Causey and Walsh TPE Causey and Walsh TPE 0.01 0.1 1 0.01 0.1 1 600 800 400 800 D /B e O /B e Figure 8 Comparison of D/Be levels in beryllium codeposits with the O/Be levels in the same codeposits. Reproduced with permission from Baldwin, M. J.; Schmid, K.; Doerner, R. P.; Wiltner, A.; Seraydarian, R.; Linsmeier, Ch. J. Nucl. Mater. 2005, 337–339, 590–594. 636 Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices retention amounts and the release behavior of hydro- gen isotopes from beryllium codeposits. In this sec- tion, a ‘codeposit’ includes both the codeposition (where a BeD or BeD2 molecule is deposited on a surface) and co-implantation (where deposited layers of beryllium are bombarded with energetic hydrogen isotopes) processes. Initial interpretation of studies of beryllium code- posits were made difficult by relatively high oxygen impurity content within the codepositing surface.102,103 Subsequent measurements104 with lower oxygen con- tent seemed to indicate that the oxygen level within the codeposit was correlated to the level of hydrogen isotope retention in the codeposit. The other variable that was identified to impact the retention level in these studies was the temperature of the codepositing surface. Measurements seriously questioning the impor- tance of oxygen on the retention level in beryllium codeposits were made by Baldwin et al.105 In this data set, the oxygen content throughout the codeposit was measured by depth profiled X-ray photoelectron spectroscopy and the oxygen content did not corre- late with the deuterium retention level (Figure 8), although the temperature of the codepositing surface was still a dominating term in determining the deu- terium retention level. Later, more detailed measure- ments confirmed that the presence of a beryllium oxide surface layer was not correlated with an increase in retention in beryllium.106 A systematic study of beryllium codeposition fol- lowed,107 identifying three experimental parameters that seemed to impact the retention level in a code- posit. Along with the surface temperature, the inci- dent deuterium energy and the beryllium deposition rate were determined to be influential scaling para- meters. The previously reported data in the literature was also evaluated using the derived scaling and found to agree with the predictions of the retention levels measured under the various experimental conditions present in the different machines. Later the derived scaling was revised108 to use the ratio of the fluxes of the codepositing species, rather than the deposition rate to permit more accurate extrapo- lation to conditions expected in the edge of confine- ment devices. The ability to predict the level of tritium retention in beryllium codeposits is an important aspect of a safety program; however, developing techniques to remove the trapped tritium from codeposits is a more important issue. The deuterium release behavior during thermal heating of beryllium code- posits has been investigated.109 The results show that the maximum temperature achieved during a bake- out is the figure of merit for determining the amount of deuterium release from beryllium. Increasing Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices 637 the time spent at lower baking temperatures did not increase the amount of deuterium released from the beryllium codeposits. These results, along with the retention level predictions, should make it possi- ble to design baking systems for different areas of a confinement device to control the accumulation rate of tritium to a desired level. 4.19.3.3 Mixed-Material Effects A recent review of mixed-material effects in ITER62 provides background information on mixed-material formation mechanisms and plasma–surface interac- tion effects. Here, the focus is on beryllium- containing mixed-material surfaces (i.e., Be/C and Be/W) and the conditions when one might expect these surfaces to dominate the observed plasma– surface interactions. In addition to plasma interac- tions with mixed-material surfaces, which will be discussed here, other aspects such as changes to ther- mal conductivity, material strength, and ductility, the impact of impurities on material joints, etc., must also be carefully evaluated. 4.19.3.3.1 Be–C phenomena Beryllium and carbon have been observed to begin thermally interdiffusing at a temperature of around 500 �C,56 resulting in the formation of a beryllium carbide layer. However, beryllium carbide has also been observed to form during energetic carbon ion bombardment of beryllium surfaces at room temper- ature.110 As mentioned in Section 4.19.3.1.2, the change in the binding energy of the carbide molecule affects the sputtering yield of both the beryllium and carbon atoms. In addition, the formation of beryllium carbide also has a dramatic effect on the chemical erosion properties of a carbon surface bombarded with energetic beryllium ions.67,68,111 The presence of beryllium carbide on the surface of a carbon sample exposed to deuterium plasma has been shown to correlate with the reduction of chemical erosion of the carbon surface.70 The speculation for the cause of this effect is that the carbide enhances the recombination of deuterium in the surface, thereby lessening the amount of deuterium available to interact with carbon atoms on the surface. This is similar to the impact of small amounts of boron carbide in a graphite surface affecting chemical erosion.112 However, the difference here is that instead of obtaining the carbide through an expensive production technique, the car- bide forms naturally as beryllium ions in the plasma interact with the carbon surface. A systematic study of the time necessary to sup- press chemical erosion of a graphite surface due to the interaction with beryllium-containing plasma has been carried out.69 Increasing the surface tempera- ture of the graphite was seen to have the biggest impact on reducing the suppression time. Increasing the beryllium content of the plasma also reduced the suppression time in a nonlinear fashion. An increase of the incident particle energy was observed to increase the time necessary to suppress the chemical erosion of the surface, presumably due to an increase in the removal of the carbide-containing surface layer. A subsequent study showed that applying heat pulses to a graphite surface interacting with beryllium-containing plasma, to simulate surface heating due to intermittent events, acted to reduce the time necessary for the carbide surface to form and suppress the chemical erosion of the surface.113 4.19.3.3.2 Be–W alloying The existence of tungsten beryllide alloys (i.e., Be2W, Be12W, and Be22W) is an excellent example of the importance of mixed-material surface formation in plasma-facing components.114 Figure 9 shows the tungsten–beryllium phase diagram. Each of the ber- yllides shown in the figure exhibits a lower melting temperature than one would expect from a tungsten plasma-facing surface. If plasma containing beryllium impurities interacts with a tungsten surface, there is a possibility of these lower melting temperature beryllide alloys being formed. In thermodynamic equilibrium, various beryllide alloys of tungsten have been observed to form,115 and their reaction rates have beenmeasured,116 at tempera- tures in excess of 800 �C. However, as was seen with beryllium carbide forming during plasma bombard- ment at lower temperature than expected thermody- namically, the concern exists that tungsten beryllide could form at temperatures below 800 �C as well. Well controlled laboratory measurements in vac- uum117 and in plasma simulators118 have shown that although thin, nanometer scale, Be2W layers form at the interface between beryllium and tungsten surfaces, their growth below 800 �C is negligible. In addition, above 800 �C, rapid beryllium sublimation from surfaces can act to limit the amount of beryl- lium available for reacting with tungsten and thereby also limit the growth rate of the alloys. In the present low wall temperature confinement devices, modeling shows that the divertor strike point locations are the only areas where significant beryllide growth might be expected and in these regions there does not 500 1000 1500 2000 2500 3000 3500 0 0 10 20 Be W W 3422 ºC 2100 ± 50 ºC Table 2 Physical properties of beryllium Atomic number 4 Atomic weight 9.013 Crystal structure Hexagonal close- packed Density (kgm�3) 1830–1850 Melting temperature (�C) 1283–1287 Thermal conductivity (Wm�1 �C�1) �200 (RT) �82 (800 �C) Specific heat (J kg�1 �C�1) �1900 Latent heat of fusion (kJ kg�1) �1300 Latent heat of vaporization (kJ kg�1) �3.66 104 Electrical resistivity (mO cm) �4.4 (RT) Thermal expansion coefficient 10�6 �C�1 �11.6 (RT) �14.96 (400 �C) Emissivity 0.1–0.5* (at 300 �C) Source: ITER MPH, ITER Final Design Report 2001 (internal project document distributed to the ITER participants). RT, room temperature. * Depending on quality of surface Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices 639 Several reviews have been published recently related to use of beryllium in tokamaks and the status of the investigations of the Be properties for the fusion application.3,121–126 Various production and proces- sing methods of beryllium metal fabrication have been reviewed in Dombrowski.127 The majority of methods are based on powder metallurgy and include powder preparation from cast product by grinding (i.e., attrition milling, impact grinding, ball mill grinding); further powder consolidation (i.e., by cold pressing (CP), cold isostatic pressing (CIP), vacuum hot pressing (VHP), hot isostatic pressing (HIP)); and possible additional mechanical treatment (e.g., extru- sion, rolling, forging). Beryllium protective armor can also be produced by plasma spray (see Section 4.19.4.3) and vapor deposition. Several proposals were made at the beginning of the ITER Research Programme during the ITER Engineering Design Phase to develop a fusion grade beryllium with high ductility, high resistance to heat flux, and high radiation resistance. However, it was recognized that this development would require sig- nificant efforts and could not be supported only by requests from the fusion community. There are various beryllium grades, which have been developed for different applications.These grades differ by chemical composition (BeO content, impuri- ties), by method of powder preparation, by method of consolidation, etc. The nonexhaustive list of various beryllium grades from the US and the Russian Feder- ation is presented in ITER Materials Properties Handbook (MPH).128 Grades with similar composi- tion are under production in Kazakhstan and inChina. We briefly discuss below some of the most relevant physical and mechanical properties of beryllium, in relation to its application as armor for PFCs. 4.19.4.1.1 Physical properties The physical properties of beryllium are summarized in Table 2, which is taken from ITER MPH.128 These properties have been used for design and performance assessments. In addition to its low atomic number, beryllium has several excellent ther- mal properties that make it well-suited for heat removal components. The thermal conductivity is comparable with that of graphite or CFC at low and high temperatures but, in contrast to C-based mate- rials, is not significantly degraded as a result of neutron-irradiation. The specific heat of beryllium exceeds that of C-based materials typically by a factor of 2 over the temperature range of interest for operation. However, Be has poor refractory properties, such as low melting temperature and high vapor pressure. The high heat capacity and good thermal conductivity of Be can be used to maintain low surface temperatures in PFCs during normal operation, but its low melting temperature and high vapor pressure cause great design difficul- ties from the standpoint of survivability from off- normal events such as vertical displacement event (VDE), ELMs, disruptions, and runaway electron impact (see Section 4.19.6.2). For the beryllium hexagonal close packed crystal structure, the main physical properties, such as the coefficient of thermal expansion, elastic modulus etc. have some anisotropy. However, for the polycrystal- line grades these properties could be, in the first approximation, considered as isotropic. Some anisot- ropy is also typical for the highly deformed grades. The physical properties (thermal conductivity, spe- cific heat, elastic modulus, etc.) in first approximation are the same for beryllium grades with similar BeO and other impurity content and they are produced by the same fabrication method. 4.19.4.1.2 Mechanical properties Beryllium is known to be a brittle material, with a typical elongation to failure in room temperature tensile tests of roughly �0.8–6%. For material with strong anisotropy (e.g., rolled plate or sheet), elonga- tion in the rolling direction could be higher, but in the transverse direction the elongation is Table 3 Candidate grades of beryllium Producer Grades Brush Wellman, US S-65C VHP, S-65C HIP, S-65 CIP, SR-200 VHP S-200F HIP, S-200F VHP, I-400 VHP Russian Federation DShG-200, TShG-56, TR-30, TGP-56 TShGT, DIP-30, TShG-200 VHP, vacuum hot pressing; HIP, hot isostatic pressing; CIP, cold isostatic pressing. 640 Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices typically significantly lower than 1%. Recently, the mechanical properties of beryllium have been summarized in ITER MPH128 and ITER Materials Assessment Report (MAR).129 The mechanical properties of beryllium depend on the production method used and they are sensitive to a variety of factors including BeO and impurity content (which varies from less than 1% to 2–3% for various grades), method of powder preparation (impact grinding, attrition grinding), method of con- solidation, and further treatments. The main problem in using beryllium is its low ductility related to the hexagonal-close-packed structure. There is limited slip in directions not parallel to the basal planes, resulting in very small ductility perpendicular to the basal direction. Depending on the production method, ductility of beryllium can be severely anisotropic. The grain size is an important factor in determining the ductility of various beryllium com- ponents. Much of the fine grain size present in the starting powder is retained during hot pressing at 1060 �C. Without an oxide network, grain growth occurs at a much lower temperature, about 800 �C. Among various beryllium grades, it was found that grade S-65C VHP (production of Brush Wellman, US) has the highest guaranteed fracture elongation at room temperature (minimum 3%; typical is more than 4–5%). This grade is produced using impact grinding powder and has a guaranteed BeO content 850 �C. This last feature is important for the selection of the joining technology for manufacturing of the PFCs. Further details on mechanical properties, such as creep and fracture toughness, can be found else- where (see, e.g., ITER MAR129). 4.19.4.2 Selection of Beryllium Grades for ITER Applications For ITER PFC applications, various commercially available beryllium grades from the United States (Brush Wellman Inc.) and from the Russian Federa- tion, listed in Table 3, were evaluated more than a decade ago as potential candidates during the ITER Engineering Design Activity (EDA). The selection of the optimum grade for ITER PFC applications is driven mainly by the require- ments of ITER operation for structural integrity and stability against various thermal loads, and in partic- ular, the absence or minimization of macrodamage. It is believed that ion-induced and thermal erosion at elevated temperatures is very similar for various grades of Be. However, performance under high heat fluxes, especially under transient thermal loads such as disruptions, VDE, and ELMs resulted in different behavior and damage mechanisms. It is considered that the ease of joining beryllium to copper alloys (see Section 4.19.5) is not so sensitive to BeO con- tent, impurity levels, and method of consolidation, which are the parameters defining the grade of beryl- lium material. It should be noted that for tokamak applications (see Section 4.19.6) beryllium is used in the form of tiles. Some surface cracking of the tiles could be acceptable, if there is no macrodamage or delamina- tion along the surface of tiles, which leads to the loss of macroscopic particles. The resistance to thermal fatigue is the most important factor that affects the material selection Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices 641 because cracking could lead not only to enhanced armor erosion, delamination, and loss of particles, but also potentially to crack propagation to the heat sink structure. Neutron irradiation resistance is another factor to be taken into account because it may affect the thermal performance and structural integrity. Because of some of the uncertainties in the ITER thermal loads, especially during transient events, preference is given to beryllium grade(s) with poten- tially higher resistance to transient thermal loads. The selection of the reference grades was made on the basis of comprehensive assessment of the results of various tests carried out during the ITER EDA. The detailed analysis is presented in ITER MAR.129 Among the various studies, the following shall be mentioned: � Screening low cycle fatigue test of 21 different beryllium grades was performed in the past.130 It was shown that the grades with the best thermal fatigue resistance are S-65C VHP, DShG-200, TShG-56, and TShGT. Figure 10 shows the results of the comparative low cyclic thermal fatigue study of different grades of beryllium. � Various grades of beryllium were also tested in conditions simulating the disruption heat loads.131 The tests show that crack formation and behavior after surface layer melting in different grades are quite different. For Be S-65C, all cracks stopped in 0 0 500 1000 C yc le s to c ra ck in iti at io n Side crack prop Be/60% BeO Be/30% BeO SR-200 S- S S-65C (T) S-65C (L) DShG-200 (T) 94 98% S-65 Extruded (T) TShGT(T) TShG-56 (T)1500 2000 2500 0.5 Figure 10 Results of low cycle thermal fatigue tests of differe crack propagation depth. Reproduced with permission fromWat of the 2nd IEA International Workshop on Beryllium Technology the molten zone, whereas for some grades the cracks propagated to the bulk of the sample. � Results of VDE simulation tests have been reported in Linke et al.132,133 Severe melting of Be was observed for energy densities of 60MJm�2 (�1 s pulse duration); however, no cracks were observed between molten and unmolten material and in the bulk of unmolten parts for S-65C VHP grade. On the basis of the available data, Be S-65C VHP (Brush Wellman, US) was selected as the reference material on the basis of excellent thermal fatigue and thermal shock behavior, and for the good available database on materials properties, including neutron irradiation effects. DShG-200 (produced in the Rus- sian Federation) was proposed as a backup, but this grade is no longer commercially available. Recently, China and the Russian Federation, that are two of the seven International Parties engaged in the construction of ITER, have proposed the fabrica- tion of additional first-wall grades as part of their ITER contribution. The Russian Federation proposes to use beryllium grade TGP-56-FW. This grade is produced by VHP in almost the final form of the tiles foreseen for the first wall. The recent results on development of this grade have been reported in Kupriyanov et al.134 China proposes instead to use a grade called CN-G01135 that is produced from agation depth (mm) I-400 (T) 200F-H -200F (T) Grades with best fatigue performance S-65-H TGP-56 S-200F (L) % S-65 Extruded (L) 1 1.5 2 nt Be grades: number of cycles to crack initiation versus son, R.; et al. Low cyclic fatigue of beryllium. In Proceedings for Fusion, Wyoming, Sept 6–8, 1995; p 7. 642 Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices impact ground powder (similar to powder used for S-65C grade) by VHP. The grade is produced by Ningxia Orient Non-Ferrous Metals Co. Ltd. In order to accept these newly proposed beryllium grades a specific qualification program is underway. 4.19.4.3 Considerations on Plasma-Sprayed Beryllium In the past, plasma spraying was considered as a high deposition rate coating method, which could offer the potential for in situ repair of eroded or damaged Be surfaces. Development work was launched during the early phase of the ITER R&D Program in the mid-1990s.136 In the plasma spray process, a powder of the material to be depos- ited is fed into a small arc-driven plasma jet, and the resulting molten droplets are sprayed onto the target surface. Upon impact, the droplets flow out and quickly solidify to form the coating. With recent process improvements, high quality beryllium coat- ings ranging up to more than 1 cm in thickness have been successfully produced. Beryllium deposition rates up to 450 g h�1 have been demonstrated with 98% of the theoretical density in the as-deposited material. Several papers on the subject have been published.136–138 A summary of the main achieve- ments can be found in Table 4. However, based on the results available, the initial idea of using plasma-sprayed beryllium for in situ (in tokamak) repair was abandoned for several rea- sons. First was the complexity of the process and requirements to control a large number of para- meters, which affect the quality of the plasma sprayed Table 4 Main achievements of ITER-relevant plasma-spraye together) Parameter Value/results Comme Residual porosity (%) �2 Could b Thermal conductivity (WmK�1) Up to 160 at RT Depend T� 60 Bond strength (MPa) 100–200 Reason Substrate temperature (�C) >450 Very imp condu higher Substrate preparation Negative transfer arc Needed Deposition rate (kg h�1) 4.5 Reason Thickness (mm) >10 Reason Deposition efficiency (%) >90 It means Thermal fatigue (MWm�2/ number of cycles) 5/680; 1/3000 For first coatings. Some of the most important parameters include plasma spray parameters such as (1) power, gas composition, gas flow-rate, nozzle geometry, feed, and spray distance; (2) characteristics of the feedstock materials, namely, particle size distribution, morphol- ogy, and flow characteristics; (3) deposit formation dynamics, that is, wetting and spreading behavior, cooling and solidification rates, heat transfer coeffi- cient, and degree of undercooling; (4) substrate conditions, where parameters such as roughness, temperature and thermal conductivity, and cleanli- ness play a strong role; (5) microstructure and properties of the deposit, namely, splat characteris- tics, grain morphology and texture, porosity, phase distribution, adhesion/cohesion, and physical and mechanical properties; and (6) process control, that is, particle velocity, gas velocity, particle and gas temperatures, and particle trajectories. Second, plasma-sprayed beryllium needs (1) inert gas pres- sure, (2) reclamation of the oversprayed powder (more than 10%), and (3) strict control of the sub- strate temperature. The higher the temperature the higher the quality of the plasma-sprayed coating, but unfortunately, an easy and reliable method to heat the first wall to allow in situ deposition was not found. Finally, tools to reliably measure the quality of the coating and its thickness are not available today and a strict control of the coating parameters is difficult to achieve. Thus, it was concluded that plasma-sprayed beryl- lium for in situ repair is too speculative for ITER without further significant developments. Neverthe- less, this method still remains attractive and could be used for refurbishment of damaged components in d technology (summary of best results, not always achieved nts e more than 5% s on temperature of substrate, maximum achieved at 0–800 �C with addition of H able ortant for good strength, adhesion, and thermal ctivity. Keep in mind that CuCrZr temperature should not be than 500 �C for several hours due to overageing of CuCrZr , but very difficult to do in situ able able that more than 10% of powder will be lost in chamber -wall conditions tested Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices 643 hot cell, albeit it may be cheaper to replace a dam- aged component with a new one. 4.19.4.4 Neutron-Irradiation Effects Several authors have reviewed the properties of neutron-irradiated beryllium for fusion applications in the past.139–141 Neutron irradiation leads to com- plex changes in the microstructure, such as the radiation-induced change of volume in beryllium, which is dominated by the nucleation and growth of He bubbles. There are two important pathways for gas produc- tion. One is the (n, 2n) reaction in which the 9Be is reduced to 8Be, which then splits into two 4He atoms. The second is the (n,a) reaction where the 9Be absorbs a neutron and then splits to form a 4He and a 6He. The 6He rapidly undergoes a b� decay to become 6Li. The 6Li then reacts with a thermal neutron to produce 4He and 3H. These processes have been incorporated into the inventory code FISPACT,142 which is used (see, e.g., Forty et al.143) to estimate the generation rates of gas and other reaction products in a tokamak. Helium generation has significant effects on the properties of materials, especially at elevated tempera- tures. Helium is initially trapped within the beryllium lattice in submicroscopic clusters. At higher neutron fluence massive helium-bubble-induced swelling occurs, especially at elevated irradiation or postanneal temperatures. Because of the atomistic nature of the helium bubble nucleation and growth, porous beryl- lium microstructures, such as from powder metallurgy or plasma spray technology, were not found to be effective in releasing significant amounts of helium under fusion reactor conditions.2 The maximum neutron-induced damage and helium production expected in Be for ITER first- wall applications (fluence of 0.5MWam�2) are �1.4–1.7 dpa and �1500 appm, respectively and the expected irradiation temperatures are in the range of 200–600 �C. The maximum temperature is on the surface of beryllium tile and depends on thick- ness and heat flux. Tritium production in beryllium is expected to be about 16 appm. Recently, Barabash et al.144 have analyzed the specific effects of neutron- induced material property changes on ITER PFCs foreseen during ITER operation. Typically, property changes induced by neutron irradiation are investigated by exposing samples/ mock-ups in fission reactors. However, the differ- ences between the fission and fusion neutron spectra are important to interpret and predict the effects. The key difference is transmutation production, which needs to be considered for the correct predic- tion of the material performance.145 During irradia- tion in fission reactors, for example, the typical value of the ratio (appm He per dpa) is 100–250, whereas for a fusion neutron spectrum this value is �1000. Depending on operational temperature, the dpa or He transmutation must be used as a reference neu- tron damage parameter. For beryllium, during low- temperature irradiation ( 644 Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices high-temperature irradiation (more than�400 �C) or after high-temperature annealing. The maximum values of swelling could reach approximately tens of percent at temperatures more than 600 �C and helium content more than several thousand atomic parts per million. Swelling depends on the structure of the beryllium: beryllium grades with small grain size (�8–10 mm) and high BeO content (�3–4wt%) have a higher resistance to swelling than conventional Be grades.141 As concluded in ITER MAR,129 for an irradiation temperature Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices 645 with areas subjected to higher power densities during each pulse. In such cases, other options for the heat exhaust technology are being considered,156 using thinner Be tiles. 4.19.5.1.1 Be/Cu alloy joining technology 4.19.5.1.1.1 Background information The main problem of bonding Be to Cu alloys is that Be reacts with almost all possible metals (except Al, Si, Ag, and Ge) and forms brittle intermetallic phases.157–159 Such bond joints have poor mechanical integrity. More robust joints use metal interlayers to act as either diffusion barriers and/or strain accom- modating compliant layers to avoid the formation of deleterious phases and to assist in the accommoda- tion of thermal cycling-induced strains.160 R&D has been performed over a number of years to develop the design and manufacturing techniques required to meet the demanding design require- ments. Significant experience has been gained with these manufacturing techniques and the associated inspection techniques. It must be noted that in the 1990s the best joining technology developed for manufacturing the Be/Cu actively cooled compo- nents was brazing with Ag-base alloys (e.g., InCuSil with �41.75% Ag) which was successfully used in JET. However, the use of Ag base brazing alloy was not allowed in ITER mainly because of the transmu- tation by neutron irradiation to Cd (�5wt% Cd will be produced in Ag–Cu eutectic alloy at 1MWam�2) whose presence would (1) reduce the melting tem- perature of the braze; (2) lead to the formation of highly radioactive isotopes; and (3) affect the pump- ing system in case of Cd release to the vacuum chamber and codeposition in the cryopumps panels. During the early stage of the ITER first-wall design development, dispersion-strengthened copper (DS-Cu) alloys (e.g., Glidcop Al25) were considered as the first option because (1) the stresses were within the design allowable, and (2) they had better thermal stability under the manufacturing route, which required a first wall to be integrated with a 4 t shield. The main developments for fabrication of joints between Be and DS-Cu alloys are reported in ITER MAR129 and Lorenzetto et al.161 However, as a result of a design change that took place from an integrated first-wall panel to a separated first-wall panel design, a precipitation-hardened copper–chromium–zirconium alloy (CuCrZr), was subsequently chosen. This was because the fracture toughness of DS-Cu is very low above 200 �C even for unirradiated material. Fracture toughness of the unirradiated and irradiated CuCrZr alloy decreases with increasing temperature, but it remains at a rela- tively high level in the ITER working temperature range and it is significantly higher than fracture tough- ness of DS-Cu. The use of separable first-wall panels makes it possible to perform heat treatments with fast cooling rates, which are mandatory to adequately retain the mechanical properties of precipitation- hardened materials. Thus, extensive studies were then performed during the last 10 years to develop reliable silver free Be/CuCrZr alloy joining techniques and to modify the joining conditions to minimize the mechanical strength loss of the CuCrZr alloy. Dif- ferent methods have been considered and investi- gated. Some of these were eliminated because of bad results (e.g., explosive bonding, inertial welding, joint rolling, and some types of brazing). Two meth- ods gave good results and were kept for further investigations: HIPping and brazing. Good results were achieved with the HIP joining technique by lowering the HIP temperature as close as possible to the CuCrZr alloy ageing temperature (about 480 �C) and with the brazing technique in the development of a fast brazing method to minimize the holding time at high temperature. The latter was achieved by induction brazing in Europe and by fast heating and cooling using an e-beam test facility in the Russian Federation. 4.19.5.1.1.2 HIP joining technique An extensive development programme performed especially in Europe has enabled the production of very good Be/CuCrZr alloy joints by HIPping. The progress on the fabrication of Be/CuCrZr joints in Europe is described in Lorenzetto et al.155,161,162 and Sherock et al.163 The HIP joining temperatures ranged from 540 to 580 �C. Different interlayers such as Cr, Ti, and Cu were tested and reported. The selection of the joining conditions used for the fabri- cation of representative first-wall mock-ups was done on the basis of mechanical test results performed with guillotine shear test specimens. The best results were obtained with Ti and Cu interlayers at 580 �C, in which the shear strength exceeded the yield strength of the parent materials as creep of the CuCrZr part or rupture of the Beryllium part was observed. Per- formance achieved with representative first-wall mock-ups exceed the present ITER design require- ments, namely, 30 000 cycles at 0.5MWm�2 peak heat flux plus transient events up to 1.4MWm�2 for about 1000 cycles (see Section 4.19.5.1.2). 646 Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices A neutron irradiation programme is still in progress to complete the full characterization with irradiated mock-ups (see Section 4.19.5.2). In the frame of the ITER Programme, in addition to Europe, the countries that in the past were inter- ested to supply the ITER first wall were China, Korea, the Russian Federation, and the United States. The underlying development work which has recently been performed in these countries to estab- lish the necessary fabrication capability is summar- ized elsewhere (see, e.g., Sherock et al.,163 Nishi et al.,164 Hong et al.,165 Youchison et al.,166 Lee et al.,167 Park et al.,168 Liu,169 and Chen170). 4.19.5.1.1.3 Fast brazing techniques A development programme has been launched to develop a fast brazing technique to minimize the holding time at high temperature and consequently retain adequate mechanical properties of the CuCrZr alloy. This was achieved by induction brazing in Europe and by fast heating and cooling using an e-beam test facility in the Russian Federation. Induction brazing tests were done using the only appropriate silver free braze alloy available in the market, the STEMET 1108 procured from the Russian Federation. It was found that this braze alloy had poor wetting properties and the quality of the product was variable. Difficulties were met for brazing Be tiles of dimensions representative of the Be tiles of first-wall panels. A few first-wall mock-ups were fabricated with inductively brazed Be tiles but showed thermal fatigue performance well below HIPped mock-ups, with detachment of Be tiles between 1.5 and 2MWm�2.162,171 This result has been considered unsatisfactory. The development work on fast brazing equipment has been stopped in Europe and the effort is being concentrated on the development of a new silver- free braze alloy. The fast brazing development in the Russian Federation has also been done using the STEMET 1108 alloy but fast heating was performed using an e-beam test facility. First-wall mock-ups were heated on the beryllium side by the e-beam and tempera- tures as high as 780 �C were achieved at the Be/CuCrZr joints for a very short time, followed by fast active cooling of the mock-ups, minimizing therefore the production of brittle intermetallic com- pounds.172 Good results were achieved on hyper- vapotron type mock-ups, developed at the early stage of the ITER design, for Be coated divertor components.129 4.19.5.1.2 High heat flux durability of unirradiated Be/Cu joints Actively cooled first-wall mock-ups have been tested under relevant heat fluxes in electron beam facilities (accelerated fatigue tests) or in low-heat flux facil- ities, which are capable of delivering moderate heat flux (10 000) of larger dura- tion to investigate the thermomechanical performance (including integrity and possible creep phenomena) of the bonded interfaces and fatigue lifetime. Because of intrinsic limitations, cost, or the limited availability of testing facilities, electron beam facilities are used with the shortest duration of the heating cycle, com- patibly with the conditions of near thermal steady- state. In many cases, to reduce the testing time, the tests are performed at a heat flux well above the nomi- nal one. These tests are often used to select the best joining conditions and then confirmation and final selection are done including tests at representative moderate heat fluxes but for the specified large number of cycles. Because of the health hazards linked to the use of berylliummaterial, only a limited number of test facilities were available or built for testing beryllium components. As far as Be-compatible electron beam facilities are concerned, two are available in Europe, the Juelich divertor test equipments in hot cells ( JUDITH 1 and JUDITH 2) at the Forschungszen- trum Juelich (FZJ) (Germany), one in the Russian Federation at the Efremov Institute in St. Petersburg and one in the United States at Sandia National Labo- ratory in Albuquerque (New Mexico). For the perfor- mance of thermal fatigue tests at moderate heat fluxes and large number of cycles, three facilities were built in Europe, one at the Ispra Joint Research Center in Italy, one at the ENEA Research Center of Brasimone in Italy, and one at the Nuclear Research Institute (NRI) of Rěz close to Prague in the Czech Republic. From these three facilities, only the last one is still under operation. The best high heat flux test results from represen- tative first-wall mock-ups, namely, 10-mm beryllium tiles joined onto 10mm CuCrZr heat sink layer with embedded 10/12mm diameter stainless steel pipes and the assembly joined onto a stainless steel backing plate (see Figure 11), were achieved with Ti and Cu interlayers and HIP temperatures of about 580 �C as described in Section 4.19.5.1.1.2. The per- formance limit of this assembly is presently at about 3MWm�2.162 As far as the thermal fatigue tests are concerned, representative first-wall mock-ups with Be/CuCrZr Figure 11 First-wall mock-up for high heat flux tests. Figure 12 First-wall mock-up for thermal fatigue tests. Reproduced with permission from Lorenzetto, P.; et al. Fusion Eng. Des. 2008, 83, 1015–1019. Figure 13 ITER first-wall qualification mock-up, EU mock-up before testing. Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices 647 joints made as above have been successfully tested up to 30 000 cycles at about 0.6MWm�2 (see Figure 12). Ultrasonic testing of the Be/CuCrZr joints after testing did not show any indication of defects. These mock-ups will be tested at higher heat fluxes representative of transient and off-normal events to check the available performance margins. Originally, six ITER Domestic Agencies were candidates for the procurement of the ITER first wall: China, European Union (EU), Japan, South Korea, Russian Federation, and United States. (The seven members of the international ITER project have all created Domestic Agencies to act as the liaison between national governments and the ITER Organization. The Domestic Agencies’ role is to han- dle the procurement of each member’s in-kind con- tributions to ITER.) However, as stated in the Final Report of Negotiations on ITER Joint Implementa- tion of 1 April 2006, a prequalification ‘‘. . . will be needed for the critical procurement packages shared by multiParties. . .’’, such as the blanket first wall. Well in advance of the assumed start of the procure- ment, each ITER Domestic Agency shall first dem- onstrate its technical capability to carry out the procurement with the required quality, and in an efficient and timely manner. For the first wall system, this is achieved via a two-stage qualification process: a mock-up qualification stage and a semiprototype qualification stage. Each stage is also split into two phases: a manufacturing acceptance phase and a heat flux testing acceptance phase. The successful manufacturing and testing of two first wall mock- ups for stage I (see Figure 13) demonstrating in particular the know-how to assemble beryllium (Be) tiles on a CuCrZr alloy and stainless steel bimetallic structure is the prerequisite to be eligible for stage II. The qualification tests for stage I have been split between the United States and EU in three test facilities: at the SNL for the US and at the NRI in the Czech Republic and the FZJ in Germany for the EU. At the SNL facility, the qualification test programme consists of the performance of 12 000 cycles at 0.88MWm�2 for 1.6min followed by 1000 cycles at 1.4MWm�2, while in the EU test facilities it consists of the performance of 12 000 cycles at 0.62MWm�2 for 5min (at NRI) followed by 1000 cycles at 1.75MWm�2 (at FZJ). To be qual- ified, a Domestic Agency shall fabricate two mock- ups which pass both tests. The first wall mock-ups fabricated by the EU Domestic Agency have successfully achieved the above test programme conditions without any indi- cation of failure. Additional tests were also performed on these mock-ups to assess the limit and tests were performed for 200 cycles at 1.7MWm�2 at the SNL and for 100 cycles at 2.25MWm�2 plus 100 cycles at 2.75MWm�2 at FZJ without any indication of fail- ure. Tests were stopped so as not to exceed the maximum acceptable Be temperature in the test facilities. The progress on the fabrication and thermal tests is described elsewhere.163–171 Figure 14 Be/CuAl25 mock-up after postirradiation high-heat flux test at 4.25MWm�2. Reproduced with permission from Lorenzetto, P.; et al. Fusion Eng. Des. 2006, 81, 355–360. 648 Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices 4.19.5.2 Thermal Tests on Neutron-Irradiated Joints Thermal fatigue durability is a necessary but not sufficient prerequisite for the Be/CuCrZr alloy joint. Neutron irradiation is also expected to affect the high heat flux durability of these joints. This is particularly true for the first-wall components, where the joint will experience rather high neutron fluence. To investigate the effects of neutron irradiation on the joint properties there are two possible ways. The first method is postirradiation testing, that is, irradia- tion of small scale mock-ups in fission reactors and postirradiation testing in heat flux test facilities. The second method is to try to reproduce more faithfully the situation in ITER and to apply a cyclic heat flux during irradiation. This consists of in-pile heat flux testing of small mock-ups in fission reactors. Given the technical difficulty of achieving simulta- neously representative values of heat flux and neu- tron irradiation and to determine exactly what the heat fluxes are within a reactor, it is suggested that the effect of neutron irradiation on the mock-ups be determined by pre- and postirradiation heat flux tests on mock-ups. This method is probably conser- vative as the postirradiation tests are performed on materials and joints having accumulated damage corresponding to the total neutron irradiation dose. However, this should be supported by analysis and material testing. Most of the tests conducted in the past were done for DS-Cu as a heat sink. Studies of neutron irradia- tion effects on the durability of the Be/Cu-alloy joints have been performed in at least two of the ITER parties: Europe,150,162 and the Russian Federation.173 For the Russian experiment, the irradiation con- ditions were 0.3 dpa at 350 �C. The irradiated Be/Cu mock-ups were then tested in the JUDITH facility. The results of the postirradiation high heat flux test- ing of the different Be/Cu mock-ups are presented elsewhere.174 On the basis of the results of these tests, it was concluded that the effect of the neutron irradi- ation is not critical for the joints. Metallographic inspections did not show any significant changes in the braze joint after neutron irradiation. For the CuMnSnCe, a small intermetallic phase was observed in the middle of the braze layer for unirra- diated and irradiated samples. No crack formation was found in the intermetallic. In addition to thermal fatigue tests, shear tests were conducted and it was found that for CuMnSnCe the shear strength decreased after neutron irradiation from 200 to 155MPa, whereas for the InCuSil braze no irradia- tion influence was observed (�300MPa). Neverthe- less, the thermal performance of the joints during high heat flux tests was very similar to the perfor- mance of unirradiated mock-ups. For the European experiments, the first irradia- tion campaign (named Paride 1 and Paride 2) of small scale Be/Cu mock-ups and Be/Cu joints took place in 1996–1999. Mock-ups were fabricated from a sin- gle 10-mm thick Be tile (grade S-65C) of dimensions 22mm� 60mm, HIPped on a 20-mm thick CuAl25 substrate (grade IG1) with a drilled 10-mm diameter cooling channel (Figure 14). HIPping was done at 830 �C for 2 h with the use of a 50 mm Ti interlayer. One mock-up was neutron irradiated in the test reactor high flux reactor (HFR) and then high-heat flux tested in JUDITH. The neutron irradiation was done at about 200 �C up to a neutron dose of about 0.6 dpa in the Be material. The neutron dose expected at the end of life in the Be armor of the first-wall panels is about 1 dpa (Be). The irradiated mock-up was tested for 1000 cycles at 1.6MWm�2 plus 100 cycles at 1.9MWm�2 plus 100 cycles at 2.4MWm�2 plus 1000 cycles at 2.8MWm�2 plus 100 cycles at 3.3MWm�2 without any sign of failure. It failed dur- ing the first cycle at 4.25MWm�2 with a partial detachment of the Be tiles on one end (Figure 14). An unirradiated mock-up was tested for 1000 cycles at 1.5MWm�2 plus 1000 additional cycles at 3MWm�2 without any sign of failure but a detachment of the Be tile occurred during the first cycle at 4.5MWm�2. It was therefore concluded that no significant degrada- tion of the Be/CuAl25 joint was observed up to a neutron dose of about 0.6 dpa. Neutron irradiation test experiments are ongoing or in preparation with Be coated first-wall mock-ups made from CuCrZr alloy to confirm the above result with this Cu alloy. Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices 649 The first experiment was a joint European/Rus- sian irradiation test campaign. It was prepared by the Efremov Institute of St. Petersburg. The original objective was to perform thermal fatigue testing of two first-wall mock-ups at about 0.5MWm�2 simul- taneously to neutron irradiation. A failure of the surface heating system made from graphite after about 5000 cycles resulted in the discontinuation of the thermal fatigue testing and a continuation of the campaign with only neutron irradiation. The irradia- tion campaign has been stopped with the achievement of a neutron dose of 0.75 dpa. The first wall mock-ups, one made with Be tiles HIPped at 580 �C and a Cu interlayer (Figure 15) and another with brazed Be tiles with STEMET 1108 braze alloy, will then be high heat flux tested together with unirradiated reference first wall mock-ups. Two other test cam- paigns are in preparation at the NRI of Rěz (Czech Republic) and at Petten (The Netherlands) with the objective of thermal fatigue testing three first wall mock-ups in parallel to neutron irradiation. The question as to whether the correlation between fusion and fission neutron spectra assumed in many of the above measurements is valid or not needs to be discussed. Comparison of changes in the mechanical properties, especially at low temperature, needs to be made with the same He to dpa ratio to ensure that the results will be valid for ITER. Be/Cu alloy mock-ups have been tested in an electron beam for 1.5 s under a deposited energy density of 60MJm�2(132) to simulate Be damage dur- ing a VDE. For a 6mm thick Be tile, the melt layer was �1.5mm, while that calculated for the same Figure 15 First-wall mock-up for irradiation test experiments. Reproduced with permission from Lorenzetto, P.; et al. Fusion Eng. Des. 2008, 83, 1015–1019. condition is �1.25mm. A cross section of the Be tile after the simulated VDE event is shown in Figure 16. The embrittlement of the Be due to neutron irradi- ation increases the loss of material particles, espe- cially at low irradiation temperature. A clear pore formation (which is expected to be He filled) has been observed in the melt layer of all neutron irra- diated specimens after thermal shock loading (see Figure 16). An area of possible concern is the small surface cracks that form when molten metals reso- lidify. These resolidification cracks could serve as thermal fatigue crack initiation sites and accelerate this type of damage. While this effect has not been extensively studied because of the difficulty of simulating disruptions in the laboratory, it may not be a critical issue as thermal fatigue cracks form after a few hundred cycles in most materials and they grow only to depths where the thermal stress level is above the yield stress.175 High heat flux tests of neutron irradiated mock- ups conducted in the past did not reveal any damage in Be and in Be/Cu joints,176 although the irradiation conditions were not fully ITER relevant (damage dose of �0.3 dpa instead 1 dpa for the end of life and also lower He per dpa). An increase in crack formation and erosion rate has been observed in the surface of irradiated Be at 350 and 700 �C.177 The S-65C grade presented the lowest damage after CuCrZr S-65C 500 mm Figure 16 Micrograph of S-65C armor on CuCrZr (CuMnSnCe braze) after vertical displacement event simulation. The actively cooled modules have been loaded with energy densities of 60MJm�2 (effective pulse duration: 1 s). Reproduced with permission from Linke, J.; Duwe, R.; Gervash, A.; Qian, R. H.; Roedig, M.; Schuster, A. J. Nucl. Mater. 1998, 258–263, 634–639. 650 Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices irradiation. High heat flux tests (at more severe con- ditions than needed for the first wall) of the cracked unirradiated Be did not reveal any detrimental behav- ior and loss of material due to cracking.178 From an engineering point of view, to avoid possible crack formation and delamination of the brittle Be, it is recommended to use Be tiles without any stress concentrations. 4.19.6 Tokamak PFC Design Issues and Predictions of Effects in ITER During Operation 4.19.6.1 PFC Design Considerations Robust engineering solutions are needed for the com- ponents that directly surround the plasma in order to withstand the thermal and mechanical loads during normal and off-normal operation and contribute to protect the outer components from the effect of neu- trons, especially the vacuum vessel and the magnets. Considerations here are limited to the portion of the PFC surface that protects the main chamber (the so-called first wall), which for both JET and ITER is made of beryllium. The strategy and the criteria adopted to design the PFCs of JET and ITER are substantially different. One of the main differences is the longer plasma pulse duration foreseen in ITER that has required the use of water-cooled structures to handle the power in steady-state conditions. This translates into design solutions that consist of tiles bonded to an actively cooled copper-alloy substrate that must withstand the thermomechanical loads with a good engineering margin and achieve the desired fatigue lifetime. An additional problem in ITER is the deg- radation, albeit limited, of thermal and mechanical properties and the properties of the joints due to neutron irradiation (see Section 4.19.4.4). In contrast to ITER, technical constraints pre- vented the use of actively cooled first-wall struc- tures in JET and the first-wall protection relies on a series of discrete poloidal limiters. The tiles are inertially cooled, and their power handling perfor- mance is driven by the need to (1) avoid surface melting and (2) reduce thermally induced stresses to give an adequate fatigue lifetime. At 40mm typi- cal thickness, the tiles are thermally thick for a typical 10 s pulse and will handle up to about 60MJm�2 without melting. Another important driver for designing PFCs is the magnitude of the electromagnetic loads associated with plasma disruptions. In tokamaks, dis- ruptions produce large changes in magnetic field, dB/dt, which induce eddy currents in the conducting materials. The currents interact with the local mag- netic field, B, to produce a torque, which is strongly dependent on the geometry. The design of the PFCs has to manage this torques via a combination of the castellations of the tiles along with cuts, which will interrupt the eddy current loops. The tile assembly must also withstand electromagnetic forces due to halo currents which, during disruptions, pass between plasma and vacuum vessel via the tiles. Enormous attention has been paid at JET and ITER during the design phase to address this problem. The problems associated with the design of the first wall at JETand ITER are briefly described in Section 4.19.6.1.1 and 4.19.6.1.2, respectively. 4.19.6.1.1 Design of the beryllium ITER-like wall at JET JET has completed in 2011 a large enhancement programme that includes, among other things, the installation of a beryllium wall and a tungsten divertor. An overview of the status of the JET ITER-like wall project is presented in Matthews et al.179 The material combination chosen for the wall and the divertor is that chosen for the DT phase of ITER and experiments in JETwith the new wall configura- tion will provide the first fully representative test of material migration, material mixing, and consequent tritium retention under ITER relevant conditions.180 Equally important is the opportunity to develop fully integrated scenarios and control schemes for protect- ing the wall. The project will therefore provide essen- tial information for interpreting material behavior in ITER and a sound technical basis for guiding the development of ITER scenarios. The design layout, the main engineering challenges, and the operational limits of the JET ITER-like wall are discussed elsewhere (see, e.g., Nunes et al.,181 Thompson et al.,182 Riccardo,183 and Riccardo et al.184). Figure 17 describes the design layout and the planned material layout. It must be noted that the existing JET wall relies on a series of discrete poloidal limiters whereas at the moment ITER relies on a plasma conforming wall. The elec- trical resistivity of Be� 0.08 mOm is more than a 100 times lower than that of CFC (�10 mOm). Therefore, after replacing the CFC tiles in JET, the mechanical loads due to eddy currents associated with disrup- tions, which were negligible in the case of carbon tiles, have become dominant for Be tiles and this Dump plate Saddle coil protection Inner wall guard limiter Outer poloidal limiter Bulk Be W or W-coating Be-coated inconel Divertor (a) (b) Divertor B & C tiles Re-ionization protections IW cladding for pellets Inner wall cladding Bulk W Inner wall guard limiters Saddl prote Saddle c protectio Mus Poloidal limiters Upper dump plate Beryllium W-coated CFC Inconel + 8 mm Be Bulk W Restraint ring protections Inner wall g Figure 17 (a) View of the present Joint European Torus main chamber with indications of how the carbon-fiber composite tiles will be replaced (reproduced with permission from Riccardo, V. J. Nucl. Mater. 2009, 390–391, 895–899). (b) Cross-section with allocations of materials. Image courtesy of EFDA-JET. Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices 651 has posed a significant engineering challenge. Calcu- lations for the main limiter tile types clearly show that a tile of the size of the existing CFC tiles would give unacceptable eddy torques, leading to the inves- tigation of several slicing options.181 The chosen design has vertical slices with a large central block and one to three side slices, depending on the toroidal extent of the tile assembly, supported on a carrier via pins (see Figure 18). The design is defined by the balance between conflicting requirements of eddy cur- rents (avoidance of large low resistance loops) and power handling (minimum number of vertical cuts to be shadowed). The problems associatedwith the design of the JET beryllium tiles (power handling capacity and disruption induced eddy currents) are discussed in detail elsewhere (see, e.g., Thompson et al.182). The installation of the new ITER-like wall and the NB enhancement has been completed by the mid of 2011 and operation is now restarting to provide important information for ITER. 4.19.6.1.2 Design of the beryllium ITER wall At the present time, the ITER first wall and shielding blanket is still undergoing a major redesign to over- come some of the main design shortcomings that were identified in the context of design review conducted in 2007; for example, the thermal load requirements were updated, in light of experimental experience.16 Most important in this respect was the recognition that the upper X-point region would see much higher loads during burn than 0.5MWm�2; long transients (approximately up to 5–10 s) of plasma contact with the wall would have to be with- stood. In addition, NB shine-through at low densities would necessitate high heat flux first-wall protection, and a new requirement has been introduced to pro- vide remote maintainability of the first-wall panel to be done in situ and independently of the shield mod- ule (which would also have to be maintainable). The rationale for the ongoing effort is described by Lowry et al.156 Proposed design modifications are being developed while trying to avoid and minimize changes to other components which are on the criti- cal fabrication path, especially the vacuum vessel, which is under fabrication. The main features of the proposed design are the following: (1) to abandon the port-limiters and to exploit the first wall for plasma startup by relying on more benign plasma start-up scenarios, including an early X-point formation; (2) to use suitably shaped plasma-facing surfaces to hide edges such that there is no illumination of component surfaces by Be slices ‘Toast rack’ carrier Support pins Halo current path through the tile Halo current from plasma JG06.336-7c Figure 18 Inner wall guard limiter tile (exploded view, top, and prototype, bottom). The five castellated Be slices have interslice and outer slice internal toroidal edges ski-slope shadowed. The slices are held on an inconel carrier by pins which allow bowing under thermal load. The RH bolts are designed to be shadowed by the next installed tile. Reproduced with permission from Riccardo, V. J. Nucl. Mater. 2009, 390–391, 895–899. 652 Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices normal or near-normal field lines, emanating in the near SOL; (3) provide power load capability of 4–5MWm�2, in order to be able to use the first wall as a limiter for startup and termination; and (4) with- standing transients is still subject to discussion. In particular, it must be noted that the vulnerability to damage induced by thermal transients is recognized and linked to the feasibility and efficiency of all processes required for full remote maintenance of the first-wall panels, which is yet to be demonstrated. In practical terms, the approach adopted is to pro- vide a shadowed poloidal band in the center of the first-wall panel, the two sides being shaped in a form typical for limiters both to provide the shadowing of the band and to ensure that the toroidally facing edge of the first-wall panel is shadowed. Because of the port regions on the low-field side, which contain a variety of structures with varying power handling capabilities, and because the toroidal field ripple is variable with the toroidal field, it is not possible to exploit the entire wall surface in this location. For this reason, the first- wall panels on the low-field side have the poloidal bands between the ports advanced with respect to those in line with the ports (see Figure 19). The amount of set back required at the edges of the first wall is determined by the penetration angle of the field lines and the power scrape-off length, with the optimi- zation taking into account the differing power handling capability of the front face and the edge of the first wall. Considerations discussed here are limited to some problems associated with the design of the beryllium tiles and prediction of PWI effects during operation in ITER. An important design driver for the first wall in the past was the specification of the thermal load during off-normal transient events.3 In particular, the thick- ness of the beryllium tiles had to be such as to prevent overheating of the joints and possible damage of the coolant pipes (see Section 4.19.6.2.2). Also, the thickness of the tile determines the temperature Be wall (a) (b) CFC strikepoints W elsewhere Figure 19 (a) View of the low field side first-wall surface showing how the panels in line with the port openings are recessed with respect to those between. It also shows the shielded central section of the panels allowing for access to the mechanical and hydraulic connections. Reproduced from Hawryluk, R. J.; et al. Nucl. Fusion 2009, 49, 15, 065012; with permission from IAEA. (b) Allocation of armor materials. Reproduced from Hawryluk, R. J.; et al. Nucl. Fusion 2009, 49, 15, 065012; Federici, G.; Loarte, A.; Strohmayer, G. Plasma Phys. Contr. Fusion 2003, 45, 1523–1547, with permission from IOPP. Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices 653 gradient and the thermal stress under a prescribed thermal load during steady-state. Limits on the tile temperature during operation arise as a result of many processes including melting, excessive vapori- zation, thermal fatigue, reduced mechanical integrity, and chemical reactions during accidental exposure of armor or structure to air or steam. The last one of the above processes is important as explosion of hydro- gen liberated from the steam–Be reaction is a major concern. In the past, a tile thickness of 10mm was adopted. This corresponded to a Be maximum tem- perature limit of �650–750 �C, roughly the level at which the relevant Be material properties (including mechanical, embrittlement, thermal fatigue, and swelling effects) start to degrade considerably. Because of the differences in the product of the elastic modulus and the coefficient of thermal expan- sion (E) between beryllium and copper or copper alloys (EBe/ECuCrZr¼ 2.4), large thermal stresses are set up around the bond between the beryllium tile and copper allot heat-sink. The difficulty to successfully join low thermal expansion armormaterials such as beryllium and tung- sten to high thermal expansion heat sink materials has been a major problem and has been discussed in Section 4.19.5. Thermomechanical modeling has shown the desirability of using very small tiles of brush like structure for PFC armor because of the reduction of the stress at the armor–heat sink interface. The proper selection of the size of beryllium tile is an important issue which impacts all aspects of compo- nent manufacturing such as increased cost of machin- ing, nondestructive examination features, reliability and repair of unbonded tiles, etc. In general, the fatigue life issue is difficult to quantify because of a number of factors. The thermal stresses depend on the temperature profile and the degree of constraint in the tiles. Tile castellations must be introduced to further relieve the constraints, and these have been sized following an extensive program of coupled thermal and mechanical analyses using finite elements codes such as ANSYS185 and ABAQUS.186 4.19.6.2 Predictions of Effects on the ITER Beryllium Wall During Operation 4.19.6.2.1 Safety issues in ITER 4.19.6.2.1.1 In-vessel tritium inventory Estimates of the tritium inventory and of permeation in the PFCs of a magnetic fusion device are impor- tant for assessing the radiological hazards from rou- tine operation and from potential accidents, for the design of the water detritiation system, and for pre- dicting the tritium supply requirements. In addition, these estimates have contributed to the decisions involving the choice of different armor materials in ITER options, which have a strong impact on tritium retention. In spite of the experimental and modeling progress which has taken place in the recent past, understanding of the subject of tritium–wall interac- tions is still far from complete and quantification of the tritium inventory in ITER is highly uncertain. The retention and permeation of implanted tritium in ITER PFCs have been widely studied in the past (see Section 4.19.3 and the example of calculations found elsewhere).9,187–189 On the basis of the results of these calculations, it can be concluded that the inventory of tritium in the beryllium first wall of a device like ITER, due to implantation, diffusion, trapping, and neutron-induced transmutation, will be on the order of 100 g rather than the kilogram quantities estimated previously70,100 and most of that will come from neutron-induced transmutations in the Be itself. The dominant process for long-term retention of tritium in beryllium for ITER is expected to be 654 Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices codeposition (see Section 4.19.3.2.2) with eroded wall material (i.e., the incorporation of tritium in the deposited layers where impurity atoms or mole- cules are deposited together with eroded material and a flux of energetic or thermal atoms). The inven- tory of this potentially volatile tritium must be kept as low as reasonably achievable (�1 kg tritium), in order to minimize the impact on the environment in case of an accidental release, in particular to avoid the evacuation of the neighboring population. The rate of formation of the codeposited material depends on the energy of the incident particles and on the substrate temperature during the deposition. In ITER, the total amount of tritium trapped in the codeposited layers will strongly depend on whether carbon is retained in the divertor during DT opera- tion. But even in a full metal ITER configuration (e.g., with Be wall and W divertor) there is evidence for potential tritium accumulation for ITER in deposited Be layers.9 In contrast to carbon, tritium codeposition in beryllium layers is expected to be released at relatively low temperature and there are provisions to periodically bake the divertor in ITER at 350 �C to release tritium trapped in the codepos- ited Be layers (see Section 4.19.3.2.2). 4.19.6.2.1.2 Chemical reactivity of beryllium dust with steam in ITER Although not a concern in present day tokamaks, in-vessel dust and tritium inventories have been recognized as a safety and operational issue for next step devices such as ITER.190–193 In particular, accident scenarios that result in water or steam exposure of hot plasma-facing mate- rials are one of the greatest concerns for ITER, because steam interacts with hot beryllium leading to the production of hydrogen, and hydrogen in the presence of air can lead to an explosion. The steam-chemical reactivity of different grades of Be has been studied extensively in the past.194–200 The amount of hydrogen produced depends on the specific material, temperature, exposure time, and especially the effective surface area. Because of the large surface area of dust, its chemical reactivity is an issue. Dust is expected to be produced inside the vac- uum vessel of a tokamak by interaction of the plasma with the components of the first wall and the divertor. A detailed discussion of the mechanisms of dust pro- duction and of the influence of parameter variations is beyond the scope of this contribution, but it should be noted that the processes and the production rate of dust are not fully understood and the extrapolation of knowledge from existing tokamaks to ITER is difficult. Research into dust production mechanisms and rates, the appropriate dosimetric limits for per- sonnel exposure, and methods of removal has only recently begun.201,202 The location where the dust settles will determine its temperature, and consequently, its chemical reac- tivity. At the moment about 6 kg of C, 6 kg of W, and 6 kg of Be dust are allowed ‘on hot surfaces’ in ITER, with these limits set by the H production risk. This corresponds to the maximum allowable quantity of H (2.5 kg) for the vessel integrity to be guaranteed in case of explosion. A complete oxidation of Be at 400 �C and C at 600 �C is assumed for the calculation. If no C is present in the machine, the limits are relaxed to 11 kg for Be, or 230 kg for W. These quan- tities are set such that the overall hydrogen combus- tion limit is not exceeded.9 It must be recognized that a limit on the order of �10 kg for beryllium dust on ‘hot-surfaces’ is very restrictive, and in particular, the development of diagnostics techniques that can determine from local measurements the global inventory in the machine could prove to be very challenging.203 How- ever, it is also likely that dust in ITER produced by Be eroded from the wall and deposited on the divertor will not survive on plasma-facing surfaces exposed to heat fluxes and will tend to accumulate in grooves or castellations in the armors of PFCs. They are an essential feature of the design of PFCs to relieve stresses during cyclic high heat flux load- ing, thus maximizing the fatigue lifetime of the armor to heat-sink joint. Some reduction in reaction rates is expected because the steam supply is not unlimited and steam must diffuse through the dust in the grooves. Experiments have been carried out in the Russian Federation, both in the Bochvar Institute of Moscow and the Efremov Institute of St. Petersburg.204 Although not conclusive, the main results summarized in Figure 20, show a reduction of the measured Be steam reactivity, particularly at high temperatures (more than a factor of 20). How- ever, further experimental and modeling work is needed to clarify if the observed slower kinetics at high temperatures (800–900 �C) eliminates the risk of explosion in the event of an accident. 4.19.6.2.2 Erosion/damage of the ITER Be wall The erosion mechanisms that affect the erosion/ damage of the first wall in ITER are (1) sputtering erosion by D–T ions and charge-exchange neutrals Inverse temperature ((1 K- 1) ´ 10 000) H 2 ge ne ra tio n ra te (l m - 2 s- 1 ) 6 1200 1000 800 700 600 500 400 ºC 1E - 08 1E - 07 1E - 06 1E - 05 1E - 04 1E - 03 1E - 02 1E - 01 1E + 00 1E + 01 1E + 02 7 8 9 10 11 12 13 14 15 16 Courtesy of V.Filatov (Efremov) 17 (1) Pressed powder in grooves (Efremov Institute). (2) Nonpressed powder in grooves (Bochvar Institute). (3) Powder on open surfaces (Efremov Institute). (1) (2) (3) INEL 91-GV-P INEL 92-GV-P INEL 91-GV-PS INEL 92-WG-DD INEL 92-GV-D,D INEL 96-WG-D,D INEL 96-GMS-D,D INEL 96-WG-C,D INEL 96-GMS-C,D INEL 96-GMS-I INEL 96-WG-I INEEL 97-PSA-G INEEL 97-PSA-WG INEEL 97-PSB-G INEEL 97-PSB-WG INEEL 97-RA-G2 INEEL 97-RA-WG2 INEEL 97-RA-G1 INEEL 97-RA-WG1 Figure 20 Initial reactivity of Be powder in grooves with steam in recent experiments carried out at Efremov and Bochvar Institute in the Russian Federation. (Be powder: BET-0.38m2 g�1, average partial size¼15mm, free (nonpressed) dust density¼0.7 g cm�3). Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices 655 during normal operation; and (2) evaporation and loss of melt layers during off-normal transient events such as thermal quench disruptions, ELMs, VDEs, and runaway electrons impact. There are additional localized erosion phenomena such as arcing, over- heating with evaporation, and, possibly, loss of melt layer on exposed edges, but it is very difficult to make predictions of these effects for ITER. Special design attention has been given to avoid the misalignments of PFCs and avoid thermal overloading with possible localized damage. 4.19.6.2.2.1 Erosion of Be wall during normal operation Calculations have been done to compute erosion of the first wall (due to fuel charge-exchange neutrals and ions, and impurity ions).205,206 It was found that about 20–40 g of Be per 400 s discharge are eroded from the wall with a beryllium peak erosion rate of the order of 0.1 nm s�1. These predictions are con- firmed by extrapolation of experimental data from JET.180 This erosion rate would be acceptable from a component lifetime standpoint, especially during the low duty-factor operation of ITER. However, the total amount of eroded material may be significant. This material will most likely go to the divertor, and this will affect the composition of the divertor sur- face; therefore, it will affect the divertor performance and contribute to tritium codeposition and dust inventories. Modeling of the influx of the eroded beryllium on the divertor is in progress to extrapolate from present machines and, in particular, to account for effects arising from material mixing including codeposition as expected in ITER. Several studies have been recently published on this subject (see, e.g., Kirschner et al.207,208). 4.19.6.2.2.2 Erosion of the beryllium wall during ELMs Depending on the actual energy flux on the Be PFCs in ITER during ELMs, melt damage may or may not occur. For Type I ELMs, which are compatible with the ITER divertor lifetime (�10MJ ELMs16,18), the expected energy flux on the main chamber in ITER will be in the range of 2–3MJ. The area of the wall over which this flux will be distributed is �30–60m2, for a toroidally symmetric energy deposition. This leads to ELM energy fluxes �0.02–0.08MJm�2 on the main chamber wall, which will cause no Be melting at all. If toroidal asymmetries and/or poloi- dal structures dominate the ELM energy deposi- tion on the first wall, a substantial reduction of the first-wall effective area for energy deposition is expected (by a factor of �5). In this case, the ELM energy fluxes on the first wall would be 0.1–0.4MJm�2, which can cause up to 18mm of 656 Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices melting, lasting �300ms.209 Figure 21 shows the results of an analysis carried out with the code described in Raffray and Federici.210,211 The erosion lifetime, expressed in number of ELMs or corresponding ITER full power pulses (approxi- mately 700 ELMs/pulses for a Be target initially 10mm thick) is found to sharply decrease above a certain ELM energy threshold. Depending on the duration of the ELM event, the threshold energy density varies between 0.2 and 0.7MJm�2. For com- parison, the results of a W wall are also shown. More recently, analysis has been carried out using more sophisticated modeling tools and the results are described elsewhere.212 From the JET Be divertor experience, we expect that only a very small part of the melt layer produced during each ELM will be mobilized (typically 120 100 80 60 40 20 0 0 5 10 15 20 0 (a) (b) 0.5 Be(1) Be(1) W(1) W(1) Be(2) Be(2) W(2) W(2) ´ 10 Deposited energy (MJ m-2) Deposited energy (MJ m-2) M el t th ic kn es s (m m ) E va p or at ed t hi ck ne ss (m m ) 1 21.5 0 0.5 1 21.5 Figure 22 Calculated amount of material (a) melted and (b) evaporated during (1) a 0.1ms and (2) a 1 ms plasma disruption for beryllium and tungsten. Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices 657 the energy deposited by the plasma and cooling by vaporization of beryllium. The vaporization temper- ature of beryllium is variously reported as 2480– 2979 �C, as compared to over 5630 �C for tungsten. Similarly, the latent heats of melting and vaporization of Be are also lower than the corresponding tungsten values. This explanation is consistent with the results in Figure 22(b), which shows the amount of material that is vaporized for a thermal quench time of (1) 0.1ms and (2) 1ms. At an energy density of 1MJm�2 and for a time of 0.1ms the thickness of vaporized material is 10mm for beryllium and 2.5mm for tungsten. It must be noticed that vapor shielding is not included in these calculations and that the results therefore should be considered conservative. High-pressure noble-gas-jet injection, for exam- ple, of neon and argon, has shown to be a simple and robust method to mitigate the deleterious effects of disruptions in tokamaks.215 The gas jet penetrates the central plasma at its sonic velocity. The deposited species dissipate >95% of the plasma energy by radiation and substantially reduce mechanical stress on the vessel caused by poloidal halo currents. Nevertheless, there remains some concern that even mitigated disruptions could damage the Be wall in ITER. Preliminary calculations show that even during a mitigated disruption in which the plasma energy is intentionally dissipated by radiation in �1ms by disruption mitigation techniques, the entire first wall of beryllium can melt to a depth of roughly 20–50mm.212,216 The fate of this melted layer is uncertain. If the melt layer resolidifies, it provides a means of removing the oxide layer and creating a clean Be layer for oxygen gettering. On the other hand, if significant j �B forces associated with the plasma termination mobilize the melt layer within the vessel, it will likely lead to operational difficulties. Another area of possible concern is the small surface cracks that form when molten metals resolid- ify. These resolidification cracks could serve as ther- mal fatigue crack initiation sites and accelerate this type of damage. While this effect has not been exten- sively studied because of the difficulty of simulating disruptions in the laboratory, it may not be a critical issue as thermal fatigue cracks form after a few hundred cycles in most materials and they grow to depths only where the thermal stress level is above the yield stress. 4.19.6.2.2.4 Erosion of the beryllium wall during VDEs At some locations (mostly at the upper inboard region and the lower region of the first wall), the Be armored PFCs must withstand a certain number of ‘slow’ thermal transients resulting from loss-of- control of plasma position during VDEs. Typical parameters for these events are 60MJm�2 over 0.3 s. In contrast to thermal quench disruptions, VDEs lead not only to significant erosion or melting, but also to high heat fluxes and a subsequent temper- ature increase at the armor/heat sink interfaces that can result in a failure of the armor/heat sink joints.217 As a matter of fact, because of their short duration ( 658 Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices The erosion due to VDEs in a device like ITER has been modeled by various authors,221 whereas the results of analysis carried out to quantify the effects on PFCs resulting from runaway electrons can be found in Raffray et al.222 As an example, Figure 23 shows the surface temperature of a 5mm copper 10 mm W, Be, or CFC on a 5 mm Cu substrate Tmett W CFC Be 0 C op p er s ur fa ce t em p er at ur e (K ) 400 600 800 1000 1200 1400 2 4 Time (s) 6 8 10 VDE 60 MJ m-2 0.3 s Plasma-facing component design (Carbon, beryllium, or tungsten coating on copper structure) Graphite, Be, or W Copper substrate 5–10 mm 5 mm Figure 23 Interface copper surface temperature rise during a vertical displacement event for different surface coating materials. Reproduced with permission from Federici, G.; Skinner, C. H.; Brooks, J. N.; et al. Plasma-material interactions in current tokamaks and their implications for next-step fusion reactors. Nucl. Fusion 2001, 41, 1967–2137 (review special issue), with permission from IAEA. Table 5 Vaporization and melting thickness (mm) and temp shielding effect) Material/thick (mm) Beryllium/5mm Energy density (MJm�2) 30 60 Time (ms) Vap. (mm) Melta (mm) Tjoim b (�C) Vap. (mm) Melta (mm) 10 4.03 309.9 163 8.44 648.2 100 3.03 821.6 217.1 7.5 812.3 aOnly stationary melt thickness (without splashing). bTemperature at the interface armor/ heat sink. substrate at its interface with a tungsten, beryllium or carbon tile of 10mm thickness during a typical VDE releasing about 60MJm�2 to the surface in 300ms.7 Tungsten and carbon armors of similar thickness usually result in a similar and higher cop- per surface temperature than that of beryllium armor of the same thickness. This is because most of the incident plasma energy is removed by the beryllium’s higher surface vaporization rate, which leaves little energy to be conducted through the structural material. In order to reduce the tempera- ture at the copper interface, thicker tiles would be required. Only beryllium tiles of reasonable thickness (20mm) can withstand the acceptable temperature rise in the copper structure for the conditions shown. The coolant flux and, conse- quently, the Be/Cu interface temperature increase with decreasing Be thickness. The evaporated and melting thickness and temperature at the Be/Cu alloy interface during each VDE is shown in Table 5 for Be tiles (5 and 10 mm thick); for two values of the VDE energy density (30 and MJ m�2) and for two VDE durations (10 and 100 ms). 4.19.6.2.2.5 Erosion of the beryllium wall during runaway impact Compared to disruptions, the thermal effects from runaway electrons are confined to a much smaller area, but the localized damage is expected to be more severe and can cause severe melting/vaporization in virtually all materials and can lead to surface spallation. These events have been observed to cause severe damage to graphite tiles in present day tokamaks. While the beryllium in the strike region will probably be severely melted, the most critical issues for runaway electron damage and VDE are damage of coolant pipes with resulting risk of water spillage. Because of the deep erature at the Be/Cu alloy during each VDE (w/o any vapor Beryllium/10mm 30 60 Tjoim b (�C) Vap. (mm) Melta (mm) Tjoim b (�C) Vap. (mm) Melta (mm) Tjoim b (�C) 163 4.04 318.6 169.6 8.46 652 169.6 233.1 3.08 861.5 170.2 7.56 870.2 170.5 Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices 659 penetration and large spatial dispersion of the high- energy electrons, a thick armor may be required to avoid overheating of the coolant channels with subsequent coolant leakage. As thicker armor implies higher surface temperatures, the best solution may be local regions that are either uncooled or with thick armor that receives low heat flux during normal opera- tions. Typically a runaway electron energy deposition transient of 50MJm�2 over 0.3 s on the Be first-wall modules results in a maximum heat flux to the coolant of �7.4MWm�2, a maximum Cu alloy temperature �640 �C, and a Be melt layer thickness �1.8mm.222 A reduction of the armor thickness will lead to an increase in the maximum Cu alloy temperature and could lead to the damage of joints. 4.19.6.3 Prospect of Using Beryllium in Beyond-ITER Fusion Reactors The main differences between a future power reactor and ITER are the much longer operation time (e.g., >108 s vs. >107 s), high duty cycle, and the higher temperature of the fluid to cool the PFCs to maintain a high plant energy conversion efficiency. The higher surface temperature of the PFC will affect particle recycling, tritium uptake, chemical erosion, and material-mixing effects. Erosion rates at the divertor target are very diffi- cult to predict in conventional fusion power plant concepts with solid high-Z targets because the net erosion or deposition is strongly dependent on plasma parameters. The fraction of ions arriving above the sputtering threshold is crucial, as is the efficiency of the prompt local redeposition. ELMs have not really been considered in this context but we can see from the analysis in Section 4.19.6.2.2.2 that ELMs in power plant systems will have to be extremely small – much smaller than will be allow- able in ITER, which still has a relatively low duty cycle. The ideal would in fact be a quiescent ELM free high density steady state edge plasma. Calculations of the minimum erosion rate for the main wall are somewhat more robust as there has to be a hot plasma in the main chamber and the rate of leakage of neutrals into the main chamber from the divertor can be calculated using Monte-Carlo codes. In a recent study, Be, C, Fe, Cu, Mo, and W walls were compared,206 with the conclusion that in all cases the erosion rate was 1–2 t per year of continuous operation. A Be or CFC wall will erode too rapidly in a reactor and the large amount of eroded material might give rise to deleterious problems as far as control of the tritium and dust inventories are concerned. A medium-Z material, such as Fe, does not seem to be acceptable purely from the standpoint of erosion lifetime. As molybdenum is unfavorable for long-term activation problems, W is the best and only solution we have available for a reactor. Effects of plasma contamination from Mo and W at the wall of tokamaks are being addressed in Alcator C-Mod, ASDEX-Upgrade and in the near-future at JET. There is considerable gross erosion by sputtering for all materials. The contributions of ions and neutrals from the plasma to this erosion are of the same order of magnitude. The integrated total erosion due to ions and the energetic neutrals for the different wall mate- rials (Figure 24) show that because of the larger sput- tering yields for the low-Z materials, the number of atoms eroded for these materials is a factor of 10–20 larger than for high-Z materials such as W. However, the total mass loss is similar for all materials, up to several kilograms per day or about 1 t per year. The maximum wall thinning for the low-Z mate- rials is about 3.5mmyear�1, while for high-Z materi- als, such as W, it is 0.22mmyear�1, that is, lower by about a factor of 15. These values are in reasonable agreement with erosion measurements at the JET vessel walls.223 With respect to wall thinning, W is favorable for the use at the vessel walls because it has the longest ‘erosion lifetime’ (Figure 24(b)). With respect to plasma contamination, the probability of the eroded atoms entering into the plasma core, their lifetime in the plasma core, and the tolerable concen- tration of these ions in a burning fusion plasma all have to be taken into account.206 The tolerable con- centration of W in the plasma is nearly three orders of magnitude lower than for low-Z atoms, such as Be and C. However, recent observations have shown that W can be effectively removed from the plasma center by central heating.224 As this central heating is natural for burning plasmas, W may be a possible plasma- facing material, even from the viewpoint of plasma contamination. The ion and neutral flux densities on the vessel walls are of the order of 1020m�2 s�1, which may be critical with respect to the tritium implantation, accumulation in and permeation through the vessel walls. 4.19.7 Concluding Remarks Beryllium is a low-density metal that is used in a number of industries, including the nuclear, automo- tive, aerospace, defense, medical, and electronics Be Be Fe Fe Cu Mo W W Mo Cu Ctot C Cphys Cchem 1020 1021 1022 0 (a) (b) 10 Atomic number of the material Ti m e (y ea r) t o er od e 5 m m W al l e ro si on (a to m s s- 1 ) Atomic number of the material 20 30 40 50 60 70 80 0 0 5 10 15 25 20 10 20 30 40 50 60 70 80 Figure 24 (a) Integrated gross erosion due to ions and the energetic neutrals. (b) Upper estimate for the time until a thickness of 5mm is eroded by sputtering at the area of largest erosion, that is, at the reference distance of about 11m. Reproduced with permission from Behrisch, R.; Federici, G.; Kukushkin, A.; Reiter, D. J. Nucl. Mater. 2003, 313–316, 388–392. 660 Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices industries, for various applications because it is exceptionally strong, is light in weight compared with other metals, has high heat-absorbing capability, and has dimensional stability in a wide range of temperatures. Beryllium has been considered for many years as a primary candidate for protection of PFCs in toka- maks because it offers distinct advantages when com- pared with alternative materials such as carbon and tungsten. It has a low atomic number and is an excel- lent oxygen getter. The interaction of beryllium with tritium is also significantly weaker than that of carbon, leading to potentially reduced tritium inven- tory. Beryllium does not form stable hydrides above 300 �C, so there should be very little trapping expected in codeposited layers formed at such tem- peratures in the divertor after sputtering, although work is still underway to clarify this problem. How- ever, beryllium has a relatively high physical sputter- ing rate and a relatively low melting temperature and as such is more susceptible to melting damage that may occur in a tokamak during thermal transients. In addition, because of its toxicity, special precautions are needed for working with beryllium, either for manufacturing or research investigation purposes. Beryllium has been used with success in various tokamaks in the past mainly because of its ability to getter oxygen and to improve plasma performance. In particular, its successful deployment in JET that started in 1989 and is continuing today with the installation of a completely new beryllium wall is the main rationale for the selection of beryllium as a plasma-facing material for the first wall of ITER, on the basis of a combination of plasma compatibility and design considerations. This paper reviewed the properties of beryllium that are of primary relevance for plasma protection applications in magnetic fusion devices (i.e., PWIs, thermal and mechanical properties for power handling, fabricability and ease of joining, chemical reactivity, etc.) together with the available knowledge on performance and operation in existing fusion machines. Special attention was given to beryllium’s erosion and deposition, formation of mixed materials, and its hydrogen retention and release characteristics. These phenomena have a profound impact on component design, machine operation, and safety. Extensive data on the behavior of Be with plasmas have been col- lected from existing tokamaks and simulators during the last two decades and this has enabled great strides to be made in our understanding of the PWI pro- cesses involved. However, there are many issues for which there are still uncertainties and we will only learn from operating the next two major experiments that foresee the use of large amounts of Be ( JET and ITER). Much work remains to be done in this area and more machine operational time and diagnostics dedicated to PWIs are required. Initiatives on these fronts, together with modeling of the results, are essential to advance the understanding of PWIs. This includes (1) the possible surface damage (melt- ing) during transients such as ELMs and disruptions and its implications for operations and (2) the prob- lem of beryllium mixing with other armor materials and in particular the issue of codeposition of tritium with Be, which is eroded from the first wall and deposited at the divertor targets. Such material may also be locally redeposited into shadowed areas of the shaped ITER first wall. Both issues are part of Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices 661 ongoing research, the initial results of which are being taken into account in the ITER design so that the influence of these two factors on ITER operation and mission are minimized. For example, ITER will very likely employ, ELM control systems based on pellets and RMP coils, disruption mitigation systems, and increased temperature baking of the divertor to release T from Be-codeposits. Dust generation is still a process which requires more attention. Conversion from gross or net erosion to dust and the assessment of dust on hot surfaces need to be investigated. At the time of writing this paper, the ITER first wall and shielding blanket is undergoing a major rede- sign effort to overcome some of the main shortcomings that were identified in the context of a recent design review scrutinizing the internal components. Complex and interrelated materials, manu- facturing, and design issues were briefly reviewed in this paper together with the progress of the manufacturing technologies being used and tested to demonstrate the durability of the joints. A critical feature of the ITER first-wall design is the beryllium to copper alloy bond. The joints must withstand the thermal, mechanical, and neutron loads and the cyclic mode of operation, and operate under vacuum, while providing an acceptable design for lifetime performance and reliability. The availability of reli- able joining technologies has a large impact on the design of the PFCs and on the overall cost of these components. The status of the available techniques presently considered to join the Be armor to the heat sink material of Cu alloys for the fabrication of Be-clad actively cooled components for the ITER first wall was discussed. During earlier ITER design phases, the feasibility of manufacturing reliable Be–CuCrZr joints was demonstrated. The results of the perfor- mance and durability heat flux tests conducted in the framework of the further ITER first-wall qualifica- tion program were described. This program has been launched and is in progress in the ITER parties in order to qualify the design and manufacturing routes. The integrity of this bond must be assured for reliable ITER performance whatever process is used to fabri- cate joints. The original procurement sharing that assigned the fabrication of first-wall panels up to six parties was seen as a risk and the number of parties supplying these critical components has now been reduced to three, Europe, the Russian Federation, and China. The selection of specific grades of specific beryl- lium for the ITER first wall was described. The effects of neutron irradiation on the degradation of the properties of beryllium itself and on the joints were also analyzed. Some of the changes are impor- tant while others are not significant for the ITER conditions. Change of thermal conductivity and swelling are not important because of the low fluence. The bulk tritium retention in neutron irradiated Be is expected to be significantly less than tritium reten- tion in the codeposited layers. The most critical consequence of neutron irradiation under ITER con- ditions is embrittlement. This is typical of all grades of beryllium. The structural integrity of neutron irradiated brittle Be is a key issue. Embrittlement of neutron-irradiated Be could lead to increased ther- mal erosion and crack formation, which is also observed to occur for unirradiated beryllium under severe transient heat loads. These cracks could serve as thermal fatigue crack initiation sites and accelerate this type of damage. While this effect has not been extensively studied because of the difficulty of simu- lating disruptions in the laboratory, it may not be a critical issue as thermal fatigue cracks form after a few hundred cycles in most materials and they grow only to depths where the thermal stress level is above the yield stress. On the basis of the information available from exist- ing fusion machines, we discussed the problems that are still at issue in the design and operation of ITER. This includes, in particular, the problem of erosion/ damage and the problem of up-take and control of tritium in the beryllium-based codeposited films. Finally, on the basis of these results some tentative and speculative consideration of the limited prospects that beryllium has in future reactors was offered. The worldwide fusion energy research over the last four decades has developed a tremendous amount of knowledge on plasma physics and related technolo- gies. From this point of view, collecting the latest information from a wide range of studies is important in order to help the fusion community to recognize the critical issues and the status. That has been the intent of this chapter. (See also Chapter 4.17, Tungsten as a Plasma-Facing Material and Chapter 4.18, Carbon as a Fusion Plasma-Facing Material). Acknowledgments The views expressed in this publication are the sole responsibility of the authors and do not necessarily reflect the views of Fusion for Energy. Neither Fusion for Energy nor any person acting on behalf 662 Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices of Fusion for Energy is responsible for the use which might be made of the information in this publication. The authors from the ITER Organization wish to acknowledge that this paper was prepared as an account of work by or for the ITEROrganization. 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Zinkle Oak Ridge National Laboratory, Oak Ridge, TN, USA Published by Elsevier Ltd. 4.20.1 Introduction 667 4.20.2 Copper and High-Strength, High-Conductivity Copper Alloys 668 4.20.2.1 Pure Copper 668 4.20.2.2 PH Copper Alloys 668 4.20.2.2.1 CuCrZr alloy 669 4.20.2.2.2 CuNiBe alloy 670 4.20.2.2.3 CuNiSi 670 4.20.2.3 DS Copper Alloys 670 4.20.3 Physical Properties of Copper and Copper Alloys 671 4.20.4 Mechanical Properties of Copper and Copper Alloys 671 4.20.4.1 Tensile Properties 671 4.20.4.2 Fracture Toughness 673 4.20.4.3 Creep 674 4.20.4.4 Fatigue and Creep–Fatigue 674 4.20.5 Irradiation Effects in Copper and Copper Alloys 675 4.20.5.1 Effect of Irradiation on Physical Properties of Copper and Copper Alloys 676 4.20.5.2 Effect of Irradiation on Mechanical Properties of Copper and Copper Alloys 676 4.20.5.2.1 Tensile properties 676 4.20.5.2.2 Fracture toughness 678 4.20.5.2.3 Fatigue and creep–fatigue 678 4.20.5.2.4 Irradiation creep and void swelling 678 4.20.5.3 Effect of Irradiation on Microstructure of Copper and Copper Alloys 681 4.20.5.3.1 Defect structure in irradiated copper and copper alloys 681 4.20.5.3.2 Dislocation channeling 684 4.20.6 Joining 685 4.20.7 Summary 687 References 688 Abbreviations CW Cold worked DS Dispersion strengthened FFTF Fast Flux Test Facility G-P Guinier–Preston HIP Hot isostatic pressing IACS International Annealed Copper Standard JET Joint European Torus MOTA Materials Open Test Assembly OFHC Oxygen-free, high conductivity PH Precipitation hardened SAA Solution annealed, and aged condition SFT Stacking fault tetrahedral TCH Tension and compression hold 4.20.1 Introduction Copper alloys are prime candidates for high heat flux applications in fusion energy systems. High heat flux is a major challenge for various fusion devices because of the extremely high energy density required in controlled thermonuclear fusion. The removal of a large amount of heat generated in the plasma through 667 668 Physical and Mechanical Properties of Copper and Copper Alloys the first wall structure imposes a major constraint on the component design life. Materials with high con- ductivity are needed to assist heat transfer to the coolant and to reduce the thermal stress for pulsed mode of operation. A number of issues must be considered in the selection of materials for high heat flux applications in fusion reactors. While high conductivity is the key property for such applications, high strength and radiation resistance are also essential for the effective performance of materials in a high heat flux, high irradiation environment. In addition, fatigue behavior is a major concern for many high heat flux applica- tions because of planned or inadvertent changes in the thermal loading. Pure copper has high thermal con- ductivity but rather low strength, and therefore its application as heat sinks is limited. The strength of copper can be improved by various strengthening mechanisms. Among them, precipitation hardening and dispersion strengthening are the two most viable mechanisms for improving the strength of copper while retaining its high electrical and thermal con- ductivities. A number of precipitation-hardened (PH) and dispersion-strengthened (DS) copper alloys are commercially available, and have been evaluated for fusion applications, for example, PH CuCrZr, CuNiBe, CuNiSi, and DS GlidCop® Al15, Al25, Al60, MAGT-0.2, etc. Two copper alloys that are most appealing are PH CuCrZr and DS CuAl25. Surveys of copper alloys for fusion applications were conducted by Butterworth and Forty1 and Zinkle and Fabritsiev.2 In this chapter, a brief description of pure copper and several copper alloys of interest to fusion appli- cations is presented, followed by a summary of their physical and mechanical properties. The radiation effects on the physical and mechanical properties of copper and copper alloys as well as their irradiated microstructure are then discussed. Joining techniques for plasma facing components in fusion reactors are also discussed. 4.20.2 Copper and High-Strength, High-Conductivity Copper Alloys 4.20.2.1 Pure Copper Copper is widely used where high electrical or ther- mal conductivity is required. Pure copper is defined as having a minimum copper content of 99.3%. Copper with oxygen content below 10 ppm is called ‘oxygen- free.’ ‘Oxygen-free, high conductivity’ (OFHC) grade copper has room temperature electrical conductivities equal to or greater than 100% International Annealed Copper Standard (IACS),where 100% IACS¼ 17.241 nOm at 20 �C.3 Copper grades with the ASTM/SAE unified number system (UNS) designation C10100, C10200,C10400, C10500, andC10700 are classified as OFHC copper. Grades C10400, C10500, and C10700 have significant silver content, which creates activa- tion hazards. OnlyC10100 andC10200 are considered for fusion systems. The use of unalloyed copper is often limited by its low strength. Copper can be strengthened by various processes, for example, cold working, grain refine- ment, solid solution hardening, precipitation hard- ening, dispersion strengthening, etc. While these approaches can significantly increase the strength, they can also lead to a pronounced reduction in con- ductivity. The challenge is to design a material with the best combination of strength and conductivity. Cold work can significantly increase the strength of pure copper and has a relatively moderate effect on conductivity.4 However, cold-worked copper can be softened at relatively low temperatures (�200 �C) because of its low recrystallization temperature.5 A recent study has shown that ultrahigh-strength and high-conductivity copper can be produced by introducing a high density of nanoscale twin bound- aries.6 The tensile strength of the nano-grained cop- per can be increased by a factor of 10 compared to conventional coarse-grained copper, while retaining a comparable conductivity. The potential of high- strength, high-conductivity bulk nano-grained cop- per in nuclear energy systems, however, has not been widely explored. Alloying in copper can significantly improve mechanical strengths and raise the softening tempera- tures. However, additions of alloying elements also reduce electrical and thermal conductivity. Among the three alloying strengthening mechanisms, namely, solid solution hardening, precipitation hardening, and dispersion strengthening, solid solution hardening has the most detrimental effects on the conductivity4 and is the least favored mechanism to obtain high- conductivity, high-strength copper alloys. 4.20.2.2 PH Copper Alloys PH copper alloys are heat-treatable alloys. The high strength of PH copper alloys is attributed to the uniform distribution of fine precipitates of second- phase particles in the copper matrix. PH copper alloys are produced by conventional solution treatment Physical and Mechanical Properties of Copper and Copper Alloys 669 and aging treatment. Solution treatment produces a homogeneous solid solution by the heating of an alloy to a sufficiently high temperature to dissolve all solutes. The alloy is then quenched to a lower temperature to create a supersaturated condition. A subsequent aging treatment heats the alloy to an intermediate temperature below the solvus tem- perature, to precipitate fine second-phase particles. Precipitates not only give rise to high strength, but also reduce the solute content in the matrix, main- taining good conductivity. The strength of a PH alloy depends on particle size, particle shape, volume frac- tion, particle distribution, and the nature of the inter- phase boundary.7 Despite their ability to develop significant strength, PH copper alloys may be soft- ened substantially as a result of precipitation coars- ening (overaging) at intermediate to high service temperatures or because of recrystallization during brazing or diffusion bonding. Therefore, heat treat- ment and thermal processing histories can have a large influence on the strength and conductivity of this class of alloys. A number of commercial PH copper alloys have been investigated for applications in fusion design, for example, CuCrZr, CuNiBe, and CuNiSi. 4.20.2.2.1 CuCrZr alloy PH CuCrZr alloy is commercially available under several trade names, for example, Elbrodur® CuCrZr from KME Germany AG, Outokumpu Oy CuCrZr, Zollen CuCrZr, C18150®, Trefimetaux CuCrZr, MATTHEY 328® from Johnson Matthey Metals, and YZC® from Yamaha Co, Ltd. The chem- ical compositions of these alloys differ by a small amount, with Cr varying from 0.4 to 1.5% and Zr 0.03–0.25%. Low Cr content is to prevent the forma- tion of coarse Cr precipitates. The element, Zr, 50 nm (a) ( Figure 1 Representative weak-beam dark-field images showi water quenched, and aged, and (b) hot isostatic pressed, soluti improves the hardening by the enhancement of fine homogeneous precipitation and improves the ductility of the alloy by inhibiting intergranular fracture.8–10 CuCrZr-IG is the ITER grade with tighter specification for composition and heat treat- ment. CuCrZr alloys are available in different forms, for example, bars, tubes, wires, foils, sheets, and plates. Hot forming, brazing, and inert gas welding are applicable for component manufacturing. CuCrZr alloys are used in the conventional aged condition. The reference ITER heat treatment in- cludes solution annealing at 980–1000 �C for 1 h, water quench, and aging at 450–480 �C for 2–4 h.11 Typical microstructure of the prime-aged CuCrZr is shown in Figure 1(a). The alloy contains an equiaxed grain structure and uniformly distributed fine Guinier–Preston (GP) zones exhibiting primarily black dot contrasts and a small number of precipitates with lobe–lobe contrast. The number density of precipitates is on the order 1022m�3, with a mean diameter of �3 nm. A low density of micron-size Cr particles and grain boundary precipitate-free zones were also observed.12–18 CuCrZr is susceptible to overaging and recrystallization during prolonged exposure at elevated temperatures. Overaging of CuCrZr causes significant coarsening of grain struc- ture and fine precipitates. Li et al.14 reported a lower number density (�1.9� 1022m�3) of larger (�9 nm in diameter) precipitates with a mixture of coherent and incoherent particles after CuCrZr was hot iso- static pressing (HIP) treated at 1040 �C for 2 h at 140MPa followed by solutionizing at 980 �C for 0.5 h with a slow cooling rate of 50–80 �Cmin�1 between 980 and 500 �C, and final aging at 560 �C for 2 h (Figure 1(b)). The average grain size was >500 mm in comparison with �27 mm grain size in the prime-aged alloy. b) 50 nm ng precipitates in unirradiated CuCrZr (a) solutionized, onized, slow-cooled, and aged. 670 Physical and Mechanical Properties of Copper and Copper Alloys 4.20.2.2.2 CuNiBe alloy Copper–beryllium ( Physical and Mechanical Properties of Copper and Copper Alloys 671 MAGT 0.2 is a Russian alloy produced by SPEZS- PLAVCompany. It contains 0.17% Al, 0.05% Hf, and 0.09% Ti in the form of oxide particles.25,26 GlidCop contains Al-oxide particles only, while in MAGT alloy, there are Al-, Ti-, and Hf-oxide particles, and mixed Al- and Ti-oxide particles. A majority of the oxide particles in MAGT 0.2 are spherical in shape with a small fraction in the form of circular disks, with an average particle size of 6 nm.25,26 4.20.3 Physical Properties of Copper and Copper Alloys Physical properties of pure copper and copper alloys are quite similar in terms of the melting point, the density, the Young’s modulus, and the thermal expan- sion coefficient. Table 1 compares the room tem- perature physical properties of pure copper, PH CuCrZr, and DS CuAl25.2,27–29 Because PH copper alloys and DS copper alloys contain only a small amount of fine second-phase particles, the physical properties of these copper alloys closely resemble those of pure copper. The conductivity of copper and copper alloys is the most important physical property for their applications. The electrical conductivity of copper can be reduced by thermal vibration of atoms and crystal imperfections, for example, solute atoms, vacancies, dislocations, and grain boundaries. These different mechanisms have additive contributions to the increase in resistivity. As with other metals, the thermal conductivity of copper, kth, is proportional to the electrical conductivity, l, described by the Wiedemann–Franz law, that is, kth ¼ lLT ½1� where T is the absolute temperature and L is the Lorentz number. The electrical conductivity of pure copper is sensitive to temperature, and less sensitive to the amount of cold work and the grain size. The linear temperature coefficient for electrical Table 1 Physical properties of pure copper, PH CuCrZr, and DS CuAl25 Cu CuCrZr CuAl25 Melting point (�C) 1083 1075 1083 Density (g cm�3) 8.95 8.90 8.86 Thermal conductivity (Wm-K�1) 391 314–335 364 Elastic modulus (GPa) 117 123 130 resistivity in copper is dr/dT¼ 6.7�10–11OmK�1.30 Severe cold work can reduce the electrical conductiv- ity of copper by only 2–3% IACS. All alloying elements in copper reduce the elec- trical conductivity, and the amount of degradation depends on the type of element, the concentration, and microstructural form (e.g., solid solution, pre- cipitation, or dispersion). Figure 2 compares the strength and conductivity of copper and several types of copper alloys.31 4.20.4 Mechanical Properties of Copper and Copper Alloys 4.20.4.1 Tensile Properties The influence of test temperature, strain rate, and thermal–mechanical treatments on the tensile prop- erties of copper and copper alloys has been studied extensively. Figure 3 illustrates the effect of test temperature on the yield strength of pure copper (in the annealed condition), PH CuCrZr and CuNiBe alloys, and DS CuAl25.15–18,28,32–39 The strength of copper alloys decreases with increasing test tempera- ture. The decrease in strength is moderate up to 200 �C. Significant drops in strength occur at higher temperatures, except that the CuNiBe ATalloy shows a relatively small reduction in strength even up to 400 �C. Pure copper has the lowest yield strength. The tensile properties of pure copper strongly depend on the thermal–mechanical treatment and the impurity content.15–18,32,33 CuNiBe alloy has the highest strength over the entire temperature range.34 The tensile properties of PH copper alloys are sensi- tive to the thermal–mechanical treatments. CuCrZr in the solution-annealed, cold-worked, and aged con- dition (SAþCWþA) has superior yield strength at low temperatures relative to CuCrZr in the solution- annealed, and aged condition (SAA). However, the strength of CuCrZr SAþCWþA alloy drops more rapidly with increasing temperature.29,34–39 The yield strength of CuNiBe can be quite different, depending on the processing techniques. The tensile ductility of copper alloys also shows strong temperature depen- dence. The uniform elongation of the CuAl25 alloy decreases considerably as the test temperature in- creases, but increaseswith increasing test temperature above 400 �C. The CuNiBe AT alloy shows a moder- ate drop of uniform elongation below 200 �C, but a sharp drop in ductility at higher temperature.34 The uniform elongation of the CuCrZr alloy shows the smallest sensitivity to test temperature. Among 0 200 400 600 800 1000 1200 0 0 100 200 300 400 500 25 50 75 100 125 150 175 0. 2% y ie ld s tr en gt h (M P a) 0. 2% y ie ld s tr en gt h (k si ) Thermal conductivity (W m-K–1) Cu–2% Be (cold worked and aged) Cu–Ni–Be (thermomechanical treated) Cu–Ni–Be (cold worked and aged) Cu–Ni–Be (solutionized and aged) Cu–Al2O3 (cold worked) Cu–Cr–Zr (cold worked and aged) Cu–Al2O3 (wrought) Cu–Cr–Zr (solutionized and aged) Cu (cold worked) Cu (annealed) Cu–2% Be (cast and aged) Figure 2 Strength and conductivity of copper and copper alloys. After Li, G.; Thomas, B. G.; Stubbins, J. F. Metall. Mater. Trans. A 2000, 31A, 2491. 0 0 100 100 200 200 300 300 400 400 500 500 600 600 700 700 800 800 900 1000 CuCrZr, SAA (Zinkle and Eatherly,34 Zinkle,38 Singh et al. 39) CuCrZr, SA+CW+A (Piatti and Boerman,29 Fabritsiev et al.35 Fabritsiev and Pokrovsky36,37) OFHC Cu, (Singh et al.,32, Singh et al.,15–18 Singh and Toft33) CuAL25, (Zinkle and Eatherly34) CuNiBe, HT1, HT2, AT (Zinkle and Eatherly34) Y ie ld s tr en gt h (M P a) Temperature (�C) Figure 3 The yield strength of copper alloys as a function of temperature. 672 Physical and Mechanical Properties of Copper and Copper Alloys the three copper alloys, the CuCrZr alloy has the best ductility over the temperature range, and the ductility of the CuNiBe alloy is the lowest. Because of the sensitivity of mechanical properties to thermal–mechanical treatments in PH copper alloys, the strength of large components made of these alloys can be significantly lower. For example, during component manufacturing, CuCrZr often experiences additional thermal cycles, such as braz- ing, welding, or HIPing. While solution annealing Physical and Mechanical Properties of Copper and Copper Alloys 673 can be conducted during or after a brazing or HIPing process, rapid quenching is not feasible for large com- ponents, and a much slower cooling rate (e.g., furnace cooled or gas cooled) is applied in the manufacturing cycle. Significant reduction in strength due to slow cooling rates has been reported in CuCrZr.30,40–42 A slow cooling rate (50–80 �Cmin�1) and overaging at 560 �C/2 h significantly reduced theyield stress and the ultimate tensile strength, and tensile elongations of CuCrZr relative to prime-aged CuCrZr.14 Cooling rates>1200 �Cmin�1 are required to fully quench the Cu–Cr solid solution.43–45 The effect of strain rate on tensile properties for pure copper and PH CuCrZr and CuNiBe alloys as well as DS CuAl25 alloy was studied at temperatures of 20 and 300 �C.14,34,46 All three copper alloys are relatively insensitive to strain rate at room tem- perature. The strain rate sensitivity parameter of m (where sy ¼ Ce_mand C is a constant) is �0.01 for the CuAl25 alloy at room temperature. The strain rate sensitivity of this alloy increases significantly with increasing temperature as reflected by a strain rate sensitivity parameter of m� 0.07 at 300 �C. Stephens et al.47 reported a strain rate sensitivity parameter of m� 0.1 in the temperature range of 400–650 �C for CuAl25. A similar effect of strain rate on ultimate tensile strength was also observed on these materi- als.34,46 Edwards46 investigated the strain rate effect of copper alloys in air and vacuum, and found that 0 0 50 100 100 150 200 250 300 350 400 450 500 J Q (k J m −2 ) Tempe Figure 4 Fracture toughness data of PH CuCrZr, CuNiBe and testing in air or vacuum did not appear to change the strain rate dependence of the CuAl25 alloy, but that testing the CuNiBe alloy in air shifted the embrittlement to a lower temperature. 4.20.4.2 Fracture Toughness Fracture toughness data for PH copper alloys, CuCrZr and CuNiBe, and DS copper alloys, CuAl15 and CuAl25, are summarized in Figure 4.14,48–50 CuCrZr has the highest toughness, and CuNiBe the lowest among these alloys. The large scatter in mea- sured fracture toughness values for CuCrZr in differ- ent studies is likely due to different heat treatments, specimen geometryanddimensions, and testingmeth- ods. The temperature dependence of the fracture toughness in CuCrZr, while difficult to accurately define, shows an initial decrease with increasing tem- perature, and then a slight recovery at temperatures above 250 �C.The effect of thermal–mechanical treat- ment on fracture toughness of CuCrZr is insignificant in comparison with its effect on tensile properties.14 The minimum value of the JQ for unirradiated CuCrZr is as high as �100 kJm�2. The fracture toughness of DS CuAl15 and CuAl25 is significantly lower than that of CuCrZr, and shows a strong directional dependence. The toughness is higher in the L-T orientation than in the T-L orientation. The fracture toughness decreases 200 300 400 Black = CuAl15 or CuAl25 Red = CuCrZr Green = CuNiBe T-L, L-T ---------------------------------------------- : Tahtinen et al.50 : Alexander et al.48 : Alexander et al.48 : Alexander et al.48 : Li et al.14 : Li et al.14 ------------------------------------------------ : Suzuki et al.49 rature (�C) DS CuAl15, CuAl25. 674 Physical and Mechanical Properties of Copper and Copper Alloys rapidly with increasing temperature. The JQ value for CuAl25 is only 7 kJm�2 at 250 �C in the T-L orientation.48 4.20.4.3 Creep Thermal creep of copper and copper alloys can be significant at relatively low temperatures, because of copper’s low melting point (0.3Tm¼�134 �C, Tm is the melting point). Nadkarni51 and Zinkle and Fabritsiev2 compared the 100-h creep rupture strength of copper and several PH and DS copper alloys at elevated temperatures. Copper alloys have significantly higher creep rupture strength than pure copper. Creep rupture strength decreases drasti- cally as temperature increases in PH alloys such as CuCrZr, as well as in pure copper, between 200 and 450 �C. DS alloys such as CuAl25 have superior creep rupture strength even above 400 �C because of their thermal stability at high temperatures. Li et al.31 summarized steady-state thermal creep data for pure copper and several copper alloys, as shown in Figure 5. Pure copper can suffer significant creep deformation at high temperature even with a very low applied stress. The creep rate of pure cop- per can be as high as �10–4 s�1 at �100MPa at 400 �C. The creep resistance of copper alloys is con- siderably higher than that of pure copper. The creep 10−4 10−6 10−8 10−10 10−12 0.01 0 50 100 150 20 0 8 16 24 C re ep r at e (1 s– 1 ) Applied str Pure copper at 400 �C (Nix et al., 1985) GlidCop Al-25 at 350 �C (Solomon et al., 1995) Ag–Cu at 193.3 �C (Thomas, 1993) Cu (Go GlidCop Al15 at 472 �C52 Applied s Figure 5 Steady-state thermal creep laws for copper alloys. A Trans. A 2000, 31A, 2491. rates of copper alloys strongly depend on the applied stress and the temperature, and can be described by the Norton power law relation; that is, e_¼ Asn expð�Q =RT Þ where e_ is creep rate, s is the applied stress, n is the stress exponent, Q is the activation energy, R is the gas constant, and T is the temperature. DS copper alloys exhibit unusu- ally high values of the stress exponent, for example, 10–21 in the temperature range of 472–721 �C for GlidCop Al15.52 Because of the time-dependent nature of creep deformation, softening behavior due to overaging and recrystallization must be considered during the creep analysis for PH copper alloys. The creep prop- erties of this class of alloys could be significantly changed during prolonged exposure at elevated temperature. 4.20.4.4 Fatigue and Creep–Fatigue Copper alloys are subjected to severe thermal cycles in high heat flux applications in fusion systems, and so, fatigue as well as creep–fatigue performance is a primary concern. Figure 6 shows the fatigue perfor- mance of OFHC Cu, PH CuCrZr and CuNiBe, and DS CuAl25.53 All three copper alloys show signifi- cantly better fatigue performance than OFHC cop- per. Among the three alloys, CuNiBe has the best 10−4 10−6 10−8 10−10 10−12 0.01 0 250 300 350 400 32 40 48 56 ess (MPa) Cu–Ni–Be at 229 �C (Thomas, 1993) Cu–Cr–Zr at 216 �C (Thomas, 1993) Cu–Cr–Zr at 300 �C5 –Cr–Zr at 300 �C rynin et al., 1992) GlidCop Al15 at 400 �C47 tress (ksi) fter Li, G.; Thomas, B. G.; Stubbins, J. F. Metall. Mater. 102 103 104 105 106 0.1 1 5 S tr ai n ra ng e (% ) Number of cycles to failure (Nf) Cu, 25 �C CuCrZr, 25 �C CuCrZr, 350 �C CuNiBe, 25 �C CuNiBe, 350 �C CuAl25, 25 �C CuAl25, 350 �C Figure 6 Fatigue performance of OFHC copper, precipitation-hardened CuCrZr and CuNiBe, and dispersion-strengthened CuAl25 in the temperature range of 25–350 �C. 0.1 1000 10 000 100 000 1 Ttest= 22 �C S tr ai n am p lit ud e (% ) Cycles to failure (Nf) OFHC Cu, no hold OFHC Cu, TCH 10 s CuAl25, no hold CuAl25, TCH 2 s CuAl25, TCH 10 s CuCrZr PA, no hold CuCrZr PA, TCH 10 s CuCrZr HT1, no hold CuCrZr HT1, TCH 10 s CuCrZr HT2, no hold CuCrZr HT2, TCH 10 s Figure 7 Hold time effect on the fatigue life of OFHC copper, DS CuAl25, and PH CuCrZr with three different heat treatments (prime aged (PA): solution annealed at 1233K for 3 h, water quenched, and then heat treated at 733K for 3 h; heat treatment 1 (HT1): PA plus an additional anneal in vacuum at 873K for 1 h and water quenched; and heat treatment 2 (HT2): PA plus an additional anneal in vacuum at 873K for 4 h (and water quenched) tested at room temperature. TCH, tension and compression hold. Physical and Mechanical Properties of Copper and Copper Alloys 675 fatigue response. The temperature dependence of fatigue behavior is stronger in CuAl25 and CuNiBe than in CuCrZr at temperatures between 25 and 350 �C. Heat treatments have an insignificant effect on fatigue life in CuCrZr.54 The fatigue life of copper and copper alloys can be significantly reduced when a hold time is applied at peak tensile and/or compressive strains during fatigue cycling. The hold time effect is evident even at room temperature and with a hold time as short as a few seconds.53,55,56 As shown in Figure 7, the fatigue life of OFHC copper is reduced significantly by the introduction of a hold time of 10 s at both tensile and compressive peak strains. The reduction in fatigue life is more severe in the high-cycle, long- life regime than in the low-cycle, short-life fatigue regime. A similar effect of the hold time was observed in copper alloys. The hold time effect appears to be more severe in CuAl25 than in CuCrZr. The effect of hold time is stronger in overaged CuCrZr (e.g., HT2 in Figure 7) than in prime-aged CuCrZr. Stress relaxation was observed during the hold periods even at room temperature where thermally activated creep processes are not expected. The reduction in fatigue life is apparently due to a change in the crack initiation mode from transgranular with no hold period to intergranular with a hold period.56,57 The fatigue life reduction under creep–fatigue load- ing could be more severe at high temperatures, particularly in PH copper alloys. Their softening behavior at elevated temperature due to overaging and recrystallization could have significant impact on the fatigue life with a very long hold time. Few studies have been performed to characterize the fatigue propagation rates of copper alloys. The fatigue crack growth rate of CuAl25 was found to be higher than that of CuCrZr at a lower stress intensity range, DK, at room temperature.58 Crack growth rates of CuCrZr and CuAl25 alloys increase with increas- ing temperature.49,59 4.20.5 Irradiation Effects in Copper and Copper Alloys The irradiation behavior of copper and copper alloys has been extensively studied up to high doses (>100 dpa) for irradiation temperatures of �400– 500 �C.60 Most of the irradiation experiments of cop- per and copper alloys have been done in mixed spectrum or fast reactors, such as HFIR, Fast Flux Test Facility (FFTF), or EBR-II. It should be noted that the accumulation rate of helium in copper in fusion reactors is significantly higher than in fission reactors (�10 appm dpa�1 in fusion reactors vs. 0.2 appm dpa�1 in fast reactors).22 Attention must be paid to transmutation effects such as helium when the irradiation data of copper and copper alloys from fission reactors are applied for fusion reactor design. 676 Physical and Mechanical Properties of Copper and Copper Alloys 4.20.5.1 Effect of Irradiation on Physical Properties of Copper and Copper Alloys Neutron irradiation leads to the formation of trans- mutation products and of irradiation defects, dis- location loops, stacking fault tetrahedra (SFT), and voids. All these features result in reduction of electrical and thermal conductivities.36,37,61–63 At irradiation temperatures between 80 and 200 �C, the electrical resistivity is controlled by the forma- tion of dislocation loops and stacking fault tetra- hedra and transmutation products. The resistivity increase from radiation defects increases linearly with increasing dose up to �0.1 dpa and saturates. The maximum measured resistivity increase at room temperature is about �6%. At irradiation tempera- tures above �200 �C, the conductivity change from extended radiation defects becomes less significant, and void swelling becomes important to the degrada- tion of the electrical conductivity. Fusion neutrons produce a significant amount of gaseous and solid transmutation products in copper. The major solid transmutation products include Ni, Zn, and Co. The calculated transmutation rates for copper in fusion first wall at 1MW-year m�2 are 190 appmdpa�1 Ni, 90 appm dpa�1 Zn, and 7 appm dpa�1 Co.2 Ni is the main transmutation element that affects the thermal conductivity of copper. It should be noted that water-cooled fission reactors would produce significantly higher transmutation rates of copper to Ni and Zn (up to �5000 and 2000 appm dpa�1, respectively) because of thermal neutron S tr es s (M P a) 300 0.3 dpa 0.2 dpa Ttest= 373 K Tirr= 373 K 0.1 dpa 0.01 dpa 250 200 150 100 S tr es s (M P a) 50 0 0 10 OHFC Cu Unirradiated 20 Strain (%) 30 40 50 60 70 Figure 8 Engineering stress–strain curves for OFHC copper (l hardened CuCrZr (right) neutron irradiated at 80 �C. The plot for D. J.; Singh, B. N.; Bilde-Sørensen, J. B. J. Nucl. Mater. 2005, 3 reactions. The data from fission reactor irradiation experiments must be treated with care when they are applied for fusion design. 4.20.5.2 Effect of Irradiation on Mechanical Properties of Copper and Copper Alloys 4.20.5.2.1 Tensile properties Irradiation causes large changes in tensile properties of copper and copper alloys. Copper and copper alloys can be hardened or softened by irradiation, depending on the irradiation temperature and the amount of the cold work prior to irradiation. Irra- diation hardening of copper and copper alloys due to defect cluster formation is significant at irradia- tion temperatures 300 �C because of radiation-enhanced recrystallization and precipitate coarsening in PH copper alloys. Low-temperature neutron irradiation of pure copper leads to development of a yield drop and significant hardening. Typical stress–strain behavior of pure copper and copper alloys irradiated to low doses at low temperatures is illustrated in Figure 8. The data of irradiated copper are from the work of Edwards et al.,64 and the data of irradiated CuCrZr from Li et al.14 Irradiation significantly changes the work hardening behavior of pure copper. Work hard- ening capability is progressively reduced with increas- ing doses. Appreciable work hardening still exists at the dose of 0.1 dpa. The effect of irradiation on the tensile behavior of copper alloys can be quite different. A complete loss of work hardening capability and 600 500 400 300 0.14 dpa 1.5 dpa 200 100 0 0 5 10 15 20 25 30 35 CuCrZr SAA Unirradiated Strain (%) eft) neutron irradiated at 100 �C and for precipitation- copper is from the reference. Reproduced from Edwards, 42, 164. Physical and Mechanical Properties of Copper and Copper Alloys 677 uniform elongation occurs at 0.14 dpa in neutron- irradiated CuCrZr in the prime-aged condition. Irra- diation to 1.5 dpa further reduces the yield strength, and recovers some total elongation in CuCrZr. The dose dependence of radiation hardening in copper at irradiation temperatures of 30–200 �C is summarized by Zinkle et al., and shown in Figure 9.65,66 Radiation hardening in copper can be observed at a dose as low as 0.0001 dpa. The yield stress increases dramatically with increasing dose and saturates at �0.1 dpa. Significant radiation hardening is accompanied by loss of strain hardening capabil- ities, resulting in prompt necking upon yielding. The temperature dependence of radiation hard- ening of pure copper at different irradiation tempera- tures was summarized and discussed by Fabritsiev and Pokrovsky.67 The radiation hardening decreases with increasing irradiation temperature in copper. The magnitude of radiation hardening is �200MPa at 80 �C, while only �40MPa at 300 �C at a dose of 0.1 dpa. Annealing at temperatures higher than 0.4 Tm can effectively reduce the defect cluster den- sity in copper. Annealing at 300 �C for 50 h after irradiation of copper to 0.01–0.3 dpa at 100 �C and annealing at 350 �C for 10 h after irradiation of CuCrZr IG and GlidCop Al25 IG to 0.4 dpa at 150 �C can essentially recover the ductility of the cop- per and copper alloys.68,69 However, postirradiation Kruglov et al. (1969) EI-Shanshoury 1972) Mohamed et al. (1982) Vandermeulen (1986) Heinisch (1988) Fabritsiev et al. (1994) Singh et al.23 Zinkle and Gibson65 Singh et al.75 350 Tirr= 30–200 �C 300 250 200 1500. 2% y ie ld s tr en gt h (M P a) 100 50 0.0001 0.001 0.01 Damage level (dpa) 0.1 1 10 100 Figure 9 Radiation hardening in copper. Reproduced from Zinkle, S. J.; Gibson, L. T. Fusion Materials Semi-annual Progress Report; DOE/ER-0313/27; Oak Ridge National Laboratory, 1999; p 163. annealing also reduces the critical stress for flow localization in pure copper.70 Irradiation creates a large increase in strength and decrease in ductility in copper alloys for irradiation temperatures below 300 �C. The strengthening effect decreases with increasing temperature. The crossover to radiation softening occurs at approximately 300 �C. The radiation softening effect in CuAl25 alloy is not as strong as for CuCrZr alloy where precipitate stability may be an issue. Neutron-irradiated copper alloys exhibit low uniform elongation after low-dose, low-temperature irradiation. The uniform elongation is recovered to near unirradiated values at 300 �C. Figure 10 compiles the yield strength data for PH CuCrZr and DS copper alloys (CuAl 25, CuAl15, MAGT 0.2) as a function of dose for the irradiation temperature of �100 �C.14,71 Both alloys show signifi- cant radiation hardening at low doses and an apparent saturation at �0.1 dpa. Irradiation-induced harden- ing is accompanied by the loss of strain hardening capability and a complete loss of uniform elongation, while the total elongation remains on the level of �10% for doses up to 2.5 dpa for CuCrZr. The strain rate dependence of tensile properties in neutron-irradiated CuCrZr was investigated at room temperature by Li et al.14 The strain rate sensi- tivity is small at room temperature in unirradiated CuCrZr. The measured strain rate sensitivity param- eter, m, is 678 Physical and Mechanical Properties of Copper and Copper Alloys Al15 and m< 0.01 for MAGT 0.2. In general, the strain rate and temperature dependence of flow stres- ses is small in fcc metals. 4.20.5.2.2 Fracture toughness Fracture toughness data for irradiated copper alloys are scarce. The effect of neutron irradiation on fracture toughness has been studied in two alloys, CuCrZr and CuAl25.14,50,72 Fracture toughness data on neutron-irradiated CuAl25 are available to a dose of 0.3 dpa, and for CuCrZr, the data are available up to 1.5 dpa (Figure 11). Neutron irradiation to 0.3 dpa significantly reduced the fracture toughness of CuAl25 in the temperature range of 20–350 �C. The toughness of irradiated CuAl25 is two to three times lower than that of the unirradiated alloy. The effect of neutron irradiation on fracture toughness of CuCrZr was less pronounced, despite the significant effect on the tensile properties even at relatively low doses (0.14–0.15 dpa). Reduction of fracture toughness in irradiated CuCrZr was small, and the JQ value was still >200 kJm�2 up to 1.5 dpa (Figure 11).14 4.20.5.2.3 Fatigue and creep–fatigue The effect of irradiation on fatigue performance has been evaluated for PH CuCrZr and DS CuAl25.73 The fatigue data for unirradiated and irradiated CuAl25 and CuCrZr in the temperature range of 20–350 �C are compiled and compared in Figure 12.24,53,74–76 The effect of irradiation on the fatigue response of CuAl25 is small at low tempera- ture. However, the fatigue life is reduced significantly 0 100 200 300 400 500 Solid symbols: JQ Open symbols: Jmax< JQ CuCrZr SCA CuCrZr SAA Tahtinen et al. Singh et al. Suzuki et al. Gillian et al. Rowcliffe CuCrZr SCA J Q (k J m −2 ) Dose (dpa) 0 0.01 0.1 1 CuCrZr SAA Tirr= 80 �C; Ttest= 22 �C Figure 11 Fracture toughness of CuCrZr with two heat treatments as a function of dose. The heat treatment, SCA, was to simulate the manufacturing cycle for ITER large components. Reproduced from Li, M.; Sokolov, M. A.; Zinkle, S. J. J. Nucl. Mater. 2009, 393, 36. at 250 and 350 �C because of radiation exposure. The fatigue life of the CuCrZr alloy was reduced follow- ing irradiation at 250 and 350 �C, similar to CuAl25. The degradation in the fatigue performance of these two alloys from irradiation exposure was not as severe as that in the tensile properties. Creep–fatigue behavior of neutron-irradiated CuCrZr was investigated at a dose level of 0.2–0.3 dpa at 22 and 300 �C by Singh et al.54 Hold times of 10 and 100 s were applied during fatigue cycling. Radiation hardening at low temperatures (e.g., 60 �C) is beneficial to the fatigue performance, while irradiation at high temperatures (e.g., 300 �C) has no significant effect on the creep–fatigue life of irradiated CuCrZr. A number of in-reactor creep–fatigue experiments were per- formed on a CuCrZr alloy in the BR-2 reactor at Mol (Belgium) by Singh et al.77 The irradiation experiments were carried out at 70 and 90 �Cat the strain amplitude of 0.5% with hold times of 10 and 100 s. The key finding was that neither the irradiation nor the hold time has any significant effect on the fatigue life of CuCrZr during the in-reactor tests. 4.20.5.2.4 Irradiation creep and void swelling There is limited literature on irradiation creep of copper and copper alloys.78–82 A study by Witzig82 showed no enhancement of creep rates in copper relative to thermal creep at 260 �C and 69MPa under light ion irradiation. Jung79 studied irradiation creep of 20% cold-worked copper foils at tempera- tures of 100–200 �C and the applied tensile stress of 20–70MPa under 6.2MeV proton irradiation with displacement rates of 0.7–3.5� 10–6 dpa s�1. The irradiation creep rate showed a linear stress dependence with the irradiation creep compliance of 6.2� 10–11 Pa�1 dpa�1 at stresses 50MPa), the creep rate showed a power law relation with the stress exponent of 4. Ibragimov et al.78 investigated in-reactor creep of copper in the WWR-K water-cooled reactor at a neutron flux of 2.5� 1015m�2 s�1 (E> 0.1MeV) at 150–500 �C and 20–65MPa. The in-reactor creep rate of copper was significantly higher than the thermal creep rate at temperatures below 0.4 Tm (Tm is the melting point). The stress dependence of the in-reactor creep rate showed a power law relation with the stress exponent of �3. Pokrovsky et al.80 reported irradiation creep data for DS MAGT 0.2. The irradiation creep experi- ments were performed using pressurized tubes 0.2 100 1000 10 000 100 000 1 3 GlidCopTM CuAl25 unirradiated and irradiated To ta l s tr ai n ra ng e (% ) Cycles to failure (Nf) Unirr RT, small size, UIUC Unirr RT, standard size, UIUC T irr = 47 �C, RT, RISO Unirr RT, HT at 650 �C, longitudinal, srivatsan Unirr RT, HT at 650 �C, transverse, srivatsan Unirr RT, HT at 650 �C, longitudinal, srivatsan Unirr RT, HT at 650 �C, transverse, srivatsan Unirr 200 �C in air, UIUC Unirr 250 �C in vac, RISO T irr = Ttest = 250 �C, 0.1 dpa, RISO T irr = Ttest = 250 �C, 0.3 dpa, RISO Unirr 350 �C in air, UIUC Unirr 350 �C in vac, UIUC Unirr 350 �C in vac, RISO T irr = Ttest = 350 �C, 0.1 dpa, RISO 100 1000 10 000 100 000 0.1 1 3 CuCrZr alloy, unirradiated and irradiated To ta l s tr ai n ra ng e (% ) Cycles to failure (Nf) Unirr RT, small size, UIUC Unirr RT, standard size, UIUC Unirr, 200 �C in air, UIUC Unirr, 250 �C in vac, RISO T irr = T test = 250 �C 0.3 dpa, RISO Unirr, 350 �C in air, UIUC Unirr, 350 �C in vac, RISO T irr = T test = 350 �C 0.3 dpa, RISO Figure 12 Effect of irradiation on fatigue life of CuAl25 (top) and CuCrZr (bottom) between room temperature and 350 �C. Physical and Mechanical Properties of Copper and Copper Alloys 679 irradiated in coolant water in the core position of the SM-2 reactor to �3–5 dpa at temperatures of 60–90 �C. A creep rate as high as �2� 10–9 s�1 was observed at a hoop stress of 117MPa. Radiation-induced void swelling in copper and copper alloys has been studied extensively. Zinkle and Farrell83,84 measured the temperature- dependence of void swelling in pure copper and a dilute Cu–B alloy neutron irradiated to �1.1–1.3 dpa at a damage rate of 2� 10–7 dpa s�1 at temperatures of 180–500 �C (Figure 13). Maximum swelling occurs at �300–325 �C in pure copper under fission neutron irradiation conditions. The lower temperature limit for void swelling is �180 �C, and the higher temperature limit �500 �C. Low-dose irradiation ( 0.7 Cu-100 appm 10B 0.6 0.5 0.4 D en si ty c ha ng e (% ) 0.3 0.2 Stage V Pure Cu 0 150 200 250 300 Irradiation temperature (�C) 350 400 450 500 550 0.1 Figure 13 Swelling in pure copper and Cu–B alloy. Reproduced from Zinkle, S. J.; Farrell, K. J. Nucl. Mater. 1989, 168, 262; Zinkle, S. J.; Farrell, K.; Kanazawa, H. J. Nucl. Mater. 1991, 179–181, 994. 680 Physical and Mechanical Properties of Copper and Copper Alloys Residual impurity oxygen can have a significant effect on void swelling in copper. A number of neu- tron, ion, and electron irradiation studies have shown that voids are not formed in high-purity, low-oxygen copper over the wide range of irradiation tempera- tures.60,86 The oxygen content should be maintained below �10wt ppm to minimize void swelling in copper. The effect of helium production on void forma- tion and swelling in copper is a significant concern for its fusion applications.87 Helium effects have been studied by either dual-beam ion irradiation88,89 or neutron irradiation of Cu–B alloys.89 Significant enhancement of void formation and swelling was observed in copper under ion irradiation with simul- taneous helium implantation. Neutron irradiation of copper containing �18wppm 10B to �1.2 dpa for the irradiation temperatures of 182–500 �C showed that the peak swelling temperature and the lower swelling temperature limit shifted to lower values (Figure 13). A recent study by Xu et al.90 of materials enriched in the copper isotopes, 63Cu, 63þ65Cu, and 65Cu neutron irradiated in the Materials Open Test Assembly (MOTA) in the FFTF at irradiation temperatures of 373–410 �C to doses up to 15.4 dpa found that both H and He enhanced void swelling in copper. The H effect is important at lower temperatures when the H production is considerably higher than the He production. At 410 �C the hydrogen effect de- creases dramatically and void swelling is affected by the helium concentration. PH and DS copper alloys have superior void swelling resistance compared to pure copper under fission neutron irradiation.2,71 Both PH CuCrZr and DS CuAl25 showed Physical and Mechanical Properties of Copper and Copper Alloys 681 swelling resistance. CuNiBe in the cold-worked and aged condition showed �28% swelling, while CuNiBe in the annealed and aged condition swelled only�13% after fission neutron irradiation to 98 dpa at 450 �C.95 The susceptibility to radiation-enhanced recrystallization is more severe in a cold-worked alloy, leading to the swelling instability. 4.20.5.3 Effect of Irradiation on Microstructure of Copper and Copper Alloys 4.20.5.3.1 Defect structure in irradiated copper and copper alloys Copper is among the most extensively studied metals in terms of fundamental radiation damage. Several reviews on the effect of irradiation on the 1.0 Loops, SFT (irradiation hardening) 0.5 N or m al iz ed u ni ts 0 (a) (c) 0 100 10 nm 200 3 Temperatu Figure 14 (a) Stacking fault tetrahedra and defect clusters pro 180 �C (reproduced from Zinkle, S. J.; Matsukawa, Y. J. Nucl. M irradiated at 300 �C (reproduced from Zinkle, S. J.; Farrell, K. J. showing the temperature dependence of defect cluster formation Radiation on Materials, ASTM STF 1125, 15th International Sym Testing and Materials: Philadelphia, 1992; p 813. microstructure of copper and copper alloys are available in the literature.60,96,97 Neutron irradiation of copper at low temperatures produces small defect clusters, dislocation loops, and SFTs. At temperatures above �150–180 �C, the density of defect clusters starts to decrease with increasing temperature, accompanied by the formation of voids. This temper- ature-dependent formation of defect structures is shown in Figure 14.60 Low-temperature neutron irradiation produces a high number density of SFTs and a low number density of dislocation loops in copper. Edwards et al.64 reported a number density of SFTs, �2–4� 1023m�3 and a number density of dislocation loops, 5� 1021m�3 in OFHC copper neutron irradiated to �0.01 dpa at 100 �C. Disloca- tion loops are believed to be of interstitial type. 300 �C 100 nm Void swelling (b) 00 re (�C) 400 500 600 duced in OFHC copper during irradiation to 1.9 dpa at ater. 2004, 329–333, 88), (b) voids in copper Nucl. Mater. 1989, 168, 262). (c) Schematic drawing and void swelling (reproduced from Zinkle, S. J. In Effects of posium); Stoller, R. E., et al., Eds.; American Society for 682 Physical and Mechanical Properties of Copper and Copper Alloys The size of SFTs is small, �2–3 nm. As doses increased, the density of SFTs increased to a satura- tion level at �0.1 dpa, while the size of SFT is inde- pendent of the dose and temperature. In general, the dislocation loop density is low, and a significant dis- location network is not formed in irradiated copper.96 Radiation hardening in copper can be adequately described by Seeger’s dispersed barrier model, and the yield strength increase is due to the formation of defect clusters.98 Singh and Zinkle96 summarized the dose dependence of the TEM-visible defect cluster density in copper irradiated near room temperature with fission neutrons, 14MeV neutrons, spallation neutrons, and 800MeV protons (Figure 15)96 TEM- visible defect clusters were observed at a very low dose (10–5 dpa). The defect cluster density showed a linear dependence on irradiation dose at low doses. The dose dependence of the defect cluster density shifts to either a linear or a square root relation at intermediate doses (>�0.0002 dpa). The cluster den- sity reaches an apparent saturation (�1� 1024m�3) at �0.1 dpa. The dislocation loops range in size from �1 to 25 nm.99 Differences in the type of irradiation (fission, fusion, spallation, etc.) have no significant effect on the defect cluster accumula- tion behavior in copper. The density of defect clus- ters in irradiated copper shows strong temperature 1025 1024 1023 1022 1021 1020 1019 1020 n = 1 1021 C lu st er d en si ty (m −3 ) ft Figure 15 Dose dependence of defect cluster density in copp from Singh, B. N.; Zinkle, S. J. J. Nucl. Mater. 1993, 206, 212. dependence (Figure 16).100 The defect cluster den- sity is essentially independent of the irradiation temperature between 20 and 180 �C (upper tempera- ture limit is dependent on dose rate). At higher tem- perature, the cluster density decreases rapidly with increasing irradiation temperature. At irradiation temperatures between 182 and 450 �C, the density of defect clusters was reduced by over three orders of magnitude.83,84 The saturation dose of the defect cluster density is similar, �0.1 dpa, for all irradia- tion temperatures.96 The size distribution of visible defect clusters can be described by an exponential function101: N(d )¼N0 exp(�d/d0), where N(d ) is the number of defects of diameter d, N0, and d0 are constants, and their values depend on irradia- tion conditions and material purity. As the irradiation temperature decreases, a fraction of small clusters increases relative to large clusters. Void formation occurs above �180 �C in neutron- irradiated copper.60 The peak void swelling temper- ature in copper is about 320 �C at a dose rate of 2� 10–7 dpa s�1. Singh and Zinkle96 summarized the dose dependence of void density measured by TEM in copper irradiated with fission and fusion neutrons at 250–300 �C from several studies. The data showed a large variation (up to two orders of magnitude differences) of void density between n = 1/2 800 MeV protons Yoshida et al. (1985) Zinkle89 Satoh et al. (1988) Horsewell et al. (1990) Makin et al. (1962) Shimomura et al. (1985) Brager et al. (1981) (n m−2) 1022 1023 1024 1025 er irradiated near room temperature. Reproduced Physical and Mechanical Properties of Copper and Copper Alloys 683 experiments. One possible source could be residual gas atoms in copper that can have a dramatic effect on void swelling in copper. Zinkle and Lee86 discussed in detail the effect of oxygen and helium on the forma- tion of voids in copper. The stacking fault tetrahedron is predicted to be the most stable configuration of vacancy clusters in copper. A small amount of oxygen (�10 appm) or helium (� 1 appm) in copper is needed to stabilize voids. High-purity copper with low oxy- gen concentration ( 0.1µm Figure 18 Comparison of the dislocation loop microstructure in irradiated pure copper (left), Cu–5% Mn (center) and Cu–5% Ni (right) alloys. The irradiation conditions were 0.7 dpa at 90 �C (Cu), 1.6 dpa at 160 �C (Cu–5% Mn), and 0.7dpa at 90 �C (Cu–5%Ni). Reproduced from Zinkle, S. J.; Horsewell, A.; Singh, B. N.; Sommer, W. F. J. Nucl. Mater. 1994, 212-215, 132; Zinkle S. J.; Nesterova, E. V.; Barabash, V. R.; Rybin, V. V.; Naberenkov, A. V. J. Nucl. Mater. 1994, 208, 119. 684 Physical and Mechanical Properties of Copper and Copper Alloys (Figure 18).25,26 These loop densities are more than an order of magnitude larger than the highest loop density observed in pure copper. The effect of the stacking fault energy on void formation in copper alloys was also investigated. Generally speaking, the lower the stacking fault energy, the less favorable for the formation of 3D voids. For example, swelling occurred in Cu–1–2.5% Ge alloys irradiated at 250 �C, while no measurable swelling occurred in Cu–3–5% Ge that has lower stacking fault energies.97 4.20.5.3.2 Dislocation channeling Dislocation channels are frequently observed during postirradiation deformation of copper and copper alloys.102,103 Greenfield and Wilsdorf104 were the first who observed an area free of irradiation defects in the middle of a slip-line cluster by TEM in a neutron-irradiated copper single crystal. Extensive studies were conducted to establish the correlation between the deformation behavior and the slip-line structure in neutron-irradiated copper single crys- tals.104–107 Sharp108–110 studied the deformation and dislocation channels in neutron-irradiated copper single crystals in detail, and established a direct cor- relation between the surface slip steps and dislocation channels. The channels are nearly free of irradiation- produced defects, and operate parallel to the primary {111} slip plane. The cleared channels are formed by cooperative localized motion of glide dislocations that interact with and annihilate the preexisting radi- ation defect clusters. The channel characteristics have strong dependence on irradiation dose and test temperatures. The channel width decreases and the slip step height increases with increasing irradiation dose, and the channel width and the slip step height decrease with decreasing deformation temperature. Howe111 confirmed that the channel width, spacing, the slip step height, and the average shear per slip band increase with increasing test temperature in the temperature range of 4–473K. The reduction in channel width was considered to be a consequence of impeded cross-slip.108,111 Dislocation channels were also observed in neutron-irradiated copper single crystals under cyclic straining.112,113 The width and average spacing of channels changed with the number of cycles, in contrast to channels formed during tensile strain- ing where the width and spacing of channels were constant over a large range of strains.108 Dislocation channels are formed in neutron- irradiated copper alloys as well. Sharp114 observed the channeling effect in three different copper alloys neutron irradiated at ambient temperature, that is, Cu–0.8% Co, Cu–Al2O3, and Cu–4% Al single crystals. The channel spacing in the copper alloys were 1.2–1.5 mm, about half that observed in neutron-irradiated copper single crystals (2.3mm). The channel width in Cu–0.8% Co alloy is similar to that for irradiated copper crystal (0.16mm), and the chan- nels have the uniform width along the length. The presence of the second-phase particles in Cu–0.8% Co alloy has little effect on channeling. In the DS Cu–Al2O3 alloy, the channels are wider (0.24mm) and Physical and Mechanical Properties of Copper and Copper Alloys 685 more irregular in width. The channel width can vary by a factor of 2 within a few microns along the length of a channel. A high density of dislocations surrounding the particles within the channel was observed in Cu– Al2O3, indicating great difficulty of dislocations in bypassing the (nondeforming) second-phase particles. In the single-phase Cu–4% Al alloy, however, no dislocation channels were observed. Edwards13,40,64,115 studied thoroughly the deformed microstructure in neutron-irradiated CuCrZr alloys, and compared with the deformation microstructure in neutron-irradiated OFHC-Cu (Figure 19). Disloca- tion channels were observed during postirradiation deformation of the CuCrZr alloy neutron irradiated to 0.2–0.3 dpa at 100 �C. Channels were formed even before the upper yield point, and continued through- out the tensile deformation process. Some channels are completely free of defect clusters, and others contain a sizeable population of defect clusters. The width of cleared channels varied between about 100 and 250 nm. The channel formation is more pro- nounced in a higher-dose specimen than in a lower- dose specimen. In comparison with OFHC-Cu, CuCrZr showed little difference in deformation mode and channel characteristics in terms of width and size. While the channels in the OFHC-Cu were free of defects and dislocation debris, the channels in the CuCrZr alloy contained a small fraction of defects and precipitates. When the irradiated CuCrZr was annealed and deformed, deformation occurs in a much more homogeneous fashion, and no well-defined channels were observed. The formation of dislocation channels in pure copper was investigated by in situ straining experi- ments on ion-irradiated copper in an electron micro- scope.116,117 Postirradiation straining of the thin foils of polycrystalline copper irradiated with 200 keV Kr Figure 19 Dislocation channels observed in OFHC-Cu (left) an Edwards, D. J.; Singh, B. N.; Xu, Q.; Toft, P. J. Nucl. Mater. 2002, J. B. J. Nucl. Mater. 2005, 342, 164. ions to about 2� 10–4 to 0.02 dpa at room tem- perature showed that defect-free channels nucleate at grain boundaries, or in the vicinity of cracks, sug- gesting that grain boundaries and crack tips are nucleation sites for channels.117 Cross-slips were found to be responsible for channel widening and defect removal within the channel. Edwards et al.64 studied the initiation and propagation of dis- location channels in neutron-irradiated OFHC-Cu (Figure 20) and CuCrZr alloy in an interrupted tensile test. TEM observations suggested that chan- nels are initiated at boundaries, large inclusions, or existing channels. Channels are formed by interac- tions of newly formed dislocations with irradiation defects on the glide plane. Once formed, the channels propagate rapidly in the grain interior until they intercept another boundary, interface, or channel. Despite significant efforts, the exact mechanism of channel formation and evolution still remains unre- solved, and a clear connection between the slip pro- cesses, dislocation channeling, and localized flow in neutron-irradiated metals is still lacking. 4.20.6 Joining Copper and copper alloys can be joined by a variety of techniques, including mechanical coupling, weld- ing, brazing, and diffusion bonding. A comprehensive overview of joining techniques for copper and copper alloys can be found in the reference.118 The welding techniques commonly used for copper and copper alloys include arc welding, resistance welding, oxy- fuel welding, and electron beam welding. Welding is generally not recommended for joining high-strength copper alloys. PH copper alloys lose their mechanical strength because of the dissolution of precipitates d CuCrZr (right) irradiated to 0.3 dpa at 100 �C. 307–311, 439; Edwards, D. J.; Singh, B. N.; Bilde-Sørensen, Strained to 1.5%Strained to 1.5% Strained to 14.5%Strained to 14.5% (a) (b) (c) (d) Figure 20 Examples of cleared channels formed in the OFHC-Cu irradiated (to 0.3dpa) and tested at 323K to different strain levels: (a) before yield, (b) before yield, (c) 1.5%, and (d) 14.5%. Note that at 14.5% strain level the grain is subdivided by numerous channels formed on different slip planes. All images shown in this figure were taken in the STEM bright field mode. Reproduced from Edwards, D. J.; Singh, B. N.; Bilde-Sørensen, J. B. J. Nucl. Mater. 2005, 342, 164. 686 Physical and Mechanical Properties of Copper and Copper Alloys during the welding process. The welded component must be resolution annealed and aged to recover some of the initial strength in the joint. Recrystallization in the melt layer degrades the mechanical property of the weldment. DS copper alloys cannot be welded by conventional welding processes because of the loss of oxide particles and recrystallization in the weld zone. Brazing is the most common method for joining copper alloys. All conventional brazing techniques can be used to join copper and copper alloys, includ- ing furnace brazing, torch brazing, induction brazing, resistance brazing, and dip brazing. A wide range of filler metals are available, and the most common brazing filler metals are Cu–Zn, Cu–P, Cu–Ag–P, and Ag- and Au-based alloys.118 Ag- and Au-based filler metals are unacceptable in fusion reactor envir- onments because of concerns of high radioactivity from neutron-induced transmutation.119 Copper alloys are typically brazed at tempera- tures between 600 and 950 �C with hold times at the brazing temperature ranging from 10 s (torch, resistance, or induction brazing) to 10min (furnace brazing).2 The brazing process can significantly soften PH copper alloys as a result of the adverse precipita- tion process. To reduce the softening effect, a fast induction brazing technique has been developed to minimize the holding time at high temperature to retain sufficient mechanical properties.120 Alterna- tively, the brazed component can be aged following furnace brazing to restore part of its initial strength. Complete recovery of high strength after furnace brazing by heat treatment in PH alloys is rather diffi- cult in practice as the component must be heated to a temperature greater than typical brazing tempera- tures and rapidly quenched to create a supersatura- tion of solute prior to aging. Oxide DS copper has been successfully joined using torch, furnace, resis- tance, and induction brazing.2 Softening is not a seri- ous concern for the base metal of DS copper alloys because of their high recrystallization temperature. The brazed copper joints show good fatigue proper- ties and relatively low ductility.2 Diffusion bonding is a viable technique to produce joints with high mechanical strength for DS copper alloys, but cannot be used to produce high-strength Physical and Mechanical Properties of Copper and Copper Alloys 687 joints in PH alloys because of significant softening of the base metal during high-temperature exposure. The DS CuAl15 and CuAl25 alloys can be joined by diffusion bonding with acceptable bond strengths under the diffusion bonding conditions similar to the normal HIPing conditions.121 Techniques for joining copper alloys to beryllium or austenitic stainless steels have been developed for the ITER plasma-facing components. A review of the joining technology was given by Odegard and Kalin.119 Recent work has focused on small- and medium-scale mock-ups and full-scale prototypes of the ITER first wall panels.122 The first wall panels of the ITER blanket are composed of a composite Cu alloy/316L(N) SS water-cooled heat sink structure with Be tile clad. A number of joining techniques have been explored for joining copper alloys to aus- tenitic stainless steel, 316L(N), including diffusion bonding, brazing, roll bonding, explosive bonding, friction welding, and HIP.123 HIP joining is by far the most desirable technique. For the PH CuCrZr alloy, the heat treatment must be integrated with the bonding cycle, and a high cooling rate (>�50 �C min�1) is required to obtain good mechanical proper- ties of CuCrZr after subsequent aging treatments. Two alternative processes are recommended124: the HIP cycle (1040 �C and 140MPa for 2 h) followed by quenching in the HIP vessel, or a normal HIP cycle with a subsequent heat treatment in a furnace with fast cooling. Gervash et al.125 studied alternative SS/Cu alloy joining methods, for example, casting, fast brazing, and explosion bonding. Cast SS/CuCrZr joint may be suitable for some ITER applications. Brazing and diffusion bonding have been consid- ered for joining the beryllium armor to a copper alloy heat sink. The Be/DS copper alloy joints can be made by high-temperature HIPing and furnace brazing.126 Results from shear tests on small-scale specimens and from high heat flux tests of the first wall mock-ups showed good performance of joints brazed with STEMET 1108 alloy at�780 �C for less than 5min.122 The Be/Cu-Al25 solid HIPing (e.g., 730 �C and 140MPa for 1 h) showed good performance from shear tests, high heat flux tests, and neutron irradiation.122 The development of joining techniques for PH CuCrCr alloy must consider the loss of mechanical strength because of overaging at high temperatures. The HIPing temperature must be reduced to be as close as possible to the aging temperature. The best results obtained so far is for HIPing at 580 �C and 140MPa for 2 h.126 A fast induction brazing tech- nique has also been developed to minimize the holding time at high temperature. Diffusion bonding of Be/CuCrZr joints gives much better high heat flux performance than brazing, and has been selected as the reference method for the European Union ITER components.120 A low-temperature Be/Cu alloy bond- ing process has also been developed that is compatible with both DS and PH copper alloys.124,127 In the United States, several different joint assemblies for diffusion bonding a beryllium armor tile to a copper alloy heat sink have been evaluated.128 To prevent formation of intermetallic compounds and promoting a good diffusion bond between the two substrates, aluminum or an aluminum–beryllium composite (AlBeMet-150) has been used as the interfacial mate- rial. Explosive bonding was used to bond a layer of Al or AlBeMet-150 to the copper substrate that was subsequently HIP diffusion bonded to an Al-coated beryllium tile. A thin Ti diffusion barrier (0.25mm) was used as a diffusion barrier between the copper and aluminum to prevent the formation of Cu–Al intermetallic phases. The Be/Cu alloy joints showed good strength and failure resistance. 4.20.7 Summary High heat flux applications for fusion energy systems require high-strength, high-conductivity materials. Selection of materials for high heat flux applica- tions must consider thermal conductivity, strength and tensile ductility, fracture toughness, fatigue and creep–fatigue, and radiation resistance. Pure copper has excellent conductivity but poor strength. PH and DS copper alloys have superior strength and suffi- cient conductivity, and are prime candidates for high heat flux applications in fusion reactors. These two classes of alloys have their own advantages and dis- advantages with regard to fabrication, joining, and in- service performance. PH copper alloys, such as CuCrZr, are heat- treatable alloys. Their properties are strongly depen- dent on the thermomechanical treatments. They possess high strength and high conductivity in the prime-aged condition, and good fracture toughness and fatigue properties in both nonirradiated and irradiated conditions. However, this class of alloys is susceptible to softening at high temperatures because of precipitate overaging and recrystallization. Their properties can be significantly degraded during large component fabrication because of their inability to achieve rapid quenching rates. DS copper alloys such as GlidCop Al25 have excellent thermal stability, and 688 Physical and Mechanical Properties of Copper and Copper Alloys retain high strength up to temperatures near the melting point. The main disadvantages of this class of alloys are their relatively low fracture toughness and difficulty to join. The effect of neutron irradiation in copper alloys depends largely on the irradiation temperature. At irradiation temperatures below �300 �C, radiation hardening occurs along with loss of strain hardening capability and complete loss of uniform elongation. Radiation hardening saturates at about �0.1 dpa in this temperature regime. At higher temperatures, radiation-induced softening can occur. Void swelling takes place between 180 and 500 �C, and the peak swelling temperature is �300–325 �C for neutron irradiation at damage rates near 10–7 dpa s�1. PH and DS copper alloys are more resistant to void swelling than pure copper. Irradiation slightly reduces the fracture toughness of copper alloys, and the effect is stronger in CuAl25 than in CuCrZr. Irradiation has no significant effect on fatigue and creep–fatigue perfor- mance. Transmutation products can significantly change the physical properties and swelling behavior in copper alloys. Significant R&D efforts have been made to select and characterize copper alloys for high heat flux applications. The ITER Material Property Handbook provides a comprehensive database for pure copper, CuCrZr, and CuAl25. For the ITER first wall and divertor applications, CuCrZr has been selected as the prime candidate. Current focus is on fabrication, joining, and testing of large-scale components. References 1. Butterworth, G. J.; Forty, C. B. A. J. Nucl. Mater. 1992, 189, 237. 2. Zinkle, S. J.; Fabritsiev, S. A. Atomic and Plasma- Materials Interaction Data for Fusion (Supplement to the Journal Nuclear Fusion 1994, 5, 163. 3. Davis, J. R., Ed. ASM Specialty Handbook: Copper and Copper Alloys; ASM International: Materials Park, OH, 2001; p 276. 4. Atrens, A.; Nairn, J.; Fernee, H.; FitzGerald, K.; Skennerton, G.; Olofinjana, A.Mater. Forum 1997, 21, 57. 5. Taubenblat, P. W.; Smith, W. E.; Graviano, A. R. In High Conductivity Copper and Aluminum Alloys; Ling, E., Taubenblat, P. W., Eds.; Metallurgical Soc of AIM: Warrendale, PA, 1984; p 19. 6. Lu, L.; Shen, Y.; Chen, X.; Qian, L.; Lu, K. 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All rights reserved. 4.21.1 Introduction 691 4.21.2 Magnetohydrodynamic Issues and the Requirement for Insulator Coatings 691 4.21.3 Development of Insulator Coating for Liquid Li Blanket 692 4.21.3.1 In Situ Formation and Healing with CaO 693 4.21.3.2 In Situ AlN Coating 693 4.21.3.3 Er2O3 and Y2O3 as New Candidates 694 4.21.3.3.1 Scoping by bulk immersion tests 694 4.21.3.3.2 In situ coating with Er2O3 694 4.21.3.3.3 Physical coating processes 695 4.21.3.3.4 Other coating technologies 695 4.21.3.4 Two-Layer Coatings 696 4.21.3.5 Radiation Effects 697 4.21.3.6 FCI Concept as an Alternative to Insulator Coating 697 4.21.4 Summary and Remaining Issues 698 References 698 Abbreviations EB-PVD Electron beam-physical vapor deposition FCI Flow channel insert IFMIF International Fusion Materials Irradiation Facility ITER-TBM International Thermonuclear Experimental Reactor-Test Blanket Module MHD Magnetohydrodynamic MOCVD Metal-organic chemical vapor deposition MOD Metal-organic deposition RF Radiofrequency RIC Radiation-induced conductivity TBR Tritium breeding ratio TEM Transmission electron microscope 4.21.1 Introduction The use of a ceramic coating for electrical insulation is a key technology for fusion blanket systems using liquid metals as breeding and coolant material and solid metals as the structural material. Particularly for the blanket system using liquid lithium and vanadium alloys (Li–V blankets), coating develop- ment is a major feasibility issue (see also Chapter 4.12, Vanadium for Nuclear Systems for vanadium alloy and liquid Li blankets). Overviews of the coating development for liquid lithium blankets are available in recent publications.1–3 It should, however, be noted that, with the development of a more fundamental understanding of coating behavior and blanket design, there has been a paradigm shift in coating development. This chapter describes the present status of insulator coating R&D, in addition to a historical overview of its development. 4.21.2 Magnetohydrodynamic Issues and the Requirement for Insulator Coatings Breeding blankets for fusion reactors are categorized into solid breeder and liquid breeder concepts. The liquid breeder blankets have certain advantages over the solid breeder blankets such as continuous chemi- cal control of the breeding material including isoto- pic control of Li, impurity control and tritium recovery, and immunity to irradiation effects. How- ever, some issues such as compatibility of the breeder with structural materials are more serious for liquid breeder blankets4. In addition, blanket structure 691 692 Ceramic Coatings as Electrical Insulators in Fusion Blankets can be simplified significantly if the liquid breeder functions as the coolant as well (self-cooled liquid breeder blanket). At present, the major candidate liquid breeder materials are Li, Li–Pb, and molten salt Flibe (LiF–BeF2). In the cases of self-cooled liquid Li and Li–Pb blankets in magnetic confinement fusion systems, the high-speed flow of these materials perpendicular to the strong magnetic field causes an electric current, which then produces an electromagnetic force as a result of interaction with the magnetic field. This force changes the velocity profile in the cooling ducts and acts to retard the coolant flow, leading to what is called a magnetohydrodynamic (MHD) pres- sure drop. This process is schematically shown in Figure 1(a). The MHD pressure drop may result in loss of flow control and mechanical stresses exceeding the allowable limits of the structural materials. The problems arising from the MHD pressure drop are critical feasibility issues for self-cooled liquid metal breeder blanket concepts with metallic structures. The quantification of the MHD pressure drop requires a rather complex numerical analysis. How- ever, in simple cases such as straight and constant area cross-section flow in conductive ducts with a uniform magnetic field in a traverse direction, the pressure gradient along the flowing direction, dp/dx, is given as follows5: dp=dx ¼ ksUB2 where s, U, and B are the electrical conductivity, flow velocity of the liquid metal, and magnetic flux den- sity, respectively, and k is a positive function of elec- trical conductivity of the wall. The equation implies that the MHD pressure drop is an issue in the case of high magnetic field and high velocity flow of conduc- tive liquid metals. In the case of a low flow rate such as would occur in a helium-cooled Li–Pb blanket, the MHD pressure drop will not be an issue. Insulator coating Magnetic field Lorentz force Vanadium duct Liquid Li flow Figure 1 Schematic illustration of magnetohydrodynamic pressure drop (left) and the role of insulator coating (right). To reduce the MHD pressure drop, optimization of the coolant flow path by enhancing the flow fraction parallel to the magnetic field may have some effect. However, a more effective way to reduce the MHD pressure drop would be to electrically insulate the coolant flow from the surrounding walls.6 The reduc- tion of MHD pressure drop by an insulator coating is schematically illustrated in Figure 1(b). The require- ments for the coating can be summarized as follows: 1. compatibility with liquid breeder under flowing conditions with a temperature gradient, 2. high electrical resistivity under irradiation, 3. robustness and/or an effective self-healing capability, 4. potential for covering large and complex surfaces, and 5. fundamental requirements for in-vessel materials such as radiation resistance, low activation proper- ties, and low tritium inventories in blanket conditions. Quantitative evaluation of the required electrical resistance and an allowable crack fraction are subject to overall blanket design including flow channel structures. A recent model calculation showed that the ratio of electrical resistivity of the insulator to the wall needs to be ≳106 and crack areal fraction to be ≲10�6 to maintain the pressure drop within tolerable range, assuming Li wets cracks.7,8 For the Li–Pb blankets, the insulator coating should be a critical issue if a self-cooled Li–Pb blanket with metallic structural materials is to be designed. How- ever, current blanket design options with Li–Pb are (1) helium cooled with slow-flowing Li–Pb, (2) dual- coolant Li–Pb with fast-flowing Li–Pb but electrically insulated from the wall by a SiC/SiC flow channel insert (FCI), or (3) self-cooled Li–Pb using SiC/SiC as the structural material. None of these concepts needs the insulator coating. However, development of a ceramic coating, necessary mostly for tritium perme- ation reduction and possibly for corrosion protection, is still a critical issue.9 4.21.3 Development of Insulator Coating for Liquid Li Blanket Liquid Li is a strong reducing agent, and thus a coating layer of many common oxides is not stable on the wall material. Therefore, intentional coating with insulator ceramics, which are stable in liquid Li, is necessary. Figure 2 shows free energy for oxidation and nitridation for various elements. From the 200 –600 –550 –500 –450 Y2O3 Er2O3 CaO BeO Li2O MgO Al2O3 ΔG (k J p er m ol a to m ic O ) ΔG (k J p er m ol a to m ic N ) T (�C) T (�C) –300 –250 –200 –150 –100 –50 0 TiN AlN BN Si3N4 Li3N 300 400 500 600 700 800 200 300 400 500 600 700 800 Figure 2 Free energy of oxide and nitride formation for selected ceramics. M2O¥ O2– OV M¥+ MLi V-alloy (O-doped) Liquid Li (M-doped) Oxide coating Figure 3 Schematic illustration of the mass transport for in situ oxide coating in Li. Ov: oxygen in vanadium substrate; MLi: metal doped in Li for producing oxide coating. Ceramic Coatings as Electrical Insulators in Fusion Blankets 693 thermodynamic viewpoint, CaO, Y2O3 and Er2O3, and AlN are expected to be stable. (Note that the resistivity of TiN is not sufficient for the coating.) The instability, however, is a function of the impurity level of Li.10 Early development efforts focused on CaO and AlN by in situ formation during exposure to chemi- cally controlled lithium. 4.21.3.1 In Situ Formation and Healing with CaO This technology aims at using a corrosion product that forms at the interface between Li and structural materials as an insulating layer. By careful control of the corrosion reaction, the insulation layer can be formed uniformly on the inner surfaces of complex components. The corrosion layer may also be formed on the cracked area of the coating, thereby repairing insulation defects. The in situ coating with CaO and AlN have been studied in the United States11 and the Russian Federation,12 respectively. ACaO insulator layer forms during the immersion of vanadium alloys in Ca-doped Li. For enhancing this reaction and stabilizing the layer, the O level in the alloy was increased by prior doping. The pro- cess is schematically shown in Figure 3. Careful control of the Ca level in Li and the O charging condition of the vanadium alloys made in situ for- mation and healing of the CaO coating possible at 694 Ceramic Coatings as Electrical Insulators in Fusion Blankets capsule tests with Mo and vanadium alloys.15,16 Thus, unless the vanadium alloy channel walls are fully covered with AlN, the use of this coating tech- nique seems problematic. For the same reason, the use of AlN layers, sheets, or plates in Li as the insulator requires full coverage of the vanadium alloy structures with AlN. 10μm V-4Cr-4Ti Er V Figure 5 Cross-section of in situ Er2O3 coating on V–4Cr–4Ti after exposure to Li(Er) for 300h at 600 �C. Reproduced from Yao, Z.; Suzuki, A.; Muroga, T.; Katahira, K. J. Nucl. Mater. 2004, 329–333, 1414–1418, with permission from Elsevier. 4.21.3.3 Er2O3 and Y2O3 as New Candidates 4.21.3.3.1 Scoping by bulk immersion tests Exploration of new candidates having superior per- formance compared to CaO and AlN was carried out using static immersion tests. Figure 4 shows the mass loss of various insulator ceramics due to exposure to static Li at high temperatures. As predicted by ther- modynamics, Er2O3 and Y2O3 showed stability supe- rior to that of CaO.2,15 For these materials, formation of LiXO2 (X¼Er or Y) during exposure to Li was reported,17,18 although the impact of these changes on the coating properties remains to be assessed. In particular, the effects of Li flow on the stability of the corrosion products on the surfaces will be the key issue. 4.21.3.3.2 In situ coating with Er2O3 Based on the experience with CaO, in situ coating with Er2O3 and Y2O3 was attempted on V–4Cr–4Ti by doping Er or Y in Li and O into V–4Cr–4Ti. Coating with Y2O3 was shown to be difficult, proba- bly because there was almost no solubility of Y in Li. Li 1000 h 500 –10 –5 0 5 10 CaO (p-crystal) CaO (s-crystal) CaZrO3 Y2O3 Er2O3 15 20 25 M as s lo ss (m g cm –2 ) 30 600 Li temperature (�C) 700 800 Figure 4 Change of mass after exposure to static Li for 1000h for a bulk of candidate ceramics. Adapted from Pint, B. A.; Tortorelli, P. F.; Jankowski, A.; et al. J. Nucl. Mater. 2004, 329–333, 119–124, with permission from Elsevier. On the other hand, formation of Er2O3 was con- firmed.19 Because the solubility of Er in Li is much lower than that of CaO, the stability of the coating, once formed, is much higher compared with the CaO in situ coating. The cross-section of the coating with compositional profile and coating thickness with time and temperature are shown in Figures 5 and 6, respectively. In the effort to optimize the precharging condition of oxygen, the microstructural process for restoring oxygen in vanadium alloy substrate was clarified.20 Figure 7 shows the depth profile of hardness before and after oxygen charging, after heat treatment and 0 100 200 300 400 500 600 700 800 0.0 0.5 1.0 1.5 V-4Cr-4Ti (NIFS-HEAT-2) Oxidation - 700 �C, 6 h Annealing - 700 �C, 16 h Exposure - Li (Er) Th ic kn es s of E r 2 O 3 (μ m ) Exposure time (h) 500 �C 550 �C 600 �C 650 �C 700 �C Figure 6 Growth of the Er2O3 layer during the in situ coating. Reproduced from Yao, Z.; Suzuki, A.; Muroga, T.; Yeliseyeva, O. I.; Nagasaka, T. Fusion Eng. Des. 2006, 81, 951–956, with permission from Elsevier. 0 50 100 150 200 250 300 0 500 0.2μm 1500 V - 4Cr - 4Ti (NIFS-HEAT-2) Depth from surface (μm) Oxidized - 6 h, 700 �C (b) Oxidized - 6 h, annealed - 16 h, 700 �C (c) Oxidized - 6 h, annealed - 16 h, Li(Er) - 100 h, 700 �C (d) As-received (a) V ic ke rs h ar d ne ss (H V ) Ti-O [200] [020] Ti-C-O-N (a) (b) (c) (d) 1000 Figure 7 Depth distribution of hardness and transmission electron microscope microstructure near the surfaces in the V–4Cr–4Ti substrate. (a) Before oxidation, (b) after oxidation, (c) after annealing, and (d) after in situ coating (coating was removed). Reproduced from Yao, Z.; Suzuki, A.; Muroga, T.; Yeliseyeva, O. I.; Nagasaka, T. Fusion Eng. Des. 2006, 81, 951–956, with permission from Elsevier. Ceramic Coatings as Electrical Insulators in Fusion Blankets 695 after in situ coating, together with the transmission electron microscope (TEM) microstructures near the surfaces. The hardness is known to follow the approx- imate level of O in V–4Cr–4Ti.21 After oxidation for 6 h at 700 �C, the surface was covered with a complex oxide layer. After the subsequent heat treatment for 16 h at 700 �C, the matrix was composed of a high density of needle-shaped Ti–O (mostly TiO2) precipitates oriented in the h100i directions (net structure). This structure was most prominently observed after annealing at 700 �C. This is consistent with the results of the precipitation study, which showed that, although Ti interacts with impurity O already at �200 �C, Ti–O precipitates start to form at �700 �C in V–4Cr–4Ti alloys.22 Figure 7 also shows that the net structure near the surface disappeared after exposure to Li for 100 h at 700 �C because of the loss of oxygen. A recent study showed in situ healing capabilities with Er2O3, but further optimization of the process is required to obtain a reliable healing function.23 The mechanism of the net structure formation and supply of oxygen for the coating was elucidated using a kinetic model. The model successfully explained the experimental trends.24 4.21.3.3.3 Physical coating processes More conventionally deposited coatings of Er2O3 or Y2O3 were developed such as arc source plasma deposition,25 electron beam-physical vapor deposition (EB-PVD),2 and radiofrequency (RF) sputtering.26 Because of the variation in the quality of the coating and the test conditions, reported stability in liquid lithiumwas different in different tests. Er2O3 produced by EB-PVD was heavily damaged in Li already at 500 �C,27 and Er2O3 and Y2O3 coatings produced by RF sputtering were also exfoliated at 500 �C.28 How- ever, Er2O3 fabricated by arc source plasma deposition showed promising results, as shown in Figure 8. Depo- sition on a higher temperature substrate produced a highly crystalline Er2O3 coating, which was shown to be stable in Li for 1000 h at 700 �C.1 The stability in Li is known to be enhanced by improving the purity and crystallinity of the coating. An oxide layer at the coating/substrate interface may cause extensive exfoli- ation because Li introduced through cracks would preferentially attack the oxidized interface. 4.21.3.3.4 Other coating technologies The efforts in using the physical coating processes explained so far are essential for establishing the 696 Ceramic Coatings as Electrical Insulators in Fusion Blankets coating concepts. However, further efforts are also necessary to enhance the engineering feasibility of coating on large and complex surfaces. For this purpose, the sol–gel method (in other words, metal-organic deposition, MOD) and metal-organic chemical vapor deposition (MOCVD) have been explored. Er2O3 coating was formed on stainless steel by the sol–gel method. Crystallinity of the coating depended on annealing atmosphere and tempera- ture.29 MOCVD was applied to coating Er2O3 on V-alloys and other materials. Successful coating on the inner surface of a tube was demonstrated.30 4.21.3.4 Two-Layer Coatings Forming an overlayer on the ceramic coating with ductile metals is an alternative approach to in situ healing to compensate for cracks in the insulator layer. Current efforts for fabricating the two-layer Substrate temperature RT 850 K Er2O3 XRD (low angle) 15 In te ns ity (a .u .) 30 45 2q 60 Figure 8 Change of Er2O3 coating on V–4Cr–4Ti by exposure t using arc source plasma deposition method with substrate temp of Er2O3 was observed only in the case of the high substrate tem coating at high substrate temperature after exposure to Li. Repr J. Nucl. Mater. 2007, 367–370, 780–787, with permission from E Er2O3 V outer layer V-4Cr-4Ti substrate 5μm Figure 9 Cross-section of Er2O3–V two-layer coating produce two-layer coatings produced by radiofrequency sputtering on V Jankowski, A.; Hayes, J. J. Nucl. Mater. 2007, 367–370, 1165–1 Li, M.; Muroga, T.; Terai, T. Compatibility of MHD coating candi Presented at 13th International Conference of Fusion Reactor M coating use pure V, V-alloys, and pure Fe, as shown in Figure 9. In these cases, only modest compatibil- ity with Li would be required for the insulator ceramics. Instead, the compatibility of V, V-alloys, and pure Fe with Li needs to be verified. It is also necessary to avoid shorting between the over- layer and the substrate. The Er2O3–V and Y2O3–V two-layer coatings on a V-alloy substrate fabricated using the EB-PVD process showed sufficient resis- tivity in Li.31 Recently, a mono-metallic thermal convection Li loop was constructed using V–4Cr–4Ti alloy. The two-layer coatings were tested in addition to uncoated V-alloy substrates at 700 �C for >2000 h in flowing Li. The tentative results showed that the corrosion loss, degradation of the coating, and mechanical property changes to the substrate were small.32 However, a full characterization remains to be completed. Exfoliated 1000 h in liquid Li 773 K 873 K 973 K o Li at 773, 873, and 973K for 1000h. The coating was made erature at room temperature and 850K. Crystalline structure perature. Remarkable change was not observed in the oduced from Muroga, T.; Chen, J. M.; Chernov, V. M.; et al. lsevier. Fe Er2O3 V-4Cr-4Ti substrate 5μm d by electron beam-physical vapor deposition and Er2O3–Fe –4Cr–4Ti. Reproduced from Pint, B. A.; Moser, J. L.; 169, with permission from Elsevier. Suzuki, A.; Pint, B.; date materials with liquid lithium under neutron irradiation. aterials, Nice, France, Dec 10–14, 2007, ICFRM2007/488. Ceramic Coatings as Electrical Insulators in Fusion Blankets 697 4.21.3.5 Radiation Effects Radiation-induced conductivity (RIC) is the loss of insulation only during irradiation, which is a common issue for insulator ceramics in irradiation environ- ments. Historically, evaluation of RIC has been car- ried out mostly for Al2O3 33 (see also Chapter 4.22, Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors). Figure 10 shows RIC as a function of the dose rate for Er2O3, Y2O3, and CaZrO3 bulk specimens for 14MeV neutron, fission neutron, and g-ray irradiation, in comparison with data on Al2O3. 34,35 For Y2O3 and AlN, results for the coating are also shown. The RIC of the candidate materials of Er2O3, Y2O3, and CaZrO3 are comparable with Al2O3. According to these results and the expected dose rate in an Li–V fusion blanket36, the expected induced conductivity in the fusion blanket condition is much lower than the maximum allowable value of �10�2 Sm�1.37 The effects of the nuclear properties of Er on radio- activity and tritium breeding ratio (TBR) of V–Li blankets were also investigated. Figure 11 shows the contact dose rate of a V–Li blanket with and without neutron multiplier Be in the cases of (1) no coating, (2) 10mm,and (3) 1mmcoatingwithEr2O3.Without the coating, the dose rate was dominated by V–4Cr–4Ti substrates reaching the hands-on recycle limit after several decades of cooling. (Note that impurities in g-rayDT neutron Bulk Coa + Y2O3 Y2O3 10–4 10–14 10–13 10–12 10–11 10–10 10–9 10–8 10–7 10–6 10–5 10–3 10–2 10–1 100 CaZrO3 CEr2O3 Er2O AIN Dose r R ad ia tio n- in d uc ed c on d uc tiv ity (S m –1 ) Figure 10 Radiation-induced conductivity as a function of do (JMTR), and g-ray irradiations for bulk and coating of magnetohy data on Al2O3 are also shown in comparison. Adapted from Tan Muroga, T. Fusion Eng. Des. 2005, 75–79, 933–937, with permis V–4Cr–4Ti were not considered in this calculation.) With the coating, the dose rate increases because of the contribution from Er but still satisfies the remote recycling limit.38Er is a neutron-absorbing element and can reduce the TBR especially for an in situ coating where Er is doped into Li. However, because only 0.15% Er is needed in Li for the in situ coating,19 the impact of Er in Li on the TBR is not an issue.39,40 4.21.3.6 FCI Concept as an Alternative to Insulator Coating In recent years, FCIs made of ceramic materials such as silicon carbide composite (SiC/SiC) have been proposed for both electrical and thermal insulation between the liquid breeder and the channel walls. Although the electrical resistivity of SiC/SiC is lower than that of the candidate insulator coating materials, the use of an insert that is much thicker than the coating allows for sufficient reduction of induced electrical currents. An FCI is attractive for application to dual- coolant Li–Pb blanket concepts in which heat is removed by both a high flow rate of He and Li–Pb.41,42 This concept, however, may not be applied to self-cooled liquid metal blankets because these blankets need to avoid thermal insulation between the coolant and the first wall or the blanket structural components and coolants. Fission neutron First wall in V-Li blanket samples tings Variation range Past data on Al2O3 101 102 103 104 105 aZrO3 3 Er2O3 Y2O3 ate (Gy s–1) se rate measured by DT-neutron (FNS), fission neutron drodynamic insulator coating candidate ceramics. Previous aka, T.; Shikama, T.; Narui, M.; Tsuchiya, B.; Suzuki, A.; sion from Elsevier. 2 Recycling limit 22 10–4 10–3 10–2 10–1 100 101 102 103 10–10 10–8 10–6 10–4 10–2 100 102 104 D os e ra te (S v h– 1 ) Time after shutdown (year) Without coating (Li-V) Without coating (Li-Be-V) With coating (Li-V) With coating (Li-Be-V) Hands on limit Structural materials (V-4Cr-4Ti) Coating thickness Neutron wall load:1.5 MW m–2 10 years operation 10μm 1μm Figure 11 Contact dose rate of V–Li blanket with and without neutron multiplier Be in the cases of (1) no coating, (2) 10mm, and (3) 1 mm coating with Er2O3. Reproduced from Muroga, T.; Tanaka, T.; Kondo, M.; Nagasaka, T.; Xu, Q. Fusion Sci. Technol. 2009, 56, 897–901, with permission from Elsevier. 698 Ceramic Coatings as Electrical Insulators in Fusion Blankets 4.21.4 Summary and Remaining Issues Development of an insulator coating is the criti- cal feasibility issue for self-cooled Li blankets. A very limited number of ceramics are stable under long-term exposure to Li at high temperatures. The present candidates are Er2O3, Y2O3, and, under lim- ited conditions, AlN. In addition to concept verifica- tion studies using several physical coatings, chemical or reactive coatings have been explored as a potential means to cover large and complex surfaces. Considering the very low tolerable crack fraction ( Ceramic Coatings as Electrical Insulators in Fusion Blankets 699 11. Smith, D. L.; Natesan, K.; Park, J. H.; Reed, C. B.; Mattas, R. F. Fusion Eng. Des. 2000, 51–52, 185–192. 12. Vertkov, A. V.; Evtikhin, V. A.; Lyublinski, I. E. Fusion Eng. Des. 2001, 58–59, 731–735. 13. Park, J. H.; Kassner, T. F. J. Nucl. Mater. 1996, 233–237, 476–481. 14. Najmabadi, F. The ARIES Team. Fusion Eng. Des. 1997, 38, 3–25. 15. Pint, B. A.; DeVan, J. H.; DiStefano, J. R. J. Nucl. Mater. 2002, 307–311, 1344–1350. 16. Suzuki, A.; Muroga, T.; Pint, B. A.; Yoneoka, T.; Tanaka, S. Fusion Eng. Des. 2003, 69, 397–401. 17. Nagura, M.; Suzuki, A.; Muroga, T.; Terai, T. Fusion Eng. Des. 2009, 84, 1384–1387. 18. Terai, T.; Mitsuyama, T.; Yoneoka, T.; Tanaka, S. J. Nucl. Mater. 1998, 253, 219–226. 19. Yao, Z.; Suzuki, A.; Muroga, T.; Katahira, K. J. Nucl. Mater. 2004, 329–333, 1414–1418. 20. Yao, Z.; Suzuki, A.; Muroga, T.; Yeliseyeva, O. I.; Nagasaka, T. Fusion Eng. Des. 2006, 81, 951–956. 21. Heo, N. J.; Nagasaka, T.; Muroga, T.; Matsui, H. J. Nucl. Mater. 2002, 307–311, 620–624. 22. Hoelzer, D. T.; West, M. K.; Zinkle, S. J.; Rowcliffe, A. F. J. Nucl. Mater. 2000, 283–287, 616–621. 23. Chikada, T.; Suzuki, A.; Yao, Z.; Sawada, A.; Terai, T.; Muroga, T. Fusion Eng. Des. 2007, 82, 2572–2577. 24. Yeliseyeva, O.; Muroga, T.; Yao, Z.; Tsisar, V. J. Nucl. Mater. 2009, 386–388, 696–699. 25. Koch, F.; Brill, R.; Maier, H.; et al. J. Nucl. Mater. 2004, 329–333, 1403–1406. 26. Sawada, A.; Suzuki, A.; Maier, H.; Koch, F.; Terai, T.; Muroga, T. Fusion Eng. Des. 2005, 75–79, 737–740. 27. Pint, B. A.; Moser, J. L.; Tortorelli, P. F. Fusion Eng. Des. 2006, 81, 901–908. 28. Sawada, A.; Suzuki, A.; Terai, T. Fusion Eng. Des. 2006, 81, 579–582. 29. Yao, Z.; Suzuki, A.; Levchuk, D.; et al. J. Nucl. Mater. 2009, 386–388, 700–702. 30. Hishinuma, Y.; Tanaka, T.; Nagasaka, T.; et al. Development of Er2O3 coating on liquid blanket components synthesized with MOCVD process. In 14th International Conference on Fusion Reactor Materials, Sapporo, Japan, Sept 6–11, 2009. 31. Pint, B. A.; Moser, J. L.; Jankowski, A.; Hayes, J. J. Nucl. Mater. 2007, 367–370, 1165–1169. 32. Pint, B. A.; Pawel, S. J.; Howell, M.; et al. J. Nucl. Mater. 2009, 386–388, 712–715. 33. Shikama, T.; Zinkle, S. J.; Shiiyama, K.; Snead, L. L.; Farnum, E. H. J. Nucl. Mater. 1998, 258–263, 1867–1872. 34. Tanaka, T.; Shikama, T.; Narui, M.; Tsuchiya, B.; Suzuki, A.; Muroga, T. Fusion Eng. Des. 2005, 75–79, 933–937. 35. Tanaka, T.; Nagayasu, R.; Sawada, A.; et al. J. Nucl. Mater. 2007, 367–370, 1155–1159. 36. El-Guebaly, L. A.; The ARIES Team. Fusion Eng. Des. 1997, 38, 139–158. 37. Mattas, R. F.; Smith, D. L.; Reed, C. B.; et al. Fusion Eng. Des. 1998, 39–40, 659–668. 38. Muroga, T.; Tanaka, T.; Kondo, M.; Nagasaka, T.; Xu, Q. Fusion Sci. Technol. 2009, 56, 897–901. 39. El-Guebaly, L. A. Fusion Eng. Des. 2006, 81, 1327–1331. 40. Muroga, T.; Tanaka, T.; Sagara, A. Fusion Eng. Des. 2006, 81, 1203–1209. 41. Noajitra, P.; Buehler, L.; Fischer, U.; Malang, S.; Reimann, G.; Schnauder, H. Fusion Eng. Des. 2002, 61–62, 449–453. 42. Wong, C. P. C.; Malang, S.; Sawan, M.; et al. Fusion Eng. Des. 2006, 81, 461–467. 43. Levchuk, D.; Levchuk, S.; Maier, H.; Bolt, H.; Suzuki, A. J. Nucl. Mater. 2007, 367–370, 1033–1037. 4.22 Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors E. R. Hodgson Euratom/CIEMAT Fusion Association, Madrid, Spain T. Shikama Tohoku University, Sendai, Japan � 2012 Elsevier Ltd. All rights reserved. 4.22.1 Introduction 702 4.22.2 Fusion-Relevant Radiation Damage in Insulating Materials 703 4.22.3 Simulation Experiments 705 4.22.4 Degradation of Insulator Electrical Resistance 706 4.22.5 Degradation of Insulator AC/RF Dielectric Properties 712 4.22.6 Degradation of Insulator Thermal Conductivity 715 4.22.7 Degradation of Optical Properties 717 4.22.8 Concluding Remarks 720 References 721 Abbreviations AC/RF Alternating current/radio frequency BA Broader approach CDA Conceptual design activity CIEMAT Centro de Investigaciones Energéticas, Medioambientales, y Tecnológicas CVD Chemical vapor deposition DC Direct current DEMO Demonstration ECRH Electron cyclotron resonant heating EDA Engineering design activity EVEDA Engineering Validation and Engineering Design Activities FIRE Fusion ignition research experiment H&CD Heating and current drive HFIR High Flux Isotope Reactor (Oak Ridge, USA) HFR High Flux Reactor (Petten, Holland) ICRH Ion cyclotron resonant heating IEA International Energy Agency IFMIF International Fusion Materials Irradiation Facility IMR Institute for Materials Research IR Infrared ITER International Thermonuclear Experimental Reactor (Cadarache, France) JET Joint European Torus (Culham, UK) KfK Kernforschungszentrum Karlsruhe (Germany) KU1, KS-4V Russian radiation-resistant quartz glasses LAM Low-activation materials LAMPF Los Alamos Meson Physics Facility (USA) LH Lower hybrid LIDAR Light Detection and Ranging MACOR Machinable Glass Ceramic (Corning Incorporated) MI Mineral insulation/insulated NBI Neutral beam injector ORNL Oak Ridge National Laboratory (USA) OSIRIS From the Greek for Us-yri ‘the king’ (Reactor at Saclay, France) PIE Postirradiation examination RAFM Reduced activation ferritic martensitic RIA Radiation-induced absorption RIC Radiation-induced conductivity RIED Radiation-induced electrical degradation RIEMF Radiation-induced electromotive force RF Russian Federation RL Radioluminescence SCCG Subcritical crack growth TEM Transmission electron microscope UV Ultraviolet 701 702 Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors 4.22.1 Introduction It is envisaged that early in the twenty-first century ITER (International Thermonuclear Experimental Reactor) will come into operation, and it is expected that this intermediate ‘technology’ machine will help to bridge the gap between the present-day large ‘physics’ machines and the precommercial DEMO power reactor, thus paving the way for commercial fusion reactors to become available by the end of the century. Although this ‘next-step’ device will undoubtedly help to solve many of the problems, which still remain in the field of plasma confinement, it will also present additional operational and experi- mental difficulties not found in present-daymachines. These problems are related to the expected radia- tion damage effects as a result of the intense radiation field from the ‘burning’ plasma. This ignited plasma will give rise to high-energy neutron and gamma fluxes, penetrating well beyond the first wall, from which one foresees a serious materials problem that has to be solved. In the initial physics phase of opera- tion of such a machine, it is the radiation flux, which will be of concern, whereas in the later technology phase, both flux and fluence will play important roles as fluence (dose)-dependent radiation damage builds up in the materials. For structural metallic materials, radiation damage in ITER is expected to be severe, although tolerable, only near to the first wall. How- ever, the problem facing the numerous insulating components is far more serious because of the neces- sity to maintain not only the mechanical, but also the far more sensitive physical properties intact. An addi- tional concern arises from the need to carry out inspection, maintenance, and repair remotely because of the neutron-induced activation of the machine. This ‘remote handling’ activity will employ machin- ery, which requires the use of numerous standard components ranging from simple wires, connectors, and motors, to optical components such as windows, lenses, and fibers, as well as electronic devices such as cameras and various sophisticated sensors. All these components use insulating materials. It is clear, there- fore, that we face a situation in which insulating materials will be required to operate under a radiation field, in a number of key systems from plasma heating and current drive (H&CD), to diagnostics, as well as remote handling maintenance systems. All these sys- tems directly affect not only the operation, but also the safety, control, and long-term reliability of the machine. Even for ITER, the performance of some potential insulating materials appears marginal. In the long term, beyond ITER, the solution of the materials problem will determine the viability of fusion power. The radiation field will modify to some degree all of the important material physical and mechanical properties. Some of the induced changes will be flux dependent, while others will be modified by the total fluence. Clearly, the former flux-dependent pro- cesses will be of concern from the onset of operation of future next-step devices. The fluence-dependent effects on the other hand are the important para- meters affecting the component or material lifetime. The properties of concern which need to be consid- ered for the many applications include electrical resistance, dielectric loss, optical absorption, and emission, as well as thermal and mechanical proper- ties. Numerous papers have been published discuss- ing both general, and more recently, specific aspects of radiation damage in insulating materials for fusion applications, and those most relevant to the present chapter are included.1–26 In recent years, because of the acute lack of data for insulators and the recognition of their high sensi- tivity to radiation, most work has concentrated on the immediate needs for ITER. A comprehensive cera- mics irradiation program was established to investi- gate radiation effects on a wide range of materials for essentially all components foreseen for H&CD and diagnostics in ITER, and to find solutions for the problems which have been identified. A large number of relevant components and candidate materials have been, and are being, studied systematically at gradu- ally increasing radiation dose rates and doses, under increasingly realistic conditions. A considerable vol- ume of the work discussed here was carried out within the ITER framework during the CDA, EDA, and EDA extension (Conceptual and Engineering DesignActivities 1992–2002) as specific tasks assigned to the various Home Teams (T26/28 and T246; EU, JA, RF, US; T252/445 and T492; EU, JA, RF).27,28 Since these last ITER tasks, no new coordinated tasks related to insulators have been formulated. However, despite the lack of an official framework in which to develop and assign further common tasks following the end of the ITER-EDA extension, col- laborative work has continued between the EU, JA, RF, and US Home Teams on both basic and applied aspects of radiation damage in insulator materials. This has resulted in considerable progress being made on the understanding of the pertinent effects of radiation on in-vessel components and materials in particular for diagnostic applications. Problemswhich have been addressed and for which irradiation testing Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors 703 has been performed include comparison of absorption and luminescence for different optical fibers andwin- dow materials, RIEMF (radiation-induced electro- motive force) and related effects for MI (mineral insulated) cables and coils, alternative bolometers to the reference JET type gold on mica, hot filament pressure gauges, and electric field effects in aluminas. One must however remember that ITER is only an intermediate ‘technology’ machine on the road to a precommercial power reactor. Such power reactors will face the same radiation flux problems as antici- pated in ITER, but the fluence problems will be far more severe. It is also important to note that the radiation flux and fluence levels will be different from one type of device to another depending on the design (e.g., in ITER and the Fusion Ignition Research Experiment (FIRE)26), and also on the spe- cific location within that device. Because of the gen- eral uncertainty in defining radiation levels, most radiation effects studies have taken this into account by providing where possible data as a function of dose rate (flux), dose (fluence), and irradiation tempera- ture. Although the task ahead is difficult, important advances are being made not only in the identifica- tion of potential problems and operational limita- tions, but also in the understanding of the relevant radiation effects, as well as materials selection and design accommodation to enable the materials lim- itations to be tolerated. Following a brief introduction to the problem of radiation damage in both metals and insulators, the important aspect of simulating the operating envi- ronment for the component or material under exam- ination will be presented, with reference to present experimental procedures. The chapter will then con- centrate on the problems facing the use of insulators, with the radiation effects on the main physical prop- erties being discussed, concentrating in particular on the dielectric properties. 4.22.2 Fusion-Relevant Radiation Damage in Insulating Materials The study of intense radiation effects in metals has been closely associated with the development of nuclear fission reactors, and as a result at the begin- ning of the 1980s when the urgent need to consider radiation damage aspects of materials to be employed in future fusion reactors was fully realized, a consid- erable amount of knowledge and expertise already existed for metallic materials.29 This was not the case for the insulating materials, mainly because of the fact that the required use of insulators in fission- type reactors is in general limited to low radiation regions, well protected from the reactor core. How- ever, despite the late start and the reduced number of specialists working in related fields at the time, together with the complexity of the mechanisms involved in radiation damage processes in insulators, considerable progress has been made not only in assessing the possible problem areas, but also in finding viable solutions. Several general reviews give a good introduction to the specific problem of radiation damage in insulators.30–36 The materials employed in the next-step fusion machine will be subjected to fluxes of neutrons and gammas originating in the ignited plasma. The radiation intensity will depend not only on the dis- tance from the plasma, but also in a complex way on the actual position within the machine because of the radiation streaming along the numerous pene- trations required for cooling systems, blanket struc- tures, heating systems, and diagnostic and inspection channels, as well as the radiation coming from the water in the outgoing cooling channels due to the 16O(n, p)16N nuclear reaction. However one-, two-, and even three-dimensional models are now available, which enable the neutron and gamma fluxes to be calculated with confidence at most, if not all, machine positions.37–40 Radiation damage is generally divided into two components: displacement damage and ionization effects. In a fusion environment, displacement dam- age, which affects both metals and insulators, will result from the direct knock-on of atoms/ions from their lattice sites by the neutrons, giving rise to vacancies and interstitials. Those primary knock-on atoms (PKAs) with sufficient energy may go on to produce further displacements, so-called cascades. The numerous point defects thus produced may either recombine, in which case no net damage results, or they may stabilize and even aggregate producing more stable extended defects. These sec- ondary processes which determine the fate of the vacancies and interstitials are governed by their mobilities. These mobilities are highly temperature dependent, and in the case of insulators even depend on the ionizing radiation level (radiation-enhanced diffusion). Displacement damage is measured in ‘dpa’ (displacements per atom) where 1 dpa is equivalent to displacing all the atoms once from their lattice sites. At the first wall of ITER, the primary displacement dose rate will be of the order of 10�6 dpa s�1. 704 Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors In contrast, ionizing radiation although absorbed by both metals and insulators, in general, only produces heating in metals. However, certain aspects of radia- tion damage in metals, such as radiation-enhanced corrosion and grain boundary modification are related to ionization. The effects of ionization on insulators are in comparison quite marked because of the exci- tation of electrons from the valence to the conduction band giving rise to charge transfer effects. Ionizing radiation is measured in absorbed dose Gy (Gray) where 1Gy¼ 1 J kg�1. At the first wall of ITER, the dose rate will be of the order of 104Gy s�1. The response of insulators to both displacement and ionizing radiation is far more complex than in the case of metals. Apart from a few specific cases (diamond for example), insulating materials are polyatomic in nature. This leads to the following: (i) We have in general two or more sublattices which may not tolerate mixing. (ii) This gives rise to more types of defects than can exist in metals. (iii) Because of the electrically insulating nature, the defects may have different charge states, and hence different mobilities. (iv) The displacement rates and thresholds, as well as themobilities, may be different on each sublattice. (v) We may have interaction between the defects on different sublattices. (vi) Defects can be produced in some cases by purely electronic processes (radiolysis); however, in the insulating materials of interest for fusion, this is generally not the case. As a consequence of these factors, while radiation damage affects all materials, the insulators are far more sensitive to radiation damage than metals. While stainless steel, for example, can withstand sev- eral dpa and GGy with no problem, some properties of insulating materials can be noticeably modified by as little as 10�5 dpa or a few kGy. Because of this, the present ongoing programs of radiation testing for diagnostics are concentrating mainly on the insulat- ing components of the systems. The results of these radiation damage processes are flux- and fluence- dependent changes in the physical and mechanical properties of the materials, which may be particularly severe for the insulators. The properties of concern which suffer modification are the electrical and thermal conductivity, dielectric loss and permittivity, optical properties, and to a lesser extent the mechan- ical strength and volume. The effects of such changes are that the insulators may suffer Joule heating because of the increased electrical conductivity or lower thermal conductivity, and absorption in windows and fibers can increase from the microwave to the optical region and they emit strong lumines- cence (radioluminescence, RL); in addition, the materials may become more brittle and may suffer swelling. Clearly, some materials are more radiation resistant than others. The organic insulators, which are widely used in multiple applications in general, degrade under purely ionizing radiation and are not suitable for use at temperatures above about 200 �C; as a result their use will be limited to superconducting magnet insulation and remote handling applications during reactor shutdown. Inorganic insulators of the alkali halide class have been widely studied and are used as optical windows; however, they are suscepti- ble to radiolysis (displacement damage induced by electronic excitation) and in general become opaque at low radiation fluences. Of the numerous insulating materials, it is the refractory oxides and nitrides, which in general show the highest radia- tion resistance, and of these the ones which have received specific attention within the fusion program include MgO, Al2O3, MgAl2O4, BeO, AlN, and Si3N4. In addition, different forms of SiO2 and materials such as diamond and silicon have been examined for various window and optical transmis- sion applications. One other aspect of radiation damage that should be mentioned is nuclear transmutation. The high- energy neutrons will produce nuclear reactions in all the materials giving rise to transmutation pro- ducts.1 These will build up with time and represent impurities in the materials, which may modify their properties. The physical properties of insulators are particularly sensitive to impurities. Furthermore, some of these transmutation products may be radio- active and give rise to the need for remote handling and hot cell manipulation in the case of component removal, repair, or replacement. For the structural materials, in the present concepts mainly steel alloys, considerable work has been carried out on the devel- opment of so-called low or reduced activation mate- rials (LAM, RAFM – reduced activation ferritic/ martensitic) for possible use in DEMO and future commercial fusion reactors.41–45 This work with the aim of reducing the amount of nuclear waste has studied not only the substitution of radiological prob- lem alloying elements such as Mo and Nb in steels, but also the viability of other materials such as vana- dium and SiC/SiC composites. In the case of the insulating materials, no equivalent study or Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors 705 development has been carried out, in part because of the small fraction of the total material volume repre- sented by the insulators, and also because the impor- tant physical properties of these materials are expected to be degraded before the transmutation products become of concern. Certainly, for a next-step machine such as ITER, transmutation products, with the possi- ble exception of hydrogen and helium, are not expected to present a serious problem. 4.22.3 Simulation Experiments Within the fusion community, there is an acute awareness of the necessity to construct a suitable irradiation testing facility for materials, which will enable both testing and development of materials for future fusion reactor devices with a fusion-like neutron spectrum. Within this context, both concep- tual and engineering design activities were under- taken during the 1990s within the IEA framework with the view of providing such a facility, the IFMIF (International Fusion Materials Irradiation Facility).46–50 This work has been recently renewed under the EU-Japan Broader Approach (BA) activ- ities with the EVEDA (Engineering Validation and Engineering Design Activities) tasks.51,52 However, at the present time no entirely suitable irradiation testing facility exists, and as a consequence experi- ments have been performed in nuclear fission reac- tors and particle accelerators, as well as g- and X-ray sources, in an attempt to simulate the real operating conditions of the insulating materials and compo- nents. The experiments required must simulate the neutron and g radiation field, that is, the displace- ment and ionization damage rates, the radiation envi- ronment, that is, vacuum and temperature, and also the operating conditions such as applied voltage, or mechanical stress. As will be seen, for the insulator physical properties, it is furthermore essential that in situ testing is carried out to determine whether or not the required physical properties of the material or component are maintained during irradiation. Examples of this include the electrical conductivity, which can increase many orders of magnitude due to the ionizing radiation, or optical windows, which may emit intense RL. Experimental nuclear fission reactors clearly have the advantage of producing a radiation field consist- ing of both neutrons and g-rays, although in most cases the actual neutron energy spectrum and the dpa to ionization and He ratios are not those which will be experienced in a fusion reactor.50 However, it is worthwhile noting that to date experimental fission reactors have mainly been used for irradiations in the metals programs where the emphasis is on the neutron flux and little consideration is given to the g field. As a result, the irradiation channels have in general been designed and installed with this criterion. However, it should be possible to select positions within the reactors which, together with suitable neutron absorber materials and neutron to g converters, provide acceptable radiation fields. The main difficulties with in-reactor experiments come from the inaccessibility of the radiation volume and are concerned with the problem of carrying out in situ measurements and achieving the correct irra- diation environment. While considerable success has been attained in the in situ measurement require- ment, with parameters such as electrical conductivity, optical absorption and emission, and even radiofre- quency dielectric loss being determined, the problem of irradiating in vacuum still remains, with most experiments being performed in a controlled He environment. Irradiation in a controlled atmosphere such as He causes an immediate problem for in situ electrical and dielectric measurements because of the radiation-enhanced electrical conductivity of the gas,53 and even in the case of irradiation in vacuum at about 10�3mbar spurious leakage currents will occur.54 Furthermore, many in-reactor experiments rely on nuclear heating to reach the required temperature, and hence have difficulty maintaining a controlled temperature, in part because of the changes in the reactor power, and also because of the problem of calculating the final sample or component tempera- ture. These aspects will be further discussed later. One additional difficulty comes from the nuclear activation of the sample or component, which gener- ally means that postirradiation examination (PIE) has either to be carried out in a hot cell or postponed until the material can be safely handled. Particle accelerators, on the other hand, are ideal for carrying out in situ experiments in high vacuum and at well-controlled temperatures because of the easy access and the very localized radiation field. High levels of displacement damage and ionization can be achieved with little or no nuclear activation. It is however in the nonnuclear aspect of the radiation field where their disadvantage is evident, and great care has to be taken to ensure that appropriate dis- placement rates are deduced to enable reliable com- parison with the expected fusion damage. A further serious disadvantage is due to the limited irradiation 706 Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors volume and particle penetration depth. This in gen- eral means that only small thin material samples or components can be tested. The present-day situation of materials and com- ponent radiation testing for fusion applications takes full advantage not only of fission reactors and particle accelerators, but also 60Co g irradiation facilities and even X-ray sources. The use of such widely different radiation sources can be justified as long as the influ- ence of the type of radiation on the physical parame- ter of interest is known. This, in certain cases, is true for radiation-induced electrical conductivity and RL for example, where for low total fluences it is the ionizing component of the radiation field which is important. In situ measurements can now be made during irradiation of the important electrical, dielec- tric, and optical properties. In addition other aspects such as mechanical strength and tritium diffusion are being assessed during irradiation. Undoubtedly, suc- cessful modeling could be of help to address this diverse use of irradiation sources; however, general modeling for the insulators has hardly got off the ground because of the difficulties associated with describing radiation effects in polyatomic band- structured materials. As a result, in contrast to the extended activity for metallic structural materials, to date there has been no coordinated activity for the insulators, with only specific models for aspects such as electrical and thermal conductivity being developed. 4.22.4 Degradation of Insulator Electrical Resistance Electrical resistance, more generally discussed in terms of the electrical conductivity (the inverse of the resistance), is an important basic parameter for numerous systems and components including the NBI (neutral beam injector) heating system, ICRH (ion cyclotron resonant heating) windows and sup- ports, magnetic coils, feedthroughs and standoffs, MI cables, and wire insulation. Any reduction in the electrical resistance of the insulator material in these components may give rise to problems such as increased Joule heating, signal loss, or impedance change. The main candidate material for these applications is Al2O3 and is also the one which has been most extensively studied, both in the polycrys- talline alumina form and as single crystal sapphire. To a lesser extent, MgO, BeO, MgAl2O4, AlN, and SiO2 have also been studied. At the present time, three types of electrical degradation in a radiation environment are recognized and have been investi- gated; these are radiation-induced conductivity (RIC), radiation-induced electrical degradation (RIED), and surface degradation. Of these types of degradation, RIC was the first to be addressed in a fusion context, as this enhance- ment of the electrical conductivity is flux dependent and hence a possible cause for concern from the onset of operation of any fusion device. Fortunately, RIC had been studied for many years, and a sound theoretical understanding already existed.55–59 The ionizing component of the radiation field causes an increase in the electrical conductivity because of the excitation of electrons from the valence to the conduction band and their subsequent trapping in levels within the band gap near to the conduction band from where they are thermally excited once again into the conduction band. Figure 1 shows sche- matically RIC as a function of irradiation time and ionizing dose rate (flux). The increase in saturation depends not only on the dose rate as indicated, but also in a complex way on the temperature and sample impurity content, as may be seen in Figure 2 for MgO:Fe.60 Nevertheless, such behavior, including the initial step, is well predicted by theory.57 However, at the dose rates of interest for fusion applications, in the range of approximately 1Gy s�1 to >100Gy s�1, saturation is reached within minutes to seconds, and it is this saturation level which is usually the value of interest. The RIC process can lead to increases in the electrical conductivity of many orders of magnitude. For example, a standard high-purity alumina has a room temperature conductivity of generally less than 10�16 Sm�1, which increases to approximately 10�10 Sm�1 for an ionizing dose rate of only 1Gy s�1.61 The first experiments carried out within a fusion application context, that is, refractory oxide materials, high-dose rates, and temperatures, gave an insight into the effects of dose rate, tem- perature, and material impurity, and established the well-known relationship at saturation, between the total electrical conductivity measured during irradia- tion and the ionizing dose rate: stotal¼ s0þ KRdwhere s0 is the conductivity in the absence of radiation, R is the dose rate, and K and d are constants.59,61–63 Although d� 1, the detailed studies found tempera- ture, dose, and dose rate dependence in this parameter, with extreme values in certain cases ranging between 0.5 and 1.5, and in addition a temperature dependence was observed for K. At the present time, extensive RIC data are available for materials irradiated with X-rays, g-rays, electrons, protons, positive ions, and fission and 2.5 2 1.5 1 0.5 0 0 20 40 Irradiation time (a.u.) In cr ea si ng d os e ra te R IC (a .u .) Schematic RIC 60 80 100 Figure 1 Schematic RIC. Saturation is reached more rapidly at higher dose rate. For fusion applications, it is generally the saturation level that is of interest. 4 3 14 �C, 180 ppm 14 �C, 650 ppm 136 �C, 650 ppm 172 �C, 180 ppm 2 1 0 0 50 100 Irradiation time (min) R IC (p A ) MgO:Fe 0.1 Gy s–1 150 200 Time R IC (3,1) (100,0.1) (e = 0,h = 0) 0 1 0 0.2 0.4 n n( t= ¥ ) 0.6 0.8 1.0 2 gG tN T 3 4 (6,0.2) (10,0.1) (10,1) Figure 2 RIC for single crystal MgO, doped with 180 and 650ppm Fe. g irradiation at 0.1Gy s�1 for different temperatures (14, 136, and 172 �C).60 Theoretical predictions are shown inset. Reproduced from Huntley, D. J.; Andrews, J. R. Can. J. Phys. 1968, 46, 147. Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors 707 14MeV neutrons. Many of the additional results, although in some cases limited to one temperature, and/or one dose rate, add confirmation to the earlier extended studies, but more importantly show that RIC is essentially a function of the ionization, independent of the irradiating particle or source. With very few exceptions, all the data taken together over a range of dose rates from 10-6 1.8 MeV e- 450 �C 700 Gy s–1 10-10 dpa s–1 Al2O3 MgO MgAl2O4 BeO 10-7 10-8100 1000 Impurity content (ppm) R IC (S m −1 ) 104 Figure 4 RIC for different single and polycrystalline materials measured during 1.8MeV electron irradiation at 700Gy s�1, 450 �C, plotted as a function of the estimated total impurity content. The line is of slope �1. Reproduced from Hodgson, E. R. J. Nucl. Mater. 1998, 258–263, 226. 10-2 10-4 10-6 10-8 10-10 10-12 10-14 10-4 10-2 100 102 104 Dose rate (Gy s–1) E le ct ric al c on d uc tiv ity (W -m )- 1 3 Superconducting magnet First wall 2 1 4 Van Lint et al.64 Klaffky et al.59 Pells et al.65 Al2O3 MgO MgAl2O4 Pells et al.65 Hodgson and Clement60 106 108 1010 Figure 3 Representative data for RIC as a function of dose rate for different oxide materials. Irradiation with electrons, protons, and neutrons. Reproduced from Shikama, T.; Pells, G. P. J. Nucl. Mater. 1994, 212–215, 80. 708 Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors where 14MeV neutron results are given together with a small selection of other RIC data. For all the RIC data available, because of the different experimental conditions, it is difficult to draw any conclusions as to the reason for the spread in RIC values at any given dose rate. However, data obtained from electron irradiations of different aluminas and other materials under identical conditions of dose rate and temper- ature give an indication that the RIC is inversely proportional to the sample impurity content.19 From these results (Figure 4), two general conclusions/ indications may be drawn: Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors 709 RIC ðsingle crystalÞ > RIC ðpolycrystalÞ and RIC ðpureÞ > RIC ðimpureÞ However, the indication on the impurity dependence needs to be qualified, as certain impurities intro- duce levels near to the conduction band, and increase the RIC.59,60 This would imply therefore that the vast majority of the impurities in the materials act as recombination centers for the electrons and holes, thereby reducing the free charge carrier life- times, and do not introduce electron levels near to the conduction band. The reduction of the electron lifetime in the conduction band has important con- sequences for the RIED effect in different materials, as discussed below. From all the data available, at the present time one can safely say that RIC is sufficiently ‘well understood’ to allow this type of electrical degrada- tion to be accommodated by the design, and that materials exist which give rise to electrical conduc- tivities �10�6 Sm�1 for ionizing dose rates of up to >103Gy s�1. One only expects possible problems or influence near the first wall. Unfortunately, this is precisely the region where magnetic coil diagnostics that can tolerate only very low leakage conductivity will be employed. It is important to remember that RIC is a flux-dependent effect andwill be present from the onset of operation of the next-step machines. Hence, devices which are sensitive to impedance changes, which will occur for example in MI cables, 2 1.5 1 0.5 0 0 20 40 Irradiatio R IC + R IE D (a .u .) RIC dominates RI Figure 5 Schematic RIED. Initially, during irradiation RIC dom conductivity increases because of permanent degradation. must take RIC into account. Furthermore, as RIC is strongly affected by impurity content, the buildup of transmutation products will modify the RIC with irra- diation time (fluence), although this is not expected to be of serious concern for ITER. In contrast to RIC, RIED is a more serious prob- lem because it has been observed under certain con- ditions to permanently increase, that is, degrade, the electrical conductivity with irradiation dose. Figure 5 shows a schematic RIED-type degradation. The ini- tial increase in the conductivity corresponds to the RIC as described above. Following a certain irradia- tion time, or accumulated dose, the conductivity again begins to increase as s0 degrades. In Al2O3 for which most work has been performed, RIED is observed as a permanent increase or degradation of the electrical conductivity (s0) when a small electric field (�100 kVm�1) is applied during irradi- ation at moderate temperatures (�450 �C). At con- siderably higher temperatures and voltages, but without an irradiation field,67 or for irradiations per- formed without an applied electric field,68 no degra- dation occurs. Even at the present time, this type of degradation is still not fully understood; nor is there general agreement as to whether RIED is a real degradation in the volume. Following the first report of RIED effect in electron-irradiated sapphire (Al2O3) and MgO, 8 numerous experiments were carried out to assess its possible relevance to fusion insulator applications. These addressed the effect of the applied electric 60 n time (a.u.) ED influence Permanent degradation 80 100 inates, but with irradiation time (dose) the measured 710 Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors field, DC or AC/RF69 and voltage threshold,70 the irradiation temperature,71,72 and the ionizing dose rate,73 as well as observations that in addition to electrons, RIED occurred with protons (Figure 674), as,75 and neutrons,76–78 and the observation of RIED effects in other materials, for example, MgAl2O4. 74 In addition, further experiments were performed in which RIED-like effects were also observed in sapphire that was electron irradiated in air,79 for thin Al2O3 films, 80 and MgO insulated cable.81 In contrast, some experiments did not observe any RIED effect, with some reporting enhanced surface conductivity or even cracking of the material.82–88 This led to suggestions that the RIED degradation is not a real volume effect, but is caused by surface contamina- tion.82,86 Because of the potential importance of elec- trical degradation and the uncertainty, extensive discussions on RIED were held at several IEA Workshops,89,90 including the experimental techni- ques employed in the irradiations to separate and identify volume degradation from surface effects. It was pointed out at an early stage of the discussions that important factors such as dose rate, and in partic- ular material-type differences, and irradiation temper- ature, all of which could cause RIEDnot to be observed were not being taken into account.73 For example, under identical conditions RIED was observed in Vitox alumina but not in Wesgo AL995 alumina,75 strongly suggesting a material (possibly impurity and/or grain size) dependence, and further 0 -5 -4.5 -4.0 -10 -15 8.5 9.0 Log10 disp Log10 ion Lo g 1 0 el ec tr ic al c on d uc tiv ity (W -1 m -1 ) Vitox Al2O3 Figure 6 RIED observed in alumina during 18MeV proton irrad permission from Pells, G. P. J. Nucl. Mater. 1991, 184, 177. observations showed that the low purity, large grain size Wesgo AL995 material was highly susceptible to surface degradation when irradiated in high vacuum.91 The in-reactor RIED experiment in HFIR at ORNL also threw light on the complex RIED problem.92,93 Initial results indicated no significant increase in elec- trical conductivity for 12 different samples. However, moderate to substantial electrical degradationwas later reported for some of the sapphire and alumina samples, so material type is an important parameter.94 One of the major difficulties for in-reactor experiments is the determination of s0, the conductivity in the absence of radiation, and its temperature behavior. The use of nuclear heating and the residual reactor radiation level mean that changes in this parameter with temperature and its corresponding activation energy are not gener- ally measured, although these are the main indicators for the onset of degradation; hence, RIED only becomes measurable when the material conductivity in the absence of radiation is larger than the RIC; that is, s0� KRd. Furthermore, some experiments were performed at temperatures either near room tempera- ture85 or above 600 �C,95 considerably outside the expected effective temperature range for RIED of approximately 400–500 �C. In an attempt to clarify the situation, work was performed to identify possible basic causes of RIED. These experiments detected specific volume effects in Al2O3 that are observed only for irradiations carried out with an applied electric field. A marked -3.5 -3.0 -2.5 9.5 10.0 400 �C 500 �C lacements per atom ization dose (Gy) iation, with an applied field of 0.5MVm�1. Reproduced with Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors 711 enhancement of the well-characterized Fþ-center (oxygen vacancy with one trapped electron) was observed,71 and TEM identified large regions of g-alumina within the bulk of RIED degraded Al2O3. 96 The increase in Fþ-center production gave rise to enhanced oxygen vacancy mobility, and led to vacancy aggregation and aluminum colloid for- mation, as may be seen in Figure 7.97 This clarified the observed close similarity between the RIED effect and colloid production in the alkali halides,68 and helped to explain the formation of g-alumina and associated bulk electrical and mechanical degrada- tion.96 The combined work led to a RIED model being formulated, which successfully explained the role of the electric field (both DC and AC/RF), the ionization, and the anion (oxygen) vacancies.98 The model predicted a threshold electric field for degradation depending on the impurity/defect con- centration which, as mentioned above in the discus- sion of RIC, reduces the free electron lifetime. This helps to explain the negative RIED results for Wesgo AL995 alumina where the applied experimental field was below the predicted value of >0.6MVm�1.75,87 It also highlighted the importance of the ionization, in agreement with earlier conclusions.73,84 Additional support for the model, and RIED as a volume effect, came with the TEM identification of aluminum colloids, as well as previously observed g-alumina, in Al2O3 irradiated with an electric field applied. 99 At that time, an alternative model based on charge buildup and breakdown was also developed, but 2 C 1 0 3 4 Ener O p tic al a b so rp tio n (O D cm -1 ) Figure 7 Aluminum colloid band in sapphire irradiated with 1. an electric field of 0.2MVm�1 applied. Reproduced from Moroñ was not extended to explain many of the important observations.100 During the intense activities related to RIED during the 1990s, two important factors emerged, one concerned with surface electrical degradation, and the other related to the importance of the exper- imental irradiation environment. For insulating com- ponents in future fusion devices, surface electrical degradation may prove to be more serious than the RIC and RIED volume effects. At that time, two types of surface degradation were reported, a con- tamination caused by poor vacuum, sputtering, or evaporation,83,88 and a real surface degradation related to radiation-enhanced surface vacuum reduc- tion and possibly impurity segregation.101,102 Both forms are affected by the irradiation environment and ionizing radiation. However, the real surface degradation effect is strongly material dependent, and occurs in vacuum but not in air or helium.102 This stresses the extreme importance of a represen- tative irradiation environment for material testing. Most insulating materials required for fusion applica- tions in ITER and beyond must indeed operate in high vacuum, and in consequence accelerator experiments to study electrical conductivity have been performed in vacuum, whereas to date, with few exceptions,76–78,103,104 in-reactor experiments for technical reasons have been performed in helium. Another significant aspect of in-reactor experiments performed in helium is the radiation-induced leakage current in the gas,53 which makes it difficult to olloid band 5 270 �C 290 �C 310 �C gy (eV) 6 8MeV electrons at different temperatures with o, A.; Hodgson, E. R. J. Nucl. Mater. 1997, 250, 156. 712 Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors determine volume conductivity.81,104 One should also mention that severe electrical surface degradation has recently been observed when oxide insulator materials are bombarded with keV H and He ions.105 The mechanism giving rise to such surface degrada- tion is believed to be the loss of oxygen from the vacuum insulator surface region due to preferential radiolytic sputtering. Similarly, in future fusion devices such as ITER ceramic insulators and win- dows may also degrade, as they will be bombarded by energetic H isotope and He ions because of ionization of the residual gas by g radiation and acceleration by local electric fields.54 At the present time, the role of the irradiation environment in electrical degradation clearly requires further study. Additional difficulties experienced in performing in-reactor experiments include temperature control and also component testing.104,106–108 It is also impor- tant to note that several in-reactor experiments have suffered from electrical breakdowns related to the difficulty of maintaining high voltages in a radiation field, precisely what is required for some H&CD and diagnostics systems in a next-step device. Whether or not these are due to RIED, temperature excursions, He gas breakdown, or problems with the MI cables, terminations, and feedthroughs remains unexplained. 4.22.5 Degradation of Insulator AC/RF Dielectric Properties As with the DC electrical properties, it soon became apparent, even before ITER CDA, that data for radi- ation effects on the AC/RF dielectric properties (dielectric loss and permittivity) of suitable insulating materials for fusion applications were almost nonex- istent. Such materials will be needed for both H&CD and diagnostic applications, where they will be required to maintain their dielectric properties from kHz to GHz under a radiation field in high vacuum. Initial work concentrated on the characterization of candidate materials (Al2O3, MgAl2O4, BeO, AlN, and Si3N4), and also PIE of neutron- and proton- irradiated materials.109–114 In general, changes in permittivity were observed to be small (�5%) and considered to be acceptable for fusion applications. However, results for dielectric loss (loss tangent mea- surements) showed orders of magnitude variation for similar materials (�10�5–10�2 for different forms of alumina at 100MHz) even before irradiation. To address this problem, a standard material (MACOR) was distributed and measured by the main laboratories involved (EU, JA, US) to check the dif- ferent measuring systems used. However, the results showed good agreement,115 and the large variation in reported loss tangent values was later shown to be real, in part because of the effect of the different impurity contents of the materials.116,117 This may be clearly seen in Figure 8, where loss tangent data for different aluminas over a wide frequency range are given, showing marked absorption band struc- tures due to polarizable defects (impurities).116 During the early postirradiation loss tangent measurements, there was an indication of recovery, suggesting that loss during irradiation could be signif- icantly higher.65,109–111 This implied that the already difficult measurements should be made in situ during irradiation. In a simple way, dielectric loss can be considered as being due to two contributions: Loss a ðDC conductivityÞ=Frequency þ Polarization term Clearly, both terms can be modified by the radiation. RIC and RIED will increase the DC conductivity and give rise to dose rate (flux) and dose (fluence) effects, although the contribution will decrease with frequency. The polarization term depends on the defects in the material, which exist as, or can form, dipoles through charge transfer processes due to ionization (impurities, vacancies), and produces the absorption band structure observed in the loss as a function of frequency (Figure 8). This term also gives rise to both flux and fluence effects. Furthermore, defects which are modified by radiation-induced charge transfer processes, for example, levels in the band gap occupied by electrons from the conduction band, are unstable and decay after irradiation. This process is responsible for the slowdecrease in electrical conductivity observed at the end of RIC experiments, and will similarly cause a slow decrease in the polari- zation term. Hence, the initial observations of recovery in dielectric loss are to be expected, and the effort required to make measurements during irradiation fully justified. Following the earlier measurements made during X-ray and proton irradiation,65,109,118 work concen- trated on the needs for ICRH at about 100MHz with the first measurements being made during pulsed neutron irradiation (Figure 9).119,120 These pulsed neutron experiments with ionizing dose rates >104Gy s�1 found increases in loss of only about a factor 4. Such a small increase is not compatible with the PIE results, which indicated that the order of 7 6 5 4 3 2 1 0 0.1 0.14 0.18 Time (s) AIN Sapphire (´10) Lo ss t an ge nt (´ 10 -3 ) 0.22 0.26 0.3 Figure 9 The first in-reactor loss measurements at 100MHz during a narrow neutron pulse (14ms FWHM), showing the slow recovery for AlN. Reproduced from Stoller, R. E.; Goulding, R. H.; Zinkle, S. J. J. Nucl. Mater. 1992, 191–194, 602, with permission from Zinkle. 10-2 10-4 10-6 102 105 Frequency (Hz) ICRH LH ECRH Lo ss t an ge nt 108 1011 Figure 8 Loss tangent versus frequency for different aluminas and sapphire (lowest loss). Reproduced from Mollá, J.; Heidinger, R.; Ibarra, A. J. Nucl. Mater. 1994, 212–215, 1029, with permission from Molla. Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors 713 magnitude increases during irradiation. This discrep- ancy may be related to the pulsed nature of the irradiation; although the peak dose rate was high, the integrated dose is only about 500Gy per pulse, far too low for RIC to reach saturation.59–63 However, recent results indicate that for low dose (fluence), that is, at the beginning of operation, the influence of the DC conductivity term (RIC) is small for frequencies above about 1MHz even for dose rates>1 kGy s�1.121 Furthermore, in these pulsed experiments, the dpa per 714 Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors pulse (�10�7 dpa) is too small to affect either the DC conductivity (RIC) or the polarizable defects, even though this term at these dose rates becomes important even down to 100 kHz. Candidate RF heating systems for ITER (IC, ion cyclotron; LH, lower hybrid; EC, electron cyclotron) operating at about 100MHz, 5GHz, and 200GHz will require insulators (feedthroughs, stand- offs, windows) to operate with large electric fields in a radiation field. In general, the in situ experiments employed low-voltage RF, and the question then arises as to whether RIED could possibly affect the dielectric loss.120 At a time of intense RIED activity, two quite different theoretical models were presented in an attempt to explain why the application of a relatively small electric field during irradiation can substantially modify the damage production pro- cess and lead to volume electrical degradation.98,100 The earlier model was based on charge buildup and breakdown, that is, a DC mechanism, but failed to explain many of the results observed during RIED experiments.100 The later model however ex- plained the role of the ionization taking into account the production of highly unstable Fþ-centers,122 the electric field threshold, as well as g-alumina and colloid production, but more importantly predicted that RIED could occur for applied fields at frequen- cies >100GHz.98 This was in agreement with early observations of RIED from DC to >100MHz, and indications for RIED at frequencies above 1GHz.69 Dielectric loss measurements at 15GHz, made during electron irradiation at 2 kGy s�1, and postirradiation from 1 kHz to 15GHz, for sapphire, alumina, BeO, and MgAl2O4, show very varied results. 123,124 Sap- phire, the purest alumina grade, and BeO showed no prompt increase in loss, nor with a dose up to 50MGy. However, the 999 and 997 alumina grades showed significant prompt and dose-dependent increases in loss, consistent with a modification in the polarization term. Furthermore, these in situmea- surements show postirradiation recovery similar to the early reports for proton- and neutron-irradiated materials.65,109–111 In addition, sapphire samples, which had been preirradiated to 7MGy, 10�6 dpa at 450 �C with a DC electric field (210 kVm�1) to pro- duce RIED showed a significant increase in the loss (2� increase), and also in the prompt dielectric loss (�5� increase). Similar increases have only been observed for sapphire neutron irradiated, with- out an electric field applied, to >10�3 dpa.9 In this context, one should also mention recent work concerned with RF ion sources for NBI systems, where in situ measurements of dielectric loss during and following electron irradiation of alumina (Dera- nox 999) to 110MGy with a 1MHz RF voltage (0.8MVm�1) applied indicate a permanent increase in loss for irradiation at 240 �C, but not at 120 �C, as expected from previous RIED studies.125 While various alumina and BeO grades were available with adequate initial properties (dielectric loss, thermal conductivity, and mechanical strength) before irradiation for NBI, IC, and even LH applica- tions, and with potential to withstand the expected ITER radiation levels, this was not the case for ECRH windows. Sapphire or high-purity alumina, the initial ECRH window reference materials with low dielectric loss in the MHz to GHz range,116,126–128 exhibit increasing loss with increasing frequency reaching �10�4 (loss tangent) by 100GHz. Hence, to transmit the megawatts of RF power that will be required,9 these materials would have to be employed at cryogenic temperatures, and furthermore with a very low neutron tolerance level, �1020 nm�2.128 However, in recent years, considerable progress has been made with CVD diamond, a material with the required combination of low dielectric loss, high ther- mal conductivity, and mechanical strength.19,25,129–134 In this context, initial work began to examine both high-purity silicon and diamond homopolar crystal- line materials which as a result of their decreasing loss with increasing frequency offered the possibility for operation at frequencies above 150GHz with loss tangents �10�5, at room temperature.129 These two materials required development in completely oppo- site directions. The initial high-resistivity silicon had very low loss but extreme radiation sensitivity. Because of its perfection, electrons excited into the conduction band by purely ionizing radiation had very long life- times (no defect recombination sites) leading to high dielectric loss through the high electrical conductiv- ity. In contrast, the CVD diamond, initially almost black in color, had high loss because of the numerous defects in the material giving rise to polarization losses, but was almost insensitive to ionizing radia- tion because of the extremely short lifetime of the conduction band electrons. Although the radiation sensitivity of silicon could be notably reduced by electron irradiation and also by Au doping because of the introduction of recombination defects, the main limitation for silicon comes from its small 1.1 eV band gap. This allows electrons to be readily thermally excited into the conduction band at tem- peratures only slightly above room temperature, 0.9 Window grade CVD diamond 145 GHz 0.6 0.3 0.0 150 200 250 Temperature (K) D ic l. lo ss t an ge nt (1 0- 4 ) 300 Unirr. 10–21n m−2 10–22n m−2 Figure 10 Diamond dielectric loss at 145GHz while being unirradiated, and neutron irradiated at 320K (pool temperature) to 1021 and 1022 nm�2. Reproduced with permission from Thumm, M.; Arnold, A.; Heidinger, R.; Rohde, M.; Schwab, R.; Spoerl, R. Fusion Eng. Des. 2001, 53, 517. Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors 715 which rapidly increases the dielectric loss.135–138 In the case of CVD diamond, the progress has been remarkable, available samples going from black and irregular in shape to almost transparent 2mm thick 100mm diameter disks, with room temperature loss�1� 10�5 at 145GHz, comparable with sapphire at 77K, and furthermore increasing only to about 5� 10�5 by 450 �C.130,132 Loss measurements during electron and X-ray irradiation at 18 and 40GHz, respectively of the developed CVD diamond, show almost negligible contributions of conductivity (RIC) and polarizable defects, and successful high-power transmission tests have now been carried out.132,133 As may be seen in Figure 10, PIE loss tangent mea- surements of neutron-irradiated ‘window grade’ CVD diamond indicate that even by 1022 nm�2 (10�3 dpa), the room temperature loss only increases to 5� 10�5 at 145GHz (6� 10�5 at 190GHz).134 During the intense activity to find suitable mate- rials for the high-power IC, LH, and EC heating applications, work was also being carried out on materials for diagnostic systems. In particular, KU1 quartz glass provided by the Russian Federation within the ITER-EDA task sharing agreement was shown to be highly radiation resistant with respect to its optical properties for use in both diagnostic and remote handling applications, and became the main reference material not only for optical windows, but also fibers.26,139,140 In view of this, the material was also examined for possible use in DC and RF appli- cations. Both RIC and RIED, together with dielectric loss and permittivity, have been determined for as-received, as well as electron and neutron irra- diated material. A large number of different experi- mental setups were employed to obtain the dielectric spectrum of KU1 over a very wide frequency range (10mHz to 145GHz), and where possible, values were obtained during electron irradiation. In addition, data have been obtained for samples neutron irra- diated to 10�4 dpa. The results indicate that for low radiation doses the electrical and dielectric properties are only slightly degraded, and in particu- lar the use of KU1 for electron cyclotron emission (ECE) windows and low-loss DC applications is feasible.134,141 4.22.6 Degradation of Insulator Thermal Conductivity Work began at an early stage to assess the thermo- mechanical properties of candidate insulating materi- als for fusion applications. In an attempt to determine the best combination of mechanical, thermophysical, and dielectric properties for the demanding H&CD applications, Al2O3 (both alumina and sapphire), AlN, Si3N4, BeO, and MgAl2O4 in numerous differ- ent grades were examined ‘as-received’ and following irradiation.142–149 At room temperature, the unirradi- ated thermal conductivity of a typical alumina is of the order of 30Wm�1 K�1, and that of BeO about 280Wm�1 K�1. These values are sufficiently high 716 Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors for IC and LH heating systems to ensure adequate cooling in most cases; however, the thermal conduc- tivity in ceramics is reduced because of increased phonon scattering, by the presence of point defects and to a lesser extent by extended defects or aggre- gates. Hence, one expects a reduction in thermal conductivity on irradiation, together with a notable influence of the irradiation temperature, that is, irradiation above temperatures at which the radiation- induced defects become mobile and can either recombine or aggregate should lead to a lower degra- dation of the thermal conductivity, while low- temperature irradiation should have a marked effect because of the increased point defect stability. The expected general behavior was confirmed by the early data (Figure 11), and indicated that a maxi- mum reduction to about one-third of the room temperature thermal conductivity value could be expected.142–145 This will occur for a neutron fluence value (dpa), which strongly depends on the irradia- tion temperature. For near room temperature irradi- ation (300 K), reduction to the lower saturation level was observed by about 1023 nm�2 (0.01 dpa), whereas at 600 K this lower saturation level was only reached following a fluence of above 1024 nm�2. Within rea- sonable margins, these values applied for Al2O3, AlN, and MgAl2O4. Similar PIE results were obtained at a later date for reactor irradiations at different tem- peratures of a wide range of ceramic materials.150 Because of the importance of point defects in the reduction of thermal conductivity, it is reasonable 0.35 0.30 lµT-1 lµT-0.9 lµT-0.7 lµT-0.4 0.25 0.20 0.15 0.10 0.05 0 0 200 400 600 800 Tempe Th er m al c on d uc tiv ity (W cm -1 K ) Figure 11 Effect of neutron and a particle irradiations at differ thermal conductivity (KfK 700K 0.001dpa, LAMPF 600K 0.5 dp Reproduced from Rohde, M.; Schulz, B. J. Nucl. Mater. 1990, 1 to expect that postirradiation measurements may un- derestimate the effect due to possible postirradiation annealing. An attempt tomeasure thermal conductivity in situ during reactor irradiation, although unable to quantify the degradation, did highlight a very rapid decrease in thermal conductivity by �1022 nm�2 (0.001 dpa) at the startup of irradiation, followed by saturation.151 Finally, one should mention the specific case of sapphire and CVD diamond, the original and the present reference materials for ECRH. For sapphire, the need for low-temperature ( 2000 1800 1600 1400 1200 1000 800 600 400 200 0 102210211020Unirradiated Window grade Scale-up grade Torus window disc Refinement grade Refinement grade Th er m al c on d uc tiv ity (W m K -1 ) Fluence (n m-2) Figure 12 Thermal conductivity reduction for different diamond grades as a function of neutron dose. Reproduced from Heidinger, R.; Rohde, M.; Spörl, R. Fusion Eng. Des. 2001, 56–57, 471, with permission from Heidinger. Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors 717 4.22.7 Degradation of Optical Properties Within the fusion program, another area of concern is related to the effects of radiation on the optical prop- erties of the dielectric materials to be used as trans- mission components (windows, lenses, and optical fibers) for the UV, visible, and near-IR wavelength diagnostic systems needed for control and safety, as well as maintenance (remote handling).21,26,140,154,155 Radiation-induced optical absorption (RIA) and light emission or RL impose severe limitations on the use of any optical material within an intense radiation field. For remote handling applications, the optical components will be expected to maintain their trans- mission properties under high levels of ionizing radi- ation (�1Gy s�1) during hundreds of hours. For such applications, RIA imposes the main limitation, but can be tolerated. However, in the case of diagnostic applications, in addition to a higher level of ionizing radiation (tens to hundreds Gy s�1), the materials will also be subjected to atomic displacements �10�9 dpa s�1. It soon became clear that both RIA and RL would impose severe limitations on the main candidate materials (sapphire and silica). Of these two materials, sapphire is by far the most resistant to ionizing radiation. Although ionizing radiation can cause an increase in optical absorption because of trace impurities and vacancy defects present in the material, it is in general the displacement damage mechanism which induces absorption at first in the UV region as a result of oxygen vacancy-related defects.30,33,156–158 This fluence (dose) effect reduces the transmission in the UV region to essentially zero for doses above about 10�4 dpa, and more slowly in the visible as the tails of the absorption bands begin to overlap into this region. Although sapphire shows more radiation resistance than SiO2 in terms of opti- cal absorption, the material was found to be unsuit- able for many diagnostic applications because of its intense RL, as will be seen below. As with RIC, RL is ionizing flux (dose rate) dependent and hence will be a problem from the onset of operation of future fusion devices. Further- more, to assess RL clearly requires in situ measure- ments during irradiation. While many studies had been carried out on luminescence phenomena in SiO2 and sapphire, the problem was only addressed in a quantitative way because of fusion application requirements.159–164 Sapphire was quickly excluded from high-dose rate applications when it was shown that the photon emission for a typical diagnostic window dose rate would be comparable with the photon emission from the plasma.159 In contrast, certain grades of silica show virtually no RL in the UV-visible region, the emission being limited almost to the Cherenkov background. Quantitative lumines- cence data comparing UV grade sapphire and two types of silica, both of which show low RL, are given in Figure 13, indicating that suitable materials do exist in which the RL can be reduced to a minimum, although there are limited data on RL as a function of fluence.162–164 In particular, the KU1 and KS-4V quartz glass materials, provided by the Russian 1000 100 10 1 0.1 2000 3000 4000 KU1 5000 Sapphire Anhydroguide Wavelength (Å) 10 11 p ho to ns /( s. Å .s r. cm 3 ) 6000 7000 8000 Figure 13 Quantified RL emission for sapphire and two silica grades during 1.8MeV electron irradiation at 700Gy s�1, 15 �C. Reproduced from Moroño, A.; Hodgson, E. R. J. Nucl. Mater. 1998, 258–263, 1889. 718 Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors Federation for the ITER diagnostics radiation testing program, have proved to be highly resistant to RL and RIA because of ionizing radiation and displace- ment damage, and are now reference materials.26,165–170 For ionizing radiation doses up to at least 100MGy and for temperatures at or above about 100 �C, very little absorption is induced in the KU1 material over the whole visible range; one must keep in mind however that with irradiation displacement dose the optical absorption related to oxygen vacancies in SiO2, as in all oxide materials, eventually renders them opaque in the UV and visible range.171–175 In an analogous way to the ECRH transmission windows, mention should be made of windows required for high-power laser transmission, that is, the LIDAR (light detection and ranging) system. This demanding diagnostic system being considered for ITER will require very high-quality transmission windows for the high-power laser pulses at about 500 and 1000 nm. It is estimated that transmission losses of the order of 5%may cause problems with the window integrity because of laser damage. However, such small decreases in the transmission corresponding to an optical density increase of only 0.02 are extremely difficult to measure by standard PIE of irradiated optical materials. Such measurements have to be per- formed in situ. In situ measurement is also required in order to determine possible radiation-enhanced absorption which can easily reach such small values. The possibility of radiation-enhanced dielectric breakdown due to the intense laser pulse and the ionizing radiation has also to be considered. However, such a determination requires an elaborate in situ experiment. Work on laser-induced damage in KU1 and KS-4V has confirmed the limited influence of RIA and RIC on the damage threshold for high- power laser transmission.176 However, metallic depo- sition due to sputtering or evaporation can seriously reduce the damage threshold even for a few nanome- ter thickness, as may be seen in Figure 14. The effect is strongly material dependent, and furthermore self- cleaning with subthreshold laser pulses is not effective for all deposited materials.177,178 Although in general RL is considered to be a problem for diagnostic systems in future devices, it may be employed as a detector/converter for X-ray, UV, and particle emission from the plasma. The intense RL from Al2O3:Cr, for example, has been used for many years in ceramic fluorescent screens for accelerator beam alignment,179 and is now being developed with improved radiation resis- tance and rapid decay times for fusion applications, along with other alternative luminescent materials (Figure 15).180–182 Furthermore, RL is a potentially powerful tool capable of monitoring material modifi- cation during irradiation, but has been largely neglected within the fusion materials activities, in part because of the difficulty in interpreting the resulting emission spectra. However, the technique is now being successfully employed to study insulating materials such as aluminas and silicas, as well as breeding ceramics for fusion applications.183,184 1.2 1 0.8 0.6 0.4 0.2 0 1014 1015 1016 Laser power density (W m-2) D am ag e p ro b ab ili ty KU1 + steel layer KS-4V + steel layer 1017 KU1 KS-4V polished Figure 14 Effect of a thin sputtered steel layer on the laser induced damage for KU1 and KS-4V quartz glasses. Reproduced from Martin, P.; Moroño, A.; Hodgson, E.R. J. Nucl. Mater. 2004, 329–333, 1442. 1.4 1.2 14-MeV neutron irradiation SrAl2O4:Eu 2+, Dy3+ Sr4Al14O25:Eu 2+, Dy3+ 1 0.8 0.6 0.4 0.2 0 300 350 400 450 500 Wavelength (nm) R el at iv e lu m in es ce nt in te ns ity (a .u .) 550 600 650 700 Figure 15 14MeV neutron-induced luminescence in doped aluminates. Reproduced from Toh, K.; Shikama, T.; Katsui, H.; et al. J. Nucl. Mater. 2009, 386–388, 1027. Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors 719 Finally, in connection with optical transmission components, one should note the flexibility and sim- plification in diagnostic design that the use of optical fibers would allow. However, this is not straightfor- ward; although RIA and RL are problems for optical window and lens components, in the case of optical fibers the situation is far worse because of the length of the optical path. Furthermore, because of the manufacturing techniques, fibers with characteristics as good as those observed for the KU1 quartz glass for example have not been obtained. This has prompted an extensive collaborative research program to find the most suitable types of radiation-resistant fiber. Several different optical fibers have been examined to assess RIA and light emission, the viability of high-temperature operation and annealing, jacketing 720 Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors material, and the influence of hydrogen loading. In addition, parallel work is being carried out on the possibility of photobleaching using high-intensity lasers to recover transmission, ‘holey’ fibers for improved transmission and radiation resistance, and fibers with extended blue – UV transmission.26,185–190 Irradiations have been carried out to total doses above 10MGy and 1022 nm�2, and temperatures from about 30 to 300 �C. The most promising fibers are the hydrogen loaded KU1 and KS-4V, where above 400 nm they show the lowest RIA, as may be seen in Figure 16.139 Although the KU1 is the slightly better material up to about 700 nm, the intrinsic OH band and its harmonics notably affect transmission above 800 nm, so for a fiber required to transmit in the visible and IR regions, the hydrogen loaded KS-4V may be a better choice. For silica materials up to about 10MGy, the main radiation damage mechanisms involve electron and hole- trapping; hence, the wide differences observed in induced absorption of the fibers tested are due to variations in intrinsic trapping centers (defects and impurities). In general, these trapping centers are thermally unstable, hence the effective thermal annealing for irradiation at higher temperature, or postirradiation thermal annealing. For higher doses, displacement damage leading to extensive structural damage begins to dominate, but by this time the fibers are of little use for diagnostic applications. Limited work is underway to examine the possibility 16 14 12 10 8 6 4 2 0 450 500 550 600 650 Waveleng 5.2 ´ 1012n cm-2 KU-H2GR ad ia tio n- in d uc ed a b so rp tio n (d B m -2 ) Figure 16 In-reactor radiation-induced optical absorption (RIA (KU-H2G) shows the best performance. Reproduced from Brich et al. J. Nucl. Mater. 2004, 329–333, 1456. of in situ photobleaching of the radiation-induced damage using high-intensity UV lasers, the potential of so-called ‘holey’ fibers (fibers containing an array of vacuum, air, or liquid filled holes) to improve radiation resistance, as well as fibers to extend trans- mission into the blue – UV region. 4.22.8 Concluding Remarks Since the end of the 1970s, the fusion materials com- munity has been providing the necessary insulator research and database for H&CD, and diagnostic systems for a next-step burning plasma device (ITER). As described above, the continuous research has identified and highlighted the limitations and potential problems related to electrical, dielectric, and thermal conductivity degradation, window and fiber absorption, and luminescence, while at the same time providing possible solutions for the ECRH win- dows with the development of diamond for example, identifying safer operating conditions for the insula- tors, assessing optical materials for low RL, and char- acterizing dielectric properties over a wide range of frequencies (IC to EC). In addition to channeling the necessary expertise, numerous unique experiments and installations have been developed to study can- didate materials under relevant conditions, in partic- ular during irradiation. All this has produced data of direct relevance to both ITER and future fusion 700 th (nm) KU1 , 25 MGy, 130 �C F1-doped KS4V 750 800 850 900 ) for four types of fiber. The hydrogen loaded KU1 ard, B.; Fernandez Fernandez, A.; Ooms, H.; Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors 721 devices. As a result of the in-reactor experience, research was then able to concentrate on the required irradiation testing and screening of prototype com- ponents, urgently required for ITER. These have included bolometers, Hall probes, MI cables, coated mirrors, pressure gauges, as well as the optical fibers discussed above.26,108,139,191–193 The irradiation testing required for ceramic insulator materials is far more complex than that required for structural materials. For almost all of the properties of interest, in situ testing is mandatory. The difficulties associated with in situ testing together with the high cost of reactor irradiations have meant that g sources and particle accelerators have been used to full advantage, where the behavior with irradiation temperature and dose rate can be more easily assessed. The need for irradiation in vacuum is an added difficulty; one must remember that most in-reactor irradiations are performed in He. However, despite the difficulties, several experimen- tal systems have been developed to enable testing in both static and dynamic (pumped) vacuum, thereby fully simulating the expected environment. In addi- tion, the hostile, inaccessible, and noisy environment of experimental fission reactors makes the measure- ment of often small but important effects difficult. However, over the years considerable in-reactor test- ing expertise has been gained and much needed experiments performed. One must remember that ITER is only a ‘next step’; the final goal is to provide a safe and reliable fusion reactor within a reasonable time. Undoubtedly, beyond ITER the use of insulators will be severely restricted to those essential to operation and mainte- nance, but they will be of paramount importance to the success of fusion power. Hence, future ceramic insulator research activity, while keeping in mind the short-term ‘urgent’ ITER needs, must address the expected fluence degradation effects on all the material properties and enable viable solutions to be available in time. The problems to be addressed are related to long-term degradation of the required properties because of aggregation and segregation of defects and impurities. For high fluence, not only H and He, but also other transmutation elements will begin to play a role in the modification of the material physical properties. Although not discussed here, mechanical property degradation has been studied since the beginning of the fusion materials activity. 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Intro Editor in chief Section editors Copyright Editors biographies Preface Foreword Permission Acknowledgments 4.01 Radiation Effects in Zirconium Alloys 4.01.1 Irradiation Damage in Zirconium Alloys 4.01.1.1 Damage Creation: Short-Term Evolution 4.01.1.1.1 Neutron-zirconium interaction 4.01.1.1.2 Displacement energy in zirconium 4.01.1.1.3 Displacement cascade in zirconium 4.01.1.2 Evolution of Point Defects in Zirconium: Long-Term Evolution 4.01.1.2.1 Vacancy formation and migration energies 4.01.1.2.2 SIA formation and migration energies 4.01.1.2.3 Evolution of point defects: Impact of the anisotropic diffusion of SIAs 4.01.1.3 Point-Defect Clusters in Zirconium Alloys 4.01.1.3.1 a Dislocation loops 4.01.1.3.2 a Loop formation: Mechanisms 4.01.1.3.3 c Component dislocation loops 4.01.1.3.4 c Loop formation: Mechanisms 4.01.1.3.5 Void formation 4.01.1.4 Secondary-Phase Evolution Under Irradiation 4.01.1.4.1 Crystalline to amorphous transformation of Zr-(Fe,Cr,Ni) intermetallic precipitates 4.01.1.4.2 Irradiation effects in Zr-Nb alloys: Enhanced precipitation 4.01.2 Postirradiation Mechanical Behavior 4.01.2.1 Mechanical Behavior During Tensile Testing 4.01.2.1.1 Irradiation hardening: Macroscopic behavior 4.01.2.1.2 Irradiation hardening: Mechanisms 4.01.2.1.3 Post-yield deformation: Macroscopic behavior 4.01.2.1.4 Post-yield deformation: Mechanisms 4.01.2.2 Effect of Postirradiation Heat Treatment 4.01.2.3 Postirradiation Creep 4.01.3 Deformation Under Irradiation 4.01.3.1 Irradiation Growth 4.01.3.1.1 Irradiation growth: Macroscopic behavior 4.01.3.1.2 Irradiation growth: Mechanisms 4.01.3.2 Irradiation Creep 4.01.3.2.1 Irradiation creep: Macroscopic behavior 4.01.3.2.2 Irradiation creep: Mechanisms 4.01.3.3 Outlook References 4.02 Radiation Damage in Austenitic Steels 4.02.1 Introduction 4.02.2 Basic Damage Processes 4.02.2.1 Atomic Displacements 4.02.2.2 Transmutation 4.02.3 Differences in Neutron Spectra 4.02.4 Transmutation Issues for Stainless Steels 4.02.5 Evolution of Radiation-Induced Microchemistry and Microstructure 4.02.6 A Cross-Over Issue Involving Radiation-Induced Microstructural Evolution and Transmutation 4.02.7 Radiation-Induced Changes in Mechanical Properties 4.02.8 Radiation-Induced Changes in Dimension 4.02.8.1 Precipitation-Related Strains 4.02.8.2 Void Swelling and Bubble Swelling 4.02.8.3 Parametric Dependencies of Void Swelling 4.02.8.3.1 Stress state 4.02.8.3.2 Elemental composition 4.02.8.3.3 Alloy starting state 4.02.8.3.4 Irradiation temperature 4.02.8.3.5 Influence of dpa rate on swelling 4.02.8.3.5.1 Category I of dpa rate effects 4.02.8.3.5.2 Category 2 of dpa rate effects 4.02.9 Irradiation Creep 4.02.9.1 Introduction 4.02.9.2 Stages of Irradiation Creep 4.02.9.3 Examples of Creep Behavior 4.02.9.4 Creep Disappearance 4.02.9.5 Recent Revisions in Understanding of Irradiation Creep 4.02.9.5.1 Dependence of irradiation creep on dpa rate 4.02.9.5.2 Dependence of creep and creep relaxation on neutron spectra 4.02.9.5.3 Dependence of creep modulus on hydrostatic stress 4.02.9.6 Stress Relaxation by Irradiation Creep 4.02.9.7 Stress Rupture 4.02.9.8 Fatigue 4.02.10 Conclusions References 4.03 Ferritic Steels and Advanced Ferritic–Martensitic Steels 4.03.1 Introduction 4.03.2 Basic Metallurgy of Ferritic-Martensitic Steels 4.03.3 Radiation Damage of Core Components in Fast Reactors 4.03.4 Development of Ferritic Steels for Fast Reactor Core 4.03.4.1 Influence of Composition and Microstructure on Properties of Ferritic Steels 4.03.4.2 Void Swelling Resistance 4.03.4.3 Irradiation Hardening in Ferritic Steels 4.03.4.4 Irradiation Creep Resistance of Ferritic Steels 4.03.4.5 Irradiation Embrittlement in Ferritic Steels 4.03.4.5.1 GBE to reduce embrittlement in ferritic steels 4.03.5 Development of Advanced ODS Ferritic Steels 4.03.6 Ferritic Steels for Out-of-Core Applications: Improvements in Joining 4.03.7 Summary References 4.04 Radiation Effects in Nickel-Based Alloys 4.04.1 Introduction 4.04.2 Void Swelling 4.04.2.1 Compositional Dependence of Void Swelling 4.04.2.2 Void-Swelling Models 4.04.2.3 Swelling Behavior of Neutron-Irradiated Nimonic PE16 4.04.3 Irradiation Creep 4.04.4 Microstructural Stability 4.04.4.1 Dislocation Structures 4.04.4.2 Precipitate Stability 4.04.5 Irradiation Embrittlement 4.04.5.1 Fast Neutron Irradiation Experiments 4.04.5.2 Helium Implantation Experiments 4.04.6 Concluding Remarks Acknowledgments References 4.05 Radiation Damage of Reactor Pressure Vessel Steels 4.05.1 Introduction 4.05.2 Pressure Vessel Steels 4.05.3 Effect of Neutron Irradiation on Bulk Properties 4.05.3.1 Summary 4.05.4 Development of Mechanistic Insight of Factors Controlling Current Plant Lifetimes 4.05.4.1 Introduction 4.05.4.2 Radiation Damage Mechanisms 4.05.4.3 Mechanistic Framework 4.05.4.4 Cluster Development Under Irradiation: Cu 4.05.4.4.1 Introduction 4.05.4.4.2 CEC characterization 4.05.4.4.3 Development with increasing fluence 4.05.4.5 Development of Matrix Defects 4.05.4.5.1 Introduction 4.05.4.5.2 Nature 4.05.4.5.3 Development with flux and fluence and irradiation temperature 4.05.4.5.4 Effect of alloying 4.05.4.5.5 MD and hardening 4.05.4.6 Effect of Radiation Damage on Hardening 4.05.4.7 Segregation to Grain Boundaries 4.05.4.8 Summary 4.05.5 Development of Mechanistically Based DDRs 4.05.5.1 Introduction 4.05.5.2 DDRs for CMn Steels 4.05.5.3 US Mechanistically Guided DDRs 4.05.5.4 Japanese Embrittlement Correlations 4.05.5.5 Summary 4.05.6 Current Issues in the Development of DDRs References 4.06 Radiation Effects in Refractory Metals and Alloys 4.06.1 Introduction 4.06.2 Niobium and Nb-Base Alloys 4.06.2.1 Introduction and History of Nb and Nb Alloys 4.06.2.2 Radiation-Induced Swelling of Nb and Nb-Base Alloys 4.06.2.3 Mechanical Properties of Irradiated Nb and Nb Alloys 4.06.3 Tantalum and Ta-Base Alloys 4.06.3.1 Introduction and History of Ta and Ta Alloys 4.06.3.2 Irradiation-Induced Swelling of Ta and Ta-Base Alloys 4.06.3.3 Mechanical Properties of Irradiated Ta and Ta-Base Alloys 4.06.4 Molybdenum and Mo-Base Alloys 4.06.4.1 Introduction and History of Mo and Mo Alloys 4.06.4.2 Irradiation-Induced Swelling and Physical Property Changes in Mo and Mo-Base Alloys 4.06.4.3 Mechanical Properties of Irradiated Mo and Mo Alloys 4.06.5 Tungsten and W-Base Alloys 4.06.5.1 Introduction and Irradiated Properties Database for W and W Alloys 4.06.5.2 Irradiation-Induced Swelling and Physical Property Changes in W and W Alloys 4.06.5.3 Irradiated Mechanical Properties of W and W Alloys 4.06.6 Outlook References 4.07 Radiation Effects in SiC and SiC–SiC 4.07.1 Introduction 4.07.2 Irradiation-Induced Swelling and Microstructure of Pure SiC 4.07.3 Irradiation-Induced Thermal Conductivity Degradation of Monolithic SiC 4.07.4 Effect of Irradiation on the Mechanical Properties of Monolithic SiC 4.07.4.1 Elastic Modulus of Monolithic SiC 4.07.4.2 Hardness of Monolithic SiC 4.07.4.3 Fracture Toughness of Monolithic SiC 4.07.4.4 Strength and Statistical Variation in Strength for Monolithic SiC 4.07.5 Irradiation Creep of SiC 4.07.6 Silicon Carbide Composites Under Irradiation References 4.08 Oxide Dispersion Strengthened Steels 4.08.1 Introduction 4.08.2 Nanoscale Oxide Particle Control 4.08.2.1 Dissociation and Precipitation 4.08.2.2 Structure and Coherency 4.08.3 Martensitic 9Cr-ODS Steels 4.08.3.1 Chemical Composition and Microstructure 4.08.3.2 Residual Ferrite Formation and Strength Characterization 4.08.3.2.1 Mechanically alloyed powder characterization 4.08.3.2.2 Pinning of α-γ interface by oxide particles 4.08.3.2.3 Strength characterization 4.08.3.3 Cladding Manufacturing 4.08.3.3.1 Continuous cooling transformation diagram 4.08.3.3.2 Manufacturing process 4.08.3.4 Creep and tensile properties 4.08.4 Ferritic 12Cr-ODS Steels 4.08.4.1 Strength Anisotropy 4.08.4.2 Recrystallization Tests 4.08.4.3 Cold-Rolling Cladding Manufacturing 4.08.4.4 Internal Creep Rupture Property 4.08.5 Al-Added 16Cr-ODS Steels 4.08.5.1 Application and Technical Issues 4.08.5.2 Thermal Aging Embrittlement Due to High Cr Content 4.08.5.3 Mechanical Properties 4.08.5.4 Cladding Manufacturing 4.08.6 Existing ODS Steel Cladding 4.08.7 Corrosion and Oxidation 4.08.7.1 Sodium Compatibility 4.08.7.2 LBE Compatibility 4.08.7.3 SCPW Compatibility 4.08.7.4 Oxidation 4.08.8 Irradiation 4.08.8.1 Simulated Irradiation 4.08.8.2 Neutron Irradiation of Materials 4.08.8.3 Fuel Pin Irradiation 4.08.8.3.1 9Cr- and 12Cr-ODS steel cladding in BOR-60 4.08.8.3.2 12Cr-ODS steel cladding in EBR-II 4.08.8.3.3 DT2203Y05 in Phénix 4.08.9 Summary References 4.09 Welds for Nuclear Systems 4.09.1 Welding Defects 4.09.1.1 Supersolidus Cracking 4.09.1.1.1 Solidification cracking 4.09.1.1.2 Liquation cracking 4.09.1.1.3 Hot tearing 4.09.1.2 Subsolidus Cracking 4.09.1.2.1 Precipitation-induced cracking 4.09.1.2.2 Segregation-induced cracking 4.09.1.3 Other Welding Defects 4.09.2 Stresses and Strains in Welds 4.09.2.1 Quantification of Residual Stresses and Strains 4.09.2.1.1 Elastic stress 4.09.2.1.2 Plastic strain 4.09.3 In-Service Performance 4.09.3.1 Environmentally Assisted Cracking 4.09.3.2 Microchemical Changes 4.09.3.3 Microstructural Changes 4.09.4 Weldability of Specific Alloy Systems 4.09.4.1 Low-Alloy Steels 4.09.4.2 Austenitic Stainless Steels 4.09.4.3 Nickel-Based Alloys 4.09.4.4 Zirconium Alloys Acknowledgments References 4.10 Radiation Effects in Graphite 4.10.1 Introduction 4.10.2 Nuclear Graphite Manufacture 4.10.3 Graphite-Moderated Reactors 4.10.4 Displacement Damage and Induced Structural and Dimensional Changes in Graphite 4.10.5 Neutron-Induced Property Changes 4.10.5.1 Wigner Energy 4.10.5.2 Mechanical and Physical Properties 4.10.6 Irradiation Creep 4.10.6.1 The Relevance of Creep to Reactor Design and Operation 4.10.6.2 The Irradiation-Induced Creep Mechanism (In-Crystal) 4.10.6.3 Review of Prior Creep Models 4.10.6.3.1 Linear viscoelastic creep model 4.10.6.3.2 The UK creep model 4.10.6.3.3 The Kennedy model 4.10.6.3.4 The Kelly and Burchell model 4.10.6.3.5 The M2 model 4.10.6.4 Deficiencies in Current Creep Models at High Neutron Doses 4.10.7 Outlook Acknowledgments References 4.11 Graphite in Gas-Cooled Reactors 4.11.1 Introduction 4.11.2 Graphite Crystal Structures 4.11.2.1 Graphite Crystal Atomic Structure and Properties 4.11.2.2 Coefficient of Thermal Expansion 4.11.2.3 Modulus 4.11.2.4 Thermal Conductivity 4.11.2.5 Microcracking (Mrozowski Cracks) 4.11.3 Artificial Nuclear Graphite 4.11.3.1 Microstructure/Property Relationships 4.11.4 Graphite Core Fast Neutron Fluence, Energy Deposition, and Temperatures 4.11.5 Dosimetry (Graphite Damage Dose or Fluence) 4.11.5.1 Early Activation Measurements on Foils 4.11.5.2 Reactor Design and Assessment Methodology: Fuel Burnup 4.11.5.2.1 Calder effective dose 4.11.5.3 Equivalent Nickel Flux 4.11.5.4 Integrated Flux and Displacements per Atom 4.11.5.4.1 DIDO equivalent flux 4.11.5.5 Energy Above 0.18MeV 4.11.5.6 Equivalent Fission Flux (IAEA) 4.11.5.7 Fluence Conversion Factors 4.11.5.8 Irradiation Annealing and EDT 4.11.5.9 Summary of Fast Neutron Dose (Fluence) 4.11.6 Graphite `Energy Deposition´ (Nuclear Heating) 4.11.6.1 The Use of Titanium for Installed Sample Holders 4.11.7 Radiolytic Oxidation 4.11.7.1 Introduction 4.11.7.2 Ionizing Radiation 4.11.7.2.1 Energy deposition 4.11.7.3 Radiolytic Oxidation Mechanism 4.11.7.4 Inhibition 4.11.7.5 Internal Porosity 4.11.7.6 Prediction of Weight Loss in Graphite Components 4.11.7.7 Weight Loss Prediction in Inhibited Coolant 4.11.8 Graphite Temperatures 4.11.9 Variation of Fluence, Temperature, and Weight Loss in a Reactor Core 4.11.9.1 Fuel End Effects 4.11.9.2 Temperature and Weight Loss 4.11.10 Distribution of Fluence Within an Individual Moderator Brick 4.11.11 Fast Neutron Damage in Graphite Crystal Structures 4.11.11.1 Stored Energy 4.11.11.2 Crystal Dimensional Change 4.11.11.3 Coefficient of Thermal Expansion 4.11.11.4 Modulus 4.11.11.5 Thermal Conductivity 4.11.11.6 Raman 4.11.12 Property Changes in Irradiated Polycrystalline Graphite 4.11.13 Averaging Relationships 4.11.14 Dimensional Change 4.11.14.1 Pile Grade A 4.11.14.2 Gilsocarbon 4.11.14.3 Effect of Radiolytic Oxidation on Dimensional Change 4.11.14.4 Dimensional Change Rate 4.11.15 Coefficient of Thermal Expansion 4.11.15.1 Pile Grade A 4.11.15.2 Gilsocarbon 4.11.15.3 Methodology for Converting Between Temperature Ranges 4.11.15.4 Effect of Radiolytic Oxidation on CTE 4.11.16 Thermal Conductivity 4.11.16.1 Pile Grade A 4.11.16.2 Gilsocarbon 4.11.16.3 Thermal Conductivity Temperature Dependence of Irradiated Graphite 4.11.16.4 Predicting the Thermal Conductivity of Irradiated Graphite for Reactor Core Assessments 4.11.17 Young´s Modulus 4.11.17.1 Relationship Between Static and Dynamic Young´s Modulus 4.11.17.2 Pile Grade A 4.11.17.3 Gilsocarbon 4.11.17.4 Separation of Structure and Pinning Terms 4.11.17.5 Effect of Radiolytic Weight Loss on Dimensional Change and Young´s Modulus 4.11.17.6 Small Specimen Strength 4.11.18 Effect of Radiolytic Oxidation on Thermal Conductivity, Young´s Modulus, and Strength 4.11.19 The Use of the Product Rule 4.11.20 Irradiation Creep in Nuclear Graphite 4.11.20.1 Dimensional Change and Irradiation Creep Under Load 4.11.20.2 Types of Irradiation Creep Experiments 4.11.20.3 The UKAEA Creep Law 4.11.20.4 Observed Changes to Other Properties 4.11.20.4.1 Coefficient of thermal expansion 4.11.20.4.2 Young´s modulus 4.11.20.5 Lateral Changes 4.11.20.6 Creep Models and Theories 4.11.20.6.1 UKAEA creep law 4.11.20.6.2 German and US creep model 4.11.20.6.3 Further modifications to the UKAEA creep law: interaction strain 4.11.20.6.4 Recent nuclear industry model 4.11.20.7 Final Thoughts on Irradiation Creep Mechanisms 4.11.21 Concluding Remark References 4.12 Vanadium for Nuclear Systems 4.12.1 Introduction 4.12.2 Vanadium Alloys for Fusion Reactors 4.12.3 Compositional Optimization 4.12.4 Fabrication Technology 4.12.5 Fundamental Study on Impurity Effects 4.12.6 Thermal Creep 4.12.7 Corrosion, Compatibility, and Hydrogen Effects 4.12.8 Radiation Effects 4.12.9 Tritium-Related Issues 4.12.10 Development of Advanced Alloys 4.12.11 Critical Issues 4.12.12 Vanadium Alloy Development for Fusion Blankets 4.12.13 Summary References 4.13 Concrete 4.13.1 Introduction 4.13.2 Concrete Longevity, NPP Safety-Related Concrete Structures, Testing and In-Service Inspection Requirements, and Operatin 4.13.2.1 Historical Perspective on Concrete Longevity 4.13.2.2 NPP Safety-Related Concrete Structures 4.13.2.2.1 Boiling water reactors 4.13.2.2.2 Pressurized water reactors 4.13.2.3 Testing and In-Service Inspection Requirements 4.13.2.4 Operating Experience 4.13.3 Aging and Long-Term Durability Considerations 4.13.3.1 Design, Construction, Material Selection, and Maintenance Considerations 4.13.3.2 Materials of Construction, Degradation Mechanisms, Damage Modeling, and Long-Term Performance of Concrete Materials 4.13.3.2.1 Materials of construction 4.13.3.2.2 Degradation mechanisms 4.13.3.2.3 Damage modeling 4.13.3.2.4 Long-term performance of concrete materials 4.13.3.3 Assessment and Repair 4.13.3.3.1 Component selection 4.13.3.3.2 In-service inspections 4.13.3.3.3 Nondestructive examinations 4.13.3.3.4 Remedial methods 4.13.4 Structural Reliability Theory 4.13.5 Summary and Potential Research Topics References 4.14 Fracture Toughness Master Curve of bcc Steels 4.14.1 Introduction 4.14.1.1 Historical Review of Fracture Toughness Determination for Ferritic Steels 4.14.1.2 Theoretical Background Leading to Use of EPFM and Data Distributions for Ferritic Steels 4.14.1.3 Master Curve Methodology as Developed by Wallin 4.14.1.3.1 History 4.14.1.3.2 Basics 4.14.1.3.3 Methodology 4.14.2 Analysis of Fracture Toughness Test Data for Master Curve Application 4.14.2.1 Standard Test and Analysis Procedure (ASTM E1921) 4.14.2.1.1 Test specimens used 4.14.2.1.2 Determination of reference temperature T0 4.14.2.1.3 Data qualification 4.14.2.1.4 Determination of lower bound curves 4.14.2.1.5 Limits of applicability 4.14.2.2 Aspects of Applying Small and Miniature Specimens 4.14.2.3 Effect of Constraint 4.14.2.4 Effect of Ductile Crack Growth 4.14.2.5 Inhomogeneous Materials 4.14.3 Application to Integrity Assessments 4.14.3.1 Transferability of Test Data 4.14.3.2 Accomplished Analyses for Specific RPV Integrity Assessments 4.14.3.2.1 PTS test by Framatome 4.14.3.2.2 PTS test by ORNL 4.14.4 Application to Lifetime Assessment 4.14.4.1 Assessment of Irradiation Embrittlement Changes 4.14.4.2 Definition of Reference Curves and Their Use 4.14.5 On Correlations Between T0 and Other Related Parameters 4.14.5.1 Crack Arrest Reference Temperature TKIa 4.14.5.2 Dynamic Versus Static Fracture Toughness 4.14.5.3 T0 Versus Charpy V-notch Transition Temperatures 4.14.6 Summary and Conclusions References 4.15 Ceramic Breeder Materials 4.15.1 Introduction 4.15.1.1 Tritium Breeding 4.15.1.2 Breeding Blankets 4.15.2 Ceramic Breeder Blankets 4.15.2.1 Pellets/Pins/Blocks 4.15.2.2 Pebble-Bed Concepts 4.15.2.3 Blanket Design Parameters 4.15.2.4 Testing of Blanket Modules in ITER 4.15.2.5 Ceramic Breeder Requirements 4.15.3 Ceramic Breeder Fabrication 4.15.3.1 Base Properties 4.15.3.2 Fabrication of Shapes 4.15.3.2.1 Pellets or blocks 4.15.3.2.2 Pebbles 4.15.4 Pebble and Pebble-Bed Thermomechanics 4.15.4.1 Introduction 4.15.4.2 Single Pebble Testing 4.15.4.3 Properties of Pebble Beds 4.15.4.3.1 Pebble-bed density and packing factor 4.15.4.3.2 Mechanical behavior of pebble beds 4.15.4.3.3 Thermal creep 4.15.4.3.4 Cyclic loading 4.15.4.4 Heat Transfer Properties 4.15.4.4.1 Thermal conductivity 4.15.4.4.2 Heat transfer 4.15.4.5 Pebble-Bed Modeling 4.15.4.5.1 Continuum models 4.15.4.5.2 Discrete-element modeling 4.15.4.6 In-Pile Behavior 4.15.5 Tritium Production and Release 4.15.5.1 Tritium Release 4.15.5.1.1 Out-of-pile 4.15.5.1.2 In-pile testing 4.15.6 Irradiation Parameters 4.15.6.1 Irradiation Damage 4.15.7 Activation and Waste Issues 4.15.8 Summary and Outlook 4.15.8.1 Microstructure 4.15.8.2 High Burnup 4.15.8.3 High Fluence 4.15.8.4 High Temperature 4.15.8.5 Effects of Transients and Off-Normal Conditions 4.15.8.6 Accident Behavior (Safety and Investment Integrity) 4.15.8.7 Development of Tools 4.15.8.8 Compatibility with Structure 4.15.8.9 Waste Management and Reuse/Recycling Acknowledgments References 4.16 Tritium Barriers and Tritium Diffusion in Fusion Reactors 4.16.1 Introduction 4.16.2 Background 4.16.2.1 Equation of State of Gases 4.16.2.2 Diffusivity 4.16.2.3 Solubility 4.16.2.4 Trapping 4.16.2.5 Permeability 4.16.2.6 Recombination 4.16.2.7 Irradiation and Implantation 4.16.3 Fusion Reactor Materials 4.16.3.1 Plasma-Facing Materials 4.16.3.1.1 Carbon 4.16.3.1.2 Tungsten 4.16.3.1.3 Beryllium 4.16.3.2 Structural Materials 4.16.3.2.1 Austenitic stainless steels 4.16.3.2.2 Ferritic/martensitic steels 4.16.3.2.3 V-Cr-Ti alloys 4.16.3.2.4 Zirconium alloys 4.16.3.2.5 Other structural metals 4.16.3.2.5.1 Aluminum 4.16.3.2.5.2 Nickel 4.16.3.2.5.3 Copper 4.16.3.3 Barrier Materials 4.16.3.3.1 Oxides 4.16.3.3.2 Aluminides 4.16.3.3.3 Nitrides 4.16.3.3.4 Carbides 4.16.3.3.5 Low permeation metals 4.16.3.3.5.1 Molybdenum 4.16.3.3.5.2 Silver 4.16.3.3.5.3 Platinum 4.16.3.3.5.4 Gold 4.16.4 Application of Barriers 4.16.4.1 Expected In-Reactor Performance 4.16.4.2 How Barriers Work and Why Radiation Affects Them 4.16.4.3 Why Barriers Are Needed for Fusion Reactors 4.16.5 Summary References 4.17 Tungsten as a Plasma-Facing Material 4.17.1 Introduction 4.17.2 Functional Requirements 4.17.3 Material Selection 4.17.3.1 Fabrication and Microstructure 4.17.3.2 Advantages and Limitations for Fusion Application 4.17.3.2.1 High atomic number: material erosion/melting 4.17.3.2.2 Recrystallization 4.17.3.2.3 Machinability, mechanical properties, and DBTT 4.17.3.2.4 Component fabrication: CTE mismatch with heat sink 4.17.3.2.5 Neutron embrittlement 4.17.3.2.6 Neutron activation and radiological hazards 4.17.3.2.7 Material availability 4.17.3.3 Tungsten Grades 4.17.4 Influence of In-Service Conditions 4.17.4.1 Thermal Shock Resistance 4.17.4.1.1 Microstructure, composition, and mechanical properties 4.17.4.1.2 Power density and pulse duration 4.17.4.1.3 Base temperature 4.17.4.1.4 Repetition rate 4.17.4.1.5 Thermal shock during off-normal events: disruptions 4.17.4.1.6 Thermal shock during normal operation: ELMs 4.17.4.2 Thermal Fatigue Resistance 4.17.4.2.1 ITER 4.17.4.2.2 Prototype and commercial reactors 4.17.4.3 Neutron Irradiation 4.17.4.3.1 Thermophysical properties and swelling 4.17.4.3.2 Mechanical properties 4.17.4.3.3 Thermal shock on irradiated W 4.17.4.3.4 Thermal fatigue on irradiated W components 4.17.4.4 Ion Irradiation and Retention 4.17.4.4.1 He-irradiation 4.17.4.4.1.1 Influence of ion energy and fluence 4.17.4.4.1.2 Influence of temperature 4.17.4.4.1.3 Influence of material´s microstructure 4.17.4.4.2 Hydrogen-irradiation and retention 4.17.4.4.2.1 Influence of ion energy, fluence, and temperature 4.17.4.4.2.2 Influence of material´s composition and microstructure 4.17.4.4.3 Combined loading conditions 4.17.5 Conclusion References 4.18 Carbon as a Fusion Plasma-Facing Material 4.18.1 Introduction 4.18.1.1 Background 4.18.1.2 Plasma-Facing Materials 4.18.1.3 Particle-Matter Interactions 4.18.2 The Advantages of Carbon as a PFC 4.18.2.1 Plasma Impurities and the Need for Graphite Materials 4.18.2.2 Thermomechanical Loading of PFMs 4.18.2.3 Transient Loading of PFMs 4.18.3 Irradiation Effects on Thermophysical Properties of Graphite and CFCs 4.18.3.1 Graphite Irradiation Damage 4.18.3.2 Surface Effects 4.18.3.3 Properties and Property Evolution of Graphite Fiber Composite 4.18.3.3.1 Irradiation-induced dimensional changes in CFCs 4.18.3.3.2 Irradiation-induced changes in strength and modulus 4.18.3.3.3 Thermal conductivity degradation 4.18.4 Plasma-Particle Interactions 4.18.4.1 Chemical Erosion 4.18.4.2 Doping of Graphite to Suppress Erosion 4.18.4.3 Physical Sputtering 4.18.4.4 Radiation-Enhanced Sublimation 4.18.4.5 Erosion of Graphite in Simulated Disruption Events 4.18.5 Tritium Retention in Graphitic Materials 4.18.6 HHF Component Technology 4.18.6.1 Joining of CFC to Heat Sink 4.18.6.2 Evaluation of HHF Joint 4.18.7 Summary and Conclusions References 4.19 Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices 4.19.1 Introduction 4.19.2 Background 4.19.2.1 Synopsis of PWIs in Tokamaks 4.19.2.2 Brief History of Plasma-Facing Materials in Fusion Devices 4.19.2.3 Experience with Beryllium in Tokamaks 4.19.3 Beryllium PWI Relevant Properties 4.19.3.1 Beryllium Erosion Properties 4.19.3.1.1 Physical sputtering of beryllium 4.19.3.1.2 Mixed-material erosion 4.19.3.1.3 Chemically assisted sputtering of beryllium 4.19.3.1.4 Enhanced erosion at elevated temperatures 4.19.3.2 Hydrogen Retention and Release Characteristics 4.19.3.2.1 Implantation 4.19.3.2.2 Beryllium codeposition 4.19.3.3 Mixed-Material Effects 4.19.3.3.1 Be-C phenomena 4.19.3.3.2 Be-W alloying 4.19.4 Main Physical and Mechanical Properties 4.19.4.1 General Considerations 4.19.4.1.1 Physical properties 4.19.4.1.2 Mechanical properties 4.19.4.2 Selection of Beryllium Grades for ITER Applications 4.19.4.3 Considerations on Plasma-Sprayed Beryllium 4.19.4.4 Neutron-Irradiation Effects 4.19.4.4.1 Thermal conductivity 4.19.4.4.2 Swelling 4.19.4.4.3 Mechanical properties 4.19.4.4.4 Thermal shock effects 4.19.4.4.5 Bulk tritium retention 4.19.5 Fabrication Issues 4.19.5.1 Joining Technologies and High Heat Flux Durability of the Be/Cu Joints 4.19.5.1.1 Be/Cu alloy joining technology 4.19.5.1.1.1 Background information 4.19.5.1.1.2 HIP joining technique 4.19.5.1.1.3 Fast brazing techniques 4.19.5.1.2 High heat flux durability of unirradiated Be/Cu joints 4.19.5.2 Thermal Tests on Neutron-Irradiated Joints 4.19.6 Tokamak PFC Design Issues and Predictions of Effects in ITER During Operation 4.19.6.1 PFC Design Considerations 4.19.6.1.1 Design of the beryllium ITER-like wall at JET 4.19.6.1.2 Design of the beryllium ITER wall 4.19.6.2 Predictions of Effects on the ITER Beryllium Wall During Operation 4.19.6.2.1 Safety issues in ITER 4.19.6.2.1.1 In-vessel tritium inventory 4.19.6.2.1.2 Chemical reactivity of beryllium dust with steam in ITER 4.19.6.2.2 Erosion/damage of the ITER Be wall 4.19.6.2.2.1 Erosion of Be wall during normal operation 4.19.6.2.2.2 Erosion of the beryllium wall during ELMs 4.19.6.2.2.3 Erosion of the beryllium wall during thermal quench disruptions 4.19.6.2.2.4 Erosion of the beryllium wall during VDEs 4.19.6.2.2.5 Erosion of the beryllium wall during runaway impact 4.19.6.3 Prospect of Using Beryllium in Beyond-ITER Fusion Reactors 4.19.7 Concluding Remarks Acknowledgments References 4.20 Physical and Mechanical Properties of Copper and Copper Alloys 4.20.1 Introduction 4.20.2 Copper and High-Strength, High-Conductivity Copper Alloys 4.20.2.1 Pure Copper 4.20.2.2 PH Copper Alloys 4.20.2.2.1 CuCrZr alloy 4.20.2.2.2 CuNiBe alloy 4.20.2.2.3 CuNiSi 4.20.2.3 DS Copper Alloys 4.20.3 Physical Properties of Copper and Copper Alloys 4.20.4 Mechanical Properties of Copper and Copper Alloys 4.20.4.1 Tensile Properties 4.20.4.2 Fracture Toughness 4.20.4.3 Creep 4.20.4.4 Fatigue and Creep-Fatigue 4.20.5 Irradiation Effects in Copper and Copper Alloys 4.20.5.1 Effect of Irradiation on Physical Properties of Copper and Copper Alloys 4.20.5.2 Effect of Irradiation on Mechanical Properties of Copper and Copper Alloys 4.20.5.2.1 Tensile properties 4.20.5.2.2 Fracture toughness 4.20.5.2.3 Fatigue and creep-fatigue 4.20.5.2.4 Irradiation creep and void swelling 4.20.5.3 Effect of Irradiation on Microstructure of Copper and Copper Alloys 4.20.5.3.1 Defect structure in irradiated copper and copper alloys 4.20.5.3.2 Dislocation channeling 4.20.6 Joining 4.20.7 Summary References 4.21 Ceramic Coatings as Electrical Insulators in Fusion Blankets 4.21.1 Introduction 4.21.2 Magnetohydrodynamic Issues and the Requirement for Insulator Coatings 4.21.3 Development of Insulator Coating for Liquid Li Blanket 4.21.3.1 In Situ Formation and Healing with CaO 4.21.3.2 In Situ AlN Coating 4.21.3.3 Er2O3 and Y2O3 as New Candidates 4.21.3.3.1 Scoping by bulk immersion tests 4.21.3.3.2 In situ coating with Er2O3 4.21.3.3.3 Physical coating processes 4.21.3.3.4 Other coating technologies 4.21.3.4 Two-Layer Coatings 4.21.3.5 Radiation Effects 4.21.3.6 FCI Concept as an Alternative to Insulator Coating 4.21.4 Summary and Remaining Issues References 4.22 Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors 4.22.1 Introduction 4.22.2 Fusion-Relevant Radiation Damage in Insulating Materials 4.22.3 Simulation Experiments 4.22.4 Degradation of Insulator Electrical Resistance 4.22.5 Degradation of Insulator AC/RF Dielectric Properties 4.22.6 Degradation of Insulator Thermal Conductivity 4.22.7 Degradation of Optical Properties 4.22.8 Concluding Remarks References