Comprehensive Nuclear Materials, Volume 4: Radiation Effects in Structural and Functional Materials for Fission and Fusion Reactors

EDITOR IN CHIEF Rudy J. M. Konings European Commission, Joint Research Centre, Institute for Transuranium Elements, Karl...

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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: [email protected]. 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 Sedmidubsky´, 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 computational 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, specifically 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.

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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 thermochemistry 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 uranium by neutron capture. After World War II, a rapid exploration of fission technology took place in the United States and the Union of Soviet Socialist Republics, and after the Atoms for Peace speech by Eisenhower at the United Nations Congress in 1954, also in Europe and Japan. A variety 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 the most 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.

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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 dynamics, 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 priesthood 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

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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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Figure 39 of Carbon as a Fusion Plasma-Facing Material Figure 40 of Carbon as a Fusion Plasma-Facing Material Table 5 of Carbon as a Fusion Plasma-Facing Material www.wiley.com The following material is reproduced with kind permission of Springer Figure 4 of Neutron Reflector Materials (Be, Hydrides) Figure 6 of Neutron Reflector Materials (Be, Hydrides) Figure 1 of Properties and Characteristics of SiC and SiC/SiC Composites Figure 3 of Properties and Characteristics of SiC and SiC/SiC Composites Figure 4 of Properties and Characteristics of SiC and SiC/SiC Composites Figure 5 of Properties and Characteristics of SiC and SiC/SiC Composites Figure 6 of Properties and Characteristics of SiC and SiC/SiC Composites Figure 7 of Properties and Characteristics of SiC and SiC/SiC Composites Figure 8 of Properties and Characteristics of SiC and SiC/SiC Composites Figure 9 of Properties and Characteristics of SiC and 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 Simulations Figure 14 of Oxide Fuel Performance Modeling and Simulations Figure 5 of Thermodynamic and Thermophysical Properties of the Actinide Nitrides Figure 51 of Phase Diagrams of Actinide Alloys Figure 6 of Thermodynamic and Thermophysical Properties of the Actinide Oxides Figure 7b of Thermodynamic and Thermophysical Properties of the Actinide Oxides Figure 9b of Thermodynamic and Thermophysical Properties of the Actinide Oxides Figure 35 of Thermodynamic and Thermophysical Properties of the Actinide Oxides Table 11 of Thermodynamic and Thermophysical Properties of the Actinide Oxides Table 13 of Thermodynamic and Thermophysical Properties of the Actinide Oxides Table 17 of Thermodynamic and Thermophysical Properties of the Actinide Oxides Figure 18 of Radiation Damage of Reactor Pressure Vessel Steels Figure 7 of Radiation Damage Using Ion Beams Figure 9b of Radiation Damage Using Ion Beams Figure 28 of Radiation Damage Using Ion Beams Figure 34 of Radiation Damage Using Ion Beams Figure 35 of Radiation Damage Using Ion Beams Figure 36d of Radiation Damage Using Ion Beams Figure 37 of Radiation Damage Using Ion Beams Table 3 of Radiation Damage Using Ion Beams


Figure 5 of Radiation Effects in UO2 Figure 9a of Ab Initio Electronic Structure Calculations for Nuclear Materials Figure 9b of Ab Initio Electronic Structure Calculations for Nuclear Materials Figure 9c of Ab Initio Electronic Structure Calculations for Nuclear Materials Figure 10a of Ab Initio Electronic Structure Calculations for Nuclear Materials Figure 23 of Thermodynamic and Thermophysical Properties of the Actinide Carbides Figure 25 of Thermodynamic and Thermophysical Properties of the Actinide Carbides Figure 26 of Thermodynamic and Thermophysical Properties of the Actinide Carbides Figure 27 of Thermodynamic and Thermophysical Properties of the Actinide Carbides Figure 28a of Thermodynamic and Thermophysical Properties of the Actinide 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 Figure 2 of Atomic-Level Dislocation Dynamics in Irradiated Metals Figure 3 of Atomic-Level Dislocation Dynamics in Irradiated Metals Figure 5 of Atomic-Level Dislocation Dynamics in Irradiated Metals Figure 10a of Atomic-Level Dislocation Dynamics in Irradiated Metals Figure 10b of Atomic-Level Dislocation Dynamics in Irradiated Metals Figure 10c of Atomic-Level Dislocation Dynamics in Irradiated Metals

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Figure 10d of Atomic-Level Dislocation Dynamics in Irradiated Metals Figure 12a of Atomic-Level Dislocation Dynamics in Irradiated Metals Figure 12b of Atomic-Level Dislocation Dynamics in Irradiated Metals Figure 12c of Atomic-Level Dislocation Dynamics in Irradiated Metals Figure 12d of Atomic-Level Dislocation Dynamics in Irradiated Metals Figure 16a of Atomic-Level Dislocation Dynamics in Irradiated Metals Figure 16b of Atomic-Level Dislocation Dynamics in Irradiated Metals Figure 16c of Atomic-Level Dislocation Dynamics in 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

4.01

Radiation Effects in Zirconium Alloys

F. Onimus and J. L. Bećhade Commissariat a` l’Energie Atomique, Gif-sur-Yvette, France

ß 2012 Elsevier Ltd. All rights reserved.

4.01.1

Irradiation Damage in Zirconium Alloys

4.01.1.1 4.01.1.1.1 4.01.1.1.2 4.01.1.1.3 4.01.1.2 4.01.1.2.1 4.01.1.2.2 4.01.1.2.3 4.01.1.3 4.01.1.3.1 4.01.1.3.2 4.01.1.3.3 4.01.1.3.4 4.01.1.3.5 4.01.1.4 4.01.1.4.1 4.01.1.4.2 4.01.2 4.01.2.1 4.01.2.1.1 4.01.2.1.2 4.01.2.1.3 4.01.2.1.4 4.01.2.2 4.01.2.3 4.01.3 4.01.3.1 4.01.3.1.1 4.01.3.1.2 4.01.3.2 4.01.3.2.1 4.01.3.2.2 4.01.3.3 References

Damage Creation: Short-Term Evolution Neutron–zirconium interaction Displacement energy in zirconium Displacement cascade in zirconium Evolution of Point Defects in Zirconium: Long-Term Evolution Vacancy formation and migration energies SIA formation and migration energies Evolution of point defects: Impact of the anisotropic diffusion of SIAs Point-Defect Clusters in Zirconium Alloys hai Dislocation loops hai Loop formation: Mechanisms hci Component dislocation loops hci Loop formation: Mechanisms Void formation Secondary-Phase Evolution Under Irradiation Crystalline to amorphous transformation of Zr-(Fe,Cr,Ni) intermetallic precipitates Irradiation effects in Zr–Nb alloys: Enhanced precipitation Postirradiation Mechanical Behavior Mechanical Behavior During Tensile Testing Irradiation hardening: Macroscopic behavior Irradiation hardening: Mechanisms Post-yield deformation: Macroscopic behavior Post-yield deformation: Mechanisms Effect of Postirradiation Heat Treatment Postirradiation Creep Deformation Under Irradiation Irradiation Growth Irradiation growth: Macroscopic behavior Irradiation growth: Mechanisms Irradiation Creep Irradiation creep: Macroscopic behavior Irradiation creep: Mechanisms Outlook

Abbreviations BWR CANDU DAD EAM EID FP-LMTO

Boiling-water reactor Canadian deuterium uranium Diffusion anisotropy difference Embedded atom method Elastic interaction difference Full-potential linear muffin-tin orbital

2

GGA hcp HVEM LDA MB MD NRT

2 2 2 2 4 4 4 6 7 7 8 9 9 10 10 10 13 14 14 14 14 16 16 17 18 19 19 19 21 24 24 25 26 27

Generalized gradient approximation Hexagonal close-packed High-voltage electron microscope Local density approximation Many body Molecular dynamics Norgett–Robinson–Torrens

1

2


PKA PWR RXA SANS SIA SIPA SIPA-AD SIPN SRA TEM Tm UTS YS

Primary knocked-on atom Pressurized water reactor Recrystallization annealed Small-angle neutron scattering Self interstitial atom Stress-induced preferential absorption Stress preferential induced nucleationanisotropic diffusion Stress preferential induced nucleation Stress-relieved annealed Transmission electron microscopy Melting temperature Ultimate tensile strength 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 crystallographic sites (depending on the energy of the incoming 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 transferred to the target atom. The interaction probability is given by the elastic collision differential cross section1,2 which depends on both the neutron kinetic energy and the transferred energy.3 For a typical fast Þ of neutron of 1 MeV, the mean transferred energy ðT 22keV. For low value of the the Zr atom is T 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 irradiations performed at low temperatures ( 1 MeV)), the number of displaced 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 Lune´ville 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 year1, 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 n m2 (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 elastic 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

performed MD computations also using different types of interatomic potentials. It is shown that, at the end of the cascade creation ( 1 MeV), at irradiation temperatures between 320 and 360 C, but saturates above 1 1024 n m2 (E > 1 MeV) and little change occurs from 1 1024 up to 1.5 1025 n m2 (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 n m2 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 1000

4.01.2.1.2 Irradiation hardening: Mechanisms

True stress (MPa)

Strain rate 2.5% min–1 0.25 0.025

800 600

0.5 400

0.025

Irradiated (~3 ⫻ 1024 n m–2) Unirradiated

200 0

0

0.02

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 hardening decreases. According to these authors, this shows that the accumulation of damage decreases as the irradiation temperature increases, presumably due to recovery during irradiation. The chemical composition seems to play a secondary role in the irradiation-induced hardening compared 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 hardening. Indeed, Adamson and Bell101 have shown using microhardness tests that the irradiation-induced hardening 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 irradiationinduced 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.

0.04 0.06 0.08 True strain (mm mm–1)

0.1

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.

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 pointdefect 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 dislocations and loops in zirconium alloys. Particularly, Carpenter111 has considered the mechanism


15

80

YS

60

24 20

40

16 12

20

Elongation

Strength (kg mm–2)

UTS

8 4 Uniform elongation 0.2 ⫻ 1021 0.4 ⫻ 1021 0.6 ⫻ 1021 0.8 ⫻ 1021 1.0 ⫻ 1021 1.2 ⫻ 1021 1.4 ⫻ 1021 1.6 ⫻ 1021 Fast fluence (E > 1 MeV)

Figure 13 The effect of fast fluence (given in n cm2, 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 Proportional limit, yield and ultimate tensile stress 700

spl

sy

19

3.2 ⫻ 10 nvt Unirrad

spl, sy and sUTS (MPa)

600

sUTS (0.1% offset) (lower yield point)

500 400 300 200 100

300

400

500

600

700

Test temperature (K) 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 1.1 104 s1. Adapted from Onchi, T.; Kayano, H.; Higashiguchi, Y. J. Nucl. Mater. 1980, 88(2–3), 226–235.

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 prismatic 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 computations114 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.

16


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 hardening 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 elongation, 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 fracture 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 rapidly with the fluence from 10% to values lower than 1% for RXA alloys at 350 C, and saturates from a fluence of 5 1024 n m2.92 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 temperatures 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 by Hirsch,110 Wechsler,126 and Luft.127 According

700 nm Figure 15 Propagating basal channels observed after tensile testing at 350 C. Adapted from Onimus, F.; Monnet, I.; Bećhade, J. L.; Prioul, C.; Pilvin, P. J. Nucl. Mater. 2004, 328, 165–179.

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 possible 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 dislocation gliding, leading to plastic strain localization 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 disagreement 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 characteristic of RXA tubing or rolled sheets, with hci axes oriented 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


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 dislocations 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 prismatic 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 progressive 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 irradiation.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 dislocation 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 instability 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 Bećhade,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

17

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 irradiation decreases the plastic anisotropy of the RXA zirconium alloys. Concerning the SRA zirconium alloys, the mechanical behavior is already more isotropic 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 various 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 coldworked zirconium alloy most of the recovery occurring 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 recovery 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 dislocations, 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 irradiationproduced 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 zirconium 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 radiationinduced hardening.101,105

18


90 UTS

80 Pile temperature (280 ⬚C)

Normal stress (psi ⫻ 10–3)

100

70 60 50 40 100

UTS ANN

YS PL YS ANN PL ANN

200 300 400 500 600 700 Postirradiation annealing temperature (⬚C)

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.

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 irradiation up to a fluence of 1.4 1024 n m2. 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 onethird 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 studied 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, explaining 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 zirconium 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 irradiation defects that harden the material. However, according to Ito et al.142 and Scha¨ffler et al.,152 irradiation 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 irradiation 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 anisotropy 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 prismatic 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


19

0.09 0.08 0.07 Diametral creep eq (–)

350 ⬚C 445 MPa

4.39 ⫻ 1024 n m–2 unirr.

8.25 ⫻ 1024 n m–2

0.06 0.05 0.04 20.80 ⫻ 1024 n m–2

0.03 0.02

45.03 ⫻ 1024 n m–2

0.01

92.41 ⫻ 1024 n m–2

0 0

50 000

100 000

150 000

200 000

t (s) 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.

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 material. In this study, it is clearly shown that in RXA zirconium alloys, the irradiation loops are progressively 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 irradiation on materials is the dimensional change without 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

20


0 8

1

2

3

dpa 5

4

7

8

9

Iodide zirconium (6,7) Zone-refined zirconium (2,5)

6 Growth strain (10–4)

6

553 K

4

Carpenter et al.164

A-axis crystals

2 0 C-axis crystals

–2 –4 0

1

2

3

4

5

6

7

8

Neutron fluence (1025 n m–2) 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.

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 h1120i or h1010i 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 rolling 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 crystal.161 This demonstrates the role of the grain boundaries on the growth mechanisms. For higher fluence, higher than 3–5 1025 n m2, a growth breakaway is observed, yielding a high growth rate. It is reported150,160,166 that for polycrystalline zirconium alloys, the grain size affects the growth rate


353 K 553 K Annealed zircaloy–2.20 mm 25% C.W. zircaloy 2.5–8 mm

55

21

FL = 0.1

50

Growth strain ⫻ 104

45 40 35 30 25 20 15 10 50 0 0

20

40

60

80

100

120

140

160

180

Fast neutron dose (⫻1024 n m–2, E > 1 MeV) 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.

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 working 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 treatment, 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 initial dislocations network in the growth mechanisms. Several authors have studied the effect of the impurity 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 zirconium single crystal obtained by using the iodine purification 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 influence 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 temperature 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 observations of the microstructure have progressed.

22


700

10–27

Temperature (K) 500

600

400

350

f = 0.10

Growth rate (m2 n–1)

Q » 150 kJ mol–1 p = 5 ⫻ 1014 m–2

10–28 Cold work

ed

10–29

Recrystallized (postbreakaway)

p = 1 ⫻ 1014 m–2 Q » 3 kJ mol–1

Recrystalliz

ed (prebre

akaway)

10–30 1.4 ⫻ 10–3

1.8 ⫻ 10–3

2.2 ⫻ 10–3

2.6 ⫻ 10–3

3.0 ⫻ 10–3

1/T(K)

Figure 20 Generalized representation of the temperature dependence of irradiation growth of Zircaloy. Adapted from Fidleris, V. J. Nucl. Mater. 1988, 159, 22–42.

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 zirconium alloys is provided by Holt.162 Some of these mechanisms are not compatible with all the observations. 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 Go¨sele174 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 chapter, several authors28,33,34,175 have shown, using atomistic 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.

(a)

(b)

(c)

(d) Figure 21 (a–c) The three phases of growth of recrystallized zirconium alloys. (d) Growth mechanisms of stress relieved zirconium alloys.

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 explains why 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 dominant sink for point defects. In the absence of hci component dislocation (as is the case in RXA zirconium alloys), calculations of DAD-induced bias


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 neighborhood of prismatic grain boundaries, or surfaces, which experience a net influx of SIAs, there will be a higher vacancy supersaturation leading to a predominance 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 stationary growth rate observed. For irradiation doses higher than 5 1025 n m2, 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 eliminated preferentially on the hci component loops and on the grain boundaries parallel to the basal plane. The SIAs are eliminated on hai type dislocations

Irradiation growth strain (%)

D3T

A2

D2L

No component dislocations

0.15

and grain boundaries parallel to the prismatic plane. This partitioning of the point defects on these different 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 hc þ ai line dislocations are already present before irradiation, under irradiation, the vacancies are preferentially eliminated on the dislocations with hc þ 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 partitioning of point defects therefore leads to the fast and continuous growth of cold-worked or SRA zirconium alloys (Figure 21 (d)). Here the growth created by the point-defect flux to the grain boundaries is relatively unimportant because they are not dominating 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, especially by Holt,162 that due to the polycrystalline nature of the material, the growth strain of the individual grains can induce strain incompatibilities between adjacent grains that exhibit different orientations. Intergranular stresses can then result from these strain incompatibilities, leading to a local irradiation creep of individual grains even without

G2

D2 D1 Many component dislocations

Some component dislocations

0.10

0.05 Growth specimens in DIDO (553 K) 0

1

23

Fuel assembly guide tubes in calvert cliffs-1 (508–583 K)

2 3 4 5 6 7 Fluence (n m–2) (E > 1 MeV) ⫻ 10–25

8

9

Figure 22 Irradiation growth in annealed (RXA) Zircaloy at 550–580 K, showing accelerating growth at 4 1025 n m2 (E > 1 MeV). Adapted from Fidleris, V. J. Nucl. Mater. 1988, 159, 22–42.

24


macroscopic applied stress on the material. This phenomenon 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 ‘thermal’ creep rate, the creep rate increasing as the neutron flux increases. The behavior under irradiation of zirconium alloys, and particularly the creep behavior, 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 deformation of the fuel assembly in the case of PWR and boiling-water reactor (BWR)169 or in-reactor structure 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 dislocation 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, causing 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 irradiation. 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 n m2 and is followed by a steady-state creep rate. Concerning

cold-worked materials, the effect of the preirradiation 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 significant 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 irradiation 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 temperature is usually described by an Arrhenius equation.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 deformation component. e_ ¼ e_ thermalcreep þ e_ irradiationcreep þ e_ growth ¼ e_ creep þ e_ growth with e_ creep ¼ K sn fp exp 1

Q RT

½1

½2

where e_ is the strain rate in s ; s is the effective stress for thermal creep in MPa; 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 K1 mol1; is the fast neutron flux in n m2 s1 (E > 1 MeV); p is the flux exponent; and K is a constant for thermal creep in s1 (MPa)n(n m2 s1)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 p ¼ 1 is commonly obtained for CANDU pressure tube deformation.163,183 For uniaxial creep tests performed 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 n ¼ 1 is obtained at 300 C for low applied stress (s 100 MPa). As the stress increases, the stress exponent increases, reaching values up to n ¼ 25 for 450 MPa applied stress for cold-worked Zr–2.5% Nb.183


10–4

350

400

Temperature (⬚C) 300

4.01.3.2.2 Irradiation creep: Mechanisms 250

200

+ 10–5

In-reactor tests, t > 6000 h Flux = 9⫻1016 n m–2 s–1, E > 1 MeV

Creep rate (h–1)

+

+ 207 MPa

+

138 MPa

10–7 Laboratory tests, t > 6000 h + +138 MPa 207 MPa 10–8 1.4

1.5

1.6

1.7 1.8 1/T ⫻ 103 (K)

1.9

2.0

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 preferential 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 dislocation glide mechanisms, which are essentially a combination of climb of dislocations due the absorption of point defects under irradiation and glide resulting in a creep deformation. For this category of mechanisms, the strain is essentially produced by glide but the strain rate is controlled by the climb.

+

10–6

25

2.1

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.

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 activation 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 thermal 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 anisotropic, 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.

Other mechanisms involving irradiation-induced loops have also to be added to these two categories of deformation mechanisms involving line dislocations. Indeed, the stress-induced preferential nucleation (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 dislocation with a Burgers vector parallel to the applied stress exhibits a higher bias toward SIAs than a dislocation 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


However, it has been seen that for growth, the anisotropic diffusion of SIAs is believed to play an important role in the deformation mechanism. Therefore, any irradiation creep model proposed for zirconium 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 irradiation creep mechanisms are not clearly identified yet. Indeed, since the initial dislocation density is very low, another deformation mechanism has to be activated. 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 density and mean loop diameter saturate at relatively low doses. Since the initial dislocation density is very low in RXA zirconium alloys, creep mechanisms involving climb of dislocations due to the SIPA mechanism or climb-plus-glide of dislocations require the generation 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 irradiation and applied stress, leading to the creation of a dislocation network that undergoes a SIPA or climbenhanced 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 overcome 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 channeling 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 dislocation climb and glide processes occur at even higher stress. For very high stress, close to the YS, dislocation 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 mechanism for irradiation creep. However, according to Holt,171 the creep anisotropy of cold-worked zirconium alloys computed from the SIPA mechanism assuming only hai type dislocations is not in agreement with the experimental anisotropy. The anisotropy computed from the climb-plus-glide mechanism assuming 80% prism slip and 20% basal slip is in good agreement with the experimental anisotropy, demonstrating 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 cluster formation, improvement in the knowledge of anisotropic diffusion of SIAs as well as better understanding of the microstructure of vacancy and interstitial hai loops and basal hci vacancy loops (origin of the loop alignment, origin of the corduroy contrast


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 precipitates, etc. Concerning the mechanical behavior of Zr alloys after irradiation, multiscale modeling of the postirradiation deformation with a better understanding of the dislocation channeling mechanism and understanding of its effects on the postirradiation mechanical behavior are needed. Moreover, better understanding of the postirradiation creep deformation mechanisms is also needed using multiscale modeling. The last point concerns the deformation mechanisms under irradiation. In that field, the basic questions are still without answers: What are the irradiation creep deformation mechanisms? What are the coupling between the deformation under irradiation and the thermal creep and growth? Progress has to be made especially using in situ deformation devices under irradiation, coupled with modeling approaches. (See also Chapter 1.01, Fundamental Properties of Defects in Metals; Chapter 2.07, Zirconium Alloys: Properties and Characteristics and Chapter 5.03, Corrosion of Zirconium Alloys).

12. 13. 14. 15. 16. 17. 18. 19. 20. 21. 22. 23. 24. 25. 26. 27. 28. 29. 30. 31. 32. 33. 34. 35. 36. 37. 38. 39.

References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11.

Beyster, J. R.; Walt, M.; Salmi, E. W. Phys. Rev. 1956, 104, 1319–1331. Guenther, P.; Smith, A.; Whalen, J. Phys. Rev. C 1975, 12, 1797–1808. Lune´ville, L.; Simeone, D.; Jouanne, C. J. Nucl. Mater. 2006, 353(1–2), 89–100. Neely, H. H. Radiat. Effects 1970, 3, 189–201. Biget, M.; Maury, F.; Vajda, P.; Lucasson, A.; Lucasson, P. Radiat. Effects 1971, 7, 223–229. Griffiths, M. J. Nucl. Mater. 1989, 165, 315–317. Ackland, G. J.; Woodings, S. J.; Bacon, D. J. Philos. Mag. A 1995, 71, 553–565. Gao, F.; Bacon, D. J.; Howe, L. M.; So, C. B. J. Nucl. Mater. 2001, 294, 288–298. Was, G. S. Fundamentals of Radiation Materials Science: Metals and Alloys; Springer: New York, 2007. Norgett, M. J.; Robinson, M. T.; Torrens, I. M. Nucl. Eng. Des. 1975, 33, 50–54. Zinkle, S. J.; Singh, B. N. J. Nucl. Mater. 1993, 199(3), 173–191.

40. 41. 42. 43. 44. 45. 46. 47. 48. 49. 50. 51.

27

Shishov, V.; Peregud, M. M.; Nikulina, A. V.; et al. J. ASTM Int. (JAI); 2005, 2(8). Woo, C. H.; Singh, B. N. Philos. Mag. A 1992, 65, 889–912. Woo, C. H.; Semenov, A. A.; Singh, B. N. J. Nucl. Mater. 1993, 206(2–3), 170–199. Kapinos, V. G.; Osetsky, Y. N.; Platonov, P. A. J. Nucl. Mater. 1991, 184(2), 127–143. Wooding, S. J.; Howe, L. M.; Gao, F.; Calder, A. F.; Bacon, D. J. J. Nucl. Mater. 1998, 254(2–3), 191–204. MacEwen, S. R.; Zee, R. H.; Birtcher, R. C.; Abromeit, C. J. Nucl. Mater. 1984, 123, 1036. Hood, G. M. J. Nucl. Mater. 1988, 159, 149–175. Frank, W. Philos. Mag. A 1991, 63, 897–913. Horvath, J.; Dyment, F.; Mehrer, H. J. Nucl. Mater. 1984, 126, 206–214. Douglass, D. L. Atomic Energy Review Supplement; International Atomic Energy Agency: Vienna, 1971; pp 311–342. Hood, G. M.; Schultz, R. J. Acta Mater. 1974, 22, 459. Hood, G. M.; Zou, H.; Schultz, R. J.; Jackman, J. A. J. Nucl. Mater. 1984, 126, 79–82. Hood, G. M. J. Nucl. Mater. 1986, 139, 179–184. Frank, W. J. Nucl. Mater. 1988, 159, 122–148. Lubbehusen, M.; Vieregge, K.; Hood, G. M.; Mehrer, H.; Herzig, C. J. Nucl. Mater. 1991, 182, 164–169. Hood, G. M.; Zou, H.; Gupta, D.; Schultz, R. J. J. Nucl. Mater. 1995, 223, 122–125. Bacon, D. J. J. Nucl. Mater. 1988, 159, 176–189. Bacon, D. J. J. Nucl. Mater. 1993, 206, 249–265. Le Bacq, O.; Willaime, F.; Pasturel, A. Phys. Rev. B 1999, 59, 8508–8515. Pasianot, R. C.; Monti, A. M. J. Nucl. Mater. 1999, 264, 198–205. Pasianot, R. C.; Monti, A. M.; Simonelli, G.; Savino, E. J. J. Nucl. Mater. 2000, 276, 230–234. Osetsky, Y. N.; Bacon, D. J.; De Diego, N. Metall. Mater. Trans. A 2002, 33A, 777–782. Woo, C. H.; Liu, X. Philos. Mag. 2007, 87(16), 2355–2369. Willaime, F. J. Nucl. Mater. 2003, 323, 205–212. Domain, C.; Legris, A. Philos. Mag. 2005, 85, 569–575. Domain, C. J. Nucl. Mater. 2006, 351(1–3), 1–19. Ve´rite´, G.; Willaime, F.; Chun, F. C. Solid State Phenom. 2007, 129, 75–81. Pe´rez, R. A.; Weissmann, M. J. Nucl. Mater. 2008, 374(1–2), 95–100. Pasianot, R. C.; Pe´rez, R. A.; Ramunni, V. P.; Weissmann, M. J. Nucl. Mater. 2009, 392(1), 100–104. Johnson, R. A.; Beeler, J. R. In Interatomic Potential and Crystalline Defects; Lee, J. K., Ed.; AIME: New York, 1981; p 165. Woo, C. H.; Huang, H.; Zhu, W. J. Appl. Phys. A: Mater. Sci. Process. 2003, 76, 101–106. Barbu, A.; Martin, G. Solid State Phenom. 1993, 30/31, 179–228. Woo, C. H. J. Nucl. Mater. 1988, 159, 237–256. Northwood, D. O.; Gilbert, R. W.; Bahlen, L. E.; et al. J. Nucl. Mater. 1979, 79, 379–394. Griffiths, M. J. Nucl. Mater. 1988, 159, 190–218. Gulden, T. D.; Bernstein, I. M. Philos. Mag. 1966, 14, 1087–1091. Kelly, P. M.; Blake, R. G. Philos. Mag. 1973, 28(2), 415–426. Northwood, D. O.; Fidleris, V.; Gilbert, R. W.; Carpenter, G. J. C. J. Nucl. Mater. 1976, 61(2), 123–130. Jostsons, A.; Kelly, P. M.; Blake, R. G. J. Nucl. Mater. 1977, 66(3), 236–256. Northwood, D. O. Atomic Energy Rev. 1977, 15, 547–610.

28 52. 53. 54. 55. 56. 57. 58. 59. 60.

61.

62. 63. 64. 65. 66. 67. 68. 69. 70. 71. 72. 73.

74. 75.

76.

Radiation Effects in Zirconium Alloys Gilbert, R. W.; Farrell, K.; Coleman, C. E. J. Nucl. Mater. 1979, 84(1–2), 137–148. Carpenter, G. J. C.; Watters, J. F. J. Nucl. Mater. 1981, 96, 213–226. Griffiths, M.; Gilbert, R. W.; Fidleris, V.; Tucker, R. P.; Adamson, R. B. J. Nucl. Mater. 1987, 150, 159–168. Hellio, C.; de Novion, C. H.; Boulanger, L. J. Nucl. Mater 1988, 159, 368–378. Griffiths, M. Philos. Mag. B 1991, 63(5), 835–847. Griffiths, M. J. Nucl. Mater. 1993, 205, 225–241. Griffiths, M.; Loretto, M. H.; Sallmann, R. E. J. Nucl. Mater. 1983, 115(2–3), 313–322. Griffiths, M.; Loretto, M. H.; Smallman, R. E. Philos. Mag. A 1984, 49(5), 613. Griffiths, M.; Mecke, J. F.; Winegar, J. E. Evolution of microstructure in zirconium alloys during irradiation. In Eleventh International Symposium on Zirconium in the Nuclear Industry; Bradley, E. R., Sabol, G. P., Eds.; American Society for testing and Materials: West Conshohocken, PA, p 580, ASTM STP 1295. Jostsons, A.; Kelly, P. M.; Blake, R. G.; Farrell, K. In 9th Conference on Neutron Irradiation-Induced Defect Structures in Zirconium, Effects of Radiation on Structual Materials; Sprague, J. A., Kramer, D., Eds.; American Society for Testing and Materials: West Conshohocken, PA, 1979; pp 46–61, ASTM STP 683. Bell, W. L. J. Nucl. Mater. 1975, 55, 14–22. Kelly, P. M.; Blake, R. G.; Jostons, A. J. Nucl. Mater. 1976, 59, 307–315. Foll, H.; Wilkens, M. Phys. Stat. Sol. (a) 1977, 39, 561–571. Kapinos, V. G.; Osetsky, Y. N.; Platonov, P. A. J. Nucl. Mater. 1992, 195(1–2), 83–101. De Diego, N.; Osetsky, Y. N.; Bacon, D. J. J. Nucl. Mater. 2008, 374(1–2), 87–94. Jostsons, A.; Blake, R. G.; Napier, J. G.; Kelly, P. M.; Farrell, K. J. Nucl. Mater. 1977, 68, 267–276. Dubinko, V.; Turkin, A.; Abyzov, A.; Griffiths, M. J. ASTM Int. 2005, 2, 157. Are´valo, C.; Caturla, M. J.; Perlado, J. M. J. Nucl. Mater. 2007, 367–370, 338–343. Holt, R. A.; Gilbert, R. W. J. Nucl. Mater. 1983, 116, 127–130. Holt, R. A.; Gilbert, R. W. J. Nucl. Mater. 1986, 137, 185–189. Holt, R. A. J. Nucl. Mater. 1988, 159, 310–338. Griffiths, M.; Gilbert, R. W. In Fidleris V Accelerated irradiation growth of zirconium alloys, In Eighth International Symposium on Zirconium in the Nuclear Industry, Philadelphia, PA, 1989; Van Swam, L. F. P., Eucken, C. M., Eds.; American Society for Testing and Materials: West Conshohocken, PA, pp 658–677, ASTM STP 1023. Griffiths, M.; Loretto, M. H.; Smallman, R. E. J. Nucl. Mater. 1983, 115(2–3), 323–330. Gilbon, D.; Simonot C Effect of irradiation on the microstructure of zircaloy-4. In Eleventh International Symposium on Zirconium in the Nuclear Industry; Garde, E. M., Bradley, E. R., Eds.; American Society for Testing and Materials: West Conshohocken, PA, 1994; pp 521–548, ASTM STP 1245. De Carlan, Y.; Re´gnard, C.; Griffiths, M.; Gilbon, D. Influence of iron in the nucleation of component dislocation loops in irradiated zircaloy-4. In Eleventh International Symposium on Zirconium in the Nuclear Industry; Bradley, E. R., Sabol, G. P., Eds.; American Society for Testing and Materials: West Conshohocken, PA, 1996; pp 638–653, ASTM STP 1295.

77.

78. 79. 80. 81. 82.

83.

84. 85. 86. 87.

88. 89. 90. 91.

92. 93.

94.

95.

Adamson, R. B. Effects of neutron irradiation on microstructure and properties of zircaloy. In Twelfth International Symposium on Zirconium in the Nuclear Industry; Sabol, G. P., Moan, G. D., Eds.; American Society for Testing and Materials: West Conshohocken, PA, 2000; pp 15–31, ASTM STP 1354. Faulkner, D.; Woo, C. H. J. Nucl. Mater. 1980, 90, 307–316. Griffiths, M.; Gilbert, R. W.; Carpenter, G. J. C. J. Nucl. Mater. 1987, 150(1), 53–66. Gilbert, R. W.; Griffiths, M.; Carpenter, G. J. C. J. Nucl. Mater. 1985, 135, 265–268. Yang, W. J. S. J. Nucl. Mater. 1988, 158, 71–80. Motta, A. T.; Lefebvre, F.; Lemaignan, C. Amorphisation of precipitates in Zircaloy under neutron and charged particle irradiation. In Ninth International Symposium on Zirconium in the Nuclear Industry; Eucken, C. M., Garde, A. M., Eds.; American Society for Testing and Materials: West Conshohocken, PA, 1991; pp 718–739, ASTM Special Technical Publication 1132. Garzarolli, F.; Goll, W.; Seibold, A.; Ray, I. Effect of in-PWR irradiation on size, structure and composition of intermetallic precipitates of Zr alloys; Bradley, E. R., Sabol, G. P., Eds.; 1996; pp 541–556, ASTM STP 1295. Motta, A. T.; Lemaignan, C. J. Nucl. Mater. 1992, 195, 277–285. Toffolon-Masclet, C.; Barberis, P.; Brachet, J. C.; Mardon, J. P.; Legras, L. J. ASTM Int. 2005, 2, 81–101. Doriot, S.; Gilbon, D.; Bećhade, J. L.; Mathon, M.; Legras, L.; Mardon, J. P. J. ASTM Int. 2005, 2, 175–201. Shishov, V.; Nikulina, A. V.; Markelov, V. A.; et al. Influence of neutron irradiation on dislocation structure and phase composition of Zr-base alloys; Bradley, E. R., Sabol, G. P., Eds.; American Society for Testing and Materials: West Conshohocken, PA, 1996; pp 603–622, ASTM-STP 1295. Cann, C. D.; So, C. B.; Styles, R. C.; Coleman, C. E. J. Nucl. Mater. 1993, 205, 267–272. Kobylyansky, G. P.; Novoselov, A. E.; Ostrovsky, Z. E.; et al. J. ASTM Int.; 2008, 5(4). Howe, L.; Thomas, W. R. J. Nucl. Mater. 1960, 2(3), 248–260. Hardy, D. G. The effect of neutron irradiation on the mechanical properties of zirconium alloy fuel cladding in uniaxial and biaxial tests. In Irradiation Effects on Structural Alloys for Nuclear Reactor Applications; Bement, A. L., Ed.; American Society for Testing and Materials: West Conshohocken, PA, 1970; pp 215–258, ASTM-STP 484. Higgy, H. R.; Hammad, F. H. J. Nucl. Mater. 1972, 44, 215–227. Rieger, G. F.; Lee, D. Strength and ductility of neutron irradiated and textured Zircaloy-2 In Zirconium in Nuclear Applications; Schemel, J. H., Rosenbaum, H. S., Eds.; American Society for Testing and Materials: West Conshohocken, PA, 1974; pp 355–369, ASTM STP 551. Baroch, C. J. Effect of irradiation at 130, 650, and 775 F on tensile properties of Zircaloy-4 at 70, 650, and 775 F. In Properties of Reactor Structural Alloys after Neutron or Particle Irradiation; Baroch, C. J., Ed.; American Society for Testing and Materials: West Conshohocken, PA, 1975; p 129, ASTM STP 570. Pettersson, K.; Vesterlund, G.; Andersson, T. Effect of irradiation on the strength, ductility, and defect sensitivity of fully recrystallized zircaloy tube. In Schemel, J. H., Papazoglou, T. P., Eds.; American


96. 97.

98.

99. 100.

101.

102. 103. 104. 105. 106. 107. 108. 109. 110. 111. 112.

113. 114. 115. 116.

Society for Testing and Materials: West Conshohocken, PA, 1979; pp 155–173, ASTM STP 681. Onchi, T.; Kayano, H.; Higashiguchi, Y. J. Nucl. Mater. 1980, 88(2–3), 226–235. Yasuda, T.; Nakatsuka, M.; Yamashita, K. Deformation and fracture properties of neutron-irradiated recrystallized Zircaloy-2 cladding under uniaxial tension. In Seventh International Symposium on Zirconium in the Nuclear Industry, Strasbourg, France, June 24–27, 1985; Adamson, R. B., Van Swan, L. F. P., Eds.; American Society for Testing and Materials: West Conshohocken, PA, 1987; pp 734–747, ASTM STP 939. Morize, P.; Baicry, J.; Mardon, J.-P. Effect of irradiation at 588 K on mechanical properties and deformation behavior of zirconium alloy strip. . In In Zirconium in the Nuclear Industry, Strasbourg, France, June 24–27, 1985; Adamson, R. B., Van Swan, L. F. P., Eds.; American Society for Testing and Materials: West Conshohocken, PA, 1987; pp 101–119, ASTM STP 939. Byun, T. S.; Farrell, K. J. Nucl. Mater. 2004, 326, 86–96. Lemaignan, C.; Motta, A. T Zirconium alloys in nuclear applications;In Materials Science and Technology Series, A Comprehensive Treatment; Cahn, R. W, Haasen, P., Kramer, E. J., Eds.; VCH-Verlag publisher: Weinheim, Germany, 1994; Vol. 10B, Nuclear Materials, Part 2, Frost, B. R. T., Ed. Adamson, R. B.; Bell, W. L. Effects of neutron irradiation and oxygen content on the microstructure and mechanical properties of zircaloy In Microstructure and Mechanical Behaviour of Materials, X’ian, Republic of China, Oct 21–24, EMAS Warley, UK, 1985; Vol. I, pp 237–246. Torimaru, T.; Yasuda, T.; Nakatsuka, M. J. Nucl. Mater. 1996, 238(2–3), 169–174. Nakatsuka, M.; Nagai, M. J. Nucl. Sci. Tech. 1987, 24, 832–838. Zu, X. T.; Sun, K.; Atzmon, M.; et al. Philos. Mag. 2005, 85(4), 649–659. Ribis, J.; Onimus, F.; Bećhade, J. L.; et al. J. ASTM Int.; 2008, 5(3). Kroupa, F.; Hirsch, P. B. Discuss. Faraday Soc. 1964, 38, 49–55. Makin, M. J. Philos. Mag. 1964, 10, 695–711. Saada, G.; Washburn, J. J. Phys. Soc. Jpn. 1963, 18 (Suppl. 1), 43. Foreman, A. J. E.; Sharp, J. V. Philos. Mag. 1969, 19, 931–937. Hirsch, P. B. In Proceedings of a Conference on Point Defect Behavior and Diffusional Processes, University of Bristol, Sept 13–16 1976. Carpenter, G. J. C. Scr. Metall. 1976, 10, 411–413. Fregonese, M.; Re´gnard, C.; Rouillon, L.; Magnin, T.; Lefebvre, F.; Lemaignan, C. Failure mechanisms of irradiated Zr alloys related to PCI: Activated slip systems, localized strains, and iodine-induced stress corrosion cracking. In 12th International Symposium on Zirconium in Nuclear Industry; Sabol, G. P., Moan, G. D., Eds.; American Society for Testing and Materials: West Conshohocken, PA, 2000; p 377, ASTM STP 1354. Onimus, F.; Monnet, I.; Bećhade, J. L.; Prioul, C.; Pilvin, P. J. Nucl. Mater. 2004, 328, 165–179. Voskoboynikov, R. E.; Osetsky, Y. N.; Bacon, D. J. Mater. Sci. Eng. A 2005, 400–401, 54–58. Onchi, T.; Kayano, H.; Higashiguchi, Y. J. Nucl. Sci. Technol. 1977, 14(5), 359–369. Onimus, F.; Bećhade, J. L.; Duguay, C.; Gilbon, D.; Pilvin, P. J. Nucl. Mater. 2006, 358, 176–189.

117.

29

Garde, A. M Effects of irradiation and hydriding on mechanical properties of Zy-4 at high fluence. In Eighth International Symposium on Zirconium in the Nuclear Industry; Van Swam, L. F. P., Eucken, C. M., Eds.; American Society for Testing and Materials: West Conshohocken, PA, 1989; pp 548–569, ASTM STP 1023. 118. Williams, C. D.; Adamson, R. B.; Olshausen, K. D. Effects of boiling water reactor irradiation on tensile properties of Zircaloy. In European Conference on Irradiation Behaviour of Fuel Cladding and Core Component Materials, Karlsruhe, Germany 1974. 119. Adamson, R. B; Wisner, S. B; Tucker, R. P; Rand, R. A. Failure strain for irradiated zircaloy based on subsized specimen testing and analysis, the use of small scale specimens for testing irradiated materials; Corwin, W. R., Rosinski, S. R., van Walle, E., Eds.; American Society for Testing and Materials: West Conshohocken, PA, 1986, p 171, ASTM STP 888. 120. Re´gnard, C.; Verhaeghe, B.; Lefebvre-Joud, F.; Lemaignan, C. Activated slip systems and localized straining of irradiated alloys in circumferential loadings. In 13th International Symposium on Zirconium in the Nuclear Industry; Moan, G. D., Rudling, P., Eds.; American Society for Testing and Materials: West Conshohocken, PA, 2002; p 384, ASTM STP 1423. 121. Coleman, C. E.; Mills, D.; van der Kuur, J. Can. Metall. Q. 1972, 11, 91–100. 122. Tomalin, D. S. Localized ductility of irradiated Zircaloy-2 cladding in air and iodine environments. In 13th International Symposium on Zirconium in the Nuclear Industry; Lowe, P., Jr, Ed.; American Society for Testing and Materials: West Conshohocken, PA, 1977; pp 557–572, ASTM STP 633. 123. Pettersson, K. J. Nucl. Mater. 1982, 105, 341–344. 124. Onimus, F.; Bećhade, J. L.; Prioul, C.; et al. J. ASTM Int.; 2005, 2(8). 125. Fournier, L.; Serres, A.; Auzoux, Q.; Leboulch, D.; Was, G. S. J. Nucl. Mater. 2009, 384(1), 38–4731. 126. Wechsler, M. S. The Inhomogeneity of Plastic Deformation; ASM: Metals Park, OH, 1973; pp 19–52. 127. Luft, A. Prog. Mater. Sci. 1991, 35, 97–204. 128. Sharp, J. V. Philos. Mag. 1967, 16, 77–96. 129. Sharp, J. V. Radiat. Effects 1972, 14, 71. 130. Makin, M. J. Philos. Mag. 1970, 21, 815–821. 131. Nogaret, T.; Robertson, C.; Rodney, D. Philos. Mag. 2007, 87(6), 945. 132. Nogaret, T.; Rodney, D.; Fivel, M.; Robertson, C. J. Nucl. Mater. 2008, 380, 22–29. 133. Lee, D.; Adamson, R. B. In Modeling of localized deformation in neutron irradiated Zircaloy-2; Lowe, P., Jr., Ed.; American Society for Testing and Materials: West Conshohocken, PA, 1977; pp 385–401, ASTM STP 633. 134. Franklin, D. G.; Lucas, G. E.; Bement, A. L. Creep of zirconium in nuclear reactors; Franklin, D. G., Lucas, G. E., Bement, A. L., Eds.; American Society for Testing and Materials: West Conshohocken, PA, 1983; pp 1–167, ASTM STP 815. 135. Onimus, F.; Bećhade, J. L. J. Nucl. Mater. 2009, 384, 163–174. 136. Nakatsuka, M.; Nagai, M. J. Nucl. Sci. Technol. 1987, 24, 832–838. 137. Delobelle, P.; Robinet, P.; Geyer, P.; Bouffioux, P. J. Nucl. Mater. 1996, 238(2–3), 135. 138. Coleman, C. E.; Chow, P. C. K.; Ells, C. E.; Griffiths, M.; Ibrahim, E. F.; Sagat, S. Rejuvenation of fracture properties of irradiated Zr-2,5 Nb by heat treatment; Stoller, R. E., Kumar, A. S., Gelles, D. S., Eds.; American

30

139. 140. 141.

142.

143. 144. 145. 146. 147. 148. 149. 150. 151.

152. 153.

154.

155. 156. 157.

158. 159.

160. 161.

Radiation Effects in Zirconium Alloys Society for Testing and Materials: West Conshohocken, PA, 1992; pp 318–336, ASTM-STP 1125. Dollins, C. C. Radiat. Effects 1972, 16, 271–280. Snowden, K. U.; Veevers, K. Radiat. Effects Defects Solids 1973, 20(3), 169–174. Carpenter, G. J. C; Watters, J. F Irradiation damage recovery in some zirconium alloys. In Zirconium in the Nuclear Industry; Schemel, J. H., Rosenbaum, H. S., Eds.; American Society for Testing and Materials: West Conshohocken, PA, 1974; p 400, ASTM STP 551. Ito, K.; Kamimura, K.; Tsukuda, Y. Evaluation of irradiation effect on spent fuel cladding creep properties. In Proceedings of the 2004 International Meeting on LWR Fuel Performance, Orlando, FL, Sept 19–22, 2004. Veevers, K.; Rotsey, W. B. J. Nucl. Mater. 1968, 27, 108–111. Williams, C. D.; Gilbert, R. W. Radiation damage in reactor In Proceedings of the IAEA Symposium, Vienna, 1969; Vol. 1, p 235. Northwood, D. O.; Causey, A. R. J. Nucl. Mater. 1977, 64, 308–312. Eyre, B. L.; Maher, D. M. Philos. Mag. 1971, 24, 767–797. Hirth, J. P.; Lothe, J. Theory of Dislocations; Wiley: New York, 1982. Burton, B.; Speight, M. V. Philos. Mag. A 1986, 53, 385. Peehs, M.; Fleisch, J. J. Nucl. Mater. 1986, 137, 190–202. Fidleris, V. J. Nucl. Mater. 1988, 159, 22–42. Mayuzumi, M.; Murai, K. In Proceedings of the 1993 International Conference on Nuclear Waste Management and Environmental Remediation. Volume 1: Low and Intermediate Level Radioactive Waste Management; American Society of Mechanical Engineers: New York, NY, 1993; Vol. 776, pp 607–612. Scha¨ffler, I.; Geyer, P.; Bouffioux, P.; Delobelle, P. J. Eng. Mater. Technol. Trans. ASME 2000, 122, 168–176. Soniak, A.; L’Hullier, N.; Mardon, J. P.; Rebeyrolle, V.; Bouffioux, P.; Bernaudat, C. Irradiation creep behavior of Zr-base alloys. In Thirteenth International Symposium on Zirconium in the Nuclear Industry, 2002; Moan, G. D., Rudling, P., Eds.; pp 837–862, ASTM STP 1423. Cappelaere, C.; Limon, R.; Gilbon, D.; et al. Impact of irradiation defects annealing on long-term thermal creep of irradiated Zircaloy-4 cladding tube. In Thirteenth International Symposium on Zirconium in the Nuclear, Industry; 2002; Moan, G. D., Rudling, P., Eds.; pp 720–739, ASTM STP 1423. Limon, R.; Lehmann, S. J. Nucl. Mater. 2004, 335, 322–334. Yasuda, T.; Nakatsuka, M.; Mayuzumi, M. Trans. Am. Nucl. Soc. 1990, 61, 77–78. Murty, K. L.; Mahmood, S. T. Effects of recrystallization and neutron irradiation on creep anisotropy of zircaloy cladding. .In In Ninth International Symposium on Zirconium in the Nuclear Industry; Eucken, C. M., Garde, A. M., Eds.; American Society for Testing and Materials: West Conshohocken, PA, 1990; p 198, ASTM STP 1132. Tsai, H.; Billone, M. C. J. ASTM Int. 2006, 3(1). Fidleris, V.; Tucker, R. P.; Adamson, R. B. An overview of microstructural and experimental factors that affect the irradiation growth behavior of zirconium alloys. In Seventh International Symposium on Zirconium in the Nuclear Industry; Adamson, R. B., Van Swan, L. F. P., Eds.; American Society for Testing and Materials: West Conshohocken, PA, 1987; pp 49–85, ASTM STP 939. Rogerson, A. J. Nucl. Mater. 1988, 159, 43–61. Carpenter, G. J. C.; Zee, R. H.; Rogerson, A. J. Nucl. Mater. 1988, 159, 86–100.

162. 163. 164. 165.

166.

167. 168.

169. 170. 171. 172. 173.

174. 175. 176. 177. 178. 179. 180.

181. 182.

183. 184.

Holt, R. A. J. ASTM Int.; 2008, 5(6). Holt, R. A. J. Nucl. Mater. 2008, 372, 182–214. Carpenter, G. J. C.; Murgatroyd, R. A.; Rogerson, A.; Watters, J. F. J. Nucl. Mater. 1981, 101, 28–37. Williams, J.; Darby, E. C; Minty, D. C. C Irradiation growth of annealed Zircaloy-2. In Sixth International Symposium on Zirconium in the Nuclear Industry; Franklin, D. G., Adamson, R. B., Eds.; American Society for Testing and Materials: West Conshohocken, PA, 1984; pp 376–393, ASTM STP 824. Garzarolli, F.; Dewes, P.; Maussner, G.; Basso, H. H. Effects of high neutron fluence on microstructure and growth of Zircaloy-4. In Eighth International Symposium on Zirconium in the Nuclear Industry, Philadelphia, PA; Van Swam, L. F. P., Eucken, C. M., Eds.; American Society for Testing and Materials: West Conshohocken, PA, 1989; pp 641–657, ASTM STP 1023. Baron, J. L.; Esling, C.; Feron, J. L.; et al. Textures Microstruct. 1990, 12(1–3), 125–140. Gilbon, D.; Soniak, A.; Doriot, S.; Mardon, J. P. Irradiation creep and growth behavior, and microstructural evolution of advanced Zr-base alloys. In Twelfth International Symposium on Zirconium in the Nuclear Industry; Sabol, G. P., Moan, G. D., Eds.; American Society for Testing and Materials: West Conshohocken, PA, 2000; pp 51–71, ASTM STP 1354. Franklin, D. G.; Adamson, R. B. J. Nucl. Mater. 1988, 159, 12–21. Ibrahim, E. F.; Holt, R. A. J. Nucl. Mater. 1980, 91(2–3), 311–321. Holt, R. A. J. Nucl. Mater. 1980, 90(1–3), 193–204. Buckley, S. N. In Properties of Reactor Materials and Effect of Radiation; Litter, W. J., Ed.; Butterworths: London, 1962; p 413. Northwood, D. O. Irradiation growth in zirconium and its alloys. In Effects of Radiation on Structural Materials; Sprague, J. A., Kramer, D., Eds.; American Society for Testing and Materials: West Conshohocken, PA, 1979; pp 62–76, ASTM STP 683. Woo, C. H.; Go¨sele, U. J. Nucl. Mater. 1983, 119, 219–228. Monti, A. M.; Sarce, A.; Smetninsky-De-Grande, N.; Savino, E. J.; Tome, C. N. Philos. Mag. A 1991, 63, 925–936. Ball, C. J. J. Nucl. Mater. 1981, 101, 147–149. Fleck, R. G.; Holt, R. A.; Perovic, V.; Tadros, J. J. Nucl. Mater. 1988, 159, 75–85. Tome´, C. N.; Christodoulou, N.; Turner, P. A.; et al. J. Nucl. Mater. 1996, 227(3), 237–250. Field, G. J. J. Nucl. Mater. 1988, 159, 3–11. Christodoulou, N.; Causey, A. R.; Holt, R. A.; et al. Modeling in-reactor deformation of Zr–2.5Nb pressure tubes in CANDU power reactors. In Eleventh International Symposium on Zirconium in the Nuclear Industry; 1996; Bradley, E. R., Sabol, G. P., Eds.; p 518, ASTM STP 1295. Matthews, J. R.; Finnis, M. W. J. Nucl. Mater. 1988, 159, 257–285. Garde, A. M.; Smerd, P. G.; Garzarolli, F.; Manzel, R. Influence of metallurgical condition on the in-reactor dimensional changes of Zircaloy fuel rods. In Sixth International Symposium on Zirconium in the Nuclear Industry; Franklin, D. G., Adamson, R. B., Eds.; American Society for Testing and Materials: West Conshohocken, PA, 1984; p 289, ASTM STP 824. Coleman, C. E.; Causey, A. R.; Fidleris, V. J. Nucl. Mater. 1976, 60, 185–194. Tinti, F. Nucl. Technol. 1983, 60(1), 104–113.

Radiation Effects in Zirconium Alloys 185. Faulkner, D.; McElroy, R. J. Irradiation creep and growth in zirconium during proton bombardment. In Effects of Radiation on Structural Materials; Sprague, J. A., Kramer, D., Eds.; American Society for Testing and Materials: West Conshohocken, PA, 1979; pp 329–345, ASTM STP 683. 186. Garner, F. A.; Gelles, D. S. J. Nucl. Mater. 1988, 159, 286–309.

187. 188. 189. 190.

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MacEwen, S. R.; Fidleris, V. J. Nucl. Mater. 1977, 66, 250–257. Nichols, F. A. J. Nucl. Mater. 1970, 37, 59. Griffiths, M.; Gilbert, R. W. J. Nucl. Mater. 1987, 150(2), 169–181. Griffiths, M.; Gilbon, D.; Re´gnard, C.; Lemaignan, C. J. Nucl. Mater. 1993, 205, 273–283.

4.02

Radiation Damage in Austenitic Steels

F. A. Garner Radiation Effects Consulting, Richland, WA, USA


4.02.1 4.02.2 4.02.2.1 4.02.2.2 4.02.3 4.02.4 4.02.5 4.02.6 4.02.7 4.02.8 4.02.8.1 4.02.8.2 4.02.8.3 4.02.8.3.1 4.02.8.3.2 4.02.8.3.3 4.02.8.3.4 4.02.8.3.5 4.02.9 4.02.9.1 4.02.9.2 4.02.9.3 4.02.9.4 4.02.9.5 4.02.9.5.1 4.02.9.5.2 4.02.9.5.3 4.02.9.6 4.02.9.7 4.02.9.8 4.02.10 References

Introduction Basic Damage Processes Atomic Displacements Transmutation Differences in Neutron Spectra Transmutation Issues for Stainless Steels Evolution of Radiation-Induced Microchemistry and Microstructure A Cross-Over Issue Involving Radiation-Induced Microstructural Evolution and Transmutation Radiation-Induced Changes in Mechanical Properties Radiation-Induced Changes in Dimension Precipitation-Related Strains Void Swelling and Bubble Swelling Parametric Dependencies of Void Swelling Stress state Elemental composition Alloy starting state Irradiation temperature Influence of dpa rate on swelling Irradiation Creep Introduction Stages of Irradiation Creep Examples of Creep Behavior Creep Disappearance Recent Revisions in Understanding of Irradiation Creep Dependence of irradiation creep on dpa rate Dependence of creep and creep relaxation on neutron spectra Dependence of creep modulus on hydrostatic stress Stress Relaxation by Irradiation Creep Stress Rupture Fatigue Conclusions

Abbreviations ATR BN-350 BN-600 BOR-60

Advanced Test Reactor in Idaho Falls, Idaho Russian acronym for Fast Neutron at 350 MW in Actau, Kazakhstan Russian acronym for Fast Neutron at 600 MW in Zarechney, Russia Russian acronym for Fast Experimental Reactor at 60 MW in Dimitrovgrad, Russia

BR-2 BR-10 BWR CAGR CANDU DFR

34 35 35 37 37 40 44 49 50 61 62 65 67 67 68 69 69 70 74 74 78 79 79 83 83 84 85 86 88 89 90 91

Belgium Research Reactor-II in Mol, Belgium Russian acronym for Fast Reactor at 10 MW in Obninsk, Russia Boiling water reactor Commercial Advanced Gas Reactor Registered trademark for Canadian Deuterium Uranium Reactor Dounreay Fast Reactor in Dounreay, Scotland

33

34


DMTR EBR-II FFTF HFIR HFR IASCC IGSCC JMTR NRU ORNL ORR PWR T/F VVER

Dounreay Materials Test Reactor in Dounreay, Scotland Experimental Breeder Reactor-II in Idaho Falls, Idaho Fast Flux Test Facility, fast reactor in Richland, WA High Flux Isotope Reactor at Oak Ridge National Laboratory High Flux Reactor in Petten, Netherlands Irradiation-assisted stress corrosion cracking Intergranular stress corrosion cracking Japan Material Testing Reactor in Oarai, Japan National Research Universal Reactor in Chalk River, Canada Oak Ridge National Laboratory: Oak Ridge Research Reactor in Oak Ridge, Tennessee Pressurized water reactor Thermal-to-fast neutron ratio Russian acronym for water-cooled, water moderated energetic reactor

4.02.1 Introduction Austenitic stainless steels are widely used as structural 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 Applications. 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 characteristics of iron-based stainless steels following neutron irradiation are also observed in nickelbased alloys. Whenever appropriate, the similarities between the two face-centered-cubic alloy systems will be highlighted. A more comprehensive treatment 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 particle studies that can strongly influence the results. Examples of processes characteristic of charged particle 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 challenging 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 temperature 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 significant levels of stress acting on the component. Stress not only influences cracking and corrosion (see Chapter 5.08, Irradiation Assisted Stress Corrosion Cracking) but can also impact the dimensional stability of stainless steel, primarily due to


thermal creep and irradiation creep, and also from the influence of stress on precipitation, phase stability, and void growth, some of which will be discussed later. However, it will be shown that neutron irradiation can strongly affect both the microstructure and microchemistry of stainless steels and high-nickel alloys, with strong consequences on physical properties, 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 sodiumcooled 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 2 MeV 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 14 MeV and especially from spallation events occurring at energies 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 significant 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

35

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 displacements can be in the form of single displacements resulting from a low-energy neutron collision with a single atom or a glancing collision with 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 displacement of atoms from their lattice site, but these are usually of second-order importance. The first of these processes involve production of energetic electrons produced by high-energy photons via the photoelectric effect, Compton Effect, or pair production.18 These electrons can then cause atomic displacements, but at a much lower efficiency than that associated with neutron-scattering events. The second 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 calculated 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 primarily from displacements and not from transmutation. The use of the dpa concept also relieves researchers 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


X MeV,’ where X is most frequently 0.0, 0.1, 0.5, or 1.0 MeV, have been used for different reactor concepts at different times in history. The threshold energy of 0.1 MeV 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 1.0 MeV. 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 ‘displacement effectiveness’ for 300 series austenitic steels that are useful for estimating dpa from >0.1 MeV fluences for both in-core or near-core spectra in most fission spectra. Examples are 7 dpa per 1022 n cm2 (E > 0.1) for most in-core light water spectra with lower in-core values of 5 dpa per 1022 n cm2 (E > 0.1) for metal fueled fast reactors and 4 dpa per 1022 n cm2 (E > 0.1) for oxide-fueled fast reactors.4 Such conversion 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 cm2 (E > 0.1) to 14 dpa per 1022 n cm2 (E > 1.0). In Russia, a threshold energy of >0.5 MeV is popular for light water

reactors with 9 dpa per 1022 n cm2 (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 sometimes 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 models (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

300 LASREF, 40 C RTNS-II, 90 C OWR, 90 C

250

Yield stress change (MPa)

Yield stress change (MPa)

300

200

150

100

50

0

1017

1018

1019

Neutron fluence, E > 0.1 MeV

1020

250

LASREF, 40 C RTNS-II, 90 C OWR, 90 C

200

150

100

50

0

10-3

10-2

dpa

Figure 1 Radiation-induced yield stress changes of 316 stainless steel versus (left) neutron fluence (n cm2 E > 0.1 MeV), and (right) displacements per atom. Reproduced from Heinisch, H. L.; Hamilton, M. L.; Sommer, W. F.; Ferguson, P. J. Nucl. Mater. 1992, 191–194, 1177, as modified by Greenwood, L. R. J. Nucl. Mater. 1994, 216, 29–44.


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 displacement of atoms there can be a chemical and/or isotopic alteration of the steel via transmutation. With the exception of helium production, transmutation 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 processes 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 relevant 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 categories 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 transmutants, and ‘isotope shifts,’ the latter involving production 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 discrete nuclear reactions to produce discrete isotopes of specific elements, the spallation-induced transmutation 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 transmutation is that the induced radioactivity per unit

37

mass is correspondingly much higher than that produced per dpa in other spectra. The majority of the spalled fragments and their daughters/granddaughters are radioactive with relatively short half-lives, leading to materials that are often much more difficult to examine than materials irradiated in fission spectra. Most importantly, there is a very strong production 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 circumstances 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 equilibrium 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 boundaries. Additionally, it is possible to modify the neutron spectra in a given irradiation capsule by including in it

38


1.E + 16

1015 HFIR

Flux/lethargy

Flux per unit lethargy

HFIR-PTP 1.E + 14

1014 ORR

HFIR-RB* 1.E + 12 ATR-ITV 1.E + 10

1013

EBRII FFTF

1.E + 08 1.E - 9

EBR II 1012

10-8 10-7 10-6 10-5 10-4 10-3 10-2 10-1 100 Neutron energy (MeV)

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).

1014 T/F ~0.15

1.E - 7

1.E - 5 1.E - 3 1.E - 1 Neutron energy (MeV)

1.E + 1

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.

Flux per unit lethargy

1013 Baffle bolt

Top of bolt head

12

10

1011 Upper core plate 10

10

109

10-8 10-7 10-6 10-5 10-4 10-3 10-2 10-1 100 Neutron energy (MeV)

101 102

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.

or enclosing it with a moderator or absorber. Metal hydrides are used in fast reactors to soften the spectrum, while in mixed-spectrum reactors the thermalto-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 twopeaked ‘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 pronounced two-peak spectral distribution where the thermal-to-fast neutron ratio can be very large. These spectral differences lead to strong variations between various reactors in the neutron’s ability to displace atoms and to cause transmutation. Depending 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 effectiveness 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


1016

Thermal flux, clean core Thermal flux, 21-day core Total nonthermal flux >0.111 MeV >0.821 MeV

5 2 1015 5 2

1013

0

4

8

12

H2O

Permanent beryllium

Removable beryllium

Control region

Outer fuel annulus

H2O annulus

2

Inner fuel annulus

5

H2O outer annulus

1014 300-g Pu target

Neutronflux (neutrons per cm2 s–1)

39

16 20 24 28 32 36 40 44 Radial distance from core center (cm)

48

52

56

60

Figure 5 Variation in fast and thermal fluxes in HFIR as a function of radial position at mid-core at 85 MW, also showing change in thermal population with burn-up (Source: ORNL website).

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 lifetime and finally, some high-flux thermal reactors that can produce 10–15 dpa year1 in stainless steels. Most importantly, fast reactors also operate in the high-flux regime, producing 10–40 dpa year1. Therefore, the largest amount of published highdpa 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 differences in neutron spectra, but only in flux. However, the very large dpa rates characteristic of fast reactors are associated with a significant difference in spectrum. 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

5.5 Row 2

5.0 dpa 1022 (E > 0.1) Row 4

4.5

4.0 -20

-10

0

10

20

Axial position (cm) Figure 6 Displacement effectiveness values of dpa per 1022 n cm2 (E > 0.1 MeV) 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 cm2 (E > 0.1 MeV). In mid-core Row 4 the radial leakage is just becoming significant.

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

40


6.0

BC

1

2

FFTF core 3 4

5.0

5

Above core 6 7 8

FFTF cycles 2 and 3

dpa 1022 (E > 0.1)

FFTF cycle 10

4.0

3.0 −100

−75

−50

−25

0

25

50

75

100

125

150

Distance from core midplane (cm) Figure 7 Values of dpa per 1022 n cm2 (E > 0.1 MeV) 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.

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 spectrum 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 0.5–0.4 MeV while light water-cooled reactors have a fast neutron peak near 1.2 MeV. One consequence of attaining successful breeding conditions is that the spectrum-averaged crosssection 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, especially 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 transmutation arising from thermal or epithermal neutrons. Whereas the properties of these metals are particularly 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 component of most 300 series stainless steels, but when included it participates in the formation of carbide


with the major alloy components. This type of reaction occurs only above high neutron threshold energies (>6 MeV). 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 1 MeV with 58 Ni 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 100 MeV as shown in Figure 9.6 Another very important helium-generation process 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

0.14 Cross-section (barns)

precipitates that change the distribution and chemical 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 important in determining the grain boundary cracking behavior, designated irradiation-assisted stress corrosion cracking (IASCC), especially with respect to the sensitization process.29 The strong loss of manganese in highly thermalized 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 radiationinduced release of these impurities to grain boundaries 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 transmutation 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, especially 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

41

0.12

Ni

0.10 0.08 0.06

Cr Ti

0.04

Fe

0.02 1

2

4 6 Energy (MeV)

8 10

20

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.

42


2500 Inconel 304L 316L 9Cr–1Mo Fe Co Ni Cu

1500 1000

1.6

500 0

0

5

10

15

Ratio to initial value

He (appm)

2000

60

Ni Natural nickel 58Ni 67.85% 60Ni 26.2%

1.2

61Ni 58

Ni

0.8

62Ni 64Ni

6.1% total

dpa

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 800 MeV 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 from Garner, 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.

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 64 Ni. There is no natural 59Ni or 63Ni at the beginning of radiation. During irradiation in a highly thermalized 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 cm2 where the 59/58 ratio peaks at 0.04 and then declines, as shown in Figure 10. This transmutation sequence in nickel is an example of the isotopic shift category of transmutation defined earlier. For other elements used to make stainless steels, there are no consequences to such a shift since the total amount of the element is unchanged

0.4

0.0 1021

59

Ni

1022 1023 Thermal fluence (n cm-2)

1024

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.

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 dpa1 can be experienced along the length of a 316 stainless baffle bolt in the baffle-former assembly of a pressurized water


43

100 Pure nickel in HFIR-PTP Percentage increase

80

60 56Fe

40

340 keV 1701 dpa

4

He 4.8 MeV 62 dpa

20

0

20

40

60

80

100

120

140

160

Displacements (dpa) neglecting 59Ni (n, a) 56Fe reaction 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 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.

reactor4,33,39 while comparable rates in fast reactors are in the order of 0.1–0.2 appm dpa1. In thermalized 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 concomitant increases in nuclear heating rates. Nuclear heating by elastic collisions with high-energy neutrons is usually too small to be of much significance. The 59Ni(n, a) reaction releases 5.1 MeV, producing a 4.8 MeV alpha particle which loses most of its energy by electronic losses, depositing significant thermal energy but producing only 62 atomic displacements per each event. However, the recoiling 56 Fe 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 (1.85 MeV), producing a total of 222 displacements per event.7,40 In addition, approximately five displaced 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 thermalto-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 highnickel alloy X-750 were examined after 18.5 years of operation far from the core and were found to be

44


completely relaxed. Calculating the 59Ni contribution, 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 calculated 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 irradiation. At the peak 59Ni level reached at 4 1022 n cm2, the nuclear heating rates from the energetic (n, a) and (n, p) reactions are 0.377 and 0.023 W g1 of nickel, significantly larger than the neutron heating level of 0.03 W g1 of natural nickel. Thus, an increase in nuclear heating of 0.4 W g1 of nickel must be added to the gamma heating rate at the peak 59Ni level. Fractions of the peak heating rates that are proportional 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 54 W g1 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 1–3 W g1. In this case, an additional 0.4 W g1 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 consequences not only on 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 sawtooth variation of gamma heating rate in the baffleformer 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 compared to nonthermalized spectra typical of fast reactors. 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 surface 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 fission 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 irradiation, 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 cladding 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 RadiationInduced Microchemistry and Microstructure When metals are subjected to displacive irradiation, especially at elevated temperatures, an intricate and coordinated coevolution of microstructure and


microchemistry commences that is dependent primarily on the alloy starting state, the dpa rate, and the temperature, and secondarily dependent on variables such as He/dpa rate and applied or internally generated stresses. In general, the starting microstructure and microchemistry 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 starting state and arrives at a destination determined only by irradiation temperature and dpa rate. This quasiequilibrium 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 disappearing by shrinkage or annihilation. The displacement process produces two types of crystalline point defects, vacant crystalline positions (vacancies) and displaced atoms in interstitial crystalline positions (interstitials). These two defect types are both mobile, but move with different diffusional modes and at vastly different velocities, with interstitials diffusing much faster than vacancies. 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 distribution of the alloy. It is important to note that interstitial agglomerations are constrained to be two-dimensional, while vacancies can agglomerate in both two-dimensional and three-dimensional forms. This dimensional disparity is the root cause of the void swelling phenomenon covered in a later section. The developing ensemble of various defect agglomerations with increasing dose induces significant 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

45

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 classifications are equally applicable to phases formed in other classes of steel. Radiation-induced alterations of microstructure and microchemistry occur because new driving forces arise that do not occur in purely thermal environments. The first of these new driving forces is the presence of very large supersaturations of point defects, especially at relatively low irradiation temperatures (250–550 C). Not only are vacancies present in uncharacteristically high levels, thereby accelerating normal vacancy-related diffusional processes, 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 segregation 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 sinkinduced vacancy gradients. This mechanism is particularly 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 radiationproduced microstructural sinks. This type of segregation 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 new driving force results from the action of the other two driving forces when operating on microstructural sinks that are produced only in irradiation environments. These are Frank interstitial loops, helium bubbles, and voids that may have developed 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

46


(a)

CW 316 SS, thimble tube 70 dpa, 315 ºC

50 nm

(b)


50 nm (c)


50 nm Figure 12 Frank loops observed in a 316 stainless flux thimble from a PWR power reactor (a) 70 dpa, 315 C and (b) 33 dpa, 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 11 dpa. 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.

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. Additionally, 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 during 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 diffusivity.57,61 On a per atom basis, phosphorus has been


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 21.7 dpa at a dpa rate of 0.84 107 dpa s1. 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.

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.

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

47

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 determining the transient duration before void swelling accelerates.1 This evolution often does not necessarily proceed by only one path but occurs in several interactive 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 radiationinduced 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 00 g 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 microstructural 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 preexisting 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

48


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.

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 frequently 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, vacuumfilled ‘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 volumeconservative 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


CW 316 SS, thimble tube 70 dpa, 330 C

49

CW 316 SS, thimble tube 70 dpa, 330 C

Bubbles on grain boundary

Matrix bubbles 1.6 1023 m−3

20 nm

−256 nm UF

20 nm

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.

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 temperature and to a lesser extent on dpa rate. The two major deformation processes, swelling and irradiation creep, do not saturate but reach steady-state deformation rates when quasi-equilibrium microstructures are attained. This coupling of saturation microstructure with steady-state behavior has been characterized 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 quasiequilibrium. These new states usually arise from either the microstructural or microchemical evolution, 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 quasiequilibrium 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 perception that such storage could not occur at reactorrelevant 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 pressurized water reactors. Figure 17 presents such a microstructure observed in a PWR flux thimble tube (cold-worked 316 stainless steel) at 70 dpa and 330 C.76 There is a very high density (>1017 cm3) of nanocavities with diameters 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.

50 20% CW 316 25% CW PCA

40 Uniform elongation (%)

Yield stress (MPa)

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800

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SA 316L SA PCA

SA 316L SA PCA 20% CW 316 25% CW PCA

30

20

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0 0

2

4

6 dpa

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10

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400 800 Yield stress (MPa)

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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 35 appm dpa1. 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.

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 differences of behavior as a function of dpa rate.

More recently, Chatani and coworkers showed that at relatively low irradiation temperatures characteristic of boiling water reactors, the radiationinduced 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 microstructure 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

54


1000

Yield strength (MPa)

365 C 800

600 Cold-worked Annealed 400 25 Ni + 0.04 P

25 Ni 200

45 Ni With 59Ni Without

~0.5 and ~15 appm He per dpa

0

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25 Ni

45 Ni

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10

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10

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20

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dpa 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.

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 proceeds, there is a progressive tendency toward flow localization followed by necking. As seen in Figure 31 the failure surface shows this evolution with increasing 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 radiationproduced 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 suppression of the steel’s ability for plastic deformation. This distinction is made because under some conditions 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 hardening. 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


55

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495 C Original series

600

400

45 Ni

25 Ni + 0.04 P

25 Ni 200

Isothermal repeat series

With 59Ni Without

~0.5 and ~5.0 appm He per dpa


0

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Original series

20

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40

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40

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60

dpa 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.

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 irradiation 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 relationship between strength and swelling at this temperature. As shown in Figure 34 failure surfaces at high swelling levels exhibit transgranular cup-cone morphology where failure proceeded by micropore coalescence 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

56


25

1024

45

700 600 MFE-4

500 400 300

2.3

2.1

400 C

2.4

1023

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1020

AD-1

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395 C

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1.9 nm

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800

Ni 35

7 Cr 15 20

1019

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AD-1 experiment in EBR-II

20

43

30 Nickel (wt%)

450 C 40

40

50

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.

800

0.2% proof stress (MPa)

Phénix Rapsodie

600

400

200

0

10

20 30 Fluence (dpa)

40

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. Pheńix operated at a displacement rate that was three times higher than that of RAPSODIE.

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%, indicating complete brittle failure, as also indicated by the fully transgranular nature of the failure surface. Concurrently, the ductility vanished and the tearing modulus plunged to zero, indicating no resistance to crack propagation. Once a crack has initiated it then propagates 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 microstructural 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 surfaces that precipitated the rather abrupt change in failure behavior.1,95 This late-term void-induced microchemical evolution induces a martensite instability in the matrix, as evidenced by the failure surface being completely coated with alpha-martensite.95


57

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6 2 4 8 Neutron fluence (E > 0.1 MeV, n cm–2)

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.

Ultimate

500

Ultimate strength

Ti, ºC 399 378 374 424 372 371

100

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 the macroscopic behavior exhibited by each alloy. In AISI 316, the flux dependence of precipitation canceled the opposite dependence of other microstructural components.

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 influence of both falling nickel level and low deformation temperature to decrease the stacking fault energy of the matrix.1 When encountered by the advancing crack tip, the epsilon-martensite is converted 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 experiencing 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 concurrent rejection of chromium can lead to strong changes in composition in the matrix during irradiation, pushing the matrix toward ferrite rather than martensite at higher temperatures, especially for steels with nickel content of 0.1 MeV)

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 and Mechanical Properties of Metals and Alloys; ASTM STP 529; 1973; pp 149–164.

1000

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 0.05 mm 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.

regained not because the steel has softened, but because it becomes exceptionally strong and hardened during deformation. As a consequence, the steel has lost the ability to neck.

Ultimate tensile strength (MPa)

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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.


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 70 mm position in Figure 33.

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 deformation in the order of 40–45%, but as the wave moves forward it leaves a relatively uniform local deformation 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

Ti = 460 C ft = 15.5 1022 n cm–2

Tt = 20 C et = 1.9%

Ti = 385 C ft = 12.8 1022 n cm–2

Tt = 205 C et = 5.2%

Ti = 460 C ft = 15.5 1022 n cm–2

Tt = 460 C et = 7.2%

100 mm

59

10 mm

Figure 35 Fractographs of failure surfaces of 20% cold-worked 316 specimens cut from an FFTF duct at high exposure. 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. Note change of fracture mode from channel fracture when tested at 205 and 460 C to brittle fracture when tested at 20 C.

60


s0.2 sUTS

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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.

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 ( 0.1 MeV) 100

KIC (MPa 冪m)

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60 Intergranular fracture

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400 500 Test temperature (C)

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700

Test temperature = 649 C

Fatigue precracked zone

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.

It should be noted that radiation-induced segregation can lead to overall changes in average lattice parameter without actually culminating in observable precipitation. Although there is no convincing evidence 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-ofphase 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’ densities (1012–1017 cm3) leads to a macroscopic increase in volume of the steel. The concentration of these cavities tends to increase with decreasing temperature 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 relatively 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 accumulation without further gas addition. It is well known that bubbles can serve as nuclei for voids, accounting for the known tendency of 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 discussed earlier in Section 4.02.6 these ‘helium-filled’ bubbles are probably pressurized with stored hydrogen as well as helium. Interestingly, the onset of this lateterm 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 systems it is void swelling that is the most important contributor to dimensional instability. In the absence of physical restraint or applied stress field void swelling distributes its strains isotropically with the most

66


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 separate 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.

20% CW 316 1 cm

Unirradiated control

Fluences beyond FFTF goal

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 cm2 (E > 0.1 MeV) 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 phenomenon, although pressure vessel embrittlement has also received a similar amount of attention. A comprehensive review on void swelling and irradiation creep was written in 1994 1 and is now being revised116 not only to incorporate new insights developed over the past decade and a half, but also to revise some earlier perceptions that have not survived 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 bodycentered cubic (bcc) metals, the void swelling process is inherently self-limiting, usually saturating at some value below 5%.9 Such saturation is usually accompanied by a process referred to as ‘self-organization’ whereby voids arrange themselves in threedimensional arrays that exhibit the same crystalline orientation as that of the crystal structure. Unfortunately, for most face-centered cubic (fcc) metals, especially stainless steels, self-organization and saturation of void swelling do not operate under most reactorrelevant conditions, and as a result swelling in austenitic stainless steels is an inherently unsaturable process. Void swelling normally exhibits a transient or incubation 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 commercial 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 discerned in Figure 50 where each irradiation temperature in this experiment is coupled with 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 conditions for both void nucleation and especially rapid void growth have been attained. The conditions for void nucleation must be favorable to end the transient. This usually requires


510 C 80

1% per dpa 538 C 482 C

60 Swelling (%)

593 C

40 427 C 650 C

20

0

454 C

400 C 0

0.2% per dpa

10 20 30 ´ 1022 Neutron fluence (n cm-2) (E > 0.1 MeV)

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.

attainment of a dislocation network to the quasiequilibrium value of 3 1010 cm2, 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

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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 posttransient swelling rate of AISI 316 is typical of all austenitic stainless steel at 1% per dpa, essentially independent of all irradiation or material variables.1,120 This terminal rate also appears to be characteristic 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 steadystate 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 boundaries 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 parametric 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 produced at a stress below yield. In effect, the steel is plastically deformed and warm-worked during the rise to power, raising the dislocation density.

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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 experiments involved only tensile stress states, some experiments suggested that both tensile and compressive stress states shortened the transient regime.1 Two recent studies have convincingly shown that the hydrostatic component of stress is relatively unimportant and that it is the deviatoric component or shear stress that accelerates swelling.124,125 This is especially evident 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 stressenhanced 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 increments 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-beforecreep and creep-before-swelling scenarios are different 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 pressurized tube, although the memory faded as irradiation continued.127 The memory is thought to reside in the stress-induced anisotropic distribution of Burger’s vectors, 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 transient duration. Increases in nickel initially increase the effective vacancy diffusion and thereby the transient, 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 therefore 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

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.


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 irradiation, often resulting in higher swelling.1 Additionally, in some alloys and metals it is difficult 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 determining 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 temperature range employed in producing cold-worked tubing by multiple drawings. It is speculated that

Double-aged to precipitate all carbon from solution

DD (%) D

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 0.009 wt%, 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, vanadium, 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.

69

Solution annealed

20% cold-worked Distance along fuel pin 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.

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 intermetallic phase evolution, and its chemical activity can be strongly affected by thermal and mechanical history, 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 versus 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 gradients, 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


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 isothermal 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 temperature 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 early data on swelling was 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 temperaturedependent 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 temperature 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 conducted 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 previously 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 components were chosen because they were made of the same steel used to construct the baffle-former-barrel assembly of PWR internals and the hexagonal subassemblies 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 temperatures (373–388 C), but covering a very large range of dpa rates (0.06–3.8 107 dpa s1), with no significant 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 temperatures 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 temperatures near 400 C at seven different dpa rates. Measurement techniques used were density change


71

3.0 U1603

Swelling (%)

2.5

U9009

U1603 Row 14 0.062–0.156 ´ 10-7

2.0

U8972

U9807

U9009 Row 10 0.38–0.96 ´ 10-7 U8972 Row 9 1.00–2.05 ´ 10-7

1.5 1.0

U9807 Row 8 1.25–3.60 ´ 10-7 dpa/sec

0.5 0.0 0

5

10

15

(a)

20

25

30

35

dpa 3.0 U9009

U9007

Swelling (%)

2.5 U9009 0.38–0.96 ´ 10-7 dpa/sec

2.0 1.5 1.0

U9007 0.44–1.12 ´ 10-7 dpa/sec

0.5 0.0 0

5

10

15

(b)

20

25

30

35

dpa

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.

and microscopy. Multiple specimens were irradiated side-by-side and the measured swelling was remarkably 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 approaching 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.

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10 dpa 0.15 ´ 10-7 dpa s–1 1.2% swelling

100 nm

14.3 dpa 1.8 ´ 10-7 dpa s–1 0.42% swelling

100 nm

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%.

Fe–15Cr–16Ni

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 producing 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 irradiated 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 5.06 107 dpa s1 and 480 C in BOR-60 and at 1.58 106 dpa s1 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 uniformly observed in the lower flux irradiation in BOR-60, regardless of alloy composition.

Fe–15Cr–16Ni–0.25Ti

30

Swelling (%)

25 0.54

20

0.78

15 10

0.31 0.09

5 0

0

10

1.70 ´ 10-6 dpa s–1 20

30

40

10 20 50 60 0 Cumulative dose (dpa)

30

40

50

60

70

Figure 55 Swelling of simple ternary and quaternary model austenitic alloys at 400 C in FFTF, showing a progressive decrease in the transient duration as the dpa rate decreases. Reproduced from Okita, T.; Sekimura, N.; Garner, F. A.; Greenwood, L. R.; Wolfer, W. G.; Isobe, Y. In Proceedings of 10th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors; 2001. Note that all swelling curves have reached or are approaching a terminal swelling rate of 1% per dpa (see dotted line).


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Swelling (%)

25 20 15 10 5 0 0

1.5

0.5 1 Silicon (wt%)

10 nm Avg. size = 8.6 nm

25 Swelling (%)

20

20 nm

15 10 5 0 Nb

Nb+Ti Solute addition BOR-60

Ti

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, 5 107 dpa s1) and BN-350 (490 C, 53 dpa, 15.6 107 dpa s1), 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.

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 powerproducing light water-cooled reactors. Such a conclusion 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 perceived 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, the maximum water temperature in PWRs. It was also thought that the lower dpa rates characteristic of PWRs would reduce vacancy supersaturations and would therefore inhibit void nucleation.

Density = 0.61 ´ 1022 m-3 Swelling = 0.20%

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.

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 transient 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 constructed 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 resolvable 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

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annealed 304 at the same temperature and dpa rate the worrisome inference is that the 304 plate surrounding 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 researchers148 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.

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 Neustroev and coworkers also found voids in a thimble tube removed from a VVER operating 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 components 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 example 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 duration decreased with lower dpa rate.157–159 Most importantly, 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 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 2 K1R1 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 50 nm in length.

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 magnitude increase in creep rate at relatively low temperatures 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


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 1.9 109 dpa s1. 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.

75

in this experiment was later calculated by Foster and coworkers to be 0.95 106 (MPa dpa)1.163 In this case, the specimen is increasing in length and decreasing in cross-section. Precipitation of carbides leads to a small densification and shrinkage of the specimen as shown in the thermal creep behavior. A similar densification process occurs during irradiation but its strains are overwhelmed by the irradiation creep strains. The accumulated damage is relatively small ( 1.0 MeV)

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.


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 temperature, 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 contribution, effectively doubling the creep rate. There are a number of consequences of the coupling between swelling, creep, and precipitationrelated strains. 1. The onset of swelling can be detected by a jump in creep modulus B long before measurable swellinginduced changes in dimension can be detected, and often before microscopy confirms the presence 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 1 106 (MPa dpa)1.

79

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 misleadingly 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 D1 or 160 MPa.1,16 4.02.9.3

Examples of Creep Behavior

Various aspects of behavior resulting in irradiation creep can be illustrated with some examples presented 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

12

1.0 PCA

10

405 C, 24.4 ´ 1022 n cm- 3 (E > 0.1 MeV)

0.8 D/De

384–406 C

8 DD (%) D0

43.3 MPa FV548 26.6 MPa

0.6

354 C, 18.4 ´ 1022

6

0.4 390 C, 15.1 ´ 1022

4

52.2 MPa 56.0 MPa 57.0 MPa

0.2

316 C, 11.6 ´ 1022

2

406 C,

0

0.0

401 C, 6.7 ´ 1022

0

50

100 150 200 Hoop stress (MPa)

4.6 ´ 1022

250

300

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%.

0

PE16

1 2 3 ´ 1024 –2 Fluence (n m ) (E > 1 MeV)

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.

80


Strain/stress (0.0001 %/MPa)

7 2 ´ 10-8 dpa s–1

6 5

20% CW 316 at 370 C

4 3

Annealed 304

2

175–200 C 300–370 C

1 0 0.0

0.1

0.2

0.3

0.4 Dose (dpa)

0.5

0.6

0.7

0.8

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 0.92 106 (MPa dpa)1. The larger transient of cold-worked 316 is due to its much higher dislocation density.

0.06 370 C, 130 MPa

316

0.04

0.02

PCA AMCR 0033

0.00 −0.02

Length change (%)

(a)

0.15 20% Cold-worked AMCR 0033 0.1 390 C, 100 MPa

420 C, 130 MPa 0.05 0.0

350 C, 100 MPa

320 C, 100 MPa

−0.05 (b)

20% Cold-worked AMCR 0033

380 C, 50 MPa 0.15 As received Aged 400 C, 1 h Aged 600 C, 1 h

0.10

0.05

0 (c)

0

1

2

3

4

dpa

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. Schu¨le, 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.

81


8 316Ti

6

4

DV (%) V0

Creep coefficient B

Swelling

6 10−6 MPa−1 dpa−1 C = 0.046 wt%

400–420 C

450–460 C

4

C = 0.006 wt%

2

316Ti+P

2 0

12 1026 4 8 −2 Fluence (n m ) (E > 0.1 MeV)

0

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 1 106 (MPa dpa)1 and D to be 0.6 102 MPa1.

400–420 C

Irradiation creep

6 Creep strain (%)

0

170 MPa 450–470 C

4 400–420 C 450–470 C

2

90 MPa

0 0 330 C 400 C

2.0 e (%)

500 C 600 C

13.1 dpa

20% CW 316

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.

12.0 dpa

1.0 BO = 2.8 10−6 MPa−1 dpa−1 0.0 300 C 400 C

2.0

500 C e (%)

600 C

13.3 dpa

25% CW PCA

12.1 dpa

1.0 BO = 3.2 10−6 MPa−1 dpa−1 0.0

0

100

200

300

400

500

Effective stress (MPa)

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 B0 ¼ 1 106 (MPa dpa)1, however, probably arising from observed high densities of helium bubbles to produce bubble-enhanced creep.

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 evidence 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

82


2.5 0 MPa

A094, T-415 C

60 MPa

Midwall creep strain/hoop stress (% per MPa)

Swelling-induced diametral strain (%)

C42, T-415 C C38, T-390 C

2.0

C39, T-390 C C40, T-390 C C44, T-390 C

1.5

83508, T-420 C

1.0

K280, T-395 C

A095, T-415 C

0.5

83508

100 MPa

0.05

140 MPa 200 MPa

Failed in next cycle

300 MPa

0.04

K280

0.03

0.02

A095

0.01 Failed in next cycle

0 0

100

50

0

150

dpa

0

50

100

150

dpa

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.

16 BEQ10-6 (MPa dpa)-1

14

56 dpa

2.

12 10 10 dpa

8

3.

6 6.3 dpa 4 2 0

0

10

20 30 Ni-equivalent (%)

40

50

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.

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%

4.

5.

per dpa or one-third of the eventual steady-state swelling rate. As swelling approaches 1% per dpa the creep rate backs down proportionately to maintain this maximum rate as shown in Figures 78–80. 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. 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. 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


83

6

Pin 32 Swelling (%)

Swelling (%)

8

4 2

Pin 31

0

Creep modulus (MPa dpa F)-1

Creep strains (%)

Pin 47 2

0

Pin 32 1.2

0.8 Pin 31 0.4

0 Bottom

Pin 1

4

Top

6´10-6

Pin 1 Pin 47

4

2

0

400

Fuel column length

500 600 Temperature (ºC)

700

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 Onset of creep disappearance Instantaneous 20 creep coefficient 1030 psi n cm-2 10

0

Swelling-enhanced creep

0

10

20

Swelling in the absence of creep 30 dpa

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.

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, irradiation creep quickly assumes all the parametric dependencies 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 perception 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

84


6

10 30 ksi

487-543 C

8

15 ksi

0.33% per dpa

6

Total

4 Diameter change (%)

DD (%) D

4 Swelling deformation

Plastic deformation

2

Hoop stress

2

= 0 ksi

0 -2

(a)

8

30 ksi Irradiation creep

0.33% per dpa

6 Stressfree swelling

4 2 0

0

2

4

6

8

10

12

Local atomic burnup (%) 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.

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 available at that time supported a flux dependence and other data sets supported an independence of dpa rate. On balance it appeared that a strong dependence of irradiation creep rate on dpa rate was the more defendable conclusion. With hindsight and additional published data supporting 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

Stress-affected swelling at 30 ksi

0 -2

(b) 0

20

40

60

80

100

dpa 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 206 MPa) does not double the deformation rate, which never exceeds 0.33% per dpa. (b) Density measurements on the 30 ksi (206 MPa) tube show that stress accelerates the rate of swelling, but also causes the creep rate to approach zero at high swelling levels.

Mosedale and Lewthwaite. The authors had not properly 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 compliance 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 produce cascades and therefore there is less in-cascade annihilation. Thus, a larger fraction of thermally produced defects are postulated to survive to contribute to creep or embrittlement.180,181


550 C Core 1

12

550 C Core 4

0.33% per dpa

0 MPa 35 MPa 70 MPa 117 MPa 163 MPa 233 MPa

8

4

DD (%) D

85

0 575 C Core 1

12

575 C Core 4

0 MPa 16 MPa 31 MPa 63 MPa 104 MPa 146 MPa

8

4

0 0

20

40

60

0 dpa

20

40

60

80

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.

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 contributions 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 transmuted 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 confirmation of the independence of creep compliance on the sign of the hydrostatic stress component can be found in some stress relaxation experiments.

86


4.8

Normalized creep rate

4.0

1.0 Stress reduction ratio

J B EN58 E FV548 Mk 1 helices 347 S.S. 240–360 C H6 helices (1) Annealed M316 T < 304 C (2) (3)

0.5

s0 = 216 MPa 3.0 0

0

0.5

1.0

2.0

1.5 dpa

1.0 0.9

Preload 23.6 N 36.5 N

1.0

0.6

0

1

2 3 4 5 Displacement rate (dpa s–1)

6 ´ 107

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.

4.02.9.6 Creep

Stress Relaxation by Irradiation

There are situations where the applied load is initially 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 experiments 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 predict 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 1.6 106 (MPa dpa)1, just slightly larger than B0 and probably

Stress reduction ratio

2.0 0.7

0.5

0.3

0

1

2 3 Neutron dose (dpa)

4

5

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.

enhanced by low levels of voids or bubbles in this high-nickel alloy. In NRU, however, the creep relaxation 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

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. 1 304 0.8 Stress ratio (a/d)

C ring (A2 = 1.2 ´ 10-6 MPa dpa–1)

0.6

0.4 Bend (A2 = 1.8 ´ 10-8 MPa dpa–1)

0.2

Irradiated temperature: - 561 K 0 0

3 2 Dose (dpa)

1

4

5

Figure 83 Stress relaxation experiments conducted on 304 stainless steel at 288 C in water-cooled JMTR at 0.82–1.7 107 dpa s1, 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.

Creep relaxation by irradiation is important in that it can reduce the opportunity for irradiationassisted 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 unloading 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 1

250 Before irradiation

200

After 3–6 dpa irradiation

sy

3W-H

0.8 Stress relaxation s/s0

150

sy (MPa)

100 50 0 -50

Irradiation temperature: 561 K

0.6

0.4

Tensile (6 mm from surface)

0.2

-100 -150

Compressive (4 mm from surface) Distance from surface, 4 mm

0

10

20 30 40 Distance from left edge (mm)

50

60

0

0

1

2 Dose (dpa)

3

4

Figure 84 Residual stresses in SA 304 associated with a one-pass weld with mechanical constraint. Stress reversals occur with depth from the surface. Reproduced from Obata, M.; Ishiyama, Y.; Nakata, K.; Sakamoto, H.; Anzai, H.; Asano, K. J. ASTM Int. 2006, 3, 15–31. Residual stresses before and after irradiation were measured by neutron diffraction. Note that B0 was determined to be 1 106 (MPa dpa)1 and is independent of the sign of the hydrostatic stress.

88


350

200 10 years 40 years

300

180

140 200

Torque (nm)

Axial stress (MPa)

A B C

160

250

150

100

50

120 100 80 60

0 0

10

20

30

40

50

60

70

Time (year)

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.

differential swelling. Figure 85 shows several calculated histories of bolt loading for PWR-relevant temperatures 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

0

10

20

30

40

Dose (dpa) 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 from Massoud, J. P.; Dubuisson, P.; Scott, P.; Ligneau, N.; Lemaire, E. In Proceedings of Fontevraud; 2002; Vol. 5; paper 62, 417.

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 produce high He/dpa ratios will decrease rupture life, especially at higher temperatures, compared to irradiation 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.


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. 103 700 C

Stress (MPa)

Annealed Thermal aged

Cold-worked 102 Irradiated in BR-2

101 1.3

1.4

1.5

1.6

1.7

1.8

T(13.5 + log tR) (K)

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-ofreactor 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-ofreactor 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 in Figure 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 unirradiated material, irradiation will result in much earlier crack formation

In-air In-sodium In-reactor

500

Hoop stress (MPa)

89

300 Irradiation effect

Sodium effect

100 80 60 14

14.5

15

15.5

16

16.5

17

17.5

18

LMP = T (14.04 + log tR) / 1000 Figure 88 Creep rupture behavior of 20% cold-worked modified 316 stainless steel, showing effect of sodium and irradiation to reduce failure lifetimes. Reproduced from Ukai, S.; Mizuta, S.; Kaito, T.; Okada, H. J. Nucl. Mater. 2000, 278, 320–327.

90


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 radiation on austenitic steels when exposed to neutron

irradiation. Structural components used in various nuclear reactors may have been constructed from alloys with carefully tailored and optimized properties, 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.

Hoop stress (MPa)

1000

100 D9I D9 575 C 605 670 750

10 12

D9

D9I 575 C 630 695 775

14 16 18 LMP, T (13.5 + log tR) ´ 10-3 (K, h)

20

Figure 89 Improvement of in-reactor [FFTF fast reactor] stress rupture properties of D9 stainless steel by controlled additions of B and P. Reproduced from Hamilton, M. L.; Johnson, G. D.; Puigh, R. J.; et al. In Proceedings ASTM Symposium on Residual Elements in Steel; ASTM STP 1042; 1989, pp 124–149.

10.0

Total strain range, D

'

1

(%)

- Unirradiated - f1 = 0.7-2 ´ 1026 n m–2 - Unirradiated, aged 115 days at 430 C

1.0

0.1 102

103

104

105 106 Cycles to failure

107

108

Figure 90 Fatigue life of 20% cold-worked AISI 316 stainless steel irradiated in HFIR to a maximum dose of 15 dpa and 900 appm He. Reproduced from Grossbeck, M. L.; Ehrlich, K.; Wassilew, C. J. Nucl. Mater. 1990, 174, 264–281.


Plastic strain range (%)

100

EC - 316L : 430 °C Nonirrad. Irrad.: 10 dpa

5.

5

6.

2

7.

10-1

8.

5

9. 10. 11.

2 10-2 103

106

12.

Figure 91 Plastic strain versus number of cycles to failure of annealed EC-316L irradiated to 10 dpa at 430 C in BR2. Reproduced from Grossbeck, M. L.; Ehrlich, K.; Wassilew, C. J. Nucl. Mater. 1990, 174, 264–281; Vandermueulen, W.; Hendrix, W.; Massault, V.; Van de Velde, J. J. Nucl. Mater. 1988, 155–157, 953–956. Using total strain rather than plastic strain, the reduction of life was only a factor of 2, relatively independent of strain range.

13.

104 105 Number of cycles to failure

14. 15. 16. 17. 18.

This evolution of properties and dimensions frequently determines the lifetime of any given structural component, a lifetime that will be very specific to each nuclear environment. It is important to recognize that all potential degradation processes may not yet have been identified and that others may lie over the current exposure horizon, especially as light water reactors are being considered for life extension to 60 or 80 years, and as fast reactors are being designed for doses well beyond 200 dpa.

References 1.

2. 3.

4.

Garner, F. A. In Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors; Vol. 10A of Materials Science and Technology: A Comprehensive Treatment, VCH Publishers, 1994; Chapter 6, pp 419–543. Garner, F. A. In Proceedings of NATO Advanced Research Workshop on; Springer: Berlin, 2007; pp 307–327. Garner, F. A. In Understanding and Mitigating Ageing in Nuclear Power Plants: Materials and Operational Aspects of Plant Life Management (PLiM); Tipping, P. G., Ed.; Woodhead Publishing Series in Energy: Number 4; Woodhead Publishing: Cambridge, UK, 2010; Chapter 10, pp 308–356. Garner, F. A.; Greenwood, L. R. In 11th International Conference on Environmental Degradation of Materials

19. 20.

21. 22. 23. 24.

25. 26. 27. 28. 29. 30. 31.

91

in Nuclear Power Systems – Water Reactors; 2003; pp 887–909. Garner, F. A.; Greenwood, L. R. Radiat. Eff. Def. Solids 1998, 144, 251–283. Garner, 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. Greenwood, L. R.; Garner, F. A. J. Nucl. Mater. 1996, 233–237, 1530–1534. Garner, F. A.; Greenwood, L. R.; Oliver, B. M. In Effects of Radiation on Materials: 18th International Symposium; ASTM STP 1325; 1999; pp 794–807. Was, G. S. Fundamentals of Radiation Materials Science; Spinger: Berlin, 2007. Garner, F. A. J. Nucl. Mater. 1983, 117, 177–197. Lee, E. H.; Mansur, L. K. J. Nucl. Mater. 1979, 85–86, 577–581. Garner, F. A.; Thomas, L. E. In Effects of Radiation on Substructure and Mechanical Properties of Metals and Alloys; ASTM STP 529; 1973; pp 303–325. Bullough, R.; Haynes, M. R. J. Nucl. Mater. 1977, 68, 286–293. Garner, F. A.; Guthrie, G. L. In Radiation Effects and Tritium Technology for Fusion Reactors; CONF-750989; 1976; Vol. 1; pp 491–518. Whitley, J. B.; Kulcinsky, G. L.; Wilkes, P.; Billen, J. J. Nucl. Mater. 1979, 85–86, 701–706. Garner, F. A.; Guthrie, G. L.; Gilbert, E. R. In Radiation Effects and Tritium Technology for Fusion Reactors; CONF-750989; 1976; Vol. I; pp 474–490. Wolfer, W. G.; Garner, F. A. J. Nucl. Mater. 1979, 85 and 86, 583–589. Garner, F. A.; Greenwood, L. R.; Roy, P. In Effects of Radiation on Materials: 18th International Symposium; ASTM STP 1325; 1999; pp 52–74. Norgett, J. J.; Robinson, J. T.; Torrens, I. M. Nucl. Engr. Design 1975, 33, p 50. Heinisch, H. L.; Hamilton, M. L.; Sommer, W. F.; Ferguson, P. J. Nucl. Mater. 1992, 191–194, 1177, as modified by Greenwood, L. R. J. Nucl. Mater. 1994, 216, 29–44. Garner, F. A.; Greenwood, L. R. Mater. Trans.; Jpn. Inst. Met. 1993, 34(11), 985–998. Garner, F. A.; Greenwood, L. R. Radiat. Eff. Def. Solids 1998, 144, 251–283. Greenwood, L. R.; Garner, F. A. J. Nucl. Mater. 2004, 329–333, 1147–1150. Edwards, D. J.; Hamilton, M. L.; Garner, F. A.; Samal, P.; Troxell, J. In Effects of Radiation on Materials: 18th International Symposium; ASTM STP 1325; 1999; pp 973–990. Edwards, D. J.; Garner, F. A.; Greenwood, L. R. J. Nucl. Mater. 1994, 212–215, 404–409. Greenwood, L. R.; Garner, F. A. J. Nucl. Mater. 1994, 212–215, 635–639. Garner, F. A.; Toloczko, M. B.; Greenwood, L. R.; Eiholzer, C. R.; Paxton, M. M.; Puigh, R. J. J. Nucl. Mater. 2000, 283–287, 380–385. Bates, J. F.; Garner, F. A.; Mann, F. M. J. Nucl. Mater. 1981, 103 and 104, 999–1003. Scott, P. J. Nucl. Mater. 1994, 211, p 101. Chung, H.; Garner, F. A. In Effects of Radiation on Materials: 18th International Symposium; ASTM STP 1325; 1999; pp 647–658. Garner, F. A.; Greenwood, L. R.; Mizia, R. E.; Tyler, C. R. Assessment of XM-19 as a Substitute for AISI 348 in ATR Service, INL/EXT-07–13530, Idaho National Laboratory, November 2007.

92 32. 33.

34.

35.

36. 37. 38. 39. 40. 41. 42. 43. 44. 45.

46. 47. 48.

49.

50. 51. 52. 53. 54. 55.

Radiation Damage in Austenitic Steels Mansur, L. K.; Lee, E. H.; Maziasz, P. J.; Rowcliffe, A. F. J. Nucl. Mater. 1986, 141–143, 633–646. Garner, F. A.; Greenwood, L. R.; Harrod, D. L. In Proceedings of the 6th International Symposium on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors; 1993; pp 783–790. 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. Griffiths, M.; Butcher, F. J.; Ariani, I.; Douglas, S.; Garner, F. A.; Greenwood, L. In Proceedings of 8th International Conference on CANDU Maintenance, Toronto, Ontario, CN, 2008. Mansur, L. K.; Grossbeck, M. L. J. Nucl. Mater. 1988, 155–157, 130–147. De Raedt, C. H. In Proceeding of the Conference on Fast, Thermal and Fusion Reactor Experiments; ANS; 1982; p 226. Greenwood, L. R. J. Nucl. Mater. 1983, 116, 137–142. 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. Garner, F. A.; Greenwood, L. R.; Oliver, B. M. In Effects of Radiation on Materials: 18th International Symposium; ASTM STP 1325; 1999; pp 794–807. Yamada, H.; Poeppel, R. B.; Sevy, R. H. Neutron-Induced Helium Implantation in GCFR Cladding; ANL-80–72, Argonne National Laboratory; 1980. Garner, F. A.; Hunter, C. W.; Johnson, G. D.; Lippincott, E. P.; Schiffgens, J. O. Nucl. Tech. 1982, 58, p 203. Hunter, C. W.; Duncan, G. D. In Proceedings International Conference on Fast Breeder Fuel Performance; Monterey CA; 1979; p 478. Duncan, D. R.; Panayotou, N. R.; Wood, W. L. ANS Trans. 1981, 38, 265. Hamilton, M. L.; Johnson, G. D.; Hunter, C. W.; Duncan, D. R. In Proceedings Conference on Dimensional Stability and Mechanical Behavior of Irradiated Metals and Alloys; British Nuclear Energy Society: London, 1981; Vol. 1; pp 211–214. Russell, K. C. Prog. Mater. Sci. 1984, 28, 229–434. Nolfi, F. W., Jr. Phase Transformations During Irradiation; Applied Science: London, 1983. Lee, E. H.; Maziasz, P. J.; Rowcliffe, A. F. In Proceedings Symposium on Phase Stability During Irradiation; The Metallurgical Society of AIME: Warrendale, 1981; pp 191–218. Yang, W. J. S.; Brager, H. R.; Garner, F. A. In Proceedings Symposium on Phase Stability During Irradiation; The Metallurgical Society of AIME: Warrendale, 1981; pp 257–269. Williams, T. M. In Effects of Radiation on Materials: 11th International Symposium; ASTM STP 782; Philadelphia, 1982; pp 166–185. Yang, W. J. S. In Radiation-Induced Changes in Microstructure: 13th International Symposium; STP 955; 1987; pp 628–646. Okamoto, P. R.; Rehn, L. E. J. Nucl. Mater. 1979, 83, 2–23. Marwick, A. D.; Piller, R. C.; Sivell, P. M. J. Nucl. Mater. 1979, 83, 35–41. Rothman, S. J.; Nowicki, L. J.; Murch, G. E. J. Phys. F. Met. Phys. 1980, 10, 383–398. Esmailzadeh, B.; Kumar, A. S. In Effects of Radiation on Materials: 12th International Symposium; ASTM STP 870; 1985; pp 468–480.

56. 57. 58. 59. 60. 61. 62. 63. 64. 65.

66. 67. 68. 69.

70. 71. 72. 73.

74.

75. 76. 77.

78. 79. 80. 81.

Esmailzadeh, B.; Kumar, A. S.; Garner, F. A. J. Nucl. Mater. 1985, 133–134, 590–594. Garner, F. A.; Kumar, A. S. In Radiation-Induced Changes in Microstructure: 13th International Symposium, Part 1; ASTM STP 955; 1987; pp 289–314. Maziasz, P. J.; Braski, D. N. J. Nucl. Mater. 1984, 122–123, 311–316. Maziasz, P. J. J. Nucl. Mater. 1984, 122–123, 472–486. Mansur, L. K.; Lee, E. H.; Maziasz, P. J.; Rowcliffe, A. F. J. Nucl. Mater. 1986, 141–143, 633–646. Garner, F. A.; Wolfer, W. G. J. Nucl. Mater. 1981, 102, 143–144. Lee, E. H.; Mansur, L. K. J. Nucl. Mater. 1986, 141–143, 695–702. Lee, E. H.; Mansur, L. K.; Rowcliffe, A. F. J. Nucl. Mater. 1984, 122–123, 299–304. Thomas, L. E. In Proceedings of 40th Annual Meeting Electron Microscopy Society of America; Electron Microscopy Society of America, 597; 1982. Itoh, M.; Onose, S.; Yuhara, S. In Proceedings of 13th International Symposium on Radiation-Induced Changes in Microstructure, Part 1; ASTM STP 955; 1987; pp 114–126. Yang, W. J. S.; Makenas, B. J. In Effects of Radiation on Materials: 12 International Symposium; ASTM STP 870; 1985; pp 127–138. Yang, W. J. S. J. Nucl. Mater. 1982, 108–109, 339–346. 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. 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. Garner, F. A.; Toloczko, M. B. Radiat. Eff. Def. Solids 1999, 148, 479–514. Garner, F. A. J. Nucl. Mater. 1993, 205, 98–117. Garner, F. A.; Toloczko, M. B. J. Nucl. Mater. 1993, 206, 230–248. Garner, F. A.; Oliver, B. M.; Greenwood, L. R.; Edwards, D. J.; Bruemmer, S. M.; Grossbeck, M. L. In 10th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors; 2001. Greenwood, L. R.; Garner, F. A.; Oliver, B. M.; Grossbeck, M. L.; Wolfer, W. G. In Effects of Radiation on Materials; ASTM STP 1447; 2004; pp 529–539. J. ASTM Int. 2004, 1(4), Paper ID JAI11365. Garner, F. A.; Simonen, E. P.; Oliver, B. M.; Greenwood, L. R.; Grossbeck, M. L.; Wolfer, W. G.; Scott, P. M. J. Nucl. Mater. 2006, 356, 122–135. Edwards, D. J.; Garner, F. A.; Bruemmer, S. M.; Efsing, P. G. J. Nucl. Mater. 2009, 384, 249–255. Conermann, J.; Shogan, R.; Fujimoto, K.; Yonezawa, T.; Yamaguchi, Y. In Proceedings of 12th International Conference on Environmental Degradation of Materials in Nuclear Power System – Water Reactors; 2005; p 277. Garner, F. A.; Hamilton, M. L.; Panayotou, N. F.; Johnson, G. D. J. Nucl. Mater. 1981, 103 and 104, 803–808. Irvin, J. E.; Bement, A. L. In The Effects of Radiation on Structural Materials; ASTM STP 426; 1967; pp 278–327. Irvin, J. E.; Bement, A. L.; Hoagland, R. G. In Flow and Fracture of Metals and Alloys in Nuclear Environments; ASTM STP 380; 1965; 236–250. Pawel, J. P.; Ioka, I.; Rowcliffe, A. F.; Grossbeck, M. L.; Jitsukawa, S. In Effects of Radiation on Materials: 18th


82. 83. 84.

85. 86. 87.

88. 89. 90. 91.

92.

93. 94. 95.

96. 97. 98. 99. 100. 101. 102.

103.

104. 105.

International Symposium; ASTM STP 1325; 1999; pp 671–688. Ehrlich, K. J. Nucl. Mater. 1985, 133–134, 119–126. Elen, J. D.; Fenici, P. J. Nucl. Mater. 1992, 191–194, 766–770. 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. Hamilton, M. L.; Okada, A.; Garner, F. A. J. Nucl. Mater. 1991, 179–181, 558–562. Garner, F. A.; Sekimura, N.; Grossbeck, M. L.; Ermi, A. M.; Newkirk, J. W.; Watanabe, H.; Kiritani, M. J. Nucl. Mater. 1993, 205, 206–218. Chatani, K.; Nishimura, S.; Kodama, M. Flux Effect on Mechanical Properties of Neutron-irradiated Type 304 Stainless Steel, Proceedings ICG-EAC meeting, Kyonju, Korea, April 2001. Neustroev, V. S.; Garner, F. A. J. Nucl. Mater. 2009, 386–388, 157–160. Neustroev, V. S.; Garner, F. A. J. Nucl. Mater. 2008, 378, 327–332. Marlowe, M.; Appleby, W. K. Trans. ANS 1973, 16, 95–96. Trantow, R. L. Ultrasonic Measurement of Elastic Properties in Irradiated 304 Stainless Steel, Hanford Engineering Development Laboratory Report HEDL-TME 73–92, 1973. Kozlov, A. V.; Shcherbakov, E. N.; Averin, S. A.; Garner, F. A. In Effects of Radiation on Materials: 22nd International Symposium; ASTM STP 1447; 2004; pp 66–77. Balachov, I. I.; Shcherbakov, E. N.; Kozlov, A. V.; Portnykh, I. A.; Garner, F. A. J. Nucl. Mater. 2004, 329–333, 617–620. Anderson, K. R.; Garner, F. A.; Stubbins, J. F. In Effects of Radiation on Materials: 15th International Symposium; ASTM STP 1125; Philadelphia, 1992; pp 835–845. 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. Neustroev, V. S.; Shamardin, V. K. Atomnaya Energiya 1990, 71(4), 345–348, in Russian. Neustroev, V. S.; Shamardin, V. K. Phys. Met. Metallogr. 1997, 83(5), 555–560. Porter, D. L. J. Nucl. Mater. 1979, 79, 406–411. Porter, D. L.; Garner, F. A.; Bond, G. M. In Effects of Radiation on Materials: 19th International Symposium; ASTM STP 1366; 2000; pp 884–893. Gusev, M. N.; Maksimkin, P.; Osipov, I. S.; Garner, F. A. J. Nucl. Mater. 2009, 386–388, 273–276. Gusev, M. N.; Maksimkin, O. P.; Garner, F. A. J. Nucl. Mater. 2010, 403, 121–125. Gusev, M. N.; Maksimkin, O. P.; Osipov, I. S.; Silniagina, N. S.; Garner, F. A. In Effects of Radiation on Nuclear Materials and the Nuclear Fuel Cycle; ASTM STP 1513; 2010; pp 210–219. J. ASTM Int. 6(7), Paper ID JAI102062. 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. 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.

106. 107. 108. 109.

110. 111. 112.

113. 114. 115. 116. 117. 118. 119. 120. 121. 122.

123. 124. 125.

126. 127. 128.

129. 130. 131.

93

Garner, F. A.; Toloczko, M. B.; Greenwood, L. R.; Eiholzer, C. R.; Paxton, M. M.; Puigh, R. J. J. Nucl. Mater. 2000, 283–287, 380–385. Ohnuki, S.; Takahashi, H.; Garner, F. A.; Pawel, J. E.; Shiba, K.; Hishinuma, A. J. Nucl. Mater. 1996, 233–237, 411–415. Weiss, B.; Stickler, R. Phase Instabilities During High Temperature Exposure of 316 Austenitic Stainless Steel, Westinghouse R&D Report 70–1D4-STABL-P1, 1970. Garner, F. A.; Cummings, W. V.; Bates, J. F.; Gilbert, E. R. Densification-Induced Strains in 20% Cold-Worked 316 Stainless Steel During Neutron Irradiation, Hanford Engineering Development Laboratory, HEDL-TME-78–9, June 1978. Puigh, R. J.; Lovell, A. J.; Garner, F. A. J. Nucl. Mater. 1984, 122 and 123, 242–245. Hales, J. W. Trans. ANS 1978, 28, 153–155. Garner, F. A.; McCarthy, J. M. In Proceedings of TMS Symposium on Physical Metallurgy of Controlled Expansion Invar-Type Alloys; Las Vegas, NV; 1989; pp 187–206. Garner, F. A.; McCarthy, J. M.; Russell, K. C.; Hoyt, J. J. J. Nucl. Mater. 1993, 205, 411–425. Brager, H. R.; Garner, F. A. J. Nucl. Mater. 1983, 117, 159–176. Straalsund, J. L.; Powell, R. W.; Chin, B. A. J. Nucl. Mater. 1982, 108–109, 299–305. Garner, F. A. In Materials Science and Technology: A Comprehensive Treatment; VCH: New York, 2011; Vol. 2(Revised); Chapter 10, currently in preparation. 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. Garner, F. A.; Black, C. A. In Effects of Radiation on Materials: 19th International Symposium; ASTM STP 1366; 2000; pp 767–777. Kumar, A.; Garner, F. A. J. Nucl. Mater. 1983, 117, 234–238. Garner, F. A. J. Nucl. Mater. 1984, 122–123, 459–471. Garner, F. A.; Brager, H. R.; Gelles, D. S.; McCarthy, J. M. J. Nucl. Mater. 1987, 148, 294–301. Garner, F. A.; Brager, H. L. In Effects of Radiation on Materials: 13th International Symposium (Part 1) Radiation-Induced Changes in Microstructure; ASTM STP 955; 1987; pp 195–206. Garner, F. A.; Gelles, D. S. J. Nucl. Mater. 1988, 159, 286–309. Hall, M. M., Jr.; Flinn, J. E. J. Nucl. Mater. 2010, 396, 119–129. Neustroev, V. S.; Makarov, E. I.; Belozerov, S. V.; Ostrovsky, Z. E. In Proceedings of Fontevraud 7, Contribution of Materials Investigations to Improve the Safety and Performance of LWRs; 26–30 September, Avignon, France, 2010. Gilbert, E. R.; Garner, F. A. J. Nucl. Mater. 2007, 367–370, 954–959. Garner, F. A.; Flinn, J. E.; Hall, M. M. J. Nucl. Mater. 2009, 386–388, 249–253. Coghlan, W. A.; Garner, F. A. In Effects of Radiation on Materials: 13 International Symposium (Part 1) RadiationInduced Changes in Microstructure; ASTM STP 955; 1987; pp 315–327. Hoyt, J. J.; Garner, F. A. J. Nucl. Mater. 1991, 179–181, 1096–1099. Porollo, S. I.; Shulepin, S. V.; Konobeev, Yu. V.; Garner, F. A. J. Nucl. Mater. 2008, 378, 17–24. Garner, F. A.; Porollo, S. I.; Konobeev, Yu. V.; Makenas, B. J.; Chastain, S. A. Trans. ANS 2010, 836–837.

94


132. Garner, F. A.; Black, C. A.; Edwards, D. J. J. Nucl. Mater. 1997, 245, 124–130. 133. Stubbins, J. F.; Garner, F. A. J. Nucl. Mater. 1992, 191–194, 1295–1299. 134. Dvoriashin, A. M.; Porollo, S. I.; Konobeev, Yu. V.; Garner, F. A. J. Nucl. Mater. 2000, 283–287, 157–160. 135. Garner, F. A.; Bates, J. F.; Mitchell, M. A. J. Nucl. Mater. 1992, 189, 201–209. 136. Garner, F. A.; Gilbert, E. R.; Gelles, D. S.; Foster, J. P. In Proceedings, ASTM 10th International Symposium on Effects of Radiation on Materials; ASTM STP 725; 1980; pp 698–712. 137. Seran, J. L.; Touron, H.; Maillard, A.; Dubuisson, P.; Hugot, J. P.; Le Boulbin, E.; Blanchard, P.; Pelletier, M. In Effects of Radiation on Materials: 14th International Symposium; ASTM STP 1046; 1990; Vol. II; pp 739–752. 138. Akasaka, N.; Yamagata, I.; Ukai, S. J. Nucl. Mater. 2000, 283–287, 169–173. 139. Garner, F. A.; Edwards, D. J.; Bruemmer, S. M.; Porollo, S. I.; Konobeev, Yu. V.; Neustroev, V. S.; Shamardin, V. K.; Kozlov, A. V. In Proceedings of Fontevraud 5, Contribution of Materials Investigation to the Resolution of Problems Encountered in Pressurized Water Reactors; September 23–27, 2002. 140. 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. 141. Garner, F. A.; Porter, D. L. In Proceedings International Conference on Dimensional Stability and Mechanical Behavior of Irradiated Metals and Alloys; April 11–13; 1983; Brighton, England; Vol. 11; pp 41–44. 142. 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. 143. Okita, T.; Sekimura, N.; Sato, T.; Garner, F. A.; Greenwood, L. R. J. Nucl. Mater. 2002, 307–31, 322–326. 144. Okita, T.; Sekimura, N.; Garner, F. A.; Greenwood, L. R.; Wolfer, W. G.; Isobe, Y. In Proceedings of 10th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors; 2001. 145. Okita, T.; Wolfer, W. G.; Sato, T.; Sekimura, N.; Garner, F. A. In Proceedings of 11th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors; 2003; pp 657–663. 146. Okita, T.; Sato, T.; Sekimura, N.; Iwai, T.; Garner, F. A. J. Nucl. Mater. 2007, 367–370, 930–934. 147. 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. 148. Etienne, A.; Radiquet, B.; Pareige, P.; Massoud, J. P.; Pokor, C. J. Nucl. Mater. 2008, 382, 64–69.S. 149. Byrne, T.; Wilson, I.; Shogan, R. In Proceedings of Fontevraud 5, Contribution of Materials Investigation to the Resolution of Problems Encountered in Pressurized Water Reactors; September 23–27, 2002. 150. Byrne, S.; Garner, F. A.; Fyfitch, S.; Wilson, I. A. In Proceedings of 10th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors; 2001. 151. Foster, J. P.; Porter, D. L.; Harrod, D. L.; Mager, T. R.; Burke, M. G. J. Nucl. Mater. 1995, 224, 207–215. 152. Fujii, K.; Fukuya, K.; Furutani, G.; Torimaru, T.; Kohyama, A.; Katoh, Y. In Proceedings of 10th

International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors; 2001. 153. Fukuya, K.; Fujii, K.; Nishioka, H.; Kitsunai, Y. J. Nucl. Sci. Technol. 2006, 43(2), 159–173. 154. Neustroev, V. S.; Dvoretzky, V. G.; Ostrovsky, Z. E.; Shamardin, V. K.; Shimansky, G. A. In Effects of Radiation on Materials: 21st International Symposium; ASTM STP 1447; 2004; pp 32–45. 155. Garner, F. A.; Porollo, S. I.; Konobeev, Yu. V.; Neustroev, V. S.; Maksimkin, O. P. In Proceedings of Fontevraud-6 Symposium on Contribution of Materials Investigations to Improve the Safety and Performance of LWRs; 2006; pp 637–648. 156. 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. 157. Neustroev, V. S.; Shamardin, V. K.; Ostrovsky, Z. E.; Pecherin, A. M.; Garner, F. A. In Effects of Radiation on Materials: 19th International Symposium; ASTM STP 1366; 2000; pp 792–800. 158. Porollo, S. I.; Konobeev, Yu. V.; Dvoriashin, A. M.; Vorobjev, A. N.; Krigan, V. M.; Garner, F. A. J. Nucl. Mater. 2002, 307–311, 339–342. 159. Porollo, S. I.; Konobeev, Yu. V.; Dvoraishin, A. M.; Krigan, V. M.; Garner, F. A. In 10th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors; 2001. 160. Maksimkin, O. P.; Tsai, K. V.; Turubarova, L. G.; Doronina, T.; Garner, F. A. J. Nucl. Mater. 2004, 329–333, 625–629. 161. Maksimkin, O. P.; Tsai, K. V.; Turubarova, L. G.; Doronina, T. A.; Garner, F. A. J. Nucl. Mater. 2007, 367–370, 990–994. 162. 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. 163. Foster, J. P.; Gilbert, E. R.; Bunde, K.; Porter, D. L. J. Nucl. Mater. 1998, 252, 89–97. 164. Gilbert, E. R.; Straalsund, J. L.; Wire, G. L. J. Nucl. Mater. 1977, 65, 277–294. 165. Gilbert, E. R.; Bates, J. F. J. Nucl. Mater. 1977, 65, 204–209. 166. Matthews, J. R.; Finnis, M. W. J. Nucl. Mater. 1988, 159, 257–285. 167. Woo, C. H.; Garner, F. A. J. Nucl. Mater. 1999, 271–272, 78–83. 168. Grossbeck, M. L.; Mansur, L. K. J. Nucl. Mater. 1991, 179–181, 130–134. 169. Grossbeck, M. L.; Gibson, L. T.; Jitsukawa, S.; Mansur, L. K.; Turner, L. J. In Effects of Radiation on Materials: 18th International Symposium; ASTM STP 1325; 1999; pp 725–741. 170. Gilbert, E. R. Reactor Technol. 1971, 14, 258–285. 171. Causey, A. R.; Carpenter, C. K. C.; MacEwen, S. R. J. Nucl. Mater. 1980, 90, 216–223. 172. 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. 173. Garner, F. A.; Puigh, R. J. J. Nucl. Mater. 1991, 179–181, 577–580. 174. Hilbert, R. F.; Kangilaski, M.; Appleby, W. K.; Craig, C. N. In Proceedings of the ANS Topical Meeting on Irradiation Experiments in Fast Reactors; 1973; pp 472–483.

Radiation Damage in Austenitic Steels 175. Ukai, S.; Ohtsuka, S. J. Nucl. Sci. Technol. 2007, 44(5), 743–757. 176. Woo, C. H.; Garner, F. A.; Holt, R. A. In Proceedings of 16th ASTM International Symposium on Effects of Radiation on Materials; ASTM STP 1175; 1992; pp 27–37. 177. Semenov, A. A.; Woo, C. H. J. Nucl. Mater. 2009, 393, 409–417. 178. Lewthwaite, G. W.; Mosedale, D. J. Nucl. Mater. 1980, 90, 205–215. 179. Garner, F. A.; Toloczko, M. B. J. Nucl. Mater. 1997, 251, 252–261. 180. Simons, R. L. J. Nucl. Mater. 1986, 141–143, 665–671. 181. Mansur, L. K.; Farrell, K. J. Nucl. Mater. 1990, 170, 236–245. 182. Foster, J. P.; Boltax, A. J. Nucl. Mater. 1980, 89, 331–337. 183. Foster, J. P.; Mildrum, C. M. J. Nucl. Mater. 1988, 151, 135–139. 184. Foster, J. P.; Karlsen, T. In 14th International Conference on Environmental Degradation of Materials in Nuclear Power Systems; 2009; pp 1355–1370. 185. Hall, M. M. J. Nucl. Mater. 2010, 396, 112–118. 186. Hall, M. M.; Flinn, J. E. J. Nucl. Mater. 2010, 396, 119–129. 187. Neustroev, V. S.; Makarov, E. I.; Belozerov, S. V.; Ostrovsky, Z. E. In Proceedings of Fontevraud 7, Contribution of Materials Investigations to Improve the Safety and Performance of LWRs, Avignon, France, 26–30 September, 2010. 188. Garzarolli, F.; Dewes, P.; Trapp-Protsching, S.; Nelson, J. L. In Proceedings of 9th International

189. 190.

191. 192. 193. 194.

195. 196. 197. 198. 199.

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Symposium on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors; TMS; 1999; pp 1027–1034. Obata, M.; Ishiyama, Y.; Nakata, K.; Sakamoto, H.; Anzai, H.; Asano, K. J. ASTM Int. 2006, 3, 15–31. 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. 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. Ukai, S.; Mizuta, S.; Kaito, T.; Okada, H. J. Nucl. Mater. 2000, 278, 320–327. Hamilton, M. L.; Johnson, G. D.; Puigh, R. J.; Garner, F. A.; Maziasz, P. J.; Yang, W. J. S.; Abraham, N. In Proceedings ASTM Symposium on Residual Elements in Steel; ASTM STP 1042; 1989; pp 124–149. Tavassoli, A. A. J. Nucl. Mater. 1998, 258–263, 85–96. Boutard, J. L. J. Nucl. Mater. 1991, 179–181, 1179–1184. Vandermueulen, W.; Hendrix, W.; Massault, V.; Van de Velde, J. J. Nucl. Mater. 1988, 155–157, 953–956. Jitsukawa, S.; Ioka, I.; Hishinuma, A. J. Nucl. Mater. 1999, 271–272, 167–172. Ioka, I.; Yonekawa, M.; Miwa, Y.; Mimura, H.; Tsuji, H.; Hoshiya, T. J. Nucl. Mater. 2000, 283–287, 440–445.

4.03 Ferritic Steels and Advanced Ferritic–Martensitic Steels B. Raj and M. Vijayalakshmi Indira Gandhi Centre for Atomic Research, Kalpakkam, India


4.03.1 4.03.2 4.03.3 4.03.4 4.03.4.1 4.03.4.2 4.03.4.3 4.03.4.4 4.03.4.5 4.03.4.5.1 4.03.5 4.03.6 4.03.7 References

Introduction Basic Metallurgy of Ferritic–Martensitic Steels Radiation Damage of Core Components in Fast Reactors Development of Ferritic Steels for Fast Reactor Core Influence of Composition and Microstructure on Properties of Ferritic Steels Void Swelling Resistance Irradiation Hardening in Ferritic Steels Irradiation Creep Resistance of Ferritic Steels Irradiation Embrittlement in Ferritic Steels GBE to reduce embrittlement in ferritic steels Development of Advanced ODS Ferritic Steels Ferritic Steels for Out-of-Core Applications: Improvements in Joining Summary

Abbreviations bcc CSL DBTT DICTRA dpa EBR EBSD fcc FFTF GBCD GBE HAADF HAZ HFIR ITER ODS steel PAGS PFR PWHT RIS SIPA SIPN TEM ▽DBTT

Body-centered cubic Coincident site lattice Ductile to brittle transition temperature Diffusion-controlled transformations Displacements per atom Experimental breeder reactor Electron back scattered diffraction Face-centered cubic Fast flux test facility Grain boundary character distribution Grain boundary engineering High angle annular dark field Heat-affected zone High flux isotope reactor International Thermonuclear Experimental Reactor Oxide dispersion strengthened steel Prior austenite grain size Power fast reactor Postweld heat treatment Radiation-induced segregation Stress-induced preferential absorption Stress-induced preferential nucleation Transmission electron microscopy Change in DBTT

97 98 101 102 103 105 106 108 110 112 114 116 119 119

4.03.1 Introduction The widespread acceptance of nuclear energy depends1 on the improved economics, better safety, sustainability, proliferation resistance, and waste management. Innovative technological solutions are being arrived at, in order to achieve the above goals. The anticipated sustainability, rapid growth rate, and economic 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 international collaborative projects (http://www.iaea.org/ INPRO/; http://www.gen4.org/) have, therefore, chosen fast reactors as one of the important constituents of the nuclear energy in the twenty-first century. The nuclear community has been constantly striving 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

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progress to achieve a target burn-up of 250 dpa, using advanced ferritic steels. The attempts by nuclear technologists to enhance the thermal efficiency have posed the challenge of improving the high temperature capability of ferritic steels. Additionally, there is an inherent disadvantage in ferritic steels, that is, their susceptibility to undergo embrittlement, 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 coolants 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 optimism with construction of the International Thermonuclear 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 reevaluate the materials technology of today for future applications. The genesis of the development of ferritic steels is, indeed, in the thermal power industry. The development 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 resistance. 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. The mechanism of irradiation hardening, embrittlement and methods to overcome the same are studied in detail. Of the different steels evaluated, 9–12% Cr ferritic–martensitic steels are the immediate future solution for fast reactor 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 973 K, by launching the next generation ferritic steels, which are currently under evaluation for nuclear applications, namely the oxide dispersion strengthened (ODS) ferritic steels (see Chapter 4.08, Oxide Dispersion Strengthened Steels). Conceptually, this series of steels combines the merits of swelling resistance 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 823 K. 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 irradiation, which will be discussed in this article. The present review begins with a brief introduction to the basic metallurgy of ferritic steels, summarizing 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 components in fast reactors. The irradiation response of ferritic steels with respect to swelling resistance, irradiation hardening, and irradiation creep are highlighted. The in-depth understanding of the damage mechanisms is explained. The main concerns of ferritic steels such as the inferior high temperature irradiation 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 circuits 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 predecessors, 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


Liquid

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523 K. 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 economically 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, vanadium, 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 austenite and ferrite stabilizing elements. A typical pseudobinary 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 temperature and the decomposition mode7–9 of high temperature austenite (Figure 1(b)) dictate the resulting microstructure at room temperature and hence, the type of steel. Accordingly, the 9CrMo family of steels can either be martensitic (9Cr–1Mo (EM10) or stabilized 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 1100 K, dictate the long term industrial performance of the steels. Evaluation of tensile and creep properties of Cr–Mo steels exposed

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Pearlite Ws Bs

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Ms Martensite

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Log {time}

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.

to elevated temperature for prolonged durations have been extensively studied.5,13,14 The following trends were established: The optimized initial alloy composition considered was 9Cr, W–2Mo ¼ 3, Si ¼ 0.5, with C, B, V, Nb, and Ta in small amounts. Higher chromium content has two effects: it increases the hardenability leading to the formation of martensite and also promotes the formation of d-ferrite thereby reducing the toughness. A reduction in the chromium

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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.

content lowers the oxidation resistance. If W þ Mo concentration is kept 0.1 MeV)

Radiation Effects in Nickel-Based Alloys

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STA DAA766 OA DAA766 STA Z260D STA Z184

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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.

precipitation of TiC and an overaged g0 structure). Swelling data derived from the density measurements 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 58 dpa at 513 C. Reproduced from Boothby, R. M. J. Nucl. Mater.1996, 230, 148–157.

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 cladding data shown in Figure 7. Void numbers are generally 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 irradiations at 600 C. Void concentrations for PE16 irradiated to fast neutron fluences (E > 0.1 MeV) of 9.4–12.3 1026 n m2 at 477–513 C in the UK-1 experiment were very similar to those determined by Sklad et al.50 for 4.0 1026 n m2 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 vacancies 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

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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 vacancies from vacancy loops which had formed during lower temperature excursions. In fact, the void concentration 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 reference wrapper material for PFR and as an alternative 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 cladding 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 70 MPa. 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 interaction 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 burnups than in PE16. The failures in PE16 cladding were

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regarded as benign and permitted continued operation, 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 mechanisms is beyond the scope of this chapter, which will instead concentrate on experimental data which enable comparisons to be made between nickelbased 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. Irradiation 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 pressurized tubes and helical springs, have been used to measure irradiation creep strains. The data are therefore generally converted to effective strain e values, using the Soderberg and effective stress s 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 components as follows: g ¼ At þ Bd t 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 5 107 dpa (N/2) s1 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 maximum observed creep rate was similar to that in austenitic steels and Harries59 represented the creep strain above a threshold dose of 8 dpa (N/2) by g ¼ 4:3 106 tðd 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 condition, 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 steadystate creep coefficients for STA and OA PE16 increased by factors of 2 at lower damage rates of 0.5–1.5 107 dpa (N/2) s1 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 (MPa1 dpa1) and creep strain rates (MPa1 s1) 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 and Mosedale63,64 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 irradiation 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


Creep coefficient, B (10–6 MPa–1 dpa–1)

6.0 M316 FV548 PE16 OA PE16 STA

5.0

4.0

3.0

2.0

1.0

0.0 0.0

1.0

2.0

3.0

4.0

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Creep strain rate (10–13 MPa–1 s–1)

9.0 8.0

M316 FV548 PE16 OA PE16 STA

7.0 6.0 5.0 4.0 3.0 2.0 1.0 0.0 0.0

1.0

2.0

Displacement rate

3.0

4.0

5.0

(10–7 dpa s–1)

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.

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 preferred absorption of interstitials at dislocations) controlled creep. It was suggested by Mosedale et al.62

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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 dependency 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 austenitic and ferritic steels, in pressurized tube experiments carried out in EBR-II at 540 C to fluences up to 4 1026 n m2 (E > 0.1 MeV). Diametral strains measured in pressurized tubes (with hoop stresses in the approximate range of 25–175 MPa) were corrected 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 precipitationhardened materials was also dependent on heat treatment, 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 B0 ft s þ 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 nonlinearity attributable to the effects of stress on swelling was observed in the solid solution alloys. An approximate value of B0 of 1.5 1028 MPa1 (n cm2)1, which is equivalent to 3 107 MPa1

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dpa1, 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 1.4 106 MPa1 dpa1 for ST PE16 to 105 MPa1 dpa1 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 precipitationhardened alloys, with the higher values being attributable 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 100 MPa for PE16 and 200 MPa for Inconel 706, were irradiated at 425, 540, and 590 C, both isothermally and with temperature steps. Diametral strains for isothermally irradiated PE16 increased with increasing fluence and temperature as expected. Following temperature changes from 540 to 590 C or 425 C, the creep rate for PE16 adjusted to the isothermal 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 tetragonal, Ni3Nb g00 phase and its effect on thermal creep resistance. Gilbert and Chin considered that the inreactor 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 subsequently 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 locations 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 150 MPa, 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 1.0–1.4 106 MPa1 dpa1 (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 material dependent (with values in the approximate range of 0.4–1.6 102 MPa1), though this variability could not be associated with any particular compositional factor. Similar results for two precipitation-hardened high-nickel alloys (with similar 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 materials 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 n m2 (E > 0.1 MeV) at hoop stresses of 0 and 167 MPa, revealed that the distribution of Frank dislocation loops was similar on all the four {111} planes in the unstressed sample but was anisotropic 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 relationship between the loop density and the normal


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 pressurized 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 n m2 at a hoop stress of 331 MPa. The results showed highly anisotropic distributions in the Burgers vectors of perfect dislocations in all the three heat-treated conditions, with dislocation densities of the various types differing by factors of up to 10–40 in each sample. The level of anisotropy produced 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 dislocations 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

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suggested that precipitates would evolve to an equilibrium size determined by competing processes affecting their growth, via the radiation enhanced and/or thermal diffusion of solutes, and their simultaneous dissolution due to damage arising in collision cascades. Two dissolution mechanisms were suggested: recoil dissolution due to the displacement of atoms from the precipitate into the matrix, and disordering 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 (dependent 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-ionirradiated 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 coarsening 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 n m2 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 irradiation-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 n m2, E > 0.1 MeV) 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 conditions. 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

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below 600 C with an activation energy that (in agreement with theoretical predictions for a process governed by point defect recombination) was approximately 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 ‘skeletal’ 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 nickelbased alloys were investigated in the 1970s and 1980s.

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 beneficial effect on (unirradiated) creep rupture life, as it segregates to grain boundaries and inhibits intergranular 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 twostage 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 irradiationinduced strengthening and grain boundary segregation and precipitation effects, have been implicated in the embrittlement of fast neutron irradiated nickel-based alloys.

4.04.5.1 Fast Neutron Irradiation Experiments

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.

Rowcliffe and Horak86 investigated the tensile properties 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 n m2 (E > 0.1 MeV). Irradiation temperatures (Ti) ranged from 450 to 735 C, with tensile tests being performed at a strain rate of 4 104 s1 at temperatures corresponding to Ti and to Ti þ 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 (>1000 MPa) yield stresses and ultimate tensile strengths (UTS) in


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 650 MPa in specimens irradiated at 625 C. The very high tensile

strengths exhibited at the lower irradiation temperatures were attributed to the instability of the (ordered body-centered tetragonal) g00 phase below 525 C and its consequent dissolution, leading to the reprecipitation of nickel and niobium as (ordered face-centered cubic) g0 on dislocation loops. At higher irradiation temperatures, both g0 and g00 were stable, but

1400

20 Yield at Ti + 110 °C Elong. at Ti + 110 °C

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Yield at Ti

18 16 14 12

800

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8 6

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Elong. at Ti

1000 Yield stress (MPa)

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4 200 2 0 400

450

500

550

600

650

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0 800

Temperature (°C) Figure 12 Yield stress and total elongation values at the irradiation temperature (Ti) and at Ti þ 110 C for Experimental Breeder Reactor-II-irradiated Inconel 706. Based on data from Rowcliffe, A. F.; Horak, J. A. Am. Nucl. Soc. Trans. 1981, 38, 266–267.

20

1400 Yield at Ti Elong. at Ti

1200

Yield at Ti + 110 °C Elong. at Ti + 110 °C

18

Yield stress (MPa)

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14 12

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16

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500

550 600 650 Temperature (°C)

700

750

0 800

Figure 13 Yield stress and total elongation values at the irradiation temperature (Ti) and at Ti þ 110 C for Experimental Breeder Reactor-II-irradiated Inconel 718. Based on data from Rowcliffe, A. F.; Horak, J. A. Am. Nucl. Soc. Trans.1981, 38, 266–267.

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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 unirradiated material. Irradiation embrittlement was generally more severe in tests at Ti þ 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 (hexagonal 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 irradiation of Inconel 718, resulting in yield strengths in excess of 1000 MPa at irradiation temperatures up to 560 C and above 800 MPa 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 Ti þ 110 C. In contrast to Inconel 706, failures in irradiated Inconel 718 were reported to be predominantly 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.

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 n m2 (E > 0.1 MeV), 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 Ti þ 110 C (to simulate reactor transients), at a strain rate of 4 104 s1, and with a small number of tests at 4 103 s1. Irradiated 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 temperature of 735 C exhibited some softening), and retained good levels of ductility with total elongation values above 10%. Yield stress and total elongation data for PE16 at higher test temperatures are shown in Figure 14 for specimens irradiated to a fast neutron fluence of 4.3 1026 n m2 (enabling direct comparison with the data for the similarly irradiated Inconel alloys shown in Figures 12 and 13). Specimens tested at the irradiation temperature again showed strengthening at temperatures in the range of 450–625 C and softening at 735 C, with good 20

1000 Yield at Ti + 110 °C Elong. at Ti + 110 °C

800

Elong. at Ti

18 16

Yield stress (MPa)

14 12

600

10 8

400

6


Yield at Ti

4

200

2 0 400

450

500

550 600 650 Temperature (°C)

700

750

0 800

Figure 14 Yield stress and total elongation values at the irradiation temperature (Ti) and at Ti þ 110 C for Experimental Breeder Reactor-II-irradiated Nimonic PE16. Based on data from Bajaj, R.; Shogan, R. P.; DeFlitch, C.; et al. In Effects of Radiation on Materials: 10th Conference; Kramer, D., Brager, H. S., Perrrin, J. S., Eds.; American Society for Testing and Materials: Philadelphia, PA, 1981; pp 326–351, ASTM STP 725. Reprinted, with permission, from ASTM STP725-Effects of Radiation on Materials, copyright ASTM International, 100 Barr Harbor Drive, West Conshohocken, PA 19428.


ductility at 450 C but total elongations reduced to 3% at 560–625 C. Tests at Ti þ 110 C showed further increases in tensile strength (consistent with 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 n m2 and to zero (i.e., failure before yield) in higher dose samples (7.1 1026 n m2). 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 samples 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 by matrix hardening with little or no change in the grain boundary 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 n m2 (E > 0.1 MeV) at 500 C and tested at strain rates from 3 105 to 3 103 s1. 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 low ductility failures were associated with an increased tendency toward intergranular fracture, and additional tests, in which samples irradiated to 4 1026 n m2 at 500 C were fractured in situ in an Auger 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

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sections of AGR (advanced gas-cooled reactor) tie bars irradiated at 512 C and above. AGR tie bars, which are approximately 10 m 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 4 m position. Stress-rupture testing at 600 C at an applied stress of 500 MPa 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 concentrations 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 contributed to the minimum in stress-rupture life, which was thought to result from the weakening of the boundaries relative to the matrix. Grain boundary helium bubbles were 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 2.5 106 s1 at temperatures approximating those of irradiation, revealed severe embrittlement with minimum elongations of 0.2% at 550 C; TEM examination 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 microstructure of developmental g0 (D21, D25, and D66) and g0 /g00 (D68) strengthened alloys were discussed

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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 n m2 (E > 0.1 MeV) in EBR-II, and were tested at Ti, Ti þ 110 C and 232 C. Severe irradiation embrittlement was evident in the ST alloys and STA D25, particularly in tests at Ti þ 110 C. Zero ductility was recorded in the lower-Ni alloy D21 (25Ni–8Cr) irradiated and tested at 550 and 600 C. Severe ductility 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 transgranular 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 compressed 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 n m2 (E > 0.1 MeV) and irradiation temperatures were in the range 460–616 C. All the materials exhibited low ductility failures at high test temperatures, particularly in tests at about Ti þ 110 C where ductilities were generally below 0.1%, though Vaidyanathan et al. found that postirradiation heat treatments (typically 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 irradiationinduced 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 precipitation 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 boundaries 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 importance. 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 failures 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 n m2 (E > 0.1 MeV). Microstructural examinations revealed a brittle intergranular fracture, with failure being attributed to a combination of matrix hardening due to g0 precipitation and grain boundary weakening due to the formation of interconnected Ni3(Ti,Nb) Z phase particles. In contrast 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 solutionannealed 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


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 cladding sections were carried out at a strain rate of 3 104 s1. Tensile tests performed at ambient temperature showed high UTS (>1000 MPa) along the full length of the pins with peak values of 1500 MPa in sections irradiated near 500 C; ductility values (uniform elongations only were given) remained low ( 1 MeV) of 1.7 1024 n m2 and a thermal fluence of 5.9 1024 n m2. 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

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out at the implantation/irradiation temperature at a strain rate of 5 104 s1. The results showed similar trends in helium-implanted and neutronirradiated 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 compared to those injected with helium. Overall, the g0 hardened alloy 7817 exhibited relatively high tensile strength (typically >700 MPa) but low ductility following helium implantation or neutron irradiation (with total elongation values generally I MeV 200 100 0.025 0.025 + 0.024% Sn

100 50

0.015

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

Charpy shift ΔT41J (⬚C)

200

0 150

A

B

C

D A Melt cast

B

C

D

0

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.

100

50

0 0

(b)

(%P) 0.003

19 19 19 19 1⫻1019 2⫻10 3⫻10 4⫻10 5⫻1019 6⫻10 7⫻1019

Fluence n cm−2 (E > 1 MeV)

Figure 3 Charpy shift (DT41 J ( C)) for (a) a US weld and a US forging containing 0.25 and 0.06 wt% Cu, respectively, and (b) US welds and a US plate containing 0.2 wt% Cu and varying levels of Ni.

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 influence 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. 1 MeV) 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.

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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 weldments 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 cm2. The authors considered that the microvoid numbers increase with dose up to 5 1018 n cm2, and then either remain constant or decrease. Analyzing positron annihilation data from annealing studies of neutronirradiated A533B plate, A508–3 forging, and welds, Carter et al.54 considered that increasing the dose from 1 1018 to 20 1018 n cm2 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 cm2s1 increased either the number density 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 components; 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 cm2 s1, E > 1 MeV).85

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Radiation Damage of Reactor Pressure Vessel Steels

Soneda65 modeled the effects of dose, dose rate, and irradiation temperature on the defect accumulation in bcc-Fe using the kinetic Monte Carlo (KMC) method.65,86 Jones and Williams83 proposed a model that describes the irradiation temperature dependence of the embrittlement of low Cu materials, DT ¼ a FT (’t)1/2, where DT, a, and ’t are the transition temperature shift (TTS), constant coefficient, and dose, respectively, and FT ¼ 1.869– 4.57 103T ( 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 conditions), 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 microstructural 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–0.05 wt% Cu alloys irradiated at 8.3 1018 n cm2, E > 1 MeVat 300 C (the irradiation time was 144 h). As a result of CDB and positron lifetime measurements on irradiated and annealed samples, the authors reported the formation of microvoids (10 vacancies), dislocation loops, and Cu-mono-vacancy-Cu complexes. 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 electronirradiated 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 (0.05 wt%) high nickel (1.26 and 1.78 wt% Ni) VVER-1000 forging and weld materials that were neutron irradiated to a total fluence of 1.38 1023 n m2 (E > 1 MeV). 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 1023 m3, while the number density of clusters in the forging was estimated to be slightly lower at 1 1023 m3. These ultrafine clusters may, or may not, be associated with vacancies. The observations 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 cm2. 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 unrelated to the change in DHv, and thus, vacancy-type defects are not solely responsible for the observed irradiation-induced hardening.


4.05.4.6 Effect of Radiation Damage on Hardening

350 WV LC LD

300 250 Dsy,meas (MPa)

The small defects formed in irradiated steels and model alloys can act as barriers to dislocation movement 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 theory, 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 dislocation 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 hardening 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 preexisting 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 precipitate. The MS simulations led to a dependence of hardening on precipitate size which differed from that predicted by the Russell–Brown model. However, 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

169

200 150 100 50 0

0

50

100

150 200 250 Dsy,pred (MPa)

300

350

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 0.4 wt% 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.

mechanical property data (particularly hardening or yield stress increase) where the MD has been subtracted 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 segregation of certain impurities to grain boundaries could cause nonhardening embrittlement. This phenomenon 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 impurity elements under irradiation.39,101–105 (Extensive experimental programs on long-term aging have permitted 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


segregation in these alloys and steels indicates that segregation is generally thermodynamically controlled, and defect gradients have no effect.) The segregation of P and C to grain boundaries in irradiated materials has received greatest attention.41,100 Overall P segregation increases with irradiation dose in all of the model alloys and steel types examined. The rate of P segregation under irradiation 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 generally 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 relatively high free C than merely to trapping of C at matrix defects. Quantifying the data has been attempted in several 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. Consistency between the models and data need not 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 elsewhere, it appears unlikely that nonhardening embrittlement will influence RPV condition during normal operation for homogeneous MnMoNi steels (i.e., A508 Class 3, A533B, 22NiMoCr37) of = or ¼ Dcopper þ Dmatrix ½4 > ; Ds 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 hardening arising from the production of point defect clusters by neutron irradiation. A further simplification was made in developing a DDR that could be applied to operational Magnox reactors. Namely, under the conditions 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. Further, the hardening from Cu clusters could be represented 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 9 DT40J > = pffiffiffiffi or ¼ B þ AFT D ½5 > ; Ds 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 materials over the years. They were revised as and when new Charpy impact energy or tensile test data became available or following revisions to the neutron 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, with manual welds, plates, and forgings

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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 surveillance 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 factor (k) needed to be introduced to modify the dose term in each material DDR. This meant that the general form of the two-term DDRs for both embrittlement and hardening (eqn [5]) became 9 DT > = pffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi ½6 ¼ B þ AFT Df þ kDt or > ; Ds In this equation, the definitions of B, A, and FT remained unchanged, but the single dose term, D, was replaced by (Df þ 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 > = < ft for f 4:39 10 n cm s 10 0:259 ðft Þe ¼ 4:39 10 > ; : ft for f < 4:39 1010 n cm2 s1 > f

300 250 200 Measured ΔT41J (⬚C)

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

173

150 100 50 0 -50 -50

0

50

100

150

200

250

300

Predicted ΔT41J (⬚C)

(a) 100

¼ effective ðflux-correctedÞ fluence

CRP ¼ B 1þ3:77Ni

1:191

80

f ðCue ; PÞ

gðCue ; Ni; fte Þ

½9

where B ¼ 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) 0 for Cu < 0:072wt% Cue ¼ min½Cuactual ; Cumax for Cu > 0:072wt% ¼ effective Cu level ½10 in which Cuactual ¼ bulk Cu level (wt%), Cumax ¼ 0.243 for typical (Ni > 0.5) Linde 80 welds, and 0.301 for all other materials. (Equations [7] and [8] are in F, reflecting units in the original reference. It is to be noted that F are employed in USNRC regulatory guides, rather than SI units.)

f ðCue ; PÞ ¼

9 8 0 for Cu 0:072 > > > > > > > > > > > = < ½Cu 0:0720:668 for Cu > 0:072 and P 0:008 > e

> > > > ½Cu 0:072 þ 1:359ðP 0:008Þ0:668 > > : for Cu > 0:072 and P > 0:008

> > > > > > ; ½11

gðCue ;Ni;fte Þ

log10 ðft Þe þ 1:139Cue 0:448Ni 18:120 1 1 ¼ þ tanh 2 2 0:629 ½12

Charpy shift ΔT41J (⬚C)

and

60

40 HSST02 SRM (0.17 wt% Cu) Prediction EONY

20

0 0

(b)

5 ⫻ 10

18

1 ⫻ 10

19

1.5 ⫻ 10

19

2 ⫻ 10

19

2.5 ⫻ 10

19

3 ⫻ 10

19

19

3.5 ⫻ 10

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 DT41 J shift for the reference plate HSST02 with the predictions of the EONY model.

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) precipitation 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

174


100 Flux > 4.4 ⫻ 1010 n cm−2 s−1 (E > 1 MeV) Flux = 1⫻1010 n cm−2 s−1 (E > 1 MeV) Flux = 1⫻109 n cm−2 s−1 (E > 1 MeV)

Predicted ΔT41J (⬚C)

80

60

40

4.05.5.4 Japanese Embrittlement Correlations

20

0 1016

1017

Predicted shift from CRP term ΔT41J (⬚C)

1019

140 Cu = 0.25 wt%, Ni = 1.0 wt%

120

−2 −1 Flux > 4.4⫻1010 n cm s Flux = 1⫻109 n cm−2 s−1

100

60

Cu = 0.25 wt%, Ni = 0.6 wt% Cu = 0.15 wt%, Ni = 1.0 wt% Cu = 0.15 wt%, Ni = 0.6 wt% −2 −1 Flux > 4.4⫻1010 n cm s

Decreasing flux Increasing Cu Decreasing Ni

40 Increasing Cu,Ni

20 1016

(b)

1018

Fluence (n cm−2) (E > 1 MeV)

(a)

80

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 correlation 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.

1017

1018

1019

1020

−2

Fluence (n cm ) (E > 1 MeV)

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 cm2 s1).

dependence of both the matrix and CRP term is broadly consistent with the understanding outlined in the previous section. The concept of fte is particularly important as it both provides a means of allowing 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 0.2–0.3 wt% 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, 0.017 wt% P),125 and concluded that out of a total hardness of DHvtot 40 6 the hardness from MD was DHvMatrix 5–10 VPN.

The first embrittlement correlation for the TTS of the Japanese RPV materials, JEAC 4201, was issued in 1991. Additional surveillance data have been compiled 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 twostep 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: mat @CSC ¼ x3 CCu þ e1 DCu þ e2 CMD @t avail 2 0 ½13 þ x8 CCu DCu 1 þ x7 CNi @CMD @CSC 0 2 ¼ x4 Ft2 x5 þ x6 CNi f @t @t mat @CCu @CSC 0 ¼ vSC vSC CSC @t @t

avail 2 vSC ¼ x2 CCu DCu tr

avail CCu

0 avail ¼ x1 CCu DCu vSC ( mat sol 0 CCu CCu mat sol mat sol CCu CCu CCu > CCu

thermal irrad thermal DCu ¼ DCu þ DCu ¼ DCu þ of

½14 ½15 ½16 ½17 ½18 ½19

where Csc and CMD are the number densities of solute mat 0 and CNi are the atom clusters and MD features, CCu

175


pffiffiffiffiffi DTSC ¼ x16 Vf qffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi mat 0 pffiffiffiffiffiffiffiffi ; CSC g CNi þ hðft Þ CSC ½20 ¼ x16 x15 f CCu mat mat C 0 CCu ; CSC ¼ x11 Cu þ x12 f CCu CSC

0 0 x14 2 g CNi ¼ 1 þ x13 CNi hðft Þ ¼ x9 ð1 þ x10 DSC Þft

DSC DCu

pffiffiffiffiffiffiffiffiffi DTMD ¼ x17 CMD DT ¼

qffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi ðDTSC Þ2 þ ðDTMD Þ2

½21 ½22 ½23 ½24 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 contribution 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

120 0.08Cu (CEC) 0.15Cu (CEC) 0.24Cu (CEC) 0.08Cu (MD) 0.15Cu (MD) 0.24Cu (MD)

100

80

DT (ºC)

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 evolution of solute atom clusters and the MD clusters, respectively (see Hiranumu et al.126 for a full description 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 irradiationinduced vacancies and thermal vacancies. Mechanical property changes are correlated with the microstructural changes using the following equations:

60

40

20

0 0.0E + 00

2.0E + 19 4.0E + 19 6.0E + 19 8.0E + 19 Fluence (n cm−2)

1.0E + 20

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.

the surveillance database of Japanese commercial reactors.126,127 The partitioning of the total embrittlement 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 calculated 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 temperature, 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 development 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 fundamental 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.

176


RG1.99r2

Stddev. 11.9 15.4

EWO

10.4

E900-02

11.7 9.4 5.4

Method

140 JEAC4201

120

100

CRIEPI CRIEPI adj

modifications in the form, or the values of, the fitting parameters. The major topics are the following:

Mean error -1.3 -1.9 2.8

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

2.3 0.7 0.1

Prediction (ºC)

80 60 w/o adjustment w adjustment

40

1:1 -2s

20

+2s

0 -20 -20

0

20

40

60

80

100

120

140

Measured value (⬚C)

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 1874 6500.

Different DDRs have been developed in different countries to describe the hardening and embrittlement of the various RPV steels. The inevitably approximate nature of the DDR expressions, the limited variation of different parameters in each surveillance database, and the limited amount of surveillance data mean that the effects of many parameters must be implicit. Different irradiation and compositional variable ranges in different surveillance schemes may contribute significantly to the forms of the DDRs and the strength of different dependences. The limitations 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 application 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

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 provides 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 embrittlement 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 cm2 s1 E > 1 MeV in the IVAR database which is not predicted by the EONY model.30 Further analysis of the IVAR database 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 (compared to lower Ni steels). At present the response of Cu-containing high Ni steels to irradiation doses of 0.1 MeV), was developed, which is as follows: DV ¼ ðFÞ0:4 f1:69 exp½ð0:018T 16:347Þ2 =ag V where a¼

14:87 þ 44:57 exp½0:09ðT 1338:71Þ 1 þ exp½0:09ðT 1338:71Þ

½1

The broader width of the swelling peak as a function of irradiation temperature for the calculation represented 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 transmutation 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 shrinkage 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 irradiations at 723 and 873 K to a fluence of 1.9 1022 n cm2 (E > 0.1 MeV), 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 cm2 (E > 0.1 MeV) at temperatures between 698 and 1323 K,46 samples of Ta–10W were included with postirradiation examination involving TEM analysis. The microstructure of the irradiated Ta–10W contained 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

and more recently by Byun and Maloy.56 In the first, irradiation to 0.13 dpa (where irradiation to 0.76 1022 n cm2, E > 0.1 MeV is 1.0 dpa in pure Ta57) at 673 K resulted in increased yield strength, though no significant loss in ductility occurred over the unirradiated control. However, work softening following the yield drop was observed. Irradiation to higher displacement doses in pure Ta by Wiffen19 showed the potential lower operating temperature limitation of Ta. Following irradiation to 1.97 dpa at 663 K, yield and ultimate tensile strengths increased to near 600 MPa with a corresponding drop in ductility to 0.1 MeV) (assumed 1.75, 5.0, and 8.5 dpa). In the 1373–1473 K 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 1773–1873 K. At 1773 K, measured swelling

220

Radiation Effects in SiC and SiC–SiC

Snead 2006 Snead 2006 Snead, unpub.

Price 1973 Blackstone 1971 Price 1969

Saturable regime point defect swelling

Amorphization regime

Price 1973,#2 Price 1973 Senor 2003 Nonsaturable regime void swelling

20

10 7 5

Swelling (%)

3

8.5 dpa

2

5 dpa

1 0.7

1.75 dpa

0.5 0.3 0.2

0.1 0

200

400

600 800 1000 1200 Irradiation temperature (°C)

1400

1600

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.

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 1773 K, 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 1273 K.3,4,18,22,34,35 Specifically, Senor et al.18 report swelling for the same type of CVD SiC irradiated in this study when irradiated in a watermoderated fission reactor (the ATR) as well. Their maximum dose, irradiation temperature, and swelling data were 1 dpa, 1373 30 K, 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 (1523 K, 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 evolution of irradiated SiC is roughly understood, at least for temperatures up to 1373 K. The swelling near the critical amorphization temperature (423 K) is classically described as the differential strain between the single interstitial, or tiny interstitial clusters, immobile vacancies, and antisite defects. As the temperature 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 1273 K, microstructural studies have noted the presence of both Frank loops and tiny voids, indicating limited mobility of vacancies.


The apparent increase in swelling with dose in the 1373–1873 K 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 2818 K where stoichiometric SiC transforms into C þ liquid phase. This value of 0.65Tm is high when viewed in comparison to fcc metal systems where void swelling typically 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. However, 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 migration 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 carbon 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 mobility measured experimentally. The silicon monovacancy has been shown to be mobile below 1273 K. Using electron spin resonance, Itoh et al.30 found the irradiation-produced T1 center in 3C–SiC disappearing above 1023 K. The T1 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 1673 K.42 Using isochronal annealing and positron lifetime analysis, Lam et al.40 have shown a carbon– silicon vacancy complex to dissociate above 1773 K 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 (1273 K. 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 n m2 (E > 0.1 MeV), 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 examination accounts for only relatively small fractions of the total swelling even in the void swelling regime. Analogous 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 irradiation 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 resistance for a material can be simply described as the summation of the individual components; that is, 1/K ¼ 1/Kgb þ 1/Ku þ 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,

222


Legend

500

Reference

N/R

Material

Note Note Single Crystal

Rohm and Haas Co.

CVD

Grain size ~5µm

Senor et al. (1996)10

CVD

Morton CVD

Graebner et al. (1998)48

CVD

Morton CVD

Pickering et al. (1990)49

CVD

Grain size ~10µm

Rohde (1991)45

Taylor et al. (1993)46 47

Highly pure and dense single-/poly-crystals

400

CVD

Grain size ~3 µm

50

CVD

Grain size >10 µm

50

CVD

Grain size 0.1 MeV) 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. Unfortunately, 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 n m2 (E > 0.1 MeV). However, it is worth noting that the Dienst work used a very limited sample population (about 10 bars.)

231

Statistically meaningful data sets on the effect of flexural strength of CVD SiC have been reported by Newsome and coworkers14 and Katoh and coworkers.58,67 Figure 19 shows a compilation Weibull plot of the flexural strength of unirradiated and irradiated Rohm and Haas CVD SiC taken from the two separate irradiation experiments carried out by Newsome 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 irradiation temperature, including data for unirradiated samples and 1.5–4.6 1026 n m2 (E > 0.1 MeV) dose range. It is likely that the Weibull modulus decreased by irradiation, appearing to be dependent on irradiation 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 573 K population. The scale parameters of flexural strength of unirradiated materials and materials irradiated at 573, 773, and 1073 K were 450, 618, 578, and 592 MPa, respectively. The Weibull modulus of the flexural strength of unirradiated materials and materials irradiated at 573, 773, and 1073 K were 9.6, 6.2, 5.5, and 8.7, respectively, with significant uncertainty. The work of Katoh, on identical material irradiated 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 773 K and 1073 K irradiation of Katoh remained almost unchanged. Given the data discussed on the effect of irradiation on the Weibull 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 determined by the internal pressurization test and the results are plotted in Figure 20. Thin-walled tubular SiC specimens of 1.22 mm outer diameter, 0.1 mm wall thickness, and 5.8 mm 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 n m2 (E > 0.1 MeV) at 1293 and 1553 K. 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

232

Radiation Effects in SiC and SiC–SiC si (MPa) 200

3

300

400

500

600

800

1000

500 ºC, 2.0 dpa m = 5.5

2

Nonirrad. m = 9.6

1

ln(ln(1/(1–Fi)))

0 -1 -2 800 ºC, 2.0 dpa m = 8.7

-3 -4 -5

300 ºC, 2.0 dpa m = 6.2

-6 5.0

5.5

6.0

(a)

6.5

7.0

ln(si) si (MPa) 3

200

300

400

600

800

1000

300 ºC, 6.0 dpa m = 5.5

2 1

500

Nonirrad. m = 9.9

ln(ln(1/(1–Fi)))

0

-1 500 ºC, 6.0 dpa m = 10.8

-2 -3 -4

800 ºC, 7.7 dpa m = 7.9

-5 -6 5.0

5.5

(b)

6.0

6.5

7.0

ln(si)

Figure 19 Weibull plots of flexural strength of unirradiated and irradiated CVD SiC in the dose range of (a) 1.5–4.6 1025 n m2 (E > 0.1 MeV) from Newsome14 and (b) 7.7 1025 n cm2 (E > 0.1 MeV) from Katoh.58

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 337 MPa (from 297 MPa) and the Weibull modulus slightly decreased to 3.9 (from 6.9) after irradiation to 1.9 1025 n m2 (E > 0.1 MeV) dpa at 1293 K. The mean fracture stresses and Weibull moduli at 4.2 1025 n m2 (E > 0.1 MeV) 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 1553 K, no discernible change in fracture stress distribution occurred. The horizontal shift indicates a simple toughening or an increase in fracture toughness alone. While the data for these surrogate TRISO samples, irradiated through internal compression, are somewhat limited, the findings indicate 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 irradiation on strength and Weibull modulus appear


3 2

200

si (MPa) 400 500

300

800

1000

Nonirrad. m = 7.6

1

ln(ln(1/(1–Fi)))

600

233

1280 ⬚C, 4.2 dpa m = 3.8

0 -1 -2

1020 ⬚C, 4.2 dpa m = 5.4

-3 1020 ⬚C, 1.9 dpa m = 4.4

-4 -5 5.0

5.5

6.0 ln(si)

6.5

7.0

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.

appropriate for application to TRISO fuel modeling. Specifically, a slight increase in the mean strength is expected (although it may be less significant 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 conditions. Irradiation creep is important for structural materials for nuclear services as it is a major contributor to the dimensional instability of irradiated 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 irradiated in a fission reactor, and the steady-state creep compliance was estimated to be in the order of 10–38 (Pa dpa m2 (E > 0.18 MeV))1 at 1053– 1403 K. Scholz and coworkers measured the transient creep deformation of SCS-6 CVD SiC-based fiber, which was torsionally loaded under penetrating proton 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 irradiation in HFIR and JMTR at 673–1353 K.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 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.

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 0.13 wt%, 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 0.08 wt%, beyond which d-ferrite is not stable. The specimens without and with 0.1 wt% Y2O3 exhibit the full martensite structure, whereas the specimens with 0.35 and 0.7 wt% Y2O3 exhibit a dual phase comprising 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 0.7 wt% 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

300 0

0.1

0.2

0.3

C content (wt%) Figure 6 Computed phase diagram with respect to carbon content for 9Cr–xC–0.2Ti–2W system without Y2O3.

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 0.7 wt% 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 measurement 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 an AC1 point of 850 C to an AC3 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 0.35 wt% Y2O3 induces an increase up to an AC3 point of 935 C. By comparing both curves, it was found that the specimen with 0.35 wt% Y2O3 exhibits a smaller degree of reduction in linear thermal expansion during the reverse transformation of the a–g-phase; this observation 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

Oxide Dispersion Strengthened Steels

1.4 1.3 1.2 Without Y2O3 AC1 1.1 1 0.9 0.8 AC3 0.7 0.6 700 750 800 850 900 950 1000 1050 1100 1.4 1.3 0.35 mass % Y O 2 3 1.2 1.1 1 AC1 0.9 0.8 AC3 0.7 0.6 700 750 800 850 900 950 1000 1050 1100 Temperature (ºC)

Figure 7 Results of linear thermal expansion measurement between 700 and 1100 C at temperature rising of 0.33 C s1 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.

Y2O3 particles are decomposed during MA, 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 suppressing 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 4 MJ m3 at 1000 C in the case of 0.13 wt% 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

12 F(0.7 mass % Y2O3)

10 Driving force (MJ m–3)

Linear thermal expansion (ΔL/L, %)

246

ΔG(0.2 mass % C) 8

F(0.35 mass % Y2O3)

6 4

F(0.1 mass % Y2O3)

2

ΔG(0.13 mass % C)

0 –2 800

900

1000

1100

1200

1300

Temperature (°C) 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.

F¼

3sfp2=3 8r

;

½3

where, s( Jm2) is the interfacial energy between a- and g-phases, and its value was selected to be 0.56 J m2.25 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 aforementioned equation, the value of pinning force F was determined for 0.1, 0.35, and 0.7 wt% 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 F Þ:

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 transformation from a- to g-phase. On the other hand, DG < F indicates that the a–g interface can be


Oxide particle AC3

γ

・

・

・

・

・

・・

・

・

・

α

・・

・

γ ・

・

・

・

・

α ・

・・

・

・

Carbide 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.

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 contents 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

Hardness (GPa)

6.0

AC1 ・

247

5.0 4.0 3.0 2.0 NT

550 ºC 1h

750 ºC 1h

800 ºC 7h

800 ºC 58 h

FC

Residual ferrite Tempered martensite Average covering residual ferrite and tempered martensite

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.

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: 1360 MPa for the residual ferrite and 930 MPa for the tempered martensite. The yield strength of the residual ferrite is 1.5 times higher than that of martensite 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 distribution 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, residual ferrite contains fine (nanosized) oxide particles which are responsible for higher strength in residual ferrite containing ODS steels. In regard to the bimodal distribution of oxide particles in martensite


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 transformation) diagram is essential to the microstructure

(a)

(b)

20 nm

20 nm

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 martensite ODS 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 normalizing heat treatment used in commercial furnaces, the cooling rate would be roughly 3000 C h1, so that

1300

1000 900

9Cr-ODS ferritic steel Containing residual-α

No residual-α

1100

Temperature, T (⬚C)

800

EM10

700

γ

α+γ 900

600 α 500

700

400

500

200 100

α+γ+m

γ+m

300

Temperature, T (K)

248

m

α+ m 18 000 (K/h) Austenite (γ) => Martensite 3000 (K/h) 101

102

103

30 (K/h) Austenite (γ) => Ferrite (α) 104

300 105

Time from 800 ⬚C, t (s) Figure 12 CCT diagram of 9Cr-ODS steel. Reproduced from Ohtsuka, S.; Ukai, S.; Fujiwara, M.; Kaito, T.; Narita, T. J. Nucl. Mater. 2006, 351, 241.

249


Low carbon steel Elemental powders Yttria powder

MA powder

9Cr–0.13C–2W–0.2Ti–0.35Y2O3

Cladding tube

Mechanical alloying (MA)

At intermediate heat treatment At final heat treatment

Cold rolling (pilger mill)

Hot extrusion (1423 K)

Figure 13 Cladding tube manufacturing process developed for 9Cr-ODS steel.

the matrix phase of 9Cr-ODS steel cladding consists of residual ferrite, martensite, and a small amount of transformed ferrite from the g-phase. 9Cr-ODS steels are promising materials to enable fast reactor fuel cladding to realize a high burnup of 200 GWd t1 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 8.5 mm in diameter by 0.5 mm in thickness by 2 m 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 hotextrusion at an elevated temperature of 1150 C to the dimensions of 32 mm in outer diameter, 5.5 mm in wall-thickness, and 4 m in length. After machining to the precise dimensions, claddings are produced at their final dimension (8.5 mm in outer diameter, 0.5 mm in thickness, and 2 m 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 manufacture 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

1st Hardness (Hv)

4.08.3.3.2 Manufacturing process

Cold rolling (Rd = 50%)

450

2nd

3rd

4th

400

350 Mother tube

1st

2nd

3rd

4th

Heat treatment

300 Figure 14 Hardness change in the process of cold rolling and intermediate and final heat treatments for cladding tube manufacturing of 9Cr-ODS steels.

rate of about 150 K h1 was applied to the intermediate heat treatment in order to induce the ferrite phase at room temperature without martensite transformation. This phase has a lower degree of hardness. Hardened cladding due to the accumulation of cold deformation can be sufficiently softened by this intermediate heat treatment, and cold rolling can then be continued with 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


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 cladding 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 120 MPa 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.

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 stainless 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 claddings can probably be ascribed to the pinning of dislocations 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

500 400

Hoop stress (MPa)

300

200

PNC316 (923 K) PNC316 (973 K) PNC316 (1023 K)

9Cr-ODS (923 K) 100 90 80 70 60 50 10

Stress range for SFR fuel cladding

HT9 (973 K)

9Cr-ODS (973 K) HT9 (923 K)

9Cr-ODS (1023 K)

HT9 (1023 K) 102

103 Time to rupture (h)

104

105

Figure 15 Creep rupture curves of 9Cr-ODS steel claddings in hoop direction by using internally pressurized specimens at temperatures of 650, 700, and 750 C, compared with those of HT9 and PNC316. Reproduced from Allen, T.; Burlet, H.; Nanstad, R. K.; Samaras, M.; Ukai, S. Mater. Res. Soc. Bull. 2009, 34(1), 20–27.


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). The manufacturing 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 7.5 mm outer

diameter, 0.4 mm thickness, and 1 m 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

1500


Tensile strength (MPa)

251

1000

500

PNC-FMS

0 200

400

600 800 1000 1200 Temperature (K)

10

5 PNC-FMS

0 200

400

600 800 1000 Temperature (K)

1200

Figure 16 Tensile strength and uniform elongation of 9Cr-ODS steel cladding in hoop direction by the ring specimens. Reproduced from Ukai, S.; Kaito, T.; Otsuka, S.; Narita, T.; Fujiwara, M.; Kobayashi, T. ISIJ Int. 2003, 43, 2038.

1000 Uni-axial Bi-axial 1 DK 1 DS 63 DSA

600

Stress (MPa)

500 400

Uni-axial

300

200

Bi-axial

MA 957 (see figure 29)

Mol-ODS (DT2203Y05, see figure 28)

100 1

10

100 1000 Time to rupture (h)

10 000

100 000

Figure 17 Creep rupture strength of 1DK, 1DS, and 63DSA claddings in hoop direction by using internally pressurized specimens at 650 C. Reproduced from Ukai, S.; Harada, M.; Okada, H.; Inoue, M.; Nishid, T.; Fujiwara, M. J. Nucl. Mater. 1993, 204, 65–73.

252


there is strong strength anisotropy, and the bi-axial creep rupture strength is considerably lower than that of the uni-axial direction. Microstructure observations 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 studied to change the substantially elongated grain structure to the equi-axed grain structure. The Y2O3 content should be 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 investigated 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 condition is expressed using applied stress s as: ½5 "_ ¼ Asn where n is the stress exponent and A is the temperature-dependent coefficient.37 In the case of

254


1000

Hoop stress (MPa)

F1 F2 F3 F4 F4 (uni-axial, this work)

Uni-axial

100

Internal, bi-axial direction

10


10 000

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.

mode of F4, and a higher strain rate is found even below a stress of 200 MPa. A transverse section of this specimen shows finely equi-axed grains of 5–10 mm (Figure 19). Apart from pinning the gliding dislocations due to oxide particle-dislocation interaction, the deformation 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 viewpoint of the interaction between Y2O3 particles and dislocations. This interaction could be formulated by the void-hardening mechanism proposed by Srolovitz,38 in which oxide particles were replaced 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=r0 Þ þ B; for screw dislocation;

10−5

Strain rate (s−1)

10−6

½6

A ¼ ð1 þ v sin2 ’Þcos ’=ð1 vÞ;

F1 F2 F3 F4 F4 (uni-axial) PNC-FMS

B ¼ 0:6 for edge dislocation; A ¼ 1 v sin2 ’=ð1 vÞ cos ’;

10−7

B ¼ 0:7 10−8

10−9

10−10 50

60

70 80 90 100

200

300

Stress (MPa)

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.

the uni-axial creep mode, a significantly high stress sensitivity of n ¼ 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 250 MPa; 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

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 dislocation 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. Further, 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 centerto-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 distribution 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


substituting l, M ¼ 3.0,41 n ¼ 0.334, b ¼ 2.48 1010 m, and G ¼ 50 600 MPa, at 700 C. Figure 22 shows the results of analyses in relation 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

400 350 F4

F4 (uni-axially longitudinal)

Hoop stress (MPa)

300

Oxide particle-hardening stress (s p) from particle distribution by TEM

250 F3

200

F2

150

F1

100 Stress at strain rate of 10−9 S−1 in the internal hoop direction

50 0

0

50 100 150 200 250 300 Face-to-face distance between particles, l (nm)

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 109s1 for F1, F3, and F4 specimens, as functions of face-to-face distance between particles. Each stress was obtained at 973 K. 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.

255

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 109 s1 in the internal hoop directional mode. For the F1 specimen, as a stress level corresponding to a strain rate of 109 s1 approaches the oxide particle-hardening stress, the strong anisotropy tends to disappear. However, for the F3 and F4 specimens with a shortened distance between particles, stress levels for a strain rate of 109 s1 in the hoop direction are degraded from the oxide particlehardening 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 5–10 mm are formed, and the grain boundaries 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

1400 Temperature (⬚C)

1200 1000 VHTR

GFR

800 600

SCWR

SFR

LFR

400 200

MSR

Generations II–III

0 0

50

100

150

200

Displacement per atom Figure 23 Temperature and dose regimes for Generation IV advanced nuclear power plants. VHTR: very high temperature reactor; SCWR: supercritical water-cooled reactor; GFR: gas fast reactor; LFR: lead fast reactor; MSR: molten salt reactor; SFR: sodium fast reactor. Reproduced from Guerin, Y.; Was, G. S.; Zinkle, S. J. Mater. Res. Soc. Bull. 2009, 34(1), 10–14.


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 SCWR and LFR is to improve their resistance to corrosion. It has been reported that the addition of chromium (>13 wt%) and aluminum (4 wt%) to ODS steels quite effectively suppresses corrosion in an SCPW and LBE environment. In general, however, an increase in the Cr content often results in increased susceptibility to thermal aging embrittlement. Furthermore, the addition of Al significantly reduces steel strength at high temperatures. 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. 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 corrosion resistance and aging embrittlement caused by increasing Cr content is one of the critical issues facing the developers of high-Cr ODS steels. The aging effects of ODS steels with different Cr content were investigated by measuring their impact fracture energy at RT after 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 >18 wt% show a significant reduction in fracture energy after aging for 100 h. In contrast, 16Cr–4Al ODS steel showed a small reduction in fracture 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 considered to be Cr-rich phases. In order to reduce susceptibility to aging embrittlement, the Cr content could be 1000 C. Tubes, sheets, and bars made from these steels are commercially used in various stationary and hightemperature 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 SCKCEN (Centre d’Etude de l’eńergie 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 hollows of 20/17 mm. 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

Basic chemical composition of ODS steels (mass %)

Steels

Cr

Mo

W

Ti

Al

Dispersoid

Fe

20 19

– –

– –

0.5 0.5

4.5 5.5

0.5Y2O3 0.5Y2O3

Bal Bal

SM/US Plansee/Austria

13

1.5

–

2.2

–

Bal

DT2906

13

1.5

–

2.9

–

0.5Y2O3, 0.9Ti2O3 1.8Ti2O3

Incoloy MA957 9Cr-ODS steel

14 9

0.3 –

– 2

1 0.2

– –

0.25Y2O3 0.35Y2O3

Bal Bal

SCK CEN Mol/Belgium SCK CEN Mol/Belgium SM/US JAEA/Japan

12Cr-ODS steel 16Cr–4Al-ODS steel

12 15.5

– –

2 2

0.3 0.1

– 4

0.23Y2O3 0.35Y2O3

Bal Bal

Turbine, combustion Incoloy MA956 PM2000 Fast reactor fuel DT2203Y05

Others

Bal 0.13C, martensite + residual ferrite 0.6Hf or 0.6Zr

SM: Special metals, former International Nickel Company; JAEA: Japan Atomic Energy Agency; KU: Kyoto University. SCKCEN: Centre d’Etude de l’eńergie Nucleaire – Studiecentrum voor Kernenergie.

Development

JAEA/Japan KU/Japan


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 nondestructively tested using eddy currents and ultrasonics 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 SCKCEN, 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 endplug 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 reactor 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 25.4 mm was gun-drilled in order to generate a tube hollow with a 4.75 mm thick wall. Extensive cladding fabrication 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

Hoop stress (N mm–2)

1000

600 ⬚C 650 ⬚C 700 ⬚C 750 ⬚C

100

10 10


259

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 significantly degraded in the hoop direction, which is essential 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 9CrODS 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 conventional 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, thereby maintaining 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 environment was observed. However, under a flowing sodium condition of 4.5 m s1, the element nickel penetrates the surface of ODS steel cladding, where an increase in nickel concentration and decrease in chromium concentration 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 microstructure change associated with nickel diffusion into the cladding surface need to be considered.53

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.

4.08.7.2

LBE Compatibility

Molten LBE has a high solubility of nickel, iron, and chromium, which are the most important alloy elements

260


10 000 STC TR PNC

STC ORT 650 ⬚C 700 ⬚C 760 ⬚C

s (MPa)

1000

650 ⬚C 100

704 ⬚C 760 ⬚C

→ → → → →

10 10

100

1000

→ →

10 000 39011045.6

tr (h)

Figure 29 Creep rupture strength of Incoloy MA957 cladding. Reproduced from Hamilton, M. L.; Gelles, D. S. PNNL-13168, Feb 2000.

80 M93: 650 ⬚C M93: 700 ⬚C F95: 650 ⬚C F95: 700 ⬚C Sodium flow < 0.001 m s–1

UTS

60

40

20 As-received 0

0

200

400

600

800

1000

1200

Sodium exposure time (h) 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.

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 considered more appropriate for LBE application. Exposure of 9Cr-ODS steels to an LBE environment 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 s1, and oxygen sensors were used to measure and maintain an oxygen concentration of about 1 106 wt%. Samples were exposed for 200, 400, and 600 h, in order to study the early stages of oxide formation and growth. A crosssectional 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


1000

Hoop stress (MPa)

Sodium flow rate; < 0.001 m s–1

261

Material: M11 700 ⬚C : In air/Ar 650 ⬚C 650 ⬚C

700 ⬚C : In sodium

650 ⬚C

700 ⬚C 100 10

100

1000 Time to rupture (h)

10 000

100 000

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

PbMb, 124

CrKa, 26

FeKa, 64

Bulk Oxide

LBE 15 kV X 2000

Diff. zone 10 μm

24 36 BEC

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.

of an outer magnetite (Fe3O4) layer and an inner Fe–Cr spinel ((Fe,Cr)3O4) layer, which is sometimes accompanied 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 temperatures 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 from magnetite to wustite 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 18 wt% Cr-ODS steel without the addition of Al dissolved markedly into LBE, while those ODS specimens containing 4 wt% 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 resistance at 650 C for 5000 h.

262


6

16Cr–4Al Weight gain (mg dm–2)

14Cr–4Al

With 4 wt% Al 1800 h

4

600 h 2 100 h 0 14

(a)

Figure 33 The appearance of Al added high Cr-ODS steel specimens after corrosion test in LBE for 1 104 h at 923 K (DO: 1 106 wt%). Reproduced from Kimura, A.; Kasada, R.; Iwata, N.; et al. In Proceedings of ICAPP ’09, Tokyo, Japan, May 10–14, 2009; Paper 9220.

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 17 wt % does not affect corrosion resistance if ODS ferritic steels contain 4 wt % Al. For 16 wt% Cr, the addition of A1 increases corrosion resistance in 16Cr-ODS steels. As shown in Figure 35,43 tested at SCPW (510 C, 25 MPa) for 600 h, the addition of 4 wt% Al did not significantly influence corrosion resistance in 19CrODS steel, though a rather dense chromia film was observed on the specimen surface. The 16 wt% 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 2 wt%. Thus, the addition of Al effectively improves corrosion resistance 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 formation 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

18

21

19Cr–4Al

1800 h Weight gain (mg dm–2)

18Cr

16 Cr content (wt%)

With 16 wt% Cr 14 600 h

7 100 h

0 0 (b)

2 Al content (wt%)

4

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 8 ppm of dissolved oxygen under a pressure of 25 MPa (10 dm = 1 m). Reproduced from Lee, J. H.; Kimura, A.; Kasada, R.; et al. In Proceedings of ICAPP ’09, Tokyo, Japan, May 10–14, 2009; Paper 9223.

controlled atmosphere of dry air. Weight measurement to evaluate the degree of oxidation was performed 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 oxidation was quite small and comparable to that of PNC316 containing 17 wt% Cr. Their weight gain is limited to below 0.1 mg mm2. On the other hand, a quite large oxidation of 0.8 mg mm2 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 17 wt% Cr containing stainless steel. The element distribution obtained by Electron probe microanalysis (EPMA) showed a scale consisting

Weight gain by oxidation (mg mm–2)


30.0 SCW 773 K.25 MPa. 600 h

20.0

Weight gain (g m–2)

10.0

SUS430 9Cr-ODS (16Cr)

1.2 16Cr-ODS

0.10 923 K ⫻ 50 h 0.08

0.04 0.02 0.00

2 Al content (mass %)

4

Weight gain by oxidation (mg mm–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.

1.0 9Cr-ODS 12Cr-ODS (fine grain) PNC316 PNC-FMS SUS430

0.8 0.6 0.4 0.2 0.0

0

500

1000 1500 Testing time (h)

2000

Cr supply through grain boundary diffusion

12Cr–ODS (fine grain)

12Cr–ODS (large grain)

PNC–FMS

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.

19Cr-ODS 0

Y2O3 effects

0.06

0.6

0

263

2500

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.

of Fe-rich oxide in the outer layers and Cr-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. Nevertheless, 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


Environmental and Molecular Science Laboratory at Pacific Northwest National Laboratory, using 5 MeV Ni ions at 500, 600, and 700 C with a damage rate of 1.4 103 dpa s1. 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 temperatures (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 in Table 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 discussed 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 temperatures 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 1.0 MeV electrons, but dissolved under 1.2 MeV 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 precipitate mass loss if vacancy absorption frees a precipitate atom. From a comparison between electron irradiation (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 5 mm-wide ring-tensile specimens with a 1.5 mm-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 reactor JOYO using the material irradiation rig at temperatures between 400 and 534 C to fast neutron fluences ranging from 5.0 1025 to 3.0 1026 nm2 (E > 0.1 MeV). 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 irradiated F94, F95, and M93 cladding, is modestly higher ( 0.1 MeV)


(1.35)

1000 (3.56)

F94 F94 unirrad. F95 F95 unirrad. M93 M93 unirrad. 1DS unirrad.

800 (0.45)

600

Y2O3(wt%)

400 Fluence (0.5)

200 600

650

(2.8)

F94 F95 M93 (3.0) (1.4, 2.5) 1DS

700 750 800 Test temperature (K)

850

0.24 0.24 0.35 0.40

900

F94 F94 unirrad. F95 F95 unirrad. M93 M93 unirrad. 1DS 1DS unirrad.


1200

8

(Fluence; ⫻ 1026 n m–2(E > 0.1MeV) (0.5) (2.8) (3.0) (1.4, 2.5)

6 4 2

Fluence (1.35)

0 600

650

(3.56)

700 750 800 Test temperature (K)

(0.45)

850

900

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.

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.

xenon and krypton tag gases was enclosed. The irradiation temperatures were 700, 725, and 750 C, and the hoop stress ranged from 45 to 155 MPa. The maximum neutron dose reached 20 dpa. It was confirmed that inpile creep rupture time is located within the out-ofpile 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 encouraging. Evidence was apparent in both MA956 and MA957

268


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.

of a0 precipitation, and in regions where recrystallization 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. 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 PRWwelded 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 integrity 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 endplug 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

Figure 42 Optical micrograph of 9Cr-ODS fuel pin after irradiation at 700 C, 5 at.% burnup and 25 dpa in BOR-60. Reproduced from Kaito, T.; Ukai, S.; Povstyanko, A. V.; Efimov, V. N. J. Nucl. Sci. Technol. 2009, 46(6), 529–533.

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 temperature 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)

100 nm

269

(c)

1 μm

1 μm

Figure 43 Precipitation occurring during the in-pile service (a) a0 -phases at 400 C, 0 dpa, (b) w-phases at 523 C to 78.8 dpa, and (c) Laves phases at 580 C to 30.5 dpa. 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.

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 and MOX 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, demonstrating some in-core ductility. This program demonstrated the viability of ODS steel as a potential cladding material for long-life advanced FRs. 4.08.8.3.3 DT2203Y05 in Pheńix

Fuel pins with DT2203Y05 cladding were irradiated in an experimental capsule placed in a special subassembly in Pheńix. The process by which they were manufactured was described in Section 4.08.6. The dose reached at midplane was 81 dpa and the temperature along the fuel pin ranged from 400 to 580 C. It was observed by TEM that the uniform distribution 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 dissolution depend on radiation-enhanced solute diffusivity and enhanced solubility under irradiation. A uniform distribution of tiny particles 68HP, which is a direct result of the decreasing alloying content of the strong melting point depressants molybdenum and niobium.7,10–15

278

Welds for Nuclear Systems

Transvarestraint test data GTAW weld buildups 180 amps, 12.7 volts, 5 ipm ~45.7 kJ in.–1

3.00

Alloy 625

2.50

Maximum crack distance (mm)

EN52MSS

2.00

EN52i 1.50 EN82H SANICRO 68HP 1.00 347 SS

0.50 308L SS Different symbols of the same color denote different heats

0.00 0

1

2

3

4

5

6 7 8 9 Applied strain (%)

10

11

12

13

14

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.

4.09.1.1.2 Liquation cracking

In contrast to solidification cracks, liquation cracks occur in the partially melted zone and the heataffected zone and can be either interdendritic or intergranular in nature. An example of an interdendritic 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 mechanism 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 secondphase particles in the heat-affected zone. The resulting 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 transgranular or interdendritic and is favored by mechanical 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.


HAZ/PMZ

279

Autogenous weld bead

Liquation cracks

500 mm

Welding direction

Liquid at the time of straining

Mushy zone at the time of straining

Solidification cracks

Solid at the time of straining

2.5 mm 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.

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 precipitation 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

280


Alloy 625

MLTS-2 (27Cr)

EN82H

68HP

347 SS

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.

loss of tensile ductility in the homologous temperature range (0.4–0.9 Tm) that corresponds to the time/temperature regime of the precipitation of a partially or fully coherent second phase. In lowstrength 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 kinetics of the detrimental phase and the macroscopic tensile ductility is shown in Figure 9, which compares 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). Ductility recovery occurs as precipitation progresses because local misfit strains decrease as chromium depletion occurs and as misfit dislocations are generated (region 3). Ductility is restored when precipitation 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 preferentially 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 stresses (e.g., from thermal strains on-cooling or applied during hot ductility testing), these cracks often link up

and form the classic ‘ductility dip’ 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 structure and are littered with (Cr,Fe)23C6-type carbides.24 Figure 10(d) illustrates how the misfit strain between the carbide and matrix increases with increasing chromium concentration in the alloy. In part, this explains why 30 wt% alloys (A690 and EN52) are more susceptible 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

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