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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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
Permission Acknowledgments
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. Be´chade 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
Radiation Effects in Zirconium Alloys
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 (<10 K). The irradiation damage was monitored in situ using electrical resistivity changes.4,5 The measured minimum displacement threshold energy transferred to the Zr atoms is Ed ¼ 21–24 eV. Measurements of Ed have also been performed using a high-voltage electron microscope (HVEM) to irradiate a Zr thin foil. The values obtained were found to be weakly orientation dependent, between 24 and 27.5 eV, with a mean Ed of 24 eV.6 The displacement energy has also been computed by molecular dynamics (MD) simulations based on various interatomic potentials. The most accurate computations have been performed using a manybody (MB) potential based on the Finnis and Sinclair formalism.7 These authors have found that the displacement energy is significantly anisotropic. Displacement energy was found to be minimum for knocking out in the basal plane, that is, in the h1120i directions, corresponding to the most favorable direction for replacement collision sequences, and to the direction of development of the basal crowdion. The corresponding displacement energy obtained (Ed ¼ 27.5 eV) is slightly above the experimental values. The value averaged over all the crystallographic directions was found to be 55 eV. The value specified in the norm reference test standard (Standard E521–89, Annual Book of ASTM Standards, ASTM, Philadelphia, PA, USA) is Ed ¼ 40 eV.8 This value is close to the spatial means obtained by MD models. 4.01.1.1.3 Displacement cascade in zirconium
The number of displaced atoms inside the cascade can be simply estimated using the Kichin– Pease formula9 or the modified Kichin–Pease formula (Norgett–Robinson–Torrens model or NRT model).10,11 According to this last model, the number of displaced atoms within the cascade in the case of a 22 keV PKA and using a displacement energy of Ed ¼ 40 eV is np ¼ 0:4ET =Ed 220. Because of the large mean free path of fast neutrons (several
3
Radiation Effects in Zirconium Alloys
centimeters), it can be considered that only one PKA is created by the incoming neutron going through the Zr cladding used in pressurized water reactors (PWRs) (with a thickness of 0.6 mm). Therefore, if the PKA creation rate per unit volume within the cladding is known for a typical fuel assembly in a PWR (with typical fast neutron flux is 5 1017 n m2 s1 (E > 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 (<2 ps), the cascade is composed of a core with a high vacancy concentration, and the self interstitial atoms (SIAs) are concentrated at the cascade periphery.14–16 The cascade creation is followed by the athermal cascade relaxation that can last for a few picoseconds. During this phase, most of the displaced atoms quickly reoccupy lattice sites as a result of prompt (less than a lattice vibration period, 0.1 ps) elastic recombination if a SIA and a vacancy are present at the same time in the elastic recombination volume (with 200
100 K
25
12 10 8 6 4 2 0 1 2 3 4 5 6 7 8 9 10 Clu
ster
size
11 12
13 14
15
0.5 25 0.5
1
1
5
5
10
rgy
ne Ae
10
(ke
20
20
Number of vacancies per cascade
Number of interstitials per cascade
14
(a)
600 K
20 15 10 5 0 1 2 3 4 5 67 8 9 1011 C
V)
luste
r siz
PK
e
1213 1415
24
0.5 30 0.5
1
1
A
PK
5
5
20 10 10 )
rgy ene
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(b)
Figure 1 Number of single and clustered (a) interstitials and (b) vacancies per cascade as a function of the PKA energy. Adapted from Gao, F.; Bacon, D. J.; Howe, L. M.; So, C. B. J. Nucl. Mater. 2001, 294, 288–298.
4
Radiation Effects in Zirconium Alloys
The form of these small clusters is also of major importance since it plays a role on the nucleation of dislocation loops. Wooding et al.16 and Gao et al.8 have shown that the small SIA clusters are in the form of dislocation loops with the Burgers vector 1=3h1120i. The collapse of the 24-vacancy cluster to a dislocation loop on the prism plane was also found to occur. 4.01.1.2 Evolution of Point Defects in Zirconium: Long-Term Evolution After the cascade formation and relaxation, which last for a few picoseconds, the microstructure evolves over a longer time. The evolution of the microstructure is driven by the bulk diffusion of point defects. For a better understanding of the microstructure evolution under irradiation, the elementary properties of point defects, such as formation energy and migration energy, have first to be examined. 4.01.1.2.1 Vacancy formation and migration energies
Concerning the vacancy, all the atomic positions are identical in the lattice and so there is only one vacancy description leading to a unique value for the vacancy formation energy. Due to the rather low a–b phase transformation temperature, the measurement of vacancy formation and migration energy in the Zr hexagonal close-packed (hcp) phase is difficult. The temperature that can be reached is not high enough to obtain an accurately measurable concentration and mobility of vacancies.18 Nevertheless, various experimental techniques (Table 1), such as positron annihilation spectroscopy or diffusion of radioactive isotopes, have been used in order to measure the vacancy formation and migration energies or the self-diffusion
Table 1 Experimental determination formation (Ef), migration (Em) and self diffusion activation (Ea) energies for vacancy (in eV) Experimental methods
Ef
Em
Ea
Reference
Semiempirical Self-diffusion Diffusion behavior of various solutes in Zr Self-diffusion
1.8–1.9 – 1.4–2.1
1.3–1.6 – 1.1–1.5
3.3 1.2–3.5 3.2–3.5
[18] [18] [19]
–
–
2.85
[20]
coefficient.18–26 The values obtained by the various authors are given in Table 1. It is pointed out by Hood18 that there is great discrepancy among the various results. It is particularly shown that at high temperature, the self-diffusion activation energy is rather low compared to the usual self-diffusion activation energy in other metals.18 However, as the temperature decreases, the self-diffusion activation energy increases strongly. According to Hood,18 this phenomenon can be explained assuming that at high temperature the vacancy mobility is enhanced by some impurity such as an ultrafast species like iron. At lower temperature, the iron atoms are believed to form small precipitates, explaining that at low temperatures the measured self-diffusion energy is coherent with usual intrinsic self-diffusion of hcp crystals. It is also shown that the self-diffusion anisotropy remains low for normal-purity zirconium, with a slightly higher mobility in the basal plane than along the hci axis.22,26,27 For high-purity zirconium, with a very low iron content, the anisotropy is reversed, with a higher mobility along the hci axis than in the basal plane.27 The vacancy formation and migration energies have also been computed either by MD methods, where the mean displacement distance versus time allows obtaining the diffusion coefficient, or by static computation of the energy barrier corresponding to the transition between two positions of the vacancy using either empirical interatomic potential7,28–34 or the most recent ab initio tools.35–38 Since the different sites surrounding the vacancy are not similar, due to the non-ideal c/a ratio, the migration energies are expected to depend on the crystallographic direction, that is, the migration energies in the basal plane == Em and along the hci direction Em? are different. The results are given in Table 2. The atomistic calculations are in agreement with the positron annihilation spectroscopy measurement but are in disagreement with the direct measurements of self-diffusion in hcp zirconium.20 As discussed by Hood,18 and recently modeled by several authors,39,40 this phenomenon is attributed to the enhanced diffusion due to coupling with the ultrafast diffusion of iron. 4.01.1.2.2 SIA formation and migration energies
In the case of SIAs, the insertion of an additional atom in the crystal lattice leads to a great distortion of the lattice. Therefore, only a limited number of configurations are possible. The geometrical description of all the interstitial configuration sites has been
Radiation Effects in Zirconium Alloys
proposed for titanium by Johnson and Beeler41 and is generally adopted by the scientific community for other hcp structures (Figure 2). T is the simplest tetrahedral site, and O is the octahedral one, with, respectively, 4 and 6 coordination numbers. BT and BO are similar sites projected to the basal plane with three nearest neighbors, but with different numbers of second neighbors. BC is the crowdion extended defect located in the middle of a segment linking two basal atoms.
Table 2 Computation determination formation (Ef), migration (Em), and self-diffusion activation (Ea) energies for vacancy (in eV) ==
Computation methods
Ef
Em
? Em
Pair potential Finnis–Sinclair MB potential Finnis–Sinclair MB potential Finnis–Sinclair MB potential EAM potential Ab initio FP-LMTO Ab initio GGA Ab initio GGA Ab initio LDA
1.59 1.79
1.21 0.93
1.10 0.93
1.79
–
–
–
[7]
1.79
0.84
0.88
2.64
[34]
1.74 2.07
0.57 –
0.59 –
2.32 –
[31] [30]
1.86 2.17 2.29
– 0.51 0.23
– 0.67 0.43
– 2.76 2.78
[36,37] [38] [38]
Ea
Reference [28] [33]
MB: many body; EAM: embedded atom method; FP-LMTO: full-potential linear Muffin-Tin orbital; GGA: generalized gradient approximation; LDA: local density approximation.
C is the interstitial atom located between two adjacent atoms of two adjacent basal planes in the h2023i direction. This direction is not a closepacked direction, and allows easier insertion of the SIA. S is the split dumbbell position in the hci direction. The only way to have access to the SIA formation energy is from atomistic computations taking into account the different configurations of the SIA given previously. In their early work on titanium, Johnson and Beeler41 found that the most stable SIA configuration was the basal-octahedral site (BO). Several other sites were also found to be metastable, like asymmetric variants of the T and C sites. As reviewed by Willaime,35 the relative stabilities of the various SIA configurations were observed to depend strongly on the interatomic potential used (Table 3). The mobility of SIAs can be estimated experimentally using electron irradiation at very low temperatures (4.2 K), followed by a heat treatment. During the recovery, the electrical resistivity is measured. The main recovery process was found around 100–120 K and analysis of the kinetics gives the SIA migration energy of Em 0.26 eV.4 Atomistic computations have also brought results (Table 3) concerning the SIA migration energy. Several authors7,28–31,33–37 have found that the mobility of SIAs is anisotropic, with low migration activation energy for the basal plane mobility (Em== 0.06 eV) and a higher migration activation energy in the hci direction (Em? 0.15 eV). In the temperature range of interest for the power reactors (T 600 K), the diffusion coefficients obtained are the following: == Di ¼ 8 109 m2 s1 (in the basal plane) and
BS S C
O
S
C
BO
BS
BC
O BO BT
5
BC
T
Figure 2 Interstitial sites configuration: (a) static localizations (adapted from Bacon, D. J. J. Nucl. Mater. 1993, 206, 249–265) and (b) relaxed configurations (adapted from Willaime, F. J. Nucl. Mater. 2003, 323, 205–212).
6
Radiation Effects in Zirconium Alloys
Table 3 eV)
Computation of SIAs formation (Ef) and migration (Em) energies in Zr by ab initio, MD, or MS (molecular statics) (in
Method
Ef
Em ==
Reference
O
BO
BS/BC
C
S
T
Em
? Em
Pair potential
–
3.83
–
4.01
–
–
BO: 0.49
[28]
EAM potential Finnis-Sinclair MB potential Finnis–Sinclair MB potential Ab initio GGA Ab initio LDA Ab initio GGA Finnis–Sinclair MB potential
2.8 – – 2.84 2.79 3.04 4.13
2.63 3.97 – 2.88 2.78 3.14 3.97
2.5 3.76 – 2.95 2.90 3.39 3.75
2.78 3.97 – 3.08 3.07 3.52 3.96
3.04 4.32 – 3.01 2.80 3.28 3.77
_ _ – 4.03 – – 3.98
BO: 0.8 – C: 0.49 0.05 – 0.06 – – – –
C: 0.29 0.14 – 0.15 – – – –
[31] [7] [33] [36,37] [35] [35] [34]
MB: many body; EAM: embedded atom method; FP-LMTO: full-potential linear Muffin-Tin orbital; GGA: generalized gradient approximation; LDA: local density approximation.
Di? ¼ 109 mm2 s1 (along the hci direction). These authors have also shown that the anisotropy depends on the temperature. Computing the effective diffusion rate of SIAs in all directions, taking into account the multiplicity of the jump configurations for each type of migration, Woo and co-workers34,42 have obtained the anisotropy for self-interstitial diffusion as a function of temperature. It is shown that the SIA mobility is higher in the basal plane than along the hci axis and that the anisotropy decreases when the temperature increases. 4.01.1.2.3 Evolution of point defects: Impact of the anisotropic diffusion of SIAs
In zirconium alloys, as in other metals, under irradiation both vacancies and SIAs (Frenkel pairs) are created within the cascade leading to an increase of the point-defect concentration with the irradiation dose. However, even at very low temperature, the Frenkel pair concentration saturates at values about 1% due to the mutual recombination of vacancies and SIAs.43 At higher temperatures, point defects migrate and can therefore disappear because of a large variety of defects/defects reactions. Three major mechanisms contribute to defect elimination: vacancy–SIA recombination, point-defect elimination on defect sinks (dislocation, grain boundaries, free surface, etc.), and agglomeration in the form of vacancy dislocation loops and interstitial dislocation loops. It has to be noted that, because of the rapid migration of SIAs compared to the slow migration of vacancies, at steady state the vacancy concentration is
several orders of magnitude higher than the SIA concentration. Because of the elimination of point defects on point-defect clusters, the clusters can grow under irradiation depending on their relative capture efficiency. In the case of cubic metals, since the relaxation volume of SIAs is usually much larger than that of vacancies, edge dislocations eliminate SIAs with a higher efficiency than vacancies (positive bias toward SIAs). Assuming an isotropic diffusion of point defects, this phenomenon leads to a preferred absorption of SIAs by dislocations, provided that there is another type of sink within the material. Because of this preferential absorption of SIAs, the interstitial loops tend to grow under irradiation and the vacancy loops tend to shrink. However, in hcp zirconium, the point-defect diffusion is usually considered to be anisotropic although there is little experimental evidence of this phenomenon. From the experimental results, it is believed that vacancy migration is only slightly anisotropic but the SIA migration is believed to be significantly anisotropic, as shown by atomistic computations. This diffusional anisotropy difference (DAD) has a strong impact on capture efficiency of point defects by sinks.44 Indeed, assuming SIAs to have a higher mobility in the basal plane than along the hci axis and that the vacancies have an isotropic diffusional behavior, it can be seen that grain boundaries perpendicular to the basal plane absorb more SIAs than vacancies. On the other hand, grain boundaries parallel to the basal plane absorb more vacancies
7
Radiation Effects in Zirconium Alloys
than SIAs. Similarly, a line dislocation parallel to the hci axis absorbs more SIAs than vacancies and a line dislocation in the basal plane absorbs more vacancies than SIAs. As discussed by Woo,44 this geometrical effect due to the DAD can overwhelm the conventional bias caused by the point-defect/sink elastic interaction difference (EID). Thus, contrary to the implications of the conventional rate theory, edge dislocations in a-zirconium are not necessarily biased toward SIAs, and grain boundaries are no longer neutral sinks. As will be described in the following, this phenomenon can explain some anomalous irradiation-induced microstructural features as well as the growth phenomenon of zirconium alloys.
0.5 mm
(b)
(a)
Figure 3 hai dislocation loops obtained in EBR-II at 700 K: 2 (a) 1.1 1025 n m2 and (b) 1.5 1026 n m . Diffracting and beam direction B ¼ 0111 vector g ¼ 1011 Griffiths, M. J. Nucl. Mater. 1988, 159, 190–218.
4.01.1.3 Point-Defect Clusters in Zirconium Alloys
It is now clearly established by numerous authors45–57 that for commercial neutron-irradiated zirconium alloys (e.g., annealed Zircaloy-2 described in Northwood et al.45) at temperatures between 250 and 400 C and for irradiation dose lower than 5 1025 n m2, the point-defect clusters that can be observed by TEM (>2 nm) consist of perfect dislocation loops, either of vacancy or interstitial nature, with Burgers vector ha i ¼ 1=3h1120i, situated in the prismatic planes with typical diameter from 5 to 20 nm, depending on the irradiation temperature (Figures 3 and 4). These loops are found in very high density, typically between 5 1021 and 5 1022 m3 depending on the irradiation temperature (Figure 5).45,51 The three hai Burgers vectors are equally represented. Thorough studies of neutron damage in zirconium using the high-voltage electron microscope (HVEM) have also been given.53,58,59
Figure 4 Typical hai loop microstructure observed on recrystallized Zy-4 irradiated at 280 C in Siloe´ up to a fluence of 6 1024 n m2.
4 N (1022m–3)
4.01.1.3.1 hai Dislocation loops
50 nm
2
(a) 0 7 d (nm)
In the case of zirconium alloys, many authors have studied the postirradiation microstructure by using transmission electron microscopy (TEM). In 1979, an international ‘round robin’ was undertaken consisting of TEM observations of neutron-irradiated recrystallized zirconium alloys45 in order to determine the nature of the point-defect clusters. A more recent compilation of observations is given by Griffiths.46 It has been now proved by numerous authors that in zirconium alloys mainly dislocation loops with hai Burgers vector can be found. Only for high fluence, the hci component dislocation loops appear. Cavities are observed only in very specific cases.
+
+ +
6 5
(b) 4
1019
+
+
+
+
Not visible
+
1021 1020 Neutron fluence (cm–2)
1022
Figure 5 Evolution with dose of the dislocation loops characteristics: (a) density and (b) mean size of defects for Zy-2 irradiated at 300 C. Adapted from Northwood, D. O. Atomic Energy Rev. 1977, 15, 547–610.
8
Radiation Effects in Zirconium Alloys
The proportion of vacancy loops to interstitial loops depends on the irradiation temperature. Indeed, it is observed that for an irradiation temperature of 350 C approximately 50% of observed loops are vacancy loops, whereas for an irradiation temperature of 400 C, 70% of loops are vacancy loops.45,46 For a low irradiation temperature (below 300 C), the majority of loops present in the material are of the interstitial type. The loop habit plane is close to the prismatic plane, but accurate determination proves that the loops are not pure edge but their habit plane is usually closer to the first-order prismatic plane f1010g. The authors have also observed that for loop diameters lower than 40 nm the loops are circular but for diameters larger than 40 nm the vacancy loops become elliptical with the great axis along the hci axis, the interstitial loops remaining circular. The hai loops also appear to be aligned in rows parallel to the trace of the basal plane.46,50 For an irradiation temperature of 300 C, no dislocation loop can be observed below a neutron fluence of 3 1023 n m2 in the case of annealed Zy-2 (Zircaloy-2) irradiated at 300 C.51 However, from this fluence, the loop density increases rapidly with increasing fluence but saturates at a density of 3 1022 m3, from a relatively low fluence of approximately 1 1024 n m2 (Figure 5). The loop density saturation has been confirmed by X-ray analysis.60 The loop size exhibits a parabolic increase with fluence but no clear saturation in the evolution of the loop size is seen even after a fluence of 1 1026 n m2.51,67 Increasing the irradiation temperature leads to a decrease in the loop density and to an increase of the loop size.45,55,61 Indeed, it was shown by Northwood et al.45 that neutron irradiation performed at 350 C of annealed Zy-2 up to a fluence of 1 1025 n m2 leads to a mean loop diameter between 8 and 10 nm and a loop density between 8 1021 and 5 1022 m3; whereas a neutron irradiation of the same alloy performed at 400 C up to a fluence of 1 1025 n m2 leads to a mean loop diameter between 16 to 23 nm and a loop density between 4 1021 and 2 1022 m3.45 Above 500 C, no irradiation damage is formed.52 The hai loop microstructure is found to be very sensitive to alloying elements such as oxygen. Indeed, for highpurity zirconium with very low oxygen content, the hai loops are large and in low density, whereas for commercial zirconium alloys (with oxygen content between 1000 and 1500 ppm) the growth speed of loops is considerably reduced yielding smaller loops in much higher density.45,55
It was also reported from TEM observations that a particular band contrast of alternative black and white was superimposed on the usual radiation damage normally visible on thin foils of irradiated materials. This phenomenon has been connected to the alignment of the loops in the same direction and is believed to be a thin-foil artifact. It has been named ‘corduroy’ contrast by Bell.62 The commonly accepted explanation of this artefact is based on the local elastic relaxation of the internal stresses in TEM thin foils, in areas where pronounced alignment of hai loops is present.63 4.01.1.3.2 hai Loop formation: Mechanisms
The origin for the stability of the hai loops in zirconium is attributed to the relative packing density of the prismatic plane compared to the basal plane, which depends on the c/a ratio of the hcp lattice. have Foll and Wilkens64 p ffiffiffi proposed that when the c/a ratio is higher than 3, loops are formed in the basal plane with Burgers vector 1=6h2023i, whereas if c/a is pffiffiffi lower than 3, then loops are formed in the prismatic plane with Burgers vector ha i ¼ 1=3h1120i. For all hcp metals, this means that loops are formed in the prismatic plane except for Zn and Cd. This is not the case for Zr, Ti, and Mg where loops are also formed in the basal planes, depending on the irradiation dose, irradiation temperature, and purity of the metal.56,57 MD computations for a-zirconium have also shown that most of the small interstitial clusters produced in the cascade have the form of a dislocation loop with Burgers vector ha i ¼ 1=3h1120i. The small vacancy clusters are also found in the prismatic plane.8,28,65 For larger point-defect clusters,66 it is shown that the point-defect clusters in the prismatic plane always relax to perfect dislocation loops with Burgers vector hai ¼ 1=3h1120i. On the other hand, vacancy clusters in the basal plane form a hexagonal loop enclosing a stacking fault with 1=2h0001i Burgers vector. The simultaneous observation of vacancy and interstitial hai loops in zirconium alloys45,48,50,54,61 is a rather surprising feature.53,57 Indeed, as discussed for usual cubic metals, interstitial loops tend to grow under irradiation and the vacancy loops tend to shrink since the edge dislocations are biased toward SIAs due to the EID. According to Griffiths,57 the coexistence of these two types of loops in zirconium can be explained by a modified SIA bias in zirconium due to (i) a relatively small relaxation volume of SIA relative to vacancy (low bias), (ii) interaction with impurities, and (iii) spatial partitioning of vacancy loops and interstitial loops as a result of elastic interactions or
Radiation Effects in Zirconium Alloys
anisotropic diffusion. Other authors53,68 think that this phenomenon is due to a subtle balance of the bias factors of the neighboring point-defect sinks that lead to an increasing bias as the loop size increases if the loop density is high. Woo44 considers that the coexistence of both types of hai loops can be explained in the frame of the DAD model, which induces a strong DAD-induced bias. Indeed, in this model, the hai type loops are shown to be relatively neutral and may therefore receive a net flow of either interstitials or vacancies, depending on the sink situation in their neighborhood. Finally, recent computations,69 using the Monte Carlo method, that take into account the large vacancy and interstitial point-defect clusters created inside the cascade as an input microstructure show that both vacancy and interstitial loops are able to grow simultaneously, the proportion of vacancy loops increasing with increasing irradiation temperature. This last phenomenon can be related to the so-called production bias discussed previously.14 4.01.1.3.3 hci Component dislocation loops
At the time of the thorough review by Northwood,51 no hci component loops had been observed yet. The ‘round robin’ work45 also established that up to an irradiation fluence of 1 1025 n m2 no hci component dislocation loop is observed. As highly irradiated Zircaloy samples became available, for fluence higher than 5 1025 n m2, evidence of hci component loops arose.46,54,70–73,189 The hci component loops have been analyzed as being faulted and of the vacancy type. They are located in the basal plane with a Burgers vector
(a)
9
1=6h2023i having a component parallel to the hci axis (Figure 6). The hci component loops are much larger than the hai loops but their density is much lower. For instance, for recrystallized Zy-2 and Zy-4 irradiated at 300 C, after 5.4 1025 n m2, hci component loops are found with a diameter of 120 nm and with a density between 3 and 6 1020 m3. Whatever the irradiation conditions, these hci component loops are always present in conjunction with more numerous and finer hai loops. The hci component loops can therefore only be observed edge-on by TEM by using the g ¼ 0002 diffraction vector, which leads to invisible hai type defects. The hci loops thus appear as straight-line segments. There is considerable evidence to show that their formation is dependent on the purity of the zirconium used (Figure 6).46,74–76,190 It is also observed that at the beginning of their formation, these dislocation loops appear to be located close to the intermetallic precipitates present in the Zircaloy samples46,76 (Figure 7). By using an HVEM on iron-doped samples, it has been possible to prove that iron enhances the nucleation of the hci loops, the loop density increasing as a function of the iron content. Moreover, iron was found to have segregated in the plane of the loops.76 4.01.1.3.4 hci Loop formation: Mechanisms
It is rather surprising that although the most stable loops are the prismatic loops, basal loops are also observed in zirconium alloys. Moreover, these loops are of the vacancy character. According to the usual rate theory, vacancy loops should not grow as a result of the bias of edge dislocation toward SIAs.
(b)
0.5 mm
Figure 6 Comparison of neutron damage in Zr at 700 K following irradiation to a fluence of 1.5 1026 n m2. (a) Crystal bar purity (500 wt ppm) with no c-component loops. (b) Sponge purity (2000 wt ppm) containing basal hci component in an edge-on orientation (arrowed). Only hci component defects are visible with diffracting vector of [0002]. The beam direction for each micrograph. Adapted from Griffiths, M. Philos. Mag. B. 1991, 63(5), 835–847. is [1010]
10
Radiation Effects in Zirconium Alloys
critical interstitial solute concentration. This volume increases as the interstitial impurity concentration is gradually supplemented by the radiation-induced dissolution of elements such as iron from intermetallic precipitates (or b-phase in the case of Zr–Nb alloys). 4.01.1.3.5 Void formation 500 nm
Figure 7 High density of c-component loops in the vicinity of the precipitates in a Zy-4 sample irradiated to 6 1025 n m2; at 585 K. The arrow shows the diffracting vector [0002]. Adapted from De Carlan, Y.; Re´gnard, C.; Griffiths, M.; Gilbon, D. Influence of iron in the nucleation of hci component dislocation loops in irradiated zircaloy-4. In Eleventh International Symposium on Zirconium in the Nuclear Industry, 1996; Bradley, E. R., Sabol, G. P., Eds.; pp 638–653, ASTM STP 1295.
The reason for the nucleation and growth of the hci component loops in zirconium alloys has been analyzed and discussed in great detail by Griffiths and co-workers.46,56,57,74 The most likely explanation for their appearance46 is that they nucleate in collision cascades, as shown recently by De Diego.66 Their stability is dependent to a large extent on the presence of solute elements, which probably lower the stacking-fault energy of the Zr lattice, making the basal hci component loops more energetically stable. It is also possible that small impurity clusters, especially iron in the form of small basal platelets, could act as nucleation sites for these loops.74,76 However, according to Griffiths,46 this cannot account for the very large vacancy hci component loops observed, since the growth of vacancy loops is not favorable considering the EID discussed previously. In order to understand the reason for the important growth of the hci component loops, another mechanism must occur. As discussed by Woo,44 the growth of hci component loops is well understood in the frame of the DAD model. Indeed, because of the higher mobility of SIAs in the basal plane rather than along the hci axis (and the isotropic diffusion of vacancies), dislocations parallel to the hci axis will absorb a net flux of SIAs whereas dislocations in the basal plane will absorb a net flux of vacancies. This can therefore explain why the basal vacancy loops can grow. The incubation period before the appearance of hci component loops can be explained, according to Griffiths et al.,73 by the fact that the hci loop formation is dependent on the volume of the matrix containing a
Early studies failed to show any cavity in Zr alloys after irradiation.77 From all the obtained data, it is seen that zirconium is extremely resistant to void formation during neutron irradiation (Figure 8).46,52 The effect of very low production of helium by (n, a) reactions during irradiation was mentioned as a possible reason for this absence of voids. But most probably, the fact that in zirconium alloys vacancy type loops are easily formed can be the reason for the absence of void.52 To favor the formation of voids, various studies performed, especially on model alloys, have shown that stabilization of voids can occur when impurities are present in the metal. Helium coming from transmutation of boron on Zr sponge67 as well as impurities located near Fe-enriched intermetallics are found to favor the stability of voids.54 Irradiations with electrons give better conditions to stabilize voids: the main reason is that irradiation doses can be very high – hundreds of displacements per atom can be reached after few hours.190 Moreover, electron irradiation on Zr samples preimplanted with He at various concentrations showed the nucleation and growth of voids only for the samples doped with at least 100 ppm of He.78 4.01.1.4 Secondary-Phase Evolution Under Irradiation 4.01.1.4.1 Crystalline to amorphous transformation of Zr-(Fe,Cr,Ni) intermetallic precipitates
In addition to point-defect cluster formation, irradiation of metals can affect the precipitation state as well as the solid solution. In the case of zirconium alloys, while investigating the effect of irradiation on corrosion, TEM observations revealed that for Zircaloy, irradiated at temperatures typical for commercial light water reactors (lower than 600 K), Zr(Fe,Cr)2 precipitates began to become amorphous after a fluence of about 3 1025 n m2. Interestingly, the other common precipitate in Zy-2, Zr2(Fe, Ni), remained crystalline up to higher irradiation doses.77 The instability of these precipitates under irradiation is of great importance since the secondary-phase precipitate plays a major role on
Radiation Effects in Zirconium Alloys
(a)
11
(b) 0.1 mm
(c)
(d)
Figure 8 Examples of radiation-induced cavities in zirconium alloys. (a) Annealed crystal-bar zirconium, prism foil, 673K, 1.21025n/m2; (b) annealed zircaloy-2, prism foil, 673K, 1.21025n/m2; (c) annealed Zr-2.5 wt% Nb, basal foil, 923K, 0.71025n/m2; (d) typical cavity attached to inclusion on a grain boundary, material (c). Adapted from Gilbert, R. W.; Farrell, K.; Coleman, C. E. J. Nucl. Mater. 1979, 84(1–2), 137–148.
the corrosion resistance of Zircaloy (see Chapter 5.03, Corrosion of Zirconium Alloys). The effect of temperature on the crystalline to amorphous transformation has been studied by various authors.75,79–83 It is shown that at low temperatures (353 K), under neutron irradiation, both Zr(Fe, Cr)2 and Zr2(Fe, Ni) undergo a rapid and complete crystalline to amorphous transformation. As the irradiation temperature increases, a higher dose is required for amorphization. It is indeed seen that, at 570 K, Zr(Fe,Cr)2 precipitates undergo only a partial amorphous transformation and Zr2(Fe,Ni) particles remain crystalline (Figure 9). It is also observed that the crystalline to amorphous transformation starts at the periphery of particles, and then the amorphous rim moves inward until the whole precipitate becomes fully amorphous. The chemical concentration profile within the precipitates also exhibits two distinct zones corresponding to the two different states: the crystalline core and the amorphous periphery. It is observed that the amorphous layer exhibits a much lower iron
(a)
(b)
0.1 mm
Figure 9 Crystalline to amorphous transformations of Zr (Cr, Fe)2 particle in Zy-4 irradiated in a BWR at 560 K: (a) 3.5 1025 n m2 and (b) 8.5 1025 n m2. Adapted from Griffiths, M.; Gilbert, R. W.; Carpenter, G. J. C. J. Nucl. Mater. 1987, 150(1), 53–66.
12
Radiation Effects in Zirconium Alloys
Counts ⫻ 101
500 400
FeKa
300 CrKa 200 100 0 350
450
650
550
750
850
Energy (eV)⫻10–1
Counts ⫻ 101
500 400 CrKa
300 200
FeKa
100
0.1 mm
0 350
450
550
650
750
850
Energy (eV)⫻10–1
Figure 10 Crystalline to amorphous transformations of Zr(Cr, Fe)2 particle in Zy-4 irradiated at 560 K at 3.5 1025 n m2. EDX spectrum shows that the amorphous volume is coincident with a depletion of Fe. Adapted from Griffiths, M.; Gilbert, R. W.; Carpenter, G. J. C. J. Nucl. Mater. 1987, 150(1), 53–66.
content than the precipitate, the iron profile showing a local drop from the standard value of 45 at.% to below 10 at.% (Figure 10). At higher temperatures (T > 640 K), amorphization was not detected and the precipitates remain crystalline, but some authors79 have nevertheless observed loss of iron and even total dissolution of Zr2(Fe, Ni) and Zr(Fe, Cr)2 precipitates and redistribution of alloying elements. The crystalline to amorphous transformation is easily understood in terms of ballistic radiation-induced disordering at a temperature where recombination of point defects or recrystallization within the intermetallic precipitate is too slow to compensate for the rate of atomic displacement (at 350 K).79 The dissolution of alloying elements remains limited at this low temperature and the amorphization is mainly due to sputtering, that is, transfer of material from the particle because of atomic displacements by neutrons. When the pointdefect concentration becomes too high and/or when the chemical disordering is too high, the crystalline structure is destabilized and undergoes a transformation to an amorphous phase.75,79 The fact that the Zr2(Fe, Ni) phase remains crystalline at intermediate temperatures (520–600 K) is presumably due to a more rapid reordering than the disordering in this structure (Zintl phase structure).
Concerning the Zr(Fe, Cr)2 (Laves phase structure), it is seen that the amorphization starts at the precipitate–matrix interface forming a front that gradually moves into the precipitate. The amorphization is believed to happen by a deviation from stoichiometry due to a ballistic interchange of iron and zirconium atoms across the precipitate–matrix interface. It also agrees with the observed kinetics of amorphization, predicting an amorphous thickness proportional to fluence and the absence of an incubation period for the transformation to start.84 The reason for the depletion of iron from the precipitates is not clearly understood yet, according to Griffiths et al.79 It is suggested that iron may be in some form of irradiation-induced interstitial state in irradiated Zr-alloys and may then diffuse interstitially out of the intermetallic particles. At high temperatures (640–710 K), corresponding to 0.3Tm, the thermal activation is sufficient to induce dynamic recrystallization impeding the amorphization of the precipitates. However, depletion and some precipitate dissolution would still occur, but the level of damage necessary for amorphization would not be reached due to the absence of cascade damage.84 Because of the high mobility of Fe and Cr, redistribution of solute can occur, leading to secondary-precipitate formation.
Radiation Effects in Zirconium Alloys
4.01.1.4.2 Irradiation effects in Zr–Nb alloys: Enhanced precipitation
In binary Zr–Nb alloys (Zr–1% Nb and Zr–2.5% Nb), the microstructure is usually in a metastable state due to the thermomechanical processing in the upper a range or in the a þ b domain. Indeed, at this relatively low temperature (around 580 C), the atomic mobility is low and the equilibrium state cannot be reached in reasonable time. After cooling, the matrix is therefore supersaturated in Nb and the composition of secondary phases (Nb rich) still corresponds to the high-temperature chemical composition. It is indeed shown by Toffolon-Masclet et al.85 that a Zr–1% Nb–O alloy that has undergone a final heat treatment at 580 C for a few hours can still evolve toward its thermodynamic equilibrium after 10000 h of heat treatment at 400 C. Under irradiation, it is observed that the microstructure of Zr–Nb alloys is not stable and very fine Nb-rich precipitates, with diameter of a few nanometers, are observed in very high density (Figure 11). This precipitation of Nb from the supersaturated matrix is observed in any type of binary alloys: in Zr–1% Nb such as M5™(86) and E110(12,87) as well as Zr–2.5% Nb.88 This needle-like precipitation has been studied mainly by TEM, and also by small angle neutron scattering (SANS) analyses.86
13
Simultaneously, a noticeable decrease of Nb content in the matrix occurs.89 This precipitation is due to an enhanced mobility of Nb atoms under irradiation due to the very high vacancy concentration created by irradiation. This enhances the Nb mobility and allows the rapid evolution of the microstructure toward its thermodynamic equilibrium, leading to precipitation of very fine Nbrich precipitates in Zr–Nb binary alloys. In Zr–Nb alloys, the Nb-rich phases also undergo chemical changes under irradiation. Indeed, it is shown that the o phase, obtained in Zr–2.5Nb by transformation of the b-Nb after extrusion, disappears and transforms into b-Nb.60 For the b-Nb phase and in the case of M5™ alloys, an evolution of the chemical composition under irradiation has also been observed, but the b-Nb precipitates still remain fully crystalline even after six PWR cycles of irradiation (70 GWd t1). Only a decrease in Nb content with a small increase in the size of the precipitates has been noticed86 (Figure 11). The same has been obtained for E110 and E635 Russians alloys, where b-Nb precipitates are altered in composition to reduce the Nb content from 85–90% to 50%.12 Moreover, for the Zr(Nb, Fe)2 Laves phases with hcp structure found in E635 and E110 alloys, it seems that a release of iron atoms into the matrix from the
(a)
(b)
(c)
(d)
(e)
(f) 100 nm
Figure 11 Micrographs of needle-like radiation-enhanced precipitation: (a) M5™ 2.1 1025 n m2, (b) Zr–1% NbO 2.8 1025 n m2, (c) M5™ 3.6 1025 n m2, (d) Zr–1% NbO 5.7 1025 n m2, (e) Zr–1% NbO 8.2 1025 n m2, and (f) M5™ 13.1 1025 n m2. Reprinted, with permission, from J. ASTM Int., copyright ASTM International, 100 Barr Harbor Drive, West Conshohocken, PA 19428.
14
Radiation Effects in Zirconium Alloys
precipitates has occurred after irradiation, leading to the transformation into b-Nb particles with bcc structure.12,89
4.01.2 Postirradiation Mechanical Behavior 4.01.2.1 Mechanical Behavior During Tensile Testing 4.01.2.1.1 Irradiation hardening: Macroscopic behavior
As for many other metals, zirconium alloys exhibit strong hardening after neutron irradiation. It is indeed observed by numerous authors90–99 and reviewed21,77,100 that the yield stress (YS), as well as the ultimate tensile strength (UTS), of both recrystallization-annealed (RXA) and stress-relieved annealed (SRA) zirconium alloys is strongly increased by neutron irradiation (Figures 12 and 13). Microhardness tests also prove this phenomenon.101–105 The irradiation-induced hardening increases rapidly for fluences below 1 1024 n m2 (E > 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
Radiation Effects in Zirconium Alloys
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
Radiation Effects in Zirconium Alloys
4.01.2.1.3 Post-yield deformation: Macroscopic behavior
Concerning the mechanical behavior beyond the YS, it is pointed out by several authors97,115,116 that for RXA zirconium alloys, the strain hardening rate is higher after irradiation at the onset of plastic flow but decreases rapidly with the plastic strain, more rapidly than before irradiation, resulting in a low strain 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.; Be´chade, 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
Radiation Effects in Zirconium Alloys
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 Be´chade,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
Radiation Effects in Zirconium Alloys
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
Radiation Effects in Zirconium Alloys
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
Radiation Effects in Zirconium Alloys
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 (<0.02%) and saturates at fluences less than 5 1024 n m2.161,164 Eventually, at higher fluence a growth breakaway (increase of the growth rate) occurs for the annealed zirconium single crystal.161 Since the deformation of the polycrystalline cladding is the result of the growth of all the grains, texture has a major influence on the growth of the polycrystalline material. A weakly textured product made of zirconium alloy, with Kearns factors close to fd 0.33 along the three directions, such as b-quenched zirconium alloys165,166 as reviewed by Fidleris,150 exhibits a very low growth. The Kearns factor fd is the resolved fraction of basal poles along the direction d. On the other hand, strongly textured products, with most grains orientated with hci axis along one given macroscopic direction (high Kearns factor, fd > 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
Radiation Effects in Zirconium Alloys
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
Radiation Effects in Zirconium Alloys
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
Radiation Effects in Zirconium Alloys
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
Radiation Effects in Zirconium Alloys
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
Radiation Effects in Zirconium Alloys
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
Radiation Effects in Zirconium Alloys
However, it has been seen that for growth, the anisotropic diffusion of SIAs is believed to play an important role in the deformation mechanism. 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
Radiation Effects in Zirconium Alloys
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.; Be´chade, 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
Radiation Effects in Zirconium Alloys
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.; Be´chade, 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.; Be´chade, 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.; Be´chade, 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.; Be´chade, 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.; Be´chade, 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.
31
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
ß 2012 Elsevier Ltd. All rights reserved.
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
Radiation Damage in Austenitic Steels
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
Radiation Damage in Austenitic Steels
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
Radiation Damage in Austenitic Steels
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.
Radiation Damage in Austenitic Steels
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
Radiation Damage in Austenitic Steels
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
Radiation Damage in Austenitic Steels
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
Radiation Damage in Austenitic Steels
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
Radiation Damage in Austenitic Steels
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
Radiation Damage in Austenitic Steels
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
Radiation Damage in Austenitic Steels
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
Radiation Damage in Austenitic Steels
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
Radiation Damage in Austenitic Steels
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
Radiation Damage in Austenitic Steels
(a)
CW 316 SS, thimble tube 70 dpa, 315 ºC
50 nm
(b)
CW 316 SS, thimble tube 33 dpa, 290 ºC
50 nm (c)
CW 316 SS, thimble tube 33 dpa, 290 ºC
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
Radiation Damage in Austenitic Steels
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
Radiation Damage in Austenitic Steels
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
Radiation Damage in Austenitic Steels
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 <3 nm in both the alloy matrix and especially on grain boundaries. The measured concentrations of 600 appm He and 2500 appm H in this specimen are thought to reside primarily within the cavities. Most importantly, these cavities are essentially invisible
50
Radiation Damage in Austenitic Steels
4.02.7 Radiation-Induced Changes in Mechanical Properties
under well-focused imaging conditions and can only be imaged using very large levels of under-focus. This implies that previous studies on similar materials may have overlooked such cavity-dominated structures. When this specimen and near-identical specimens were subjected to slow strain rate testing after irradiation, the fracture surface was indicative of 100% intergranular stress corrosion cracking (IGSCC), with lower doses and gas levels producing proportionally less IGSCC.77 As hydrogen is known to be a contributor to grain boundary cracking, it appears plausible that hydrogen storage may accelerate the cracking process and that higher exposure will lead to an increasing susceptibility to cracking. This issue may therefore become increasingly important as PWRs previously licensed for 40 years are being considered for life extension to 60 and possibly 80 years.
Long before the onset of significant phase evolution or void swelling is observed, the first manifestation of the radiation-induced microstructural/microchemical evolution appears in changes of the mechanical properties. As shown in Figure 18 the stress–strain diagrams of stainless steels begin to change significantly even at very low dpa levels. The strength of the alloy increases, the elongation decreases, and there is a progressive decrease in work-hardening. This behavior is dependent somewhat on test temperature but is not very sensitive to neutron spectrum. Movement of dislocations in metals during deformation following irradiation is impeded by the microstructural components produced by radiation (dislocations, dislocation loops, voids, bubbles, precipitates) and therefore the strength of annealed steel 1000
1000
316 SS tested at 288 C
316 SS tested at RT 800
600 Unirradiated 0.0001 dpa 0.001 dpa 0.01 dpa 0.1 dpa 0.78 dpa
400 200 0 0.0
(a)
Eng. stress (MPa)
Eng. stress (MPa)
800
0.2
0.4 0.6 Eng. strain
400 Unirradiated 0.01 dpa 0.1 dpa
200 0 0.0
1.0
0.8
600
0.2
0.4 0.6 Eng. strain
(b)
0.8
1.0
1000
Eng. stress (MPa)
800
EC316LN tested at RT 10.7 dpa 3.64 dpa 0.86 dpa 0.45 dpa
600
2.53 dpa 1.36 dpa
0.0 dpa
400 200 0 0.0
(c)
0.2
0.4 0.6 Eng. strain
0.8
1.0
Figure 18 Engineering stress–strain curves for irradiated austenitic stainless steels: (a) annealed 316 SS irradiated in HFIR mixed spectrum reactor at 60–100 C and tested at 25 C, (b) annealed 316 SS irradiated in HFIR at 350 C and tested at 288 C, and (c) annealed EC316LN irradiated in the LANSCE spallation neutron and proton spectrum at 60–100 C and tested at 25 C. Reproduced from Kim, J. W.; Byun, T. S. J. Nucl. Mater. 2010, 396, 10–19.
Radiation Damage in Austenitic Steels
increases. The strength increase usually saturates at relatively low exposure levels (<10 dpa) as shown in Figure 19, reflecting a similar saturation of microstructural densities. Since the concentration of most radiation-induced microstructural components decreases with increasing temperature above 300 C, one would expect that the saturation strength would also decrease with increasing temperature, as is shown in Figures 20–23.
1000 316LN 316
Yield strength (MPa)
800 304 600 316
304, 304L
400
316, 316LN PCA
316 200 Unirradiated 304, 316, PCA alloys 0 0.0001
0.001
0.01 0.1 1 Neutron dose (dpa)
10
100
Figure 19 Strengthening of various annealed 300 series stainless steels versus dpa in various water-cooled reactors at relatively low temperatures (280–330 C). Reproduced from Pawel, J. P.; Ioka, I.; Rowcliffe, A. F.; Grossbeck, M. L.; Jitsukawa, S. In Effects of Radiation on Materials: 18th International Symposium; ASTM STP 1325; 1999; pp 671–688. At these temperatures strengthening saturates at 10 dpa.
Frank loops
Yield strength (MPa)
1400
Cavities, precipitates
1200 1000
Defect clusters
800 600 400 200 0 0
Unirradiated 100
200 300 Temperature ( C)
400
500
Figure 20 Radiation-induced strengthening of annealed 300 series steels versus irradiation temperature and the microstructural components causing the strengthening. Note the peak strengthening at 300 C followed by a decline at higher temperatures. Reproduced from Pawel, J. P.; Ioka, I.; Rowcliffe, A. F.; Grossbeck, M. L.; Jitsukawa, S. In Effects of Radiation on Materials: 18th International Symposium; ASTM STP 1325; 1999; pp 671–688.
51
Irradiation of cold-worked steels also leads to strengthening at lower temperatures but softening can occur at higher temperatures if the saturation strength level at a given temperature is below the starting strength, as seen in Figure 21. Most importantly, both annealed and cold-worked steels converge to the same saturation level when irradiated at the same dpa rate and temperature as seen in Figure 22.78 Similar convergence behavior has been observed in the evolution of microhardness.79 Note also that radiation-induced changes in strength are roughly independent of composition within the annealed 300 series stainless steels, especially at lower irradiation temperatures, as shown by Figure 19. Such convergence behavior has been observed many times, but there are exceptions; for example, cold-worked steels converge in their notch tensile strength, but not to the level reached by annealed steels.80 Such behavior is usually observed in steels that twin heavily during deformation and were irradiated at low temperatures that resist recrystallization. Twin boundaries are not easily erased by displacements, so their hardening contribution persists. Concurrent with an increase in radiation-induced hardening is a loss of ductility,81–83 as shown in Figures 23 and 24. The concept of saturation or persistence of mechanical properties, especially with respect to temperature, applies to the most recent irradiation temperature, as demonstrated by comparing isothermal and nonisothermal histories. In Figure 25 the mechanical properties of three model alloys are seen to converge during isothermal irradiation without being affected by composition, He/dpa ratio, and mechanical starting state.84 In Figure 26, however, an early detour in temperature led to differences from isothermal behavior, but these differences disappeared when the intended isothermal temperature was reestablished.84 Previous saturation states are soon forgotten, usually by 5 dpa, but only if the hardening components are easily erased and replaced at the new temperature. If hardening arises primarily from dislocation loops and dislocations, this condition is easily met. If the primary hardening arises from a fine density of voids and especially bubbles produced at lower temperatures, then the microstructural memory cannot be easily erased, even at much higher temperatures. An example is shown in Figure 27 where a series of Fe–Cr–Ni ternary austenitic alloys were irradiated at 400 and 500 C in ORR at high He/dpa ratios
52
Radiation Damage in Austenitic Steels
1100 1000 900
371 C
Yield strength (MPa)
800
427 C
700 483 C 600 500 538 C
400
593 C 300
649 C 704 C
200
760 C
100 0
816 C 0
1
2
3
4
Neutron fluence
5
6
(n cm–2)
7
8
9
10 1022
(E > 0.1 MeV)
Figure 21 Evolution of yield strength in 20% cold-worked 316 stainless steel irradiated in EBR-II over a wide range of temperatures. Reproduced from Garner, F. A.; Hamilton, M. L.; Panayotou, N. F.; Johnson, G. D. J. Nucl. Mater. 1981, 103 and 104, 803–808.
600 650 C
20% Cold-worked
400 200
Annealed 0 538 C 20% Cold-worked
Yield strength (MPa)
600 400 200
Annealed
0 427 C 800 600
20% Cold-worked
400
Annealed
200 0 0
1
2
3
4
5
6 22
7
8
9
10
–2
Neutron fluence (10 n cm ) Figure 22 Influence of temperature and neutron exposure on evolution of yield strength in both annealed and 20% cold-worked AISI 316 irradiated in EBR-II, showing that the saturation strength level is independent of starting condition, converging at doses of 5–15 dpa. Reproduced from Garner, F. A.; Hamilton, M. L.; Panayotou, N. F.; Johnson, G. D. J. Nucl. Mater. 1981, 103 and 104, 803–808.
(27–58 appm dpa1) and 395 and 450 C in EBR-II at very low He/dpa ratios (0.7–1.2 appm dpa1).85 Note that there are very significant differences in hardening observed between the two experiments and that the differences arose primarily from a very large difference in cavity density, a difference that was too large to be explained in terms of helium content alone. It was later shown that the ORR experiment suffered a very large number (237 over 2 years) of unrecognized negative temperature setbacks of 1–2 h, with decreases varying from 50 to 500 C.86 Even though the total dpa accumulated during these setbacks was only 1% of the total dose, the frequent bloom of high densities of small Frank loops at lower temperatures provided a very large periodic increase in nucleation sites for helium bubbles on the new Frank loops that significantly strengthened the matrix. The loops could subsequently dissolve but the bubbles could not. In addition to temperature, the most prominent irradiation variable is the dpa rate and it is known that the microstructural densities, especially Frank loops and voids, are known to increase in concentration as the dpa rate increases. Various radiation-stable phases such as w0 are also known to be flux-sensitive, while other phases such as carbides and intermetallics are more time-sensitive.1 Thus, it is not surprising that some sensitivity to dpa rate might be observed in strength properties, as
Radiation Damage in Austenitic Steels
53
Test temperature = Irradiation temperature 1200
40
Yield strength (MPa)
800
400 C
450–500 C
600
525–565 C
400
Uniform elongation (%)
215–245 C
1000
20 525–565 C
10 5
2
400 C
1 585–625 C
200 0
1
2
3
215–245 C
0.5
5 ´ 1026
4
0
1
2
3
5 ´ 1026
4
Neutron fluence (n m-2) E > 0.1 MeV
Neutron fluence (n m-2) E > 0.1 MeV
Figure 23 Neutron-induced changes in tensile properties of annealed 1.4988 stainless steel irradiated in the DFR fast reactor. Reproduced from Ehrlich, K. J. Nucl. Mater. 1985, 133–134, 119–126. Ductility declines as strength increases.
50 20% CW 316 25% CW PCA
40 Uniform elongation (%)
Yield stress (MPa)
1200
800
400
0
SA 316L SA PCA
SA 316L SA PCA 20% CW 316 25% CW PCA
30
20
10
0 0
2
4
6 dpa
8
10
0
400 800 Yield stress (MPa)
1200
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
Radiation Damage in Austenitic Steels
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
Total elongation (%)
40
25 Ni + 0.04 P
25 Ni
45 Ni
30 Annealed 20 Cold-worked
10
0
0
10
20
0
10
20
0
10
20
30
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
Radiation Damage in Austenitic Steels
55
800
Yield strength (MPa)
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
Total elongation (%)
0
40
25 Ni
30
Isothermal repeat series
45 Ni
25 Ni + 0.04 P
Original series
20
10
0
0
20
40
0
40
20
0
20
40
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
Radiation Damage in Austenitic Steels
25
1024
45
700 600 MFE-4
500 400 300
2.3
2.1
400 C
2.4
1023
1.9 9.4
5.5
5.6
3.4
1022 23 nm
1021
1020
AD-1
100 400
500
600
Temperature (C)
2.9
500 C
23
395 C
35
200
0 300
MFE-4 experiment in ORR
1.9 nm
Cavity density (m-3)
Average DYS (MPa)
800
Ni 35
7 Cr 15 20
1019
25
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. Phe´nix 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
Radiation Damage in Austenitic Steels
57
1200
700
370 C
AISI 304
500
Ultimate
300 Yield 100 0.1
E, MeV
odf, dpa s–1
Ti, C
0.75 0.53 0.29 0.29 0.19 0.17
7.9 ´ 10-7 3.9 ´ 10-7 1.8 ´ 10-7 1.5 ´ 10-7 0.8 ´ 10-7 0.6 ´ 10-7
392 376 373 426 371 371
1 10 Exposure (dpa)
(a)
Strength (MPa)
Strength (MPa)
1000
100
800
600 Yield strength 400
200
0
Strength (MPa)
700
AISI 316
300 Yield
100 0.1 (b)
1
E, MeV
odf, dpa s–1
0.76 0.63 0.38 0.35 0.21 0.17
8.4 ´ 10-7 5.1 ´ 10-7 2.3 ´ 10-7 1.9 ´ 10-7 1.0 ´ 10-7 0.6 ´ 10-7
10 Exposure (dpa)
0
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 <10%. In some observations voids encased in austenite shells have been observed to exist in a pure ferrite matrix.98,99 To date, however, no significant component failure has been reported to result from this particular late-term instability. Finally, there appears to be another late-term phase instability developing at lower irradiation temperatures that involves martensite but does not appear to be due to void swelling. Gusev et al. have shown that for irradiation temperatures below 350 C a growing tendency for stress-induced martensite formation is occurring in Russian austenitic steels at doses in the range of 25–55 dpa when tested at room temperature.100–102 Surprisingly, this instability results in a restoration of engineering ductility to preirradiation levels. However, the ductility is
Radiation Damage in Austenitic Steels
2.8 ´ 1022 n cm–2
Unirradiated plastic dimpling
100
Intermediate 10.7 ´ 1022 n cm–2 mechanism
Channel fracture Uniform elongation Proportional elastic limit Yield strength Symbols Open-5C3 CRT Crossed-5A3 CRT Closed-3A1 SRT
Yield strength
Strength (ksi)
80 Proportional elastic limit
60
3
EBR-II 304 SS Irradiated at 700 F Tested at 700 F
40
2
Uniform elongation
20
1 0
0 0
1
2
3
4
5
6
7
8
9
10
Uniform elongation (%)
58
11
Fluence, ´1022 n cm–2 (E > 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)
X-strength without elongation 800
6,5 Ttest = 20 C
600
6,2
10,
18, 21,
18, 400
Ttest = Tirr. 2,7 0,6
17,
23,
0,8 25, 15, Ttest = 800 C
200 22,
11, 0 −50
26, 0
7,8
1,8
50 100 150 200 250 Position from core central plane (mm)
300
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.
Radiation Damage in Austenitic Steels
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
Radiation Damage in Austenitic Steels
s0.2 sUTS
1.0
0.8
Angle of fracture
0.6 90
Test at 20 C
60
Test at irradiation temperature 30 0
10 20 Swelling (%)
800
30
15
Test at 20 C
UE (%)
UTS (MPa)
600
400 Test at irradiation temperature
200
Test at 20 C
10
Test at irradiation temperature 5
, Nil ductility 0
0
10
20 30 Swelling (%)
40
0
0
10
20 30 Swelling (%)
40
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 (<1 dpa) where there is essentially no loss of toughness, and the second regime involves an intermediate exposure range (1–10 dpa) where toughness decreases rapidly with exposure, producing an order of magnitude reduction in Jc and two orders of magnitude degradation in tearing modulus. Finally, a saturation regime is reached, in which increasing exposure does not produce a further reduction in toughness. This saturation occurs well before any of the void-induced instabilities discussed above can occur. As shown in
Figure 47, the saturation level is remarkably independent of the original toughness level. Welds in austenitic alloys were shown by Mills to exhibit lower initial toughness values and lower saturation toughness levels as well. The fracture toughness level is sensitive to the test temperature, however, as shown in Figure 48. At high test temperatures, the fracture mode changes from transgranular to intergranular in nature, reflecting the effect of test temperature on both matrix strength and also the influence of helium embrittlement at grain boundaries.105 The level of helium needed to promote high temperature embrittlement is not very high, however, and can easily be reached after moderate neutron exposure in fast reactor-irradiated alloys with the lowest nickel level.
61
Radiation Damage in Austenitic Steels
(a)
16185
16049
(b)
422
0.1 µm
200
0.1 µm
Figure 37 (a) Bright field image of voids and deformation bands observed in a highly embrittled 20% cold-worked 316 hexagonal duct. (b) Dark field image showing a high density of thin stacking fault platelets of epsilon-martensite on one of the four sets of close-packed planes. Reproduced from Hamilton, M. L.; Huang, F. H.; Yang, W. J. S.; Garner, F. A. In Effects of Radiation on Materials: 13th International Symposium (Part II) Influence of Radiation on Material Properties; ASTM STP 956; 1987; pp 245–270.
4.02.8 Radiation-Induced Changes in Dimension
1000
100
OKh16N15M3B
UTS (MPa)
10 Some ductility retained Nil ductility 1 1000
100
10 0.01
Kh18N10T
0.1
1.0
10
Swelling (%)
Figure 38 Influence of swelling on ultimate tensile strength of 0Khl6N15M3B cladding and Kh18N10T hexagonal ducts irradiated in BOR-60. Tests were performed at the irradiation temperature, using specimens cut from regions of maximum swelling. Reproduced from Neustroev, V. S.; Shamardin, V. K. Phys. Met. Metallogr. 1997, 83(5), 555–560. The two steels develop loss of strength with swelling differently, probably reflecting the very different precipitate structures of the two steels.
One of the most challenging engineering consequences of neutron irradiation is the development of dimensional instability, whereby a structural component can shrink or grow in volume and where it can be distorted in shape, often with both processes occurring at the same time. There are two major categories of such changes: conservative of volume and nonconservative of volume. A distinction can also be made between processes that distribute the resulting strains isotropically or anisotropically. Additionally, a further distinction can be made concerning whether the process to the first-order is stress-driven or not, or whether it is stress-sensitive to the second-order. Depending on the crystal structure there are a variety of such distortion processes, some more prominent than others in a given crystal system. For austenitic stainless steels the phenomenon of radiation-induced growth (volume-conservative, anisotropic distribution of strains in the absence of stress) is not an issue, whereas for hexagonal close packed alloys based on zirconium and rhenium growth is often a dominant process.9,106 Austenitic steels also
62
Radiation Damage in Austenitic Steels
80 Test temperature = 180 C
Total elongation
20
Uniform elongation
10
0
0
5 10 Swelling (%)
Test temperature = 180 C Total absorbed energy (J cm−2)
Elongation (%)
30
60 Wrapper #1–low dose Wrapper #2–high dose
40
20
0
0
5 10 Swelling (%)
Figure 39 (Left) Correlation between ductility loss and swelling in several heats of irradiated Ti-modified steels in PHENIX. At 5% swelling the total and uniform elongations converge and by 10% no ductility remains. (Right) Correlation of swelling and embrittlement in Charpy impact tests of cold-worked Ti-modified 316 steel irradiated in PHENIX. Reproduced from Fissolo, A.; Cauvin, R.; Hugot, J. P. Levy, V. In Effects of Radiation on Materials: 14th International Symposium; STP 1046; 1990; Vol. II; pp 700–713.
Figure 41 Void-induced embrittlement of an annealed 304 steel EBR-II assembly duct after 54 dpa at 400 C. Reproduced from Flinn, J. E.; Krajcinovic, D.; Phipps, R. D.; Franklin, D. G.; Miller, S. C. Evaluation of Ex-Reactor Loading Event on High-fluence EBR-II Control-rod Thimble 5E3, ANL/EBR-068, February 1973. The duct broke during routine handling in the hot cell. Figure 40 Failure during mounting in a vise of severely void-embrittled 316 stainless steel creep tube irradiated in the EBR-II fast reactor to 130 dpa at 400 C with a hoop stress of 276 MPa. Reproduced from Porter, D. L.; Garner, F. A. J. Nucl. Mater. 1988, 159, 114–121. Swelling at the initial failure point was 14%.
are not very prone to significant transmutationinduced changes in lattice parameter as sometimes observed in alloys based on rhenium and vanadium.106,107 See also Chapter 4.01, Radiation Effects in Zirconium Alloys. Stainless steels experience three general categories of radiation-induced strain processes. These are precipitation-related strains, void swelling, and
irradiation creep. In general, these three processes are not fully independent but are interrelated and often synergistic. 4.02.8.1
Precipitation-Related Strains
Stainless steels undergo an evolution of phase structure at reactor-relevant temperatures, even in the absence of radiation. These changes involve the formation of various carbides, later followed by various intermetallic phases.1,108 This evolution is accompanied by net changes in average lattice parameter arising from differences in partial molar volume of elements when passing from one phase to another.
Radiation Damage in Austenitic Steels
By-97
63
By-92
53 dpa 27.8%
52 dpa 29.8% U-796
50 34 dpa max
50 100 200 Annealed tubes
200 MPa M
14% swelling Figure 42 Severe embrittlement and failure in three BOR60 reflector assembly ducts. The ducts were made of annealed X18H10T, the Russian equivalent of 321 steel. Reproduced from Neustroev, V. S.; Ostrovsky, Z. E.; Teykovtsev, A. A.; Shamardin, V. K.; Yakolev, V. V. In Proceedings of 6th Russian Conference on Reactor Materials Science; 11–15 September 2000, Dimitrovgrad, Russia, in Russian. The maximum swelling values (from left to right) were 27.8, 29.8, and 14%. Failure was the result of high withdrawal loads arising from both swelling and bending, the latter a consequence of radial dpa gradients in the reflector.
0
50 100 200 Cold-worked tubes
200 MPa
Figure 44 EI-847 pressurized tubes irradiated to 75 dpa, 330–342 C (top) and 82 dpa, 365 C, (bottom) in BN-350. Reproduced from Porollo, S. I.; Vorobjev, A. N.; Konobeev, Yu. V.; Dvoraishin, A. M.; Krigan, V. M.; Budylkin, N. I.; Mironova, E. G.; Garner, F.A. J. Nucl. Mater. 1998, 258–263, 1613–1617. All tubes lost pressure, either by cracking or by completely failing during removal from their canister. Before breaking, the tubes were also bent by irradiation creep due to swelling-induced interaction with the top of the canister. Swelling of 6.2% was measured in the zero stress, annealed tube and 11.2% in the cold-worked zero stress tube.
~30–35% of martensite Figure 43 Failure of 20% cold-worked D9 (Ti-modified 316) cladding during routine handling. Failure occurred where 90 dpa was attained at 460 C in FFTF, producing 32% swelling. Reproduced from Makenas, B. J.; Chastain, S. A.; Gneiting, B. C. , ‘‘Dimensional Changes in FFTF Austenitic Cladding and Ducts, Westinghouse Hanford Company Report WHC-SA-0933VA, Richland WA, 1990. Fuel was lost from the open section.
The resulting macroscopic strains are sometimes very counterintuitive, however, especially with respect to their sign. For example, formation of the less dense carbide phases leads to macroscopic densification of the alloy and shrinkage of volume,109 while the formation of denser intermetallic phases (Chi, Sigma, Laves) usually leads to an increase in volume, a form of nonvoid swelling.110,111 This counterintuitive behavior is the result of the different partial molar volumes of critical elements (C and Mo primarily) between the new
Figure 45 Deformation at room temperature of the Russian analog of AISI 321 following irradiation in BN-600 to 55 dpa at 310 C. Distortion of painted circular dots shows where the deformation wave has passed, moving toward the left. The specimen was cut from a hexagonal duct of a fuel assembly. Reproduced from Gusev, M. N.; Maksimkin, O. P.; Garner, F. A. J. Nucl. Mater. 2010, 403, 121–125.
precipitates and the alloy matrix in which they form. Both the carbide and intermetallic phase evolution appear to be accelerated and sometimes altered under irradiation. Other radiation-produced phases (w0 , G-phase) also appear to induce changes in lattice parameter
64
Radiation Damage in Austenitic Steels
Volume fraction of martensite (%)
35 30 25 20 15 10 5 0 0
2
4
6
8
10
12
14
16
Test portion length (mm)
Figure 46 Deformation-induced martensite (vol.%) produced at 20 C in the Russian analog of AISI 321 following irradiation in BN-600 to 26 dpa at 423 C. Reproduced from Gusev, M. N.; Maksimkin, O. P.; Garner, F. A. J. Nucl. Mater. 2010, 403, 121–125. The leading edge of the wave was moving from right to left and was at 2–3 mm when the test was interrupted.
1070 kJ m–2 600 Irradiated temperature = 400–427 C Test temperature = 427 C 500 316 SS 304 SS Annealed Inconel 600 Inconel 800 CF8 SS as-cast 308 SS weld
Jc (kJ m–2)
400
300
200
Base metal
100
0
Weld
0
5
10 15 Neutron exposure (dpa)
20
25
Figure 47 Irradiation-induced evolution of fracture toughness Jc in various austenitic steels and welds. Reproduced from Mills, W. J. ‘‘Irradiation Effects on the Fracture Toughness of Austenitic Fe-Cr-Ni Alloys,’’ Hanford Engineering Development Laboratory Report HEDL-TME-82–17, Richland, WA, 1982; Mills, W. J. Nucl. Technol. 1987, 82, 290–303.
but these have not been well characterized, primarily because these phases develop concurrently with void swelling that masks their contribution.1 Garner1 provides a review of precipitation-induced strains. For the current purpose it is sufficient to note that carbide-induced densification increases with carbon content and with increasing irradiation temperature. Such volume changes for the most
common carbon levels range from 0.1% to 0.4% decrease in volume. The resulting strains may or may not be isotropically distributed, depending on whether there is a pronounced starting dislocation texture on which the carbides nucleate. This process is most pronounced for titanium carbides in Ti-stabilized steels. Carbide-induced strains usually develop quickly enough to be measurable before swelling strains become dominant and therefore are relatively easy to identify compared to those of slower forming phases. The formation of intermetallic phases can generate strains in the order of 1–3%. There is insufficient evidence to support anisotropy of resulting strains, but there exists some evidence that tensile stress states may accelerate the formation of these phases.110 Additionally, there is a decrease in density and a concurrent increase in volume when ferrite is formed from austenite as a result of radiation-induced segregation of nickel. Formation of ferrite from austenite can lead to volume increases as large as 3%, but there are no available data on potential anisotropy or stress dependence. As opposed to carbide-induced strains that develop relatively quickly, ferrite and intermetallic strains develop rather slowly, and therefore are usually unrecognized, especially when other strain contributions arising from swelling and creep are present. Such precipitation-induced strains are important in that while they usually saturate in magnitude, they can be a significant portion of the total net strain at low dpa levels, thereby complicating the analysis and extrapolation of void swelling and irradiation creep data. Such strains can also affect the stress distribution and level in a structural component. For instance, a preloaded tie-rod or bolt will initially increase in load as a result of carbide-induced shrinkage even while irradiation creep proceeds to relax the load.
Radiation Damage in Austenitic Steels
65
Transgranular fracture
11.0–11.3 ´ 1022 n cm–2 (E > 0.1 MeV) 100
KIC (MPa 冪m)
80
377 C 400 C
Fatigue precracked Test temperature = 538 C zone
388 C
60 Intergranular fracture
Irradiation temperature
40
20 382 C
0 200
300
400 500 Test temperature (C)
600
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
Radiation Damage in Austenitic Steels
famous published example shown in Figure 49.115 When restrained in any direction, however, the swelling-induced stresses activate irradiation creep (to be discussed later), which then redistributes the strain in the unrestrained directions, as shown earlier in Figure 16 where fuel pins locally restrained by a spirally wrapped wire evolved into spiral fuel pins. At any given altitude on the fuel pin the interaction between wire and cladding the cross-section becomes oval in shape and the resulting deformation is called ‘ovality.’ It is important to note that, contrary to popular opinion, swelling and irradiation creep are not 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
Radiation Damage in Austenitic Steels
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
67
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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Radiation Damage in Austenitic Steels
Applied stresses have been shown to participate in the evolution of Frank loop and dislocation evolution and to produce the anisotropy of Burger’s vector distribution that is important to the operation of irradiation creep.123 Since shear stresses also assist in the unfaulting of Frank loops and in the evolution toward quasi-equilibrium network densities, it is not surprising that applied stress accelerates the onset of swelling. Although most previously reported 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.
Radiation Damage in Austenitic Steels
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
Radiation Damage in Austenitic Steels
primarily because such histories tend to accelerate the radiation-induced formation of nickel and silicon-rich phases, especially that of the g0 phase.1,136 Formation of these phases usually precedes swelling.1 Strong gradients in temperature across thin fuel cladding have also been shown to accelerate the onset of swelling, producing more swelling than what 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
Radiation Damage in Austenitic Steels
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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Radiation Damage in Austenitic Steels
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).
Radiation Damage in Austenitic Steels
73
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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Radiation Damage in Austenitic Steels
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
Radiation Damage in Austenitic Steels
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 dpa), so no swelling has yet occurred. While the irradiation creep rate is much higher than that of thermal creep at relatively low temperatures, the difference decreases between the two as the temperature increases, as shown in Figures 63 and 64.164,165 As noted earlier, the designation of creep and swelling as being separate processes appeared to be correct in the 1970s, but it is now known that once swelling begins, these two phenomena are both 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. Before swelling started, however, anisotropic distribution of mass in response to shear stresses was already in progress. This dislocation 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 in response to the local stress state. As an extreme example of
50 nm
3 dpa
6.5 dpa
22 dpa
Figure 60 Microstructure of annealed 12X18H9 specimens irradiated in the lower sections of BOR-60 reflector assembly at 320–330 C for 27 years. Lower dpa levels were reached at lower dpa rates. Reproduced from 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.
76
Radiation Damage in Austenitic Steels 0.65 dpa, 281 C
7.7 dpa, 285 C
12.3 dpa, 363 C
12.6 dpa, 380 C
8.8 dpa, 430 C
Figure 61 Void microstructure observed in a wrapper duct constructed from annealed 12X18H9 stainless steel and irradiated in the BN-350 reactor at various axial distances from the midplane. Reproduced from Maksimkin, O. P.; Tsai, K. V.; Turubarova, L. G.; Doronina, T. A.; Garner, F. A. J. Nucl. Mater. 2007, 367–370, 990–994. Lowest temperatures correspond to the bottom of the duct.
15 ´ 10-4
Temperature F
138 MPa 454 C
10
1300 1200 1100
1000
900
800
700
Thermal creep
10
5
Thermally induced densification and creep
0 0
500
1000 Time (h)
1500
In-reactor creep
1
Irradiation creep Midwall hoop strain (%)
Tensile strain
20% CW 316 SS Dose–15 dpa Time–4138 h
Hoop stress ksi/MPa 30/207 20/138 10/69
0.1
5/34
2000
Figure 62 Comparison of thermal and irradiation creep strains observed in 20% cold-worked 316 stainless steel in uniaxial creep tests during neutron irradiation in the EBR-II fast reactor or during ex-reactor thermal aging. Reproduced from Gilbert, E. R.; Kaulitz, D. C.; Holmes, J. J.; Claudsen, T. T. In Proceedings Conference on Irradiation Embrittlement and Creep in Fuel Cladding and Core Components; British Nuclear Energy Society: London, 1972; pp 239–251.
anisotropy, it has been shown that during charged particle simulation experiments on void swelling that 100% of the mass flow is directed perpendicular to the surface normal. During irradiation the swelling region experiences only a small fraction of the yield
30
0.01
20 10 5 0.001 1.0
1.1
1.2
1.3 1.4 1000/T (K-1)
1.5
1.6
Figure 63 Early comparison of thermal creep and irradiation creep in EBR-II of 20% cold-worked 316 pressurized tubes at 15 dpa and various stress levels. Reproduced from Gilbert, E. R.; Straalsund, J. L.; Wire, G. L. J. Nucl. Mater. 1977, 65, 277–294.
stress and does not produce plastic deformation.16 In effect, ion-induced swelling experiments might be better characterized as irradiation creep experiments
Radiation Damage in Austenitic Steels
700
Temperature (C) 600 500
77
400
Temperature (F) 1400
1200
1000
800
Hoop stress = 70 MPa Midwall hoop strain (%)
1.0
15 dpa
0.1
8 dpa 0.01
600 h 0.001
300 mm
2 dpa
1.0
1.1
1.2
1.3
1000 / T
(K-1)
1.4
4138 h 2400 h 1.5
1.6
Figure 64 Another comparison of thermal creep and irradiation creep in EBR-II of 20% cold-worked 316 pressurized tubes at various dpa levels at 70 MPa hoop stress. Reproduced from Gilbert, E. R.; Bates, J. F. J. Nucl. Mater. 1977, 65, 204–209.
in that two-thirds of the mass flow has been redirected along the third coordinate. One distinguishing characteristic of irradiation creep as opposed to purely thermal creep is that it is inherently nondamaging on the microstructural level. Its operation does not result in triple-point cracks or grain boundary voids and does not produce plastic deformation or martensite. In that sense irradiation creep is beneficial in that its action mitigates both microscopic and macroscopic problems arising from both externally loaded and internally generated stresses. Irradiation creep is especially effective to avoid stress concentrations anywhere in the microstructure, reducing the possibility of failure initiation. There is some evidence to support the proposal that radiation creep is not simply additive to thermal creep but that it modifies thermal creep, making it less damaging. Figure 65 shows an excellent example of the power of irradiation-modified thermal creep to produce a large amount of distortion without failure in CAGR cladding. Irradiation creep is a not completely understood phenomenon whose microstructural origins and complexity continue to evolve in our understanding, requiring additional evaluation as new data and
Figure 65 Micrographic section through a short length of Nb-stabilized Fe-20Cr-20Ni stainless steel after irradiation in a Commercial Advanced Gas Reactor, showing collapse of the cladding at 650 C into a large inter-pellet gap caused by thermal ratcheting of pellets during power level changes. The pressure difference driving the collapse is only 4 MPa but produces 55% strain without failure. The square features are cross-sections of a spiral rib on the cladding used to improve thermal performance. Reproduced from Crossland, I. G.; Roberts, G.; Jones, K. W.; Bradshaw, K. A. In Nuclear Fuel Performance; British Nuclear Energy Society: London, 1985; Vol. 1; pp 37–41.
insight are collected. As noted earlier, a comprehensive review on both void swelling and irradiation creep was written in 19941 and is now being revised,116 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. Most of the revised perceptions concern irradiation creep rather than swelling. On the microscopic level, however, there is still debate about what modes of dislocation movement and what balance of diffusion versus dislocation contributions combine to produce the observed strains under various irradiation conditions. Additional background information on this subject can be found in two other papers.166,167 While there are a very large number of creep mechanisms that have been proposed, most fall in several broad categories with respect to the microstructural components and the point defects involved. Matthews and Finnis have presented the most extensive review of these mechanisms and ranked them in terms of their relative plausibility.166 Garner and Gelles
Radiation Damage in Austenitic Steels
demonstrated that irradiation creep often leaves microstructural evidence in the alloy matrix that allows some mechanisms to be identified as strong contributors.123 A brief summary of current state of knowledge relevant to the purpose of this review is provided in the following sections. 4.02.9.2
Stages of Irradiation Creep
Irradiation creep of austenitic steels can be envisaged as having four stages. These are the transient regime, the creep regime in the absence of swelling, swellingenhanced creep, and creep disappearance. The first three contributions have traditionally been described using the following equation.1 : 2_ ¼ A½1 expðdpa=tÞ þ B0 þ DS: s The equivalent strain per unit equivalent stress, sometimes called the creep modulus B, is the sum of a transient contribution that saturates usually at a dpa or less, the creep compliance B0 in the absence of swelling, and stress-enhanced creep where the enhanced creep rate is proportional to the void swelling rate. Bubble swelling also accelerates irradiation creep, but the influence is expressed primarily in the early stages of creep.167 The coefficients A and t are empirical, experimentally determined constants that are very materialspecific (composition, thermal–mechanical treatment and texture) and sometimes stress-state dependent. For many high-exposure applications, the transient can be ignored. Even more importantly, the transient term can be obscured if significant precipitationrelated strains are developing concurrently. In general, the magnitude of the transient-induced strain increases with increasing stress, but the duration in time or dpa usually does not increase with stress. The creep transient involves both the dose needed to establish equilibrium densities of point defects, but most importantly, it involves the dose required for establishment of the quasi-equilibrium dislocation density. The transient is most pronounced for coldworked steels that start at higher dislocation density. As recombination–annihilation mechanisms reduce the dislocation density the instantaneous creep rate drops until the quasi-equilibrium dislocation density is reached and the steady-state creep rate B0 is established. For any austenitic steel it can be safely assumed that B0 is 1 106 (MPa dpa)1 and is effectively a ‘crystal constant’ similar to the 1% per dpa swelling rate of austenitic steels.
The creep compliance B0 is not only independent of composition but also starting state, dpa rate, and temperature over the range of reactor-relevant conditions. There appears to be one exception, however, in that the creep rate at temperatures somewhere below 100 C can increase significantly above B0, as shown at 60 C by Grossbeck and Mansur for various austenitic steels.168,169 At very low temperatures, vacancies are relatively immobile and cannot cancel the climb of dislocations produced by more mobile interstitials. Examples of such behavior have been seen in other studies.170,171 A good example of accelerated creep at low temperatures is shown in Figure 66. The swelling–creep coupling coefficient D was originally assumed also to be a crystal constant at 0.6 102 MPa,1 but as discussed later, it was found that D declines with increasing swelling to approximately one-third of this value or even to zero, depending on the stress state, stress history, and swelling history. This decline is an expression of the fourth stage of irradiation creep, variously designated as creep cessation, creep disappearance, or creep damping. One feature of irradiation creep that distinguishes it as different from thermal creep is that it varies with stress to power 1.0 rather than a higher power typical of thermal creep. This linearity is shown in Figure 67. The creep equation presented above also predicts
1.0
0
0.5
dpa 1.0
1.5
2.0
300 C Stress reduction ratio
78
0.5
60 C 0.2 172 < s0 < 262 MPa 0.1 0
2
4 6 8 ´ 1024 Fluence (n m-2) (E > 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.
Radiation Damage in Austenitic Steels
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
Radiation Damage in Austenitic Steels
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
Radiation Damage in Austenitic Steels
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
Radiation Damage in Austenitic Steels
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
Radiation Damage in Austenitic Steels
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
Radiation Damage in Austenitic Steels
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
Radiation Damage in Austenitic Steels
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
Radiation Damage in Austenitic Steels
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
Radiation Damage in Austenitic Steels
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
Radiation Damage in Austenitic Steels
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.
Radiation Damage in Austenitic 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
Radiation Damage in Austenitic Steels
and much earlier failure. Other researchers have reached the same conclusion.198 In general it appears that most researchers agree that helium is a contributing but not primary cause of the radiation-induced degradation in lifetime.195–199
4.02.10 Conclusions In general there are no beneficial aspects of 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.
Radiation Damage in Austenitic Steels
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
Radiation Damage in Austenitic Steels
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
Radiation Damage in Austenitic Steels
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.
95
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
ß 2012 Elsevier Ltd. All rights reserved.
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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Ferritic Steels and Advanced Ferritic–Martensitic Steels
progress to achieve a target burn-up of 250 dpa, using advanced ferritic steels. The attempts by nuclear technologists to enhance the thermal 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
Ferritic Steels and Advanced Ferritic–Martensitic Steels
Liquid
1500 Liq
Liquid + a + g
uid
a
1400
9Cr steel
1300
Temperature (⬚C)
1200
a a+g
g
1100
1000
900
a + g + (CrFe)7C3
800 a + g + (FeCr)3C 700
600
a + (CrFe)7C3 + (CrFe)4C
a + (FeCr)3C + (CrFe)7C3 a + (CrFe)7C3
a + (FeCr)3C 0
5
10
a + (CrFe)4C 15
20
25
Chromium (%)
(a)
Ae3 Ferrite Temperature (⬚C)
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
99
Pearlite Ws Bs
Widmanstatten ferrite Upper bainite
Bainite Lower bainite
Ms Martensite
(b)
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
100
Ferritic Steels and Advanced Ferritic–Martensitic Steels
(Mn+Ni)/135 Mod. 9Cr 1Mo
Mod. 9Cr 1Mo: base model
(Mn+Ni)/1.85 Mod. 9Cr 1Mo
4
(Mn+Ni)/2.32 Mod. 9Cr 1Mo
3
(Mn+Ni)/1.7 Mod. 9Cr 1Mo
1800
Mod. 9Cr 1Mo
1805
8
9
10
1795
625
(a)
1810
0.6 Si added 9Cr 1Mo
650
0.24 Si added 9Cr 1Mo
675
Plain 9Cr 1Mo
Melting point (K)
Ms, experimental (K)
1815
700
600 600
Experimental Estimated
1820
Ms/K = 904 - 474 (C + 0.46(N - 0.15Nb ) - 0.046Ta) -{17Cr + 33Mn + 21Mo + 20Ni + 39V + 5W) -45Mn2 - 25Ni2 - 100V2 + 10Co } - 44.5Ta
0.42 Si added 9Cr 1Mo
9Cr–ferritic martensitic steels 725
1W-0.23V-0.05Ta 9Cr 1Mo
750
1790 625
650
675
700
725
750
1785
Ms, empirical estimate (K)
1
2
(b)
5
6
7
11
Steel designation
1000 9Cr–ferritic steel
99
Continuous heating transformation (CHT) diagram
1248
Austenite
60
Ac3
40
1198 20
900
10
50% transformed Ac1
850
15
Ferrite+ austenite+ carbide
1
1098
Ferrite + carbide 800 100
101
(c)
102
1148
5
Temperature (K)
Temperature (⬚C)
950
103
Time (s)
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 <3%, creep strength will reduce, while higher amount promotes the formation of d-ferrite and brittle Fe2Mo Laves phase. The addition or partial replacement of molybdenum with tungsten and boron increased the stability of M23C6, and slowed down the kinetics of recovery. Lower nickel introduced d-ferrite, while its increase reduces creep strength. When Si is less than 0.3%, oxidation resistance gets lowered, while higher silicon content led to agglomeration of carbides with an increased amount of d-ferrite. On similar lines, the composition of all other elements could also be optimized, based on structure–property correlation studies. The components of the steam generators are often subjected to repeated thermal stresses as a result of temperature gradients that occur on heating and cooling during start-ups and shutdowns or during
variations in operating conditions. Steady state operation in between start-up and shutdown or transients would produce creep effects. Therefore the low cycle fatigue (LCF) and creep–fatigue interaction assume15 importance in the safe life design approach of steam generator components. The alloy exhibited a decrease in fatigue life with increasing temperature, thus limiting its upper limit of temperature up to about 773 K. The joining technologies of Cr–Mo steels have been well investigated.16,17 One of the major problems during welding of ferritic steels has been the formation of d-ferrite, if the amount of ferrite stabilizers is high. The partial substitution of Mo with W enables austenite stabilization and hence reduces the tendency to form d-ferrite. In fact, there needs to be an intricate balance between austenite and ferrite stabilizing elements in 9–12Cr–Mo steels.
Ferritic Steels and Advanced Ferritic–Martensitic Steels
This would ensure a satisfactory solidification process with a fully austenitic structure. Additionally, this enables easier hot workability during primary processing and tubemaking, without losing high temperature creep resistance. The formation of d-ferrite reduces toughness due to the notch sensitivity, promotes solidification cracking and embrittlement due to sigma-phase precipitation and reduces the creep ductility at elevated temperatures of operation. Other problems relate to solidification cracking, hydrogen cracking, and reheat cracking, which have been extensively studied.18 The Type IV cracking in ferritic steel weldments and the brittle layer formation in the dissimilar welds are discussed in detail later.
4.03.3 Radiation Damage of Core Components in Fast Reactors The core components in fast reactors include the following: clad (cylindrical tubes which house the fuel pellets) for the fuel and wrapper (a container which houses fuel elements, in between which the coolant flows) for fuel subassemblies. Figure 3 shows a schematic of clad and wrapper in a typical fuel subassembly. The necessity to develop robust technology for core component materials arises from the fact that the ‘burn-up’ (energy production from unit Head Adaptor
Shield pins, pellet stack
B&C shielding Steel shielding Top plannet
Section−MM 217 fuel pins
Fuel Bottom plannet Section−E E
Coolant entry tube
Section−HH
Section−BB Discriminator Section–XX
Fuel Pin
Figure 3 Schematic of a typical fuel subassembly.
101
quantity of the fuel) of the fuel depends on the performance of the clad materials. The higher burnup of the fuel increases the ‘residence time’ of the subassembly in the core, eventually lowering the cost. The core component materials in fast breeder reactors are exposed to severe environmental service conditions. The differences in the exposure conditions of the clad and wrapper in a fast reactor core are listed in Table 1. Under such exposure conditions, materials in the fast reactor fuel assemblies exhibit many phenomena (Figure 4), specific to fast reactor core: Void swelling, irradiation growth, irradiation hardening, irradiation creep, irradiation, and helium embrittlement. Another selection criterion, namely the compatibility of the core component materials with the coolant, the liquid sodium, has already been established. Presently, methods are known to avoid interaction of the clad material with the coolant. Detailed books and reviews19,20,21,22,23 are available on all the degradation mechanisms mentioned above, which are related to the production, diffusion, and interaction of point defects in the specific lattice of the material. Hence, a brief introduction is presented below (see also Chapter 1.03, RadiationInduced Effects on Microstructure; Chapter 1.11, Primary Radiation Damage Formation; and Chapter 1.04, Effect of Radiation on Strength and Ductility of Metals and Alloys). Void swelling in a fast reactor core can change a cube of nickel to increase (20%) its side from 1 cm to 1.06 cm, after an exposure to irradiation of 1022 n cm2. Void swelling is caused by the condensation of ‘excess vacancies’ left behind in the lattice after ‘recombination’ of point defects produced during irradiation. Void swelling is measured using the change in volume (▽V/V) of bulk components of the reactor or image analysis of voids observed using transmission electron microscope (TEM). The ‘irradiation growth’ (fluence 1020 n cm2) can increase the length of a cylindrical rod of uranium three times and reduce its diameter by 50%, retaining the same volume. This occurs mainly in anisotropic crystals, introducing severe distortion in core components. It is caused by the preferential condensation of interstitials as dislocation loops on prism planes of type (110) of hcp structures and vacancies as loops on the basal planes (0001), which is equivalent to transfer of atoms from the basal planes to prism planes, via irradiation-induced point defects. Irradiation hardening refers to the increase in the yield strength of the material with a
102
Ferritic Steels and Advanced Ferritic–Martensitic Steels
Table 1
Comparison of exposure conditions of clad and wrapper of fast reactor core
Criterion
Clad tube
Wrapper tube
Exposure conditions (only trends; exact values depend on core design)
Maximum temperature: 923–973 K
Lower temperature range than clad: 823 K Lower temperature gradient Moderate stresses from coolant pressure Flowing sodium environment Neutron environment similar
Major damage mechanisms
Selection criteria: mechanical properties
Corrosion criteria
Steeper temperature gradient Higher stresses from fission gas pressure Chemical attack from fuel Average neutron energy: 100 keV Neutron flux: 4–7 1011 n m2 s Neutron fluence: 2–4 1019 n m2 Void swelling Irradiation creep at higher temperatures Irradiation embrittlement Interactions with fuel and fission products Tensile strength Tensile ductility Creep strength Creep ductility Compatibility with sodium Compatibility with fuel Compatibility with fission products
Void swelling Irradiation creep Irradiation embrittlement Interaction with sodium Tensile strength Tensile ductility Compatibility with sodium
General common selection criteria Good workability International neutron irradiation experience as driver or experimental fuel subassembly Availability
concomitant reduction in ductility, under irradiation at temperatures <0.3Tm. The large number density of defect loops, voids, and precipitates generated during irradiation pins the mobile dislocations and acts as an obstacle to their further movement, requiring additional stress to unpin the immobile dislocations. The irradiation creep, the most important parameter for design consideration, is the augmentation of thermal creep of the material, under irradiation. This leads to premature failure of the material and restricts the service life. The mechanisms responsible for irradiation creep are identified as the ‘stressinduced preferential absorption’ (SIPA) and the ‘stress-induced preferential nucleation (SIPN)’ of point defects by dislocations, which revolve around the interaction of excess point defects generated during irradiation with dislocations. Irradiation embrittlement, another frequent observation in ferritic steels exposed to irradiation, refers to the increase in the ductile to brittle transition temperature (DBTT) during irradiation. Drastic loss in ductility at low temperatures results from a lower sensitivity of the fracture stress, sf, due to irradiation and less dependence on temperature than the yield strength sy. Materials with a high
value of the Hall-Petch constant are more prone to brittle failure. Such materials like ferritics release more dislocations into the system when a source is unlocked, causing hardening and loss of ductility. Some of the engineering materials contain nickel, an element which undergoes an (n,a) reaction, producing high concentration of helium. The binding energy of helium with a vacancy being very high 2 eV, the helium atoms stabilize the voids, enhancing their growth rate. Incorporation of helium during irradiation into voids along the grain boundaries assists grain boundary crack growth by linking voids causing ‘helium embrittlement.’ Of these many degradation mechanisms, the alloy development programmes have focused mainly on the void swelling, irradiation hardening, embrittlement, and the irradiation creep, since these are the major life limiting factors.
4.03.4 Development of Ferritic Steels for Fast Reactor Core This section begins with the optimization of chemistry and initial microstructure to develop swelling and
Ferritic Steels and Advanced Ferritic–Martensitic Steels
103
3 a axis
(a)
Growth strain (10-4)
Swelling (ÑV/V )
2 Linear swelling regime
Transient swelling regime
Threshold dose
Irradiation dose (dpa)
1 0 -1
c axis
-2 -3
(b)
Neutron fluence (E>1 MeV) 1024 n m–2
Strain (%)
Stress
Irradiated
Unirradiated
Irradiated
Unirradiated
(c)
(d)
Strain
Time (h)
Absorbed energy
Unirradiated
(e)
Irradiated
Test temperature
Figure 4 Schematic representation of major damage mechanisms in the core component materials of fast reactors: (a) The different stages of void swelling, (b) irradiation growth, (c) increase in strength with a concomitant reduction in ductility during irradiation hardening, (d) increase in creep strain and reduction in creep life after irradiation caused by irradiation creep, and (e) increase in ductile to brittle transition temperature and reduction in upper shelf energy after irradiation caused by irradiation embrittlement.
creep-resistant ferritic steels. The microstructural instability during service exposure is briefly presented. The superior swelling performance of ferritic steels is understood based on mechanisms of void swelling suppression. Following this, the irradiationinduced/-enhanced segregation/precipitation causing irradiation hardening is discussed. The irradiation creep and embrittlement, their mechanisms and methods to combat the problems are highlighted. The R&D efforts of today to reduce the severity of embrittlement in ferritic steels, using modeling methods, are outlined. Finally, typical problems in the weldments
of ferritic steels, when used for out of core applications, are presented, emphasizing the advantage of modeling in predicting the materials’ behavior. 4.03.4.1 Influence of Composition and Microstructure on Properties of Ferritic Steels Rapid strides have been made the world over, in the design and development of advanced creep-resistant ferritic or ferritic–martensitic steels. The low alloy steels can be used as either 100% ferrite–martensite
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Ferritic Steels and Advanced Ferritic–Martensitic Steels
or a mixture of both. It is possible to choose the required structure by the appropriate choice of either the chemistry or the heat treatment. For example, a completely ferrite matrix, yielding high toughness, can be obtained in steels with chromium content higher than 12%, with carbon reduced to less than 0.03%. The same steel can be used to provide higher strength by choosing the 100% martensite structure, if carbon content is increased to about 0.1%. The 9Cr steels have always been used in the 100% martensite state. Extensive studies have been carried out on phase stabilities of these steels, with changes in chemistry and heat treatment. The creep resistance of the plain Cr–Mo steels has, further, been increased by the addition of carbide stabilizers like Ti or V or Nb, leading to the modified variety of 9–12Cr–Mo steels. These
Table 2 Optimizing the constitution in the development of ferritic steels Element
Function
Cr
Basic alloying element, corrosion resistance, hardenability Solid solution strengthening Strengthening by formation of MX-carbonitride Austenite stabilizer, solid solution strengthening, carbonitride formers Grain boundary strengthening, stabilization of carbide Austenite former, inhibits d-ferrite formation
Mo, W, Re, Co V, Nb, Ti, Ta C, N B Ni, Cu, Co
Table 3
elements led24 to copious, uniform precipitation of Monte Carlo (MC) type of monocarbides, which are very fine and semicoherent. Such precipitates are very efficient in pinning the mobile dislocations, leading to improved creep behavior at higher temperatures. These carbides are stable at temperatures higher than even 1273 K and hence, do not cause deterioration of long-term mechanical properties during service exposure. The development of high creep-rupture strength 9–12% steels with various combinations of N, Mo, W, V, Nb, Co, Cu, and Ta is based on optimizing the constitution (Table 2.) and d-ferrite content, increasing the stability of the martensite, dislocation structure and maximizing the solid solution and precipitation hardening. The concentration of each element in ferritic steels has been optimized based on an in-depth understanding of the influence of the specific element on the behavior of the steel. The extensive studies related to optimization of chemistry are summarized in Table 3. Based on the strong scientific insights, large number of commercial steels have been developed (Table 4) in the later half of the last century. Most of this family of ferritic–martensitic steels is used in the normalized and tempered condition or fully annealed condition to achieve the desirable phase. The type of structure that is deliberately favored in a given steel depends on the end application. The microstructure of the steels in normalized and tempered conditions consists24 (Figure 5) of (a) martensite laths containing dislocations with a Burgers vector 1/2a0<111> with a density of approximately 1 1014 m2 (b) coarse M23C6 particles located at
Beneficial and harmful effects of different elements during design of creep-resistant ferritic steels
Element
Beneficial
Detrimental
Optimum (wt%)
Carbon
Strength
0.1
Mo
Creep strength
Ni Mn Si
S scavenger Void swelling, spheroidization
Ti, V, Nb, Ta, and W
Precipitation strengthening
Weldability Hardenability d-ferrite formation Laves phase Intermetallic Ni3P Induced radioactivity d-ferrite formation Silicide formation Undissolved carbides, low hardness of martensite Radiation embrittlement Precipitation Induced radioactivity Ductility Neutron poison
S, P, As, Sb, Sn, and Bi Cu Co N, O B
Delayed coarsening
1 0.1 0.5 0.2–0.4 <0.1 <0.001 <0.01 <0.01 <20 ppm <0.002
Ferritic Steels and Advanced Ferritic–Martensitic Steels Table 4 List25 of commercial ferritic steels, their chemistry, and properties Commercial name
Chemistry
105 h creep strength at 873 K MPa1
T22 Stab. T22 HCM2S T9 EM12 F9 T91
2.25Cr1Mo 2.25Cr1MoV 2.25Cr1MoWNb 9Cr1Mo 9Cr2MoVNb 9Cr1MoVNb 9Cr1MoVNb (optimized) 9Cr(MoW)VNb 9CrWTiV 12Cr1MoV 12Cr1MoWV 12CrMoWVNbCu 12CrWVNbCo
35 60–80 100 35 60–80 60–80 100
T92 Eurofer HT91 HT9 HCM12A SAVE12
120 120 60–80 60–80 120 180
prior austenite and ferrite grain boundaries with finer precipitates within the laths and at martensite lath and subgrain boundaries. M2X precipitates rich in Cr are isomorphous with (CrMoWV)2CN. The initial microstructure of the normalized and tempered steels described above does not remain stable during service in a nuclear reactor. Prolonged exposure at high temperature causes changes in the initial microstructure, which has been studied extensively. The M2X precipitates in the normalized and tempered stabilized 9Cr–1Mo steels are gradually replaced (Figure 6) by MX, intermetallic, and Laves phases during prolonged aging at high temperature. The high temperature and the irradiation over prolonged time of exposure introduce microstructural instabilities. These instabilities are caused mainly by the point defects caused by irradiation and complex coupling of these defects with atoms in the host lattice, their diffusion or segregation and finally the precipitation. There is a recovery of the defect structure since the irradiation-induced vacancies alter the dislocation dynamics. There are three types of processes with respect to evolution of secondary phases: irradiation-induced precipitation, irradiationenhanced transformations, and the irradiation modified phases. It is seen that the evolution of these phases depends on the composition and structure of the steel and the irradiation parameters like the temperature, dose rate, and the dose. Evolution of irradiationinduced phases and their influence on hardening and embrittlement is discussed later.
4.03.4.2
105
Void Swelling Resistance
Extensive experimental investigations found3 that the ferritic steels, whose high temperature mechanical properties are far inferior to austenitic stainless steels, displayed excellent radiation resistance. The ferritic– martensitic steels (9–12% Cr) have, therefore, been chosen for clad and wrapper applications, in order to achieve the high burn-up of the fuel. This is based26–29 (Table 5) on their inherent low swelling behavior. The 9Cr–1Mo steel, modified 9Cr–1Mo (Grade 91), 9Cr–2Mo, and 12Cr–1MoVW (HT9) have low swelling rates at doses as high as 200 dpa. For example, HT9 shows 1% swelling at 693 K for 200 dpa. The threshold dose for swelling in ferritic steels is as high as nearly 200 dpa in contrast to 80 dpa for the present generation D9 austenitic stainless steel. It is established that the void swelling depends crucially on the structure of the matrix lattice, in which irradiation produces the excess defects. Extensive basic studies have identified19,30–33 the following reasons as the origin of superior swelling resistance in ferritic steels: 1. The relaxation volume for interstitials, that is, the volumeof the matrix inwhich distortion is introduced bycreating an interstitial, in bcc ferrite is larger19 than fcc austenite. For every interstitial introduced, the lattice distortion is high and hence the strain energy of the lattice. Hence, the bias toward attracting or accommodating interstitials in the bcc lattice is less. This leaves higher density of ‘free’ interstitials in the bcc lattice than fcc lattice. As a result, recombination probability with vacancies increases significantly and supersaturation of vacancy reduces. Consequently, the void nucleation and swelling is less. 2. The migration energy of vacancies in bcc iron is only 0.55 eV, against a high value in fcc austenite, 1.4 eV. Vacancies are more mobile in bcc than fcc, increasing the recombination probabilities in bcc ferrite. Another factor is the high binding energy between carbon and vacancy in bcc iron (0.85 eV), while it is only 0.36–0.41 eV in austenite. This leads19 to enhanced point defect recombination in bcc than fcc, due to more trapping of vacancies by carbon or nitrogen. 3. In bcc iron, it is known30 that there is a strong interaction between dislocations and interstitials solutes, forming atmospheres of solutes around dislocations. The formation of ‘atmospheres’ around dislocations makes them more effective sinks for vacancies than interstitials, resulting in suppression of void growth,
106
Ferritic Steels and Advanced Ferritic–Martensitic Steels
A
B
0.5 mm
(a)
(001) (b)
V Ka
NbLa
V Kb
NbKa
FeKa 2.00
NbKa
V Ka
NbLa
4.00
6.00
8.00
10.00
12.00
14.00
16.00
(c)
1
2.00
4.00
6.00
8.00
10.00
12.00
14.00
16.00
(d)
Figure 5 Initial structure24 of normalized and tempered modified 9Cr–1Mo steel: (a) Monocarbides (MC) and M23C6 along lath boundaries in a carbon extraction replica of the sample and (b) Microdiffraction of fine particle marked B, confirming the crystal structure of MC. Energy dispersive analysis of X-rays (EDAX) identifying the MC particles (B) to be rich in (c) V and (d) Nb.
provided the following conditions are satisfied: ‘atmospheres’ comprise of either oversized substitutional atoms or interstitials, dislocations have high binding energy with solutes, and concentration of solute atoms at the core of the dislocation exceeds a critical value. On the other hand, if ‘atmosphere’ is made up of undersized atoms like Si or P, the voids can grow. The ‘atmosphere’ of interstitials reduces the dislocation bias for additional capture and inhibits dislocation climb, thus converting them to saturable sinks. Such a scenario would increase the recombination probabilities, suppressing the void growth. These fundamental differences in the behavior of solutes and point defects in bcc lattice make ferritic steel far superior to austenitic steels, with respect to radiation damage.
The challenging task for materials scientists to use ferritic steels directly in fast reactor fuel assembly was with respect to enhancing the high temperature mechanical properties of the ferritic steels, especially high temperature creep life and irradiation creep resistance. 4.03.4.3 Irradiation Hardening in Ferritic Steels The initial microstructure of the steels evolves during service, due to high temperature and irradiation for prolonged times, leading to modification of defect structure and secondary phases. These changes harden the steel, leading to concomitant embrittlement, which is discussed below.
Ferritic Steels and Advanced Ferritic–Martensitic Steels
It is reported that carbon content in 12% chromium steel is maintained high in order to use the steel as martensitic steels. The high amount of carbon in 12% chromium steel leads to copious precipitation of carbides, that is, twice as much in 9Cr steels. Both the steels have predominantly M23C6 carbides with a small fraction of monocarbides, (1 21 3)
(a)
1 mm
MoLa
Laves phase
FeKa
(b)
CrKa SiKa MoKa
2.00
4.00
6.00
8.00
10.00 12.00 14.00 16.00 18.00
24
Figure 6 Effect of prolonged exposure (823 K per 10 000 h) of modified 9Cr–1Mo steel. Transmission electron micrograph showing (a) formation of detrimental Fe2Mo Laves intermetallic phase around the M23C6. The insets show the microdiffraction pattern and magnified view of the nucleation of Laves phase (b) EDAX spectrum confirming the enrichment of iron and molybdenum. Table 5
107
eventually leading34 to deterioration of their resistance to brittle failure. The critical stress to propagate a crack is inversely proportional to the crack length. If it is assumed that fracture initiates at an M23C6 precipitate and the crack length at initiation equals the diameter of a carbide particle then the fracture stress will decrease with increasing precipitate size. The precipitates coarsen during irradiation in the range of 673–773 K, thus causing a decrease in fracture stress and an increase in DBTT even in the absence of further hardening. Additionally, Cr rich, bcc a0 precipitates formed35 in the higher chromium steel during thermal exposure and irradiation lead to hardening and embrittlement of the steel. The d-ferrite, into which there is a repartitioning of chromium, is harmful, since it promotes formation of a0 . The presence of very fine coherent particles of the w (Fe2Mo) phase has also been reported in the T91 and HT9 steels. The w phase was observed to form more rapidly in the 9Cr–2Mo type of steels, both in the d-ferrite and martensite phases. This is possibly due to the higher amount of Mo in the EM12 type of steels. The w phase is enriched in Fe, Si, and Ni and contains significant amount of Mo and P. The G phase (Mn7Ni16Si17) has been found to form very occasionally in the modified 9Cr–1Mo and HT9 (12% Cr) variety of steels. The s phase (Fe–Cr phase, enriched in Si, Ni, and P) has been observed to form around the M23C6 particles in 9–13% Cr martensitic steels after irradiation at 420–460 C in Dounray Fast Reactor. In addition Cr3P needles and MP (M ¼ Fe, Cr, and Mo) particles have also been detected in the 12 and 13Cr steels in the range of 420–615 C. The formation of these phases during irradiation may be understood in terms of the strong radiation-induced segregation (RIS) of alloying and impurity elements to point defect sinks in the steels (see Chapter 1.18, RadiationInduced Segregation). The RIS of alloying/impurity elements could lead36 to either enrichment or depletion near the sinks, depending on the size of the atom and its binding energy with iron self-interstitials.
Void swelling resistance26–29 of some commercial ferritic steels
Commercial name
Chemistry and country of origin
Reactor in which irradiation was carried out
Burn-up achieved (dpa)
FV448 EM10 1.4914 EP450 EP450
12Cr–MoVNb, UK 9Cr–1Mo, France 12CrMoVNb, Germany 12Cr–MoVNb, Russia –
PFR Phenix Phenix BN-350 BN-600
132 142 115 45 144
108
Ferritic Steels and Advanced Ferritic–Martensitic Steels
Generally, a large number of alloying elements, W, Nb, Mo, Ta, V, or Ti are dissolved into the matrix of ferritic steels, some of them being larger than the iron atom. This could lead to the expansion of the unit cell of ferrite, making an element say, chromium undersized, with a positive binding energy with iron self-interstitial. Such a situation would lead to enrichment of chromium near the sink-like grain boundary. The reverse could happen if the size of the alloying elements happen to be smaller than iron. The w, G and s phases are all enriched in Si and Ni – elements which are known to segregate to interfaces during irradiation. With the exception of G phase, all the other phases and the a0 phase are rich in Cr. In those ferritic steels, where Cr is depleted near voids and at other interfaces which act as point defect sinks, it follows that in steels containing higher than 11 or 12% Cr, the chromium enrichment within the matrix may lead to local concentrations exceeding those (14%) at which a0 forms thermally. Further, enrichment of Cr may also result from the partial dissolution of chromium rich precipitates such as M23C6 during irradiation. In addition, RIS of phosphorus can also lead to the formation of phosphides in some of the steels. The irradiation-induced point defect clusters and loops may also facilitate and enhance nucleation of these phases. Although the relatively soft d-ferrite improves the ductility and toughness of the 12Cr steel, the fracture could be initiated at the M23C6 precipitates on the d-ferrite–martensite interface. The presence of d-ferrite, extensive precipitation and radiation-induced growth of M23C6 precipitates and formation of the embrittling intermetallic phases in the 12Cr–1MoVW steel in the temperature range 573–773 K are together responsible37 for the relative change in impact behavior of 9Cr–1MoVNb and 12Cr–1MoVW between 323 and 673 K. Irradiation-induced microstructural changes are the factors that govern the creep and embrittlement behavior, which therefore, has to be minimized using appropriate chemistry and structure. 4.03.4.4 Irradiation Creep Resistance of Ferritic Steels An essential prerequisite for maximizing the ‘irradiation creep resistance’ is to ensure38 the best combination of thermal creep behavior and longterm microstructural stability at high temperature. Hence, the present section would discuss irradiation creep in the same sequence as mentioned above.
The design principles of development of creepresistant steels are as follows: Introduce high dislocation density by either transformations or cold work to increase the strength of the basic lattice; Strengthen the host lattice by either solid solution strengtheners or defects; Stabilize the boundaries created by phase transformations by precipitating carbides along the boundaries; Arrest dislocation glide and climb by appropriate selection of crystal structure, solid solution, interfaces, dislocation interactions, and crystal with low diffusivity; Resist sliding of grain boundaries by introducing special type of boundaries and anchoring the boundaries with precipitates; Ensure long-term stability of the microstructure, especially against recovery and coarsening of the fine second phase particles; In the case of 9–12 Cr steels, the martensitic lath structure (Figure 7) decorated with only MX which should39 be stable over long-term service life is the most desired structure. Thermo-Calc evaluations show39 that MX can be stabilized at the expense of M23C6 only by reducing carbon to as low a value as 0.02% in 9 Cr–1Mo steel. This value is too low to ensure acceptable high temperature mechanical behavior of the steels. In the context of fast reactor core components, the high chromium 9–12% ferritic–martensitic steels assume relevance. Hence, an extensive database25 for a large number of commercial ferritic steels has been generated and Lath boundary
Lath boundary
(a)
(b) M23C6 MX
Figure 7 The schematic39 of most undesirable (a) and desirable (b) microstructures for design of creep-resistant steels.
109
Ferritic Steels and Advanced Ferritic–Martensitic Steels stensile -0.5Mo + 0.8 W
100
0 s
T 9Cr1Mo
P91
E911
P92
T
0
0
650 ⬚C
50
T
600 ⬚C
+ 0.04N + 0.2V + 0.08Nb
(a)
0 0 0 0
+1 W
T
105 h rupture stress (MPa)
150
T
stensile 0
0
5 540 ⬚C 1 ⫻ 1023 n cm–2
4
0
(a) 0
316SS
0
DD/D (%)
0 3 D9 2
(b) Figure 9 The mechanisms of stress-induced preferential absorption (a) and stress-induced preferential nucleation (b) during irradiation creep.
HT-9 D21
1
D68 0 0
(b)
20
40
60
80
100
120
140
Hoop stress (MPa)
Figure 8 (a) Thermal creep40 of 9Cr1Mo ferritic steel. (b) irradiation creep41 of ferritics in comparison to austenitics. Reprinted, with permission, from J. ASTM Int., copyright ASTM International, 100 Barr Harbor Drive, West Conshohocken, PA 19428
Figure 8(a) shows40 the continuous improvement achieved by careful modification of alloying elements, in the thermal creep behavior of successive grades of different commercial ferritic steels. While understanding thermal creep is essential to narrow down the choice of ferritic steels for use in a fast reactor, ‘benchmarking’ the steels developed under irradiation is an essential stage before actually using the radiation-resistant steels in the reactor. The irradiation creep depends on the stress level, the temperature, and the dose. Figure 8(b) shows41 the comparison of irradiation creep of ferritic steels with competing materials like the austenitics and nickel-based alloys. It is clear that the point defects generated during irradiation act against the design principles of developing creep-resistant materials, listed earlier. The point defects accelerate the kinetics of dislocation climb, coarsen the precipitates, and generally enhance the diffusivity. In addition, the excess point defects precipitate into either interstitial or vacancy loops, but not randomly. The interaction between point defects and stress leads to the precipitation of interstitial loops parallel to the applied stress, while vacancy
loops form in planes perpendicular to the stress. This process (Figure 9(a)) called the stress-induced preferential nucleation (SIPN) results in additional creep strain solely due to irradiation. The excess point defects under temperature migrate randomly. But in the presence of an additional factor, that is, stress, the vacancies migrate preferentially to grain boundaries perpendicular to the applied stress, while the interstitials toward boundaries parallel to the stress. This is equivalent to removing material from planes parallel to the stress to those which are perpendicular to the applied stress, introducing additional creep strain. This process is called the stress-induced preferential absorption (SIPA) (Figure 9(b)). The radiation-induced defects also evolve from isolated point defect to loops and voids, which have different types of influence on irradiation creep. Most often, irradiation creep occurs19,42 simultaneously with swelling and sometimes, swelling influences irradiation creep. At very small dose levels, swelling enhances creep rates. Beyond a certain dose levels, the creep component reduces and at high dose levels, creep disappears, while swelling continues. Figure 10 shows the variation in creep coefficient at various dose levels, and the regimes where swelling has an influence. The dynamics of point defects during irradiation continuously evolve with change in structure of dislocation network and loops. At small dose levels, there is a uniform distribution of very fine voids, which act as effective pinning centers for mobile dislocations. Thus the creep rate increases. With increase in dose levels, voids grow and multiply. The chance of interstitials and
110
Ferritic Steels and Advanced Ferritic–Martensitic Steels
Instantaneous creep coefficient
Onset of disappearance of creep
Swelling without creep
Swelling enhanced creep
Dose (dpa)
Figure 10 Schematic of variation of instantaneous creep coefficient with dose, showing the interplay between irradiation creep and void swelling.
vacancies impinging on the void surface becomes more than their reaching dislocations. The number of interstitials reaching a dislocation reduces. Additionally, the defect clusters, that is, the dislocation loops also undergo ‘faulting’ contributing to the density of dislocations in the matrix. Hence, creep rate reduces, due to two factors: increased dislocation density of the matrix due to unfaulting of dislocation loops and reduced availability of interstitials to dislocations. The above process continues until complete cessation of creep, with swelling continue to take place. At very high temperatures, the point defect migration along the grain boundaries in preferential routes causes the grain boundary aided creep. This high temperature limit of ferritic–martensitic steels restricts the application of these steels to at best, wrappers of present generation fast reactors based on oxide fuel. It is necessary to develop materials with better high temperature irradiation creep properties and void swelling for clad applications. The future scenario, which envisages the development of metallic fuel to ensure sustainability by breeding, could make use of ferritic steels for both clad and wrapper. This advantage arises due to the lower value of the anticipated clad temperatures with metallic fuels (see Chapter 3.14, Uranium Intermetallic Fuels (U-Al, U-Si, U-Mo)), whose choice is mainly to ensure sustainability using high breeding ratio. 4.03.4.5 Irradiation Embrittlement in Ferritic Steels The stabilized ferritic steels in the normalized and tempered condition have a tempered martensitic structure with a preponderance of monocarbides
that impart the necessary creep strength, while the prior austenite grain and lath boundaries are decorated with Cr rich M23C6 precipitates which increase the thermal stability of the steel. It is reported that thermal aging at temperatures above 773 K causes gradual but continuous degradation in upper shelf properties in addition to increase in the DBTT. The nature of embrittlement varies for different components of the reactor. For removable components such as clad, which are subjected to high temperature and pressure, with a residence time of a few years, creep embrittlement is the issue which decides their design and performance, while for permanent support structures increase in hardening and loss in fracture toughness on irradiation are major issues. The origin of embrittlement is two-fold: segregation of tramp elements to prior austenite grain boundaries which make the grain boundaries decohesive and evolution of carbides and intermetallic phases. The latter causes progressive changes in the tempered martensitic microstructure, which deteriorate the fracture properties of the steel, by introducing irradiation hardening effects. The increase in the ductile to brittle transition temperature, DDBTT, is known to be related to irradiation hardening, which is generally observed to saturate with fluence. Evidence for a possible maximum in DBTT was observed for the 12Cr steel irradiated in the range of 35–100 dpa in fast flux test facility (FFTF). Based on observed data in a number of cases it appears that a high fluence and/or high temperature are required before a maximum is observed. This implies that the strength and impact properties are a balance between the point defect production and irradiation-induced precipitation. The precipitation during irradiation hardens the steel and irradiation accelerated recovery and aging soften the steel. The latter process is more important at high fluences and/or higher irradiation temperatures. Hence, hardening in most of these Cr–Mo steels is more than compensated for by the recovery and aging processes, leading to saturation in irradiation hardening above 723 K. For body centered cubic materials such as ferritic martensitic steels, radiation hardening at low temperatures (<0.3TM) can lead to a large increase in the DBTT and lowering of impact energy for radiation dose as low as 1 dpa (displacement per atom). The minimum operating temperature to avoid embrittlement in ferritic martensitic (F/M) steels is 473– 523 K, while the upper limit is controlled by four different mechanisms: thermal creep, high temperature helium embrittlement, void swelling, and compatibility
Ferritic Steels and Advanced Ferritic–Martensitic Steels
with fuel and coolant. Void swelling is not expected to be significant in F/M steels up to damage levels of about 200 dpa. Extensive evaluation14,15,43–58 of the embrittlement behavior of the ferritic steels for different chemistry is shown in Figure 11. The merit in focusing on chemistry around 9% chromium is very clear based on the observation of minimum shift in DBTT around this composition, under irradiation. However, higher chromium improves corrosion resistance and ease of reprocessing. Hence, chromium content has to be selected balancing these requirements. It is known44–48 that addition of phosphorous, copper, vanadium, aluminum, and silicon would increase the DBTT while sulfur reduces the upper shelf energy (USE). The 12Cr steels, HT9, show a larger shift (125 K) in DBTT as compared to modified 9Cr–1Mo steel (54 K). Hence, the balance is always between nearly nil swelling resistant 12Cr steels and 9Cr steel which is less prone to embrittlement than 12Cr steels. Microstructural parameters, like the prior austenite grain size, lath/packet size, carbides, and their distribution influence49,50 the embrittlement behavior. Studies on the effects of heat treatment and microstructure on the irradiation embrittlement in 9Cr–1MoVNb and HT9 steels are summarized below:
Prior austenite grain size (PAGS) influences51 the DBTT for the 9Cr–1MoVNb steel, but not in
250 JLF-4
DBTT shift (C)
200
150
12Cr–MoVW steel. This is attributed to the precipitates in the microstructure controlling the fracture behavior rather than the PAGS, in the 12Cr steel. The size of martensitic lath and packet, which is sensitive52 to austenitization temperature, can also affect51 the fracture behavior. Examination of the fracture surface revealed cleavage and regions of ductile tearing along prior austenite grain and lath packet boundaries. Subsurface microcracks and secondary surface cracks were found associated with large boundary carbides. It was suggested that cleavage fracture initiated in HT9 by propagation of a microcrack from a coarse carbide into the matrix. Propagation was inhibited by the intercepted boundaries, lath or grain and ductile tearing was required53 to continue propagation. The amount of tearing increased with increasing austenitization temperature. Tempering for the two normalization temperatures had very small effect on the DBTT, for the two steels. Irradiation of the two steels at 638 and 693 K resulted37 in an increase in DBTT and a decrease in USE for all conditions with the shift in DBTT for the 12 Cr steel being almost twice that for 9Cr steel. Although the 12Cr steel with the smallest grain size had55 the lowest DBTT after 20 dpa, the effect of tempering was different. In the case
2.25CrV 2.25Cr–1WV 10 dpa: 365 ⬚C 2Cr–1.5V R. Klueh et al.
JLF-6 12Cr–6Mn–1W 12Cr–6Mn–1V
36 dpa: 410 ⬚C A. Kohyama et al.
2.25Cr–2W
111
12Cr–2WV
JLF-3
100
2.25Cr–2W
F82H
5Cr–2WV
9Cr–2WV
7 dpa: 365 ⬚C 10 dpa: 365 ⬚C
50
7 dpa: 365 ⬚C R. Klueh et al.
36 dpa: 410 ⬚C
Cr–1V JLF-1 9Cr–1W
7.5Cr–2W
9Cr–2WVTa
0 0
2
4
6
8
10
12
Chromium content (wt%) Figure 11 Variation43 of shift in ductile to brittle transition temperature (DBTT) for various Cr–Mo steels with irradiation to different dose levels at around 673 K. The ferritic steel with 9Cr–1Mo has the least variation in DBTT.
112
Ferritic Steels and Advanced Ferritic–Martensitic Steels
of 12Cr steel, the higher tempering temperature causes coarsening of precipitates thus accelerating fracture. The saturation of shift in DBTT with fluence is independent54 of tempering conditions for the 9Cr steel, while for the 12Cr steel, a maximum is observed, probably due to faster growth of precipitates during irradiation. The generation of helium through (n,a) reaction in elements of structural materials is known to cause severe damage to the embrittlement behavior of core component materials. Table 6 lists the shift in DBTT, for 9 and 12CrMo steels, under reactor irradiation, with and without helium, which demonstrates56 the harmful effect of helium. These results become more pertinent in the case of fusion reactors, where the operating conditions include the generation of helium up to about 100 appm year1. The increase in the DBTT due to irradiation is a cause of serious concern for use of ferritic steels, since it makes the postirradiation operations very difficult. Several methods have been attempted57,58 to address this problem, which includes modification of the steel through alloying additions, control of tramp elements by using pure raw materials and improved melting practices, and grain boundary engineering (GBE). However, the propensity of the problem is less if the clad thickness is low, which normally is the case to ensure best heat transfer properties. For low thickness components, the triaxial stress necessary for the embrittlement does not develop, which reduces the intensity of this otherwise serious problem of embrittlement in ferritic steels. An approach to reduce shift in DBTT is an immediate concern in ferritic steels for core component applications and efforts to overcome this problem by selection of high purity metals, adoption of double or triple vacuum melting for steel making, strict control of tramp and volatile elements, and development of special processing methods, which would improve the nature of grain boundaries (GBE) are in progress. Table 6
4.03.4.5.1 GBE to reduce embrittlement in ferritic steels
GBE is an emerging field, which promises methods to improve the performance of materials, whose degradation in service is caused by the presence of high angle boundaries. The concept, first proposed59 by Prof. T. Watanabe in the early 1980s, envisages improvement of properties of materials by controlling the grain boundary character distribution (GBCD). Many processes like diffusion, precipitation, segregation, sliding, cavitation, and corrosion are kinetically faster along high angle grain boundaries. Hence, it is possible to decelerate these detrimental processes by replacing the random boundaries with low energy ones, coincident site lattice (CSL) boundaries (denoted by the ‘sigma number,’ S, which is defined as the reciprocal of the fraction of lattice points in the boundaries that coincide between the two adjoining grains on the basis of CSL model). Another prerequisite for GBE is to completely destroy the interconnectivity of random grain boundary network. The insight in the field of GBE was achieved with the advent of computer assisted EBSD (electron back scatter diffraction) technique developed during the 1980s. The embrittlement in ferritic steels is known to be caused by segregation phenomena. The kinetics of segregation can be controlled by suitable selection of the nature of grain boundaries. GBE has been applied60–63 to combat embrittlement problems in ferritic steels. The task of carrying out GBE using experimental methods is time consuming. Hence, it is prudent to resort to computational methods, which need to be validated using selected experiments. A 3D Poisson–Voronoi grain structure, simulated using MC technique was employed to study60 (Figure 12(a)) intergranular crack percolation using percolation theory. The percolation threshold was estimated to be 80%. To apply this model to specific alloys like ferritic steel, system specific characteristics need to be incorporated61 in the model. One such attempt is to define the propensity of the grain boundaries for propagation of cracks based on relative values of the grain boundary energy and the
Comparison56 of embrittlement behavior of 9 and 12Cr steels, with and without helium
Irradiation conditions
Shift in DBTT (K)
Reactor
Temperature (K)
Dose (dpa)
9Cr1Mo(VNb)
12Cr1Mo(VW)
EBR II EBR II HFIR
663 663 673
13 26 40
50 50 200 (30 appm He)
125 150 250 (110 appm He)
Ferritic Steels and Advanced Ferritic–Martensitic Steels
113
(a)
1 0.9 0.8 0.7
Failure probability
30 50 60 70 80
Fine grains (12 mm)
0.6
Coarse grains (25 mm)
% crackresistant boundaries
0.5 0.4 0.3 0.2 0.1 0 200
300
400 500 600 Critical crack length (mm)
250
Two-step normalization and tempering treatment
200 150
Conventional N&T treatment
100
67.8 J lower bound criteria
50 0
-45 ⬚C
-30 ⬚C
-80 -70 -60 -50 -40 -30 -20 -10
(c)
Fractal dimension increment
Charpy absorbed energy (J)
300
Temperature (⬚C)
0
10
20
30
0.185 0.180 0.175 0.170 0.165 0.160 0.155 0.150 0.145 0.140 0.135 0.130 0.125
0
5
800
15
10
900
20
1000
25
30
35
40
25
30
35
40
45
Average Crack initial stage
0
(d)
700
Specimen A
100
Specimen B
0
(b)
5
10
15
20
0.185 0.180 0.175 0.170 0.165 0.160 0.155 0.150 0.145 0.140 0.135 0.130 0.125
45
Charpy impact energy (J)
Figure 12 Modeling and electron back scattered diffraction studies in grain boundary engineering of ferritic steels: (a) Percolation of a crack in the 3D P–V model of grain structure generated60 using Monte Carlo methods. (b) Percolation probability for two different grain size (c) experimental confirmation63 of reduction in ductile to brittle transition temperature (DBTT) with grain size and (d) fractal analysis63 of the fracture surface revealing the tortuous path being responsible for the improvement of DBTT in fine grain size. N and T refers to normalized and tempered.
114
Ferritic Steels and Advanced Ferritic–Martensitic Steels
energy required for propagation of cracks. These calculations were carried out (Figure 12(b)) for two different grain sizes. The prediction of finer grain size being favorable to reduce embrittlement was confirmed (Figure 12(c)) experimentally. The GBCD, that is, the distribution of various grain boundary types has been evaluated62 in modified 9Cr–1Mo ferritic steel using EBSD technique. The experimental observations confirmed the reduction of DBTT by 20 K with reduction in grain size. The fractal analysis of the fracture surface demonstrated (Figure 12(d)) that the tortuous path which cracks need to follow in fine grain sample is responsible63 for the observed reduction in the propensity for embrittlement. It is shown clearly that the low energy boundaries can be introduced in engineering materials in three different methods: preferential nucleation of low angle boundaries around twins or controlled recovery or orientation relations during phase transformation, if some of the variants happen to result in CSL boundaries. Significant improvements in properties using GBE have been achieved64 in many austenitic stainless steels, in contrast to ferritic steels. The major challenges in the application of GBE to Cr–Mo ferritic steels arise from the following factors: lower twinning probability, higher stacking fault energy, and limited variants with CSL boundaries during g ! a transformation during cooling.
4.03.5 Development of Advanced ODS Ferritic Steels In recent years, an attempt to increase the high temperature creep life of ferritics to 973 K and target burn-up of the fuel to 250 dpa, has enabled a ‘revisit’ to the concept of strengthening the steel using 5 nm particles of yttria (see Chapter 4.08, Oxide Dispersion Strengthened Steels), leading to the ODS ferritic steels. ODS ferritic steels are prospective candidate materials for sodium cooled fast reactors with peak burn-up of 250 dpa as well as GenIV and fusion reactors. Earliest developments of ODS steels can be traced to the efforts65 of Belgium in 1960s, followed by Japan66 since 1987, and France67 in the last decade. The ODS steels for fast and fusion reactors68,69 are in the R&D stage. The design of ODS steels for fast and fusion reactor applications is based on Fe–Cr–W–Ti– Y2O3, either the martensitic 9 or 12Cr or the ferritic 12Cr steels. The dispersoids which confer the high temperature creep life to the ferrite matrix are70
10 nm Figure 13 Z-contrast in the high angle annular dark field (HAADF) micrograph of dispersoids in oxide dispersion strengthened (ODS)-9Cr–1M0 ferritic steels, which are responsible for the superior high temperature creep behavior.
(Figure 13) in the size range of around 5 nm with a volume fraction around 0.3%. The yttria dissolves in it some amount of titanium, leading to the formation of mixed, complex oxide, namely TiO2Y2O3. The rationale for the choice of the matrix composition is as follows: Chromium: Choice of 9% Cr and 0.1% C ensures 100% martensite, during normalization of the steel. It is possible to ensure 100% martensite in 12% chromium steel by ensuring the carbon content to be above 0.1%. Ferritic ODS steels can be obtained in 12% chromium steels by lowering the carbon content to be less than 0.03%. Higher chromium provides the corrosion and decarburization resistance in sodium at 973 K, with acceptable oxidation resistance. Carbon: Addition of 0.1% carbon ensures 100% martensite in 9% Cr steels, thus ensuring absence of anisotropy during g ! a transformation. Higher amount of carbon would promote precipitation of M23C6, thus reducing the toughness. On the other hand, M23C6 along the lath boundaries offers the long-term microstructural stability of the lath structure. Nitrogen: The solubility of nitrogen in ferrite is very low. This is useful in non-ODS ferritic steels like T91, due to enhanced creep resistance by formation of V or Nb carbides/carbonitrides. But, in ODS steels, Ti is used for refining yttria. Hence, nitrogen content is restricted to 0.01%, preventing the formation of deleterious TiN compound.
Ferritic Steels and Advanced Ferritic–Martensitic Steels
Tungsten: Tungsten is a more effective solid solution strengthener than Mo, but at the cost of ductility. Tungsten stabilizes d-ferrite and accelerates formation of Laves phase, both of which cause reduction in toughness. Hence, it is optimized to 2.0%. Yttria: The most important constituent of ODS steels is the yttria, which enhances high temperature creep strength by pinning mobile dislocations and delays void swelling by acting as sinks for point defects produced during irradiation. The strength increase is accompanied by a concomitant loss of ductility and saturates around 0.4% yttria. Hence, it is optimized to 0.35%. Titanium: The major role of titanium in ODS steels is to refine the yttria particles (20 nm after mechanical alloying) to ultra-fine (2–3 nm) particles. The complex Y–Ti–O particle imparts the necessary high temperature creep strength. The beneficial effect of titanium saturates around 0.2%. Further increase introduces manufacturing problems of the tubes and hence titanium is chosen as 0.2%. Excess oxygen: Oxygen is present during processing of the ODS steels. The oxygen present in excess of the amount required for formation of required amount of Y–Ti–O complex leads to increase in tensile and creep strength. The Y–Ti–O complex oxides requires about 0.07 þ 0.01% excess oxygen. Argon: A strict control of argon (<0.002%) during processing of ODS steels is essential to avoid embrittlement due to formation of argon bubbles during irradiation. Minor elements: Nickel and manganese are to be reduced to maintain the A1 (a ! g on heating) temperature higher than the anticipated hot spot temperature. This enables the tempering temperature to be as high as possible. Silicon, phosphorous, and sulfur undergo RIS and cause embrittlement. Silicon also accelerates formation of deleterious phases like the Laves phase. Hence, their amounts are reduced to 0.05 and less. The processing route used worldwide is the powder metallurgy route of mechanically alloying prealloyed powders of Fe–Cr–W–Ti–C þ Ti2O3, followed by hot extrusion and rolling or hipping with final heat treatments. Commercial ODS steels (Table 7) have been developed demonstrating the standardizing of fabrication technologies. The ferritic–martensitic ODS steels have been developed by adjusting the contents of chromium and carbon. The ferritic ODS steels, with 12% chromium and carbon content less than 0.02%, derive71 their high temperature creep strength basically from
115
Table 7 List of few commercial ODS ferritic steels and their chemistry Commercial name
Chemistry
MA956 MA957 M11 M92 PM2000
Fe–20Cr–4.5Al–0.5Y2O3 Fe–14Cr–0.3Mo–Ti–0.27Y2O3 Fe–9Cr–Mo–0.37Y2O3 Fe–9Cr–Mo–0.30Y2O3 Fe–20CrAlTi–0.5Y2O3
the dispersoids. The ferrite matrix offers72 superior resistance against oxidation and corrosion while the major challenge appears to be the anisotropy73 of properties. The martensitic steels based on either 12% chromium with 0.1–0.2% carbon or the 9% chromium, derive their strength from the martensitic matrix and the dispersoids. The 9% chromium steel displays isotropic properties while it suffers from inferior corrosion resistance. The conventional joining technologies pose significant problems leading to coalescence of oxide particles. Hence, solid state bonding techniques74 like the pressurized resistance welding have been developed for joining the clad tube with the end plug of a fast reactor. Postweld heat treatments (PWHT) have also been developed to match the strength levels of the clad and the end plug. The in-service performance of ODS steels in flowing sodium has been found to be satisfactory, despite the formation of austenite layer on the surface of the clad tube due to the deposition of nickel. The thickness of the oxide layer in 12%Cr ODS steel was found to be only 50% of that in 9Cr ODS steels. Reactor irradiation experiments have been performed75,76 on ODS ferritic steels. The dispersoids were found74 to be stable up to a dose of 10 dpa in JOYO. The mechanical properties at 573 K after neutron irradiation were reported to be same as conventional ferritic martensitic steels. The studies on long-term in-service behavior and postirradiation behavior are being studied. Presently, the main challenges in this variety of ODS steels are the anisotropy observed in steels with chromium more than 12%, less oxidation resistance in steels with 9% chromium, fabrication procedure with cost-effectiveness, uniformity of dispersoids in all regions of the clad tube, stability of the dispersoids under irradiation, and the joining technologies. It is hoped that the above problems would be overcome in the near future and the ODS ferritic steels can be used in fast reactor core as clad and wrapper, for burn-up
116
Ferritic Steels and Advanced Ferritic–Martensitic Steels
beyond 250 dpa, at temperatures exceeding 973 K. Additionally, ODS ferritic steels are also being considered for fusion reactor applications. The rich experience in the development of fast reactor materials would enable launching the advanced ferritic steels for fusion technology, in a shorter time span.
The HAZ comprises coarse prior-austenitic grain martensite, fine prior-austenitic grain martensite and an intercritical structure, as one traverses from the weld fusion interface toward the unaffected base metal. This is dictated by the peak temperatures experienced by the base metal during weld thermal cycle and the phase transformation characteristics of the steel. It has been established that the localized microstructural degradation in the intercritical region of HAZ is mainly responsible for the premature creep-rupture strength of Cr–Mo weld joint and can be overcome if residual stresses of the weld are adequately relieved by PWHT. The lower creep-rupture strength of weld joint than the base metal is due38,77 to the different types of cracking developed during creep exposure. Four types of cracking have been identified (Figure 15) in Cr–Mo steel weld joint. They have been categorized as Type I, Type II, Type III, and Type IV. The Type I and Type II cracks originate in the weld metal, propagate either through the weld metal itself (Type I) or cross over in the HAZ (Type II). The Type III cracking occurs in the coarse grain region of HAZ and can be avoided by refining the grain size. Type IV cracking nucleates and propagates in the intercritical/fine grain region of HAZ. Type IV failure occurs at longer creep exposure and higher test temperature, by coalescence of fine cavities leading to microcracks (Figure 16(a)) and their eventual propagation to the surface.
4.03.6 Ferritic Steels for Out-of-Core Applications: Improvements in Joining An ambitious target of increasing the temperature and pressure of steam in many power plants has provided a high impetus for the development of steels with better high temperature properties. Very often, the weld joints play a crucial life limiting role in these components. One of the recurrent problems is the frequent failure of weldments due to Type IV cracking (see below), in weldments of ferritic steels subjected to creep loading. Another problem encountered during service exposure of joints of dissimilar ferritic steels is the failure due to the formation of hard brittle zone at the heat-affected zone (HAZ). Both these issues are discussed below. The modified 9Cr–1Mo steel fusion weld joint (Figure 14) consisting of base metal, deposited weld metal, and the HAZ produces a complex heterogeneous microstructure due to thermal cycle. The base metal and weld metal consist of a tempered martensite structure, with columnar grains in the weld metal.
Interface
Weld metal
Base metal
e
rfac
Inte 9Cr–1Mo weld metal
21/4Cr–1Mo Base metal 20 mm
20 mm
Base metal
ICZ
Weld metal
FGH AZ
CG HAZ
20 mm
HAZ CGHAZ
FGHAZ
ICZ
20 mm
20 mm
Figure 14 Schematic showing different zones in a ferritic steel weldment and optical micrographs obtained from weld metal, base metal, interface, coarse grained heat-affected zone (CGHAZ), fine grained heat-affected zone (FGHAZ), and intercritical zone (ICZ) of the weldment.
Ferritic Steels and Advanced Ferritic–Martensitic Steels
Weld metal
Base metal HAZ
II
HAZ
Base metal IV
(a)
III HAZ
Base metal
I
Weld metal
HAZ
5 Cm
(b)
Figure 15 Locations77 of different types of failure in weld geometry of the ferritic steels: (a) schematic representation and (b) experimental observation in creep tested weldment of 9Cr–1Mo steels.
Cavity Cavity associated with precipitate 10 mm
(a)
(b)
M23C6
[111] Z phase Z-phase
1 mm Figure 16 Type IV cracking in same sample as in Figure 15. (a) cavities in the intercritical region and (b) Z-phase77 in creep-tested 9Cr–1Mo steel. The inset shows the microchemistry of the Z-phase.
The type IV cracking susceptibility, defined as the reduction in creep-rupture strength of weld joint compared to its base metal, depends on the type of ferritic steel. 2.25Cr–1Mo steel is most susceptible to type IV cracking; whereas the plain 9Cr–1Mo steel is the least susceptible. At higher test temperature, the type IV cracking susceptibility is higher in modified 9Cr–1Mo steel than the plain steel. This is related77 to the precipitation of Z-phase (Figure 16(b)), a complex Cr (V, Nb) N particle, in the modified steel. The Z-phase grows rapidly at elevated temperatures during long term exposure, by dissolving the beneficial
117
MX types of precipitates. This promotes the recovery of the substructure with associated decrease in strength in the intercritical region of HAZ. Although it is difficult to completely eliminate Type IV cracking, several methods are being adopted to improve type IV cracking resistance. It is more severe in thick sections due to the imposed geometrical constraint. A design modification can be adopted to decrease the variation in tensile stresses across the welded section of the component or avoid joints in critical regions having high system stresses and relocate them in the less critical region. Strength homogeneity across the weld joint can also be improved by a suitable PWHT. An increase in width of the HAZ can reduce the stress triaxiality such that the soft intercritical region deforms with less constraint with the consequence of reduced creep cavitation, to minimize type IV cracking tendency. The width of the HAZ can be increased both by changing preheat and heat-input during welding. Another contrasting approach to overcome type IV cracking is to avoid or minimize the width of the HAZ, to eliminate the intercritical zone. This is being attempted by employing advanced welding techniques such as laser welding. The resistance against intercritical softening can also be improved by increasing the base strength of the steel with the addition of solid solution hardening elements such as W, Re, and Co and also by microalloying the steel with boron. Microalloying with boron retards the coarsening rate of M23C6 by replacing some of its carbon. The boron content needs to be optimized with the nitrogen content to avoid BN formation. Addition of Cu is also found to be beneficial. Copper is almost completely insoluble in the iron matrix and when added in small amounts, precipitates as nanosize particles to impart creep resistance. A suitable adjustment of the chemical composition of steel within the specification range also reduces the large difference in creep strength between the softened HAZ, the base metal, and the coarse grain HAZ of the joint. A weld joint of modified 9Cr–1Mo steel with low carbon, nitrogen, and niobium has been reported to possess creep strength comparable to that of the base steel. It is expected that a judicious combination of changes in chemistry and process variables would reduce the failures due to type IV cracking in weldments of ferritic steels subjected to creep loading. Another frequent problem78–81 is the formation of ‘hard brittle zone’ during service exposure of dissimilar joints between ferritic steels, leading to failures. The formation (Figure 17(a)) of microscopic layer of
118
Ferritic Steels and Advanced Ferritic–Martensitic Steels
hard, brittle zone along the HAZ in dissimilar weldments of steels is known to be responsible for the cold cracking, stress corrosion cracking, and higher frequency of failures of the weldments. This is one of the cases where modeling has enabled an in-depth understanding of the problem, in addition to providing an industrial solution to prevent the formation of brittle zone. The brittle layer at the interface between 9Cr–1Mo weld and 2.25Cr–1Mo base metal is shown77 to be a manifestation of a number of synergistic factors: (a) microstructural changes in regions close to the heat source during welding (b) migration of carbon during PWHT, driven by the gradient in its activity and (c) formation (inset in Figure 17(a))
of series of fine carbides when there is a local supersaturation of carbon. It has been possible to use modeling methods like Finite Difference Methods to predict the carbon profile across the weld region of 9Cr–1Mo and the base metal of 2.25Cr– 1Mo (Figure 17(b)), which were in good agreement with the profiles obtained using electron probe microanalysis. These calculations could be refined using Thermo-Calc and diffusion-controlled transformations (DICTRA) to take into account the simultaneous precipitation of carbides and diffusion of carbon. These computational methods were instrumental in predicting the methods to prevent the formation of hard zone in dissimilar joints of ferritic steels. Three elements which would repel carbon 0.50 Temp: 1023 K
0.45 M1
B
C
A M23C5
2.25Cr–1Mo
Soft zone
A
ppt zone
9Cr–1Mo
500 nm
Carbon concentration (wt%)
(b)
0.35 0.30
21/4Cr–1Mo
9Cr–1Mo
0.25 0.20 HZ
0.15
SZ
0.10 0.05
100 m
(a)
1h 15 h
0.40
0.00 -0.1
0.1
0.0 ‘x’ (mm)
(b)
Cmaximum in hard zone (wt%)
0.5 Ni Co Cu
0.4
7 6.5 6
0.3
5.5
Z
5 4.5
0.2
4 6 5
5.0, 5.5, 5.5
0.1
(d) 0.00
(c)
0.02
0.04 0.06 d (mm)
0.08
Y
4
6 5.5
3 4.5
5
6.5
7
X
0.10
Figure 17 Modeling in preventing formation of hard zone in dissimilar joints of ferritic steels: (a) Optical micrograph78 of formation of hard zone in the heat-affected zone between 9Cr–1Mo weld and the 2.25Cr–1Mo base metal, exposed to 1023 K for 15 h. The hard zone is marked as A and the soft zone as B. The inset shows the transmission electron micrograph of individual carbides in the hard zone. (b) Finite difference method calculation78 predicting the diffusion profile of carbon in the same weld geometry. (c) Variation79 of amount of carbon in the hard zone versus thickness of the ‘diffusion barrier’ introduced between the 9Cr–1Mo and the 2.25Cr–1Mo to prevent the formation of hard zone and (d) positions81 of carbon atoms in a bcc iron lattice calculated using molecular dynamics.
Ferritic Steels and Advanced Ferritic–Martensitic Steels
atoms, that is, with the positive interaction energy were chosen for this purpose. Figure 17(c) shows78 the comparison of three different metals, Ni, Cu, and Co in preventing the formation of hard zone. Experimental confirmation was obtained79 using interlayer between the two dissimilar ferritic steels. Further insight could also be arrived81 at in the diffusion behavior (Figure 17(d)) of carbon interstitial in the lattices of bcc iron and fcc nickel using molecular dynamics. These calculations could demonstrate that the activation energy for diffusion of carbon in a fcc nickel lattice is higher than bcc iron. This sluggish diffusion kinetics is due to the repulsive potential of nickel toward carbon, which is the main reason for the choice of nickel as the most effective diffusion barrier between the two ferritic steels. Thus, an industrial solution to prevent the formation of brittle zone in joints of dissimilar ferritic steels after service exposure could be arrived at, based on an in-depth understanding of the interaction between the lattice potentials of atoms. It has been demonstrated in the above studies that modeling methods could be used most effectively to reduce the experimental time required for overcoming an industrial problem. Experimental benchmarking was required only for final confirmation of the predictions. These trends are becoming more common in almost all problems in materials technology, in recent years, be it atomistic mechanisms or fabrication technologies or prediction of life of components. It is hoped that this approach of knowledgebased design of materials would gradually replace the time consuming empirical methods of today.
4.03.7 Summary Future trends in the global fast reactor industry are toward higher operating temperatures, higher burnup (250 GWd t1), higher breeding ratios (1.4) and longer lifetime for reactor (60–100 years). These goals require several developments in materials science and technology across all components of nuclear plants, especially for core component materials. Ferritic steels have a much better void swelling resistance compared to currently used austenitic stainless steels and are capable of enhancing the burn-up of the fuel up to about 200 GWd t1. Ferritic–martensitic steels based on 9–12% Cr compositions exhibit the highest swelling resistance and a number of commercial swelling resistant materials have been marketed. The principles behind the
119
design of swelling resistant ferritic steels for core components of fast reactors have been discussed. However, their use is rendered difficult due to their poorer creep strengths at temperatures higher than 873 K. Improvement of higher temperature tensile and creep strengths in these alloys will enable us to achieve higher temperatures, in addition to higher burn-up, thus improving the economics of nuclear power production. Presently, the reduced creep strength of 9–12Cr ferritic steels at temperatures above 798 K, has restricted their use to certain low stressed components such as subassembly wrappers. Another crucial problem is ‘embrittlement’ in ferritic steels. The mechanisms and methods which are being attempted to overcome embrittlement problems are discussed. Alloy development programmes are in progress to explore ferritic–martensitic oxide dispersion strength variants, for higher target burn-up of 250 dpa, with enhanced high temperature (973 K) capability, by improving mechanical properties. Conventional alloy melting routes will have to be abandoned in favor of powder metallurgy techniques of ball-milling, hot isostatic pressing, and hot extrusion for the synthesis of these ODS steels. Process optimization for the development of 9Cr-based ferritic–martensitic steels strengthened by a fine dispersion of yttria nanoparticles has been completed. The major concerns in this family of ferritic or ferritic–martensitic steels are the anisotropy of properties in ferritic 12Cr steels or oxidation resistance in 9Cr steels, fabrication procedure, microstructural stability under irradiation, and dissolution during back-end technologies. Materials science, engineering, and technology have become an integral part of the aspiration of the nuclear community to improve the economic viability of fast reactors. One of the major concerns in the alloy development programmes has been the unacceptably long time taken to launch newer materials. It is expected that the current trends in materials development, through intense international collaborations and increased role of modeling in materials behavior, would certainly reduce the time and cost of alloy development programmes for future reactors.
References 1. 2.
Baldev Raj; Vijayalakshmi, M.; Vasudeva Rao, P. R.; Rao, K. B. S. MRS Bull. 2008, 33, 327–337. Seren, J. L.; Levy, V.; Dubuisson, P.; et al. In Effects of Radiation Materials ASTM STP 1125, Stoller, R. E.,
120
3. 4. 5.
6. 7. 8. 9. 10.
11. 12. 13. 14. 15. 16. 17. 18. 19. 20. 21. 22. 23.
24. 25. 26.
27. 28.
29.
Ferritic Steels and Advanced Ferritic–Martensitic Steels Kumar, A. S., Gilles, D. S., Eds.; ASTM: PA, 1992; pp 1209–1233. Little, E. A.; Stow, D. A. J. Nucl. Mater. 1979, 87(1), 25–39. Baldev Raj; Vijayalakshmi, M. To appear in Journal of Metals. Vijayalakshmi, M.; Baldev Raj. Phase Evolution Diagrams – A New Approach to Mean Metal Temperature of Ferritic Components; University Press: Hyderabad, India, 2006. Klueh, R. L.; Harries, D. R. High Chromium Ferritic and Martensitic Steels for Nuclear Applications; ASTM: USA, 2001. Saroja, S.; Vijayalakshmi, M.; Raghunathan, V. S. Mater. Sci. Eng. 1992, 154 A, 59–65. Saroja, S.; Vijayalakshmi, M.; Raghunathan, V. S. J. Mater. Sci. 1992, 27, 2389–2395. Parameswaran, P.; Vijayalakshmi, M.; Raghunathan, V. S. J. Nucl. Mater. 1996, 232, 226. Thomas Paul, V.; Saroja, S.; Hariharan, P.; Rajadurai, A.; Vijayalakshmi, M. Pressure Vessels and Piping: Materials and Properties; Baldev Raj, Choudhary, B. K, Anish Kumar, Eds.; Narosa Publishing House: India, 2009. Jeyaganesh, B.; Raju, S.; Murugesan, S.; et al. Met. Mat. Trans. 2009, 40, 386–397. Raju, S.; Jeyaganesh, B.; Arun Kumar Rai; et al. J. Thermophys. 2010, 31, 399–415. Chandravathi, K. S.; Laha, K.; Bhanusankara Rao, K.; Mannan, S. L. Mater. Sci. Tech. 2001, 17(5), 559–565. Moitra, A.; Sreenivasan, P. R.; Mannan, S. L.; Vakil Singh Met. Mat. Trans. 2005, 36A, 2957–2965. Choudhary, B. K.; Bhanusankara Rao, K.; Mannan, S. L.; Kashyap, B. P. Int. J. Fatigue 1992, 14(4), 219–223. Albert, S. K.; Ramasubbu, V.; Parvathavarthini, N.; Gill, T. P. S. Sadhana 2003, 28(3–4), 383–393. Moitra, A.; Parameswaran, P.; Sreenivasan, P. R.; Mannan, S. L. Mater. Characterization 2002, 48, 55–61. Albert, S. K.; Ramasubbu, V.; Gill, T. P. S. Indian Weld. J. 2001, 342, 37–43. Was, G. S. Fundamentals of Radiation Materials Science; Springer Berlin Heidelberg: New York, 2007. Bykov, V. N.; Konobeev, Y. u. V. Atomic Energy 1977, 43(1), 618–624. Klueh, R. L.; Harries, D. R. High Chromium Ferritic and Martensitic Steels for Nuclear Applications; ASTM publication, 2001. Klueh, R. L. Int. Mater. Rev. 2005, 50(5), 287–310. Baldev Raj; Vijayalakshmi, M. Lecture notes during Joint ICTP/IAEA School on Physics and Technology of Fast Reactor Systems; The Abdus Salam International Centre for Theoretical Physics, Nov 9–20, 2009; http://cdsagenda5.ictp.trieste.it. Thomas Paul, V.; Saroja, S.; Vijayalakshmi, M. J. Nucl. Mater. 2008, 378, 271–281. Francis, J. A.; Mazur, W.; Bhadesia, H. K. D. H. Mater. Sci. Technol. 2006, 22(12), 1387–1395. Brown, C.; Levy, V.; Seran, J. L.; Ehrlich, K.; Roger, R. J. C.; Bergmann, H. In Fast Reactors and Related Fuel Cycles, FR-91. Proceedings of the Conference in Atomic Energy Society, Tokyo 1991; Vol. 1, Paper 7.5. Brown, C.; Lilley, R. J.; Crittenden, G. C. Nucl. Eng. 1994, 35, 122. Khabarov, V. S.; Dvoriashin, A. M.; Porollo, S. I. In Proceedings of Technical Committee Meeting on Influence of High Dose Irradiation on Advanced Reactor Core Structural and Fuel Materials, IAEA-TECDOC-1039, IAEA, Vienna, 1998, 139. Poplavsky, V. M.; Zabudko, L. M. In Proceedings of Technical Committee Meeting on Influence of High Dose
30. 31. 32. 33. 34. 35. 36. 37. 38. 39. 40. 41. 42. 43. 44. 45. 46. 47. 48. 49. 50.
51. 52. 53. 54. 55. 56. 57. 58. 59.
60.
Irradiation on Advanced Reactor Core Structural and Fuel Materials, IAEA-TECDOC-1039, IAEA, Vienna, 1998, 7. Little, E. A. J. Nucl. Mater. 1979, 87, 11–24. Little, E. A.; Stow, D. A. J. Nucl. Mater. 1979, 87, 25–39. Little, E. A. Radiat. Effects 1972, 16, 135–140. Little, E. A.; Bullough, R.; Wood, M. H. Proc. Roy. Soc. Lond. 1980, 372A, 565–579. Vitek, J. M.; Klueh, R. L. Met. Trans. 1983, 14A, 1047–1055. Kai, J. J.; Klueh, R. L. J. Nucl. Mater. 1996, 230, 116–123. Lu, Z.; Faulkner, R. G.; Was, G.; Wirth, B. D. Scripta Mater. 2008, 58(10), 878–881. Klueh, R. L.; Alexander, D. J. J. Nucl. Mater. 1998, 258–263, 1269–1274. Bhadeshia, H. K. D. H. ISIJ Int. 2001, 41(6), 626–640. Abe, F.; Taneika, M.; Sawada, K. Int. J. Pressure Vessels Piping 2007, 84, 3–12. Ennis, P. J.; Czyrska-Filemonowicz, A. Operation Maintenance Mater. Issues 2002, 1(1), 1–28; Pub. European Technology Development Ltd., UK. Toloczko, M. B.; Garner, F. A. J. ASTM Int. 2004, 1, 4. Garner, F. A.; Toloczko, M. B. J. Nucl. Mater. 1997, 11, 252–261. Kohyama, A.; Hishinuma, A.; Gelles, D. S.; Klueh, R. L.; Dietz, W.; Ehrlich, K. J. Nucl. Mater. 1996, 233–237, 138–147. Hudson, J. A.; Druce, S. G.; Gage, G.; Wall, M. Theor. Appl. Fract. Mech. 1988, 10, 123–133. Hippsley, C. A.; Haworth, N. P. Mat. Sci. Tech. 1988, 4, 791–802. Qingfen, L.; Shanglin, Y.; Li, L.; Zheng, L.; Tingdong, X. Scripta Mater. 2002, 47, 389–392. Gelles, D. S. J. Nucl. Mater. 1994, 212–215, 714—719. Sathyanarayanan, S.; Moitra, A.; Samuel, K. G.; Sasikala, G.; Ray, S. K.; Singh, V. Mater. Sci. Eng. 2008, 488A, 519–528. Shamardin, V. K.; Golovanov, V. N.; Bulanova, T. M.; et al. J. Nucl. Mater. 1999, 271–271, 155–161. Klueh, R. L.; Alexander, D. J. Effect of Radiation on Materials: 16th International Symposium, ASTM STP 1175; Kumar, A. S, Gelles, D. S, Nanstad, R. K, Little, E. A., Eds.; ASTM: PA, 1993; pp 591–612. Karthikeyan, T. IGCAR, India, Personal Communique. Elliot, C. K.; Lucas, G. E.; Maiti, R.; Odette, G. R. J. Nucl. Mater. 1986, 141–143, 439–443. Odette, G. R.; Lucas, G. E.; Maiti, R. J. Nucl. Mater. 1987, 148, 22–27. Vitek, J. M.; Corwin, W. R.; Klueh, R. L.; Hawthorne, J. R. J. Nucl. Mater. 1986, 141–143(2), 948–953. Klueh, R. L.; Alexander, D. J. J. Nucl. Mater. 1998, 258–263(2), 1269–1274. Klueh, R. L.; Hashimoto, N.; Sokolov, M. A.; Maziasz, P. J.; Shiba, K.; Jitsukawa, S. J. Nucl. Mater. 2006, 357(1–3), 169–182. Shamardin, V. K. In Proceedings of International Conference on Radiaiton Materials Science, Alushta, USSR, 1990, 3. Baldev Raj; Saroja, S.; Laha, K.; Karthikeyan, T.; Vijayalakshmi, M.; Rao, K. B. S. J. Mater. Sci. 2009, 44(9), 2239–2246. Watanabe, T.; Tsurekawa, S.; Zhao, X.; Zuo, L. In Proceedings of International Conference on Microstructure and Texture in Steels and Other Materials, Jamshedpur, India, Feb 2008; Haldar, A., Satyam, S., Bhattacharjee, D., Eds.; Springer Verlag: London, 2009; pp 43–82. Karthikeyan, T.; Saroja, S.; Vijayalakshmi, M.; Murthy, K. P. N. Frontiers in Design of Materials; Baldev Raj, Ranganathan, S., Mannan, S. L., Rao, K. B. S.,
Ferritic Steels and Advanced Ferritic–Martensitic Steels
61. 62. 63.
64. 65. 66. 67.
68.
69. 70.
Mathew, M. D., Shankar, P., Eds.; 2006, University Press: Hyderabad, India. Karthikeyan, T. Annals of the Indian National Academy of Engineering, VI, 2009; pp 189–193. Karthikeyan, T.; Thomas Paul, V.; Mishra, S. K.; Saroja, S.; Vijayalakshmi, M.; Samajdar, I. Metal. Mater. Trans. 2009, 40A, 2030–2032. Karthikeyan, T.; Saroja, S.; Rachael Reena, S.; Vijayalakshmi, M. In Presented in Annual Technical Meeting of The Indian Institute of Metals, Kolkatta, Nov 14–16, 2009. Parvathavarthini, N.; Mulki, S.; Dayal, R. K.; Samajdar, I.; Mani, K.; Baldev Raj Corrosion Sci. 2009, 51(9), 2144–2150. Huet, J. J. Powder Metall. 1967, 10, 108–115. Ukai, S.; Nishida, T.; Okuda, T.; Toshitake, T. J. Nucl. Mater. 1998, 258–263, 1745–1749. Dubuisson, P.; Schill, R.; Hugon, M. P.; Grislin, I.; Seran, J. L. 18th International Symposium on Effects of Radiation on Materials, ASTM STP 1325; Nanstad, R. K., Hamilton, M. L., Garner, F. A., Kumar, A. S., Eds.; ASTM: West Conshohocken, 1999; pp 882–898. Huet, J. J.; Casteels, F.; De Wilde, L.; Leroy, V.; Snykers, M.; Van Asbroeck, Ph. Properties of ferritic steels strengthened for use as fuel cladding in fast reactors in fuel for fast reactors; IAEA: Vienna, 1974, 215–225; Vol. II. Odette, G. R.; Alinger, M. J.; Wirth, B. D. Ann. Rev. Mater. Res. 2008, 38, 471–503. Saroja, S.; Dasgupta, A.; Divakar, R.; et al. In Joint EC-IAEA Topical Meeting on Development of New
71. 72. 73. 74. 75. 76. 77. 78. 79. 80. 81.
121
Structural Materials for Advanced Fission and Fusion Reactor Systems, Oct 5–9, 2009; Fusion for Energy, Barcelona, Spain; J. Nucl. Mat. 2011, 409(2), 131–139. Ukai, S.; Kaito, T.; Ohtsuka, S.; Narita, T.; Fujiwara, M.; Kobyashi, T. ISI Int. 2003, 43(12), 2038–2045. Kaito, T.; Narita, T.; Ukai, S.; Matsuda, Y. J. Nucl. Mater. 2004, 329–333, 1388–1392. Ukai, S.; Okuda, T.; Fujiwara, M.; Kobayashi, T.; Mizuta, S.; Nakashima, H. J. Nucl. Sci. Technol. 2002, 39(8), 872–879. Seki, M.; Hirako, K.; Kono, S.; Kihara, Y.; Kaito, T.; Ukai, S. J. Nucl. Mater. 2004, 329–333(2), 1534–1538. Akasaka, M.; Yamashita, S.; Yoshitake, T.; Ukai, S.; Kimura, A. J. Nucl. Mater. 2004, 329–333, 1053–1056. Yamashita, S.; Oka, K.; Ohnubi, S.; Akasaka, N.; Ukai, S. J. Nucl. Mater. 2002, 307–311, 283–290. Laha, K.; Chandravathi, K. S.; Parameswaran, P.; Rao, K. B. S. Metall. Trans. 2009, 40, 386–397. Sudha, C.; Thomas Paul, V.; Terrance, A. L. E.; Saroja, S.; Vijayalakshmi, M. Weld. J. 2006, 85(3), 71s. Anand, R.; Sudha, C.; Karthikeyan, T.; Terrance, A. L. E.; Saroja, S.; Vijayalakshmi, M. Trans. Indian Inst. Met. 2008, 616, 483–486. Anand, R.; Sudha, C.; Karthikeyan, T.; Terrance, A. L. E.; Saroja, S.; Vijayalakshmi, M. J. Mater. Sci. 2009, 44(1), 257–265. Anand, R.; Karthik, V.; Vijayalakshmi, M. Proceedings of ISRS, IIT Madras, Dec 20–24, 2010.
4.04
Radiation Effects in Nickel-Based Alloys
R. M. Boothby National Nuclear Laboratory, Harwell, Oxfordshire, UK
ß 2012 Elsevier Ltd. All rights reserved.
4.04.1
Introduction
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4.04.2 4.04.2.1 4.04.2.2 4.04.2.3 4.04.3 4.04.4 4.04.4.1 4.04.4.2 4.04.5 4.04.5.1 4.04.5.2 4.04.6 References
Void Swelling Compositional Dependence of Void Swelling Void-Swelling Models Swelling Behavior of Neutron-Irradiated Nimonic PE16 Irradiation Creep Microstructural Stability Dislocation Structures Precipitate Stability Irradiation Embrittlement Fast Neutron Irradiation Experiments Helium Implantation Experiments Concluding Remarks
124 124 129 133 136 138 138 139 140 140 145 147 148
Abbreviations AGR DFR EBR-II EDX HVEM N/2 NRT OA PFR PS SIPA ST STA TEM UTS VEC
Advanced gas-cooled reactor Dounreay Fast Reactor Experimental Breeder Reactor-II Energy dispersive X-ray High-voltage electron microscope dpa calculated according to half-Nelson model dpa calculated according to Norget, Robinson, and Torrens model Overaged Prototype Fast Reactor Proof stress Stress-induced preferred absorption Solution treated Solution treated and aged Transmission electron microscope Ultimate tensile strength Variable energy cyclotron
4.04.1 Introduction Research into the effects of irradiation on nickelbased alloys peaked during the fast reactor development programs carried out in the 1970s and 1980s. Interest in these materials focused on their high resistance to radiation-induced void swelling compared to austenitic steels, though a perceived susceptibility to
irradiation embrittlement limited their application to some extent. Nevertheless, the Nimonic alloy PE16 was successfully used for fuel element cladding and subassembly wrappers in the United Kingdom, and Inconel 706 was utilized for cladding in France. Both of these materials are precipitation hardened and consequently have high creep strength, and much research and development of alternative alloys was directed toward maintaining swelling resistance and creep strength while aiming to alleviate, or at least understand, irradiation embrittlement effects. There has been some revival of interest in nickel-based alloys for nuclear applications, and various aspects of radiation damage in such materials have recently been reviewed by Rowcliffe et al.1 in the context of Generation IV reactors, and by Angeliu et al.2 in consideration of their use for the pressure vessel of the Prometheus space reactor. Nickel-based alloys are also candidate structural materials for molten salt reactors, for which resistance to corrosion by molten fluoride salts and high-temperature creep strength are prime requirements, though intergranular attack by the fission product tellurium and irradiation embrittlement due to helium production are potentially limiting factors for this application.3 This chapter focuses on the void swelling behavior, irradiation creep, microstructural stability, and irradiation embrittlement of precipitation-hardened nickel-based alloys. Fundamental to all of these effects are the basic processes of damage production 123
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Radiation Effects in Nickel-Based Alloys
by the creation of vacancies and interstitial atoms in displacement cascades, and the ways in which these point defects migrate and interact with, causing the redistribution of, solute atoms. Detailed discussions of damage processes and radiation-induced segregation are beyond the scope of this chapter but these topics will be introduced where necessary, particularly in relation to void swelling models. More detailed reviews are given in Chapter 1.01, Fundamental Properties of Defects in Metals; Chapter 1.03, Radiation-Induced Effects on Microstructure; Chapter 1.11, Primary Radiation Damage Formation; Chapter 1.12, Atomic-Level Level Dislocation Dynamics in Irradiated Metals and Chapter 1.18, Radiation-Induced Segregation. Typical compositions of nickel-based alloys and some precipitation-hardened steels, which are considered in this chapter, are shown in Table 1. Alloy compositions are generally given in weight percent throughout this chapter unless stated otherwise. Precipitation-hardened alloys may be utilized in a number of different heat-treated conditions, which are generally abbreviated here as ST (solution treated), STA (solution treated and aged), and OA (overaged). Further information on the material properties of nickel alloys is given in Chapter 2.08, Nickel Alloys: Properties and Characteristics. Neutron fluences are generally given for E > 0.1 MeV unless indicated otherwise. Atomic displacement doses (dpa) are generally given in NRT (Fe) units, although the half-Nelson (N/2) model was Table 1
widely used particularly in the United Kingdom in the 1970s6. The exact relationship between these units will vary depending on the neutron spectrum (which may differ, not only from one reactor to another, but also depending on location within a reactor), but approximate conversion factors for fast reactor core irradiations are 1026 n m2 ðE > 0:1MeVÞ ¼ 5dpa NRTðFeÞ ¼ 6:25dpa ðN=2Þ
4.04.2 Void Swelling 4.04.2.1 Compositional Dependence of Void Swelling Nimonic PE16 was first identified as a low-swelling alloy in the early 1970s. Void swelling data derived from density measurements on fuel pin cladding materials from the Dounreay Fast Reactor (DFR) were reported by Bramman et al.7 and were complemented by electron microscope examinations described by Cawthorne et al.8 Swelling in STA PE16 was found to be lower than in heat-treated austenitic steels and comparable to cold-worked steels. Comparison of data for PE16 and FV548 (a Nb-stabilized austenitic steel) irradiated under identical conditions in DFR to a peak neutron fluence of 6 1026 n m2 indicated that the lower swelling of PE16 was due to smaller void concentrations at irradiation temperatures up to 550 C and reduced void sizes at higher
Nominal compositions (wt%) of commercial and developmental nickel-based alloys
Alloy
Ni
Cr
Mo
Ti
Al
Nb
Mn
Si
C
Nimonic PE16 Inconel 706 Inconel 718 Inconel 600 Inconel 625 Incoloy 800 Hastelloy X D21a D25a D66a D66b D68a D68b PE16 matrix Incoloy DS Alloy 7817 Alloy 7818
43 41.5 52.5 75 61 34 48 25 30 45 40 45 34 36 39 40 40
16.5 16 19 16 22 20.5 21 8.4 10.5 12 11 12 12.5 20 18 15 15
1.1 – 3.0 – 9.0 – 9.0 1.0 3.7 3.0 2.0 – – 4.0 – 3.2 3.0
1.2 1.8 0.9 0.3 0.3 0.4 – 3.3 1.8 2.5 3.0 1.8 1.6 – 0.04 2.0 0.3
1.2 0.2 0.5 0.2 0.3 0.4 – 1.7 1.3 2.5 1.5 0.4 0.25 – 0.02 0.9 –
– 2.9 5.2 – 3.5 – – – – – – 3.6 2.8 – – – 3.0
0.1 0.2 0.2 0.2 0.2 0.9 0.5 1.0 1.0 – 0.2 0.3 0.2 0.1 1.0 0.2 0.2
0.2 0.2 0.2 0.2 0.2 0.5 0.5 1.0 1.0 0.5 0.5 0.4 0.4 0.2 2.0 0.5 0.5
0.05 0.03 0.04 0.08 0.05 0.07 0.10 0.04 0.04 0.03 0.04 0.03 0.02 0.07 0.08 0.02 0.02
a
Composition indicated by Yang et al.4 Composition indicated by Toloczko et al.5
b
Other
0.3Cu 0.5Cu 0.5W, 2.0Co
Fe Bal. Bal. Bal. Bal. Bal. Bal. Bal. Bal. Bal. Bal. Bal. Bal. Bal. Bal. Bal. Bal. Bal.
125
Radiation Effects in Nickel-Based Alloys
energy Ed ¼ 40 eV). In addition to precipitationhardened alloys, including PE16 and Inconel 706, this experiment included nonhardenable high-Ni alloys, such as Inconel 600 and Hastelloy X, a range of commercial steels, and Fe–Cr–Ni ternary alloys containing 15% Cr and 15–35% Ni. The alloys were preimplanted with 15 appm helium prior to ion bombardment, and the irradiation temperature was chosen as being close to the peak swelling temperature for ionirradiated austenitic steels. The extent of void swelling was determined by electron microscope examinations in low-swelling alloys, but was estimated from stepheight measurements (comparing the surfaces of irradiated and nonirradiated regions) in high-swelling materials. As illustrated in Figure 1, the results showed negligible swelling (<0.1%) in PE16, Inconel 706, Hastelloy X, and the Fe–15Cr–35Ni ternary alloy, low swelling (<1%) in other high-Ni alloys, but high swelling (generally >20%) in austenitic steels. In commercial alloys containing 18% Cr, minimum swelling occurred at Ni contents of about 40–45%. Although void diameters generally appeared to be smaller in the Ni-based alloys than in austenitic steels, the main factor accounting for reduced swelling was a much lower void concentration. In the ternary alloys, reducing the Ni content from 35% to 30% resulted in
60 Commercial alloys Fe–15Cr–Ni alloys 50 1
Commercial alloys: 1 Type 304 2 Type 321 3 Type 316 4 Type 318 5 12R72HV 6 A286 7 Incoloy 800 8 Inconel 706 9 Nimonic PE16 10 Hastelloy X 11 Inconel 625 12 Inconel 600 13 Inconel 702
2
40 Swelling (%)
temperatures. At around the same time, Hudson et al.9 compared the swelling behavior of PE16, type 316 steel, and pure nickel, using 20 MeV C2+ ion irradiations in the Harwell VEC (variable energy cyclotron). The materials were implanted with 10 appm (atomic parts per million) of helium prior to ion bombardment to peak displacement doses >200 dpa (N/2) at 525 C. Void swelling in 316 steel and nickel exceeded 10% at the highest doses examined, compared to 0.5% in PE16. Void nucleation appeared to occur earlier in nickel (at 0.1 dpa) than in PE16 or type 316 (2 dpa), but the peak void concentration was higher by a factor of about 10 in the austenitic steel than in nickel or PE16. Hudson et al.9 originally attributed the swelling resistance of PE16 to the presence of the coherent, ordered face-centered cubic, Ni3(Al,Ti) g0 precipitates, which were thought either to trap vacancies and interstitials at their surface, thereby enhancing pointdefect recombination, or to inhibit the climb of dislocations, thereby preventing them from acting as preferential sinks for interstitial atoms. In support of the first of these two suggested mechanisms, Bullough and Perrin10 argued that the surface of a coherent precipitate would be a more effective trapping site than an incoherent one where the identity of the point defects would immediately be lost (and where, as a consequence, void nucleation was likely to occur). The efficiency of point defect trapping would be expected to be greater the higher the total surface area of the g0 precipitates, that is, to be inversely proportional to the precipitate size at constant volume fraction. On the other hand, the second mechanism proposed by Hudson et al. should be most effective at an intermediate particle size where dislocation pinning is strongest. Support for the latter process was provided by Williams and Fisher11 from HVEM (high-voltage electron microscope) irradiations of PE16 at a damage rate of about 102 dpa s1 at 600 C, where the swelling rate was higher at small (3 nm) and large (70 nm) g0 particle diameters than at intermediate sizes of about 20 nm. However, it is now considered that any effect that the g0 precipitates may have on the swelling resistance of Nimonic PE16 is secondary to that of the matrix composition. The generally low-swelling behavior of Ni-based alloys compared to austenitic steels was shown by Johnston et al.12 following bombardment with 5 MeV Ni2+ ions at 625 C. The damage rate in these experiments was 102 dpa s1 and the displacement dose was originally estimated as 140 dpa but this was subsequently revised by Bates and Johnston13 to 116 dpa (based on displacement
3 4
30
5
20
10 6 0
7 8 9 10
0
10
20
30
11
40 50 60 Nickel (wt%)
12 13 70
80
90
Figure 1 Swelling versus nickel content of commercial alloys and ternary Fe–15Cr–Ni alloys bombarded with Ni2þ ions to a damage level of 116 dpa at 625 C. Reproduced from Johnston, W. G.; Rosolowski, J. H.; Turkalo, A. M.; Lauritzen, T. J. Nucl. Mater.1974, 54, 24–40.
126
Radiation Effects in Nickel-Based Alloys
an increase in overall swelling from <0.1% to 12%, although it was noted that the 35% Ni alloy showed a localized swelling of 5% in a region close to a grain boundary. Additional experiments reported by Johnston et al.12 indicated that the peak swelling temperature for PE16 irradiated with 5 MeV Ni2+ ions was 675 C, but even then, swelling at 116 dpa remained below 0.2%. Swelling data for a wider range of pure Fe–Cr–Ni austenitic alloys, with Cr contents up to 30% and Ni up to 100%, following Ni ion bombardment to 116 dpa at 675 C, were reported by Bates and Johnston.13 These results showed a strong dependence on both Cr and Ni, with the swelling increasing with increasing levels of Cr but being minimized at Ni contents of about 45–60%. Examination of the dose dependence of swelling in ternary alloys containing 15% Cr and 20–45% Ni showed that the incubation dose required for the onset of swelling increased with increasing Ni content. Furthermore, although high-swelling rates of the order of 1% per dpa were attained in 20–35% Ni alloys, the swelling rate of the 45% Ni alloy remained low even at doses above 250 dpa. Following their earlier C2+ ion irradiation experiments, Hudson and coworkers moved to the use of 46.5 MeV Ni6+ ions to investigate void swelling behavior. This was considered preferable because the recoil spectra of high-energy Ni ions provided a better simulation of fast neutron damage, and because carbon implantation encouraged the formation of carbides which acted as void nucleation sites. A summary of some of the Ni ion irradiation work carried out by the Harwell group was given by Makin et al.14 No significant differences in the swelling behavior of Nimonic PE16 were evident between ST or aged conditions. Peak swelling in Ni6+ ionirradiated PE16 (preimplanted with 10 appm He) occurred at 625 C, where a swelling of 1.5% was recorded at 120 dpa (N/2). Void concentrations in PE16 were reported to be lower by a factor of about 5 than in similarly irradiated type 316 and 321 austenitic steels. A drawback of charged particle irradiation experiments for evaluating void swelling is that the evolution of other microstructural features may differ significantly from that during neutron irradiation (see also Chapter 1.07, Radiation Damage Using Ion Beams). In the case of Nimonic PE16, for example, the precipitation and/or redistribution of the g0 phase during long-term neutron exposure might be expected to influence swelling behavior. In order
to simulate swelling in a more appropriate microstructure, Bajaj et al.15 examined the effect of 4 MeV Ni2+ ion irradiation on PE16, which had been preconditioned by exposure to neutrons in Experimental Breeder Reactor-II (EBR-II). Reactor-conditioned samples had been exposed to neutron fluences in the range of 3–6 1026 n m2 (E > 0.1 MeV) at temperatures from 454 to 593 C. Swelling rates during Ni ion irradiations at 675 C were higher by a factor of about five in reactor-conditioned material than in a nonconditioned sample. The increased swelling rate was attributed to changes in the matrix composition resulting from an increased volume fraction of g0 in the reactor-conditioned material. Early attempts to account for the effects of matrix composition on void swelling focused on the stability of the austenite phase. Harries16 suggested that the swelling behavior of austenitic steels and nickelbased alloys could be rationalized in terms of their Ni and Cr equivalent contents (i.e., the relative austenite and ferrite stabilizing effects of their constituent elements), with the composition of highswelling alloys then falling into the (g þ s) phase field in the Fe–Cr–Ni ternary phase diagram. Harries postulated that the partitioning of solute elements into the sigma phase would have a detrimental effect on the swelling resistance of austenite. Watkin17 took a similar approach, but found that an improved correlation could be obtained using the concept of electron vacancy numbers rather than Ni and Cr equivalents. The average electron vacancy number, Nv, of the matrix is calculated from the atomic fractions of its constituents, with allowance being made for the precipitation of carbides and g0 (or g00, etc.), and has been widely used to predict the susceptibility of nickel-based alloys to the formation of intermetallic phases.18Nv was calculated from: Nv ¼ 0:66Ni þ 1:70Co þ 2:66Fe þ 3:66Mn þ 4:66ðCr þ MoÞ Watkin found that void swelling in a range of alloys with Ni contents up to 43%, which were irradiated in DFR to a peak dose of 30 dpa at 600 C, remained low for Nv below about 2.5 (corresponding to low susceptibility to s phase formation), but increased approximately linearly at higher Nv. However, as was clearly argued by Bates and Johnston,13 correlations based on sigma-forming tendency could not account for the minimum in swelling observed at about 45% Ni, since higher Ni contents should continue to be beneficial.
Radiation Effects in Nickel-Based Alloys
A better understanding of the swelling behavior of Fe- and Ni-based alloys resulted from a series of fast neutron irradiation experiments which were carried out in EBR-II in the early 1980s. Irradiation temperatures in these experiments ranged from about 400 to 650 C. Initial data for a range of commercial alloys, including ferritic and austenitic steels, as well as nickel-based alloys, were reported by Bates and Powell19 and Powell et al.,20 with higher dose data (up to a peak fluence (E > 0.1 MeV) of 25 1026 n m2, corresponding to 125 dpa) being reported by Gelles21 and Garner and Gelles.22 Swelling data for Fe–Cr–Ni ternary alloys, irradiated in EBR-II to a peak fluence of 22 1026 n m2 (110 dpa), were presented by Garner and Brager.23 The extent of void swelling in these experiments was determined by density change measurements. In general, alloys with nickel contents in the range of 40–50% exhibited the lowest swelling. Swelling in commercial nickel-based alloys was generally lower in ST than in aged conditions, this being attributed to the beneficial (though temporary) effect of minor elements remaining in solution and being able to interact with point defects19; subsequent precipitation during irradiation would be expected to reduce this benefit and the resulting densification, though small, would also effectively reduce the measured swelling. Swelling data for a number of ST alloys, which were irradiated in the AA-1 rig in EBR-II, are shown in Figure 2; data are shown for two withdrawals, at peak fluences of 14.7 1026 n m2 and 25.3 1026 n m2, with measurements for Inconel 600 and Inconel 625 reported at both fluence levels, data for Nimonic PE16 and Inconel 706 at the lower level, and data for Incoloy 800 and Hastelloy X at the higher level. The nickel contents of the alloys range from about 34% in Incoloy 800 to 75% in Inconel 600. Swelling remained relatively low in the three Inconel alloys and in PE16. However, both Incoloy 800 and Hastelloy X exhibited high swelling at some temperatures, with swelling in the latter reaching 80% at 593 C. The reason for such high swelling in neutron-irradiated Hastelloy X (nickel content 48%) is unclear, but it was noted that densification up to 3% occurred at the lower irradiation temperatures – indicating microstructural instability and possibly signaling changes in the composition of the matrix which may have affected the swelling behavior. (Note that Hastelloy X was identified as a low-swelling alloy in the Ni2+ ion irradiation experiments described by Johnston et al.12) Some data for different heat-treated conditions of PE16 at the higher fluence level were reported by
127
Garner and Gelles,22 and are compared for irradiations at 538 C (more or less corresponding to the peak swelling temperature for PE16 in the AA-1 experiment) with lower fluence data from Bates and Powell19 in Figure 3. The heat-treated conditions indicated in Figure 3 are ST (ST 4 h at 1080 C), A1 (ST and aged 16 h at 705 C), A2 (ST and aged 1 h at 890 C plus 8 h at 750 C), and OA (ST and aged 24 h at 840 C). Note that the silicon content of the PE16 used in these experiments was much lower at 0.01% than the level of 0.2% typically found in UK heats of the alloy. Overall, the data appear to show little effect of initial heat treatment on the swelling of PE16, except that the OA condition exhibited the most swelling (5.2%) at the higher fluence. Although it is clear that the swelling behavior of austenitic alloys is largely dependent on nickel content, there is ample evidence to show that minor solute additions can have significant effects. Much of the work on minor solutes has focused on steels similar to type 316, but some data are available for higher nickel alloys. For example, Mazey and Hanks24 used 46.5 MeV Ni6+ ion irradiations to examine the effects of Si, Ti, and Al additions on the swelling response of model alloys with base compositions approximating that of the matrix phase in PE16. Solute additions of 0.25% Si or 1.2% Ti reduced swelling, but the addition of 1.2% Al (in the absence of Si or Ti) markedly increased it. The beneficial effect of Si was believed to arise from its high diffusivity in solution (this is discussed further in Section 4.04.2.2), whereas that of Ti appeared to be related to the formation of Z phase (hexagonal-structured Ni3Ti). The addition of Al resulted in an increase in the concentration of voids, the surfaces of which were coated in a thin layer of the g0 phase (Ni3Al). A beneficial effect of Si on the swelling response of modified Incoloy DS alloys under Ni6+ ion irradiation was also reported by Mazey et al.25 However, it should be noted that high Si contents can give rise to the formation of radiationinduced phases which are enriched with Ni and Si, such as the Ni3Si form of g0 and the silicide G-phase (M6Ni16Si7, where M is usually Ti, Nb, or Mn). G-phase particles are generally found in association with large voids and their formation may therefore give rise to an increase in the swelling rate.26,27 Swelling data derived from density measurements for neutron irradiated, modified Incoloy DS alloys, with Si contents ranging from 0.19 to 2.05% (compared to a specified level of 1.9–2.6% in the commercial alloy), are compared with data for a ‘PE16 matrix
128
Radiation Effects in Nickel-Based Alloys
5.0
4.0
Fluence
15
Swelling (%)
3.5 3.0 10 2.5 2.0 1.5 5
1.0 0.5
Fluence (1026 n m−2, E > 0.1 MeV)
In 600 In 625 PE16 In 706
4.5
0.0 −0.5 350
400
450
500
550
600
650
0 700
Temperature (°C) 100
80
In 600 In 625 In 800 Hast X
Fluence
Swelling (%)
70
25
20
60 50
15
40 30
10
20 10
5
Fluence (1026 n m−2, E > 0.1 MeV)
90
0 −10 350
400
450
500
550
600
650
0 700
Temperature (°C) Figure 2 Void swelling of nickel-based alloys irradiated in AA-1 rig in Experimental Breeder Reactor-II. Based on data from Bates, J. F.; Powell, R. W. J. Nucl. Mater.1981, 102, 200–213; Garner, F. A.; Gelles, D. S. In Effects of Radiation on Materials: 14th International Symposium; Packan, N. H., Stoller, R. E., Kumar, A. S., Eds.; American Society for Testing and Materials: Philadelphia, PA, 1990; Vol. II, pp 673–683, ASTM STP 1046.
alloy’ and Nimonic PE16 in Figure 4. The materials were all in ST condition apart from PE16 which was in an STA condition (aged 4 h at 750 C). The alloys were irradiated in the UK-1 rig in EBR-II to fluences in the range of 9–16 1026 n m2 (E > 0.1 MeV) at temperatures of 390–640 C. These data are previously unpublished except those for STA PE16 (heat DAA 766) which were reported by Boothby.28 Swelling in the modified Incoloy DS alloys generally decreased with increasing Si content. The 0.19% Si
alloy exhibited high swelling at all temperatures with indications of swelling peaks at about 440 and 640 C. Increased Si levels tended to suppress the high temperature swelling peak and reduce the magnitude of swelling at lower temperatures. The PE16 matrix alloy containing 0.24% Si exhibited a high temperature swelling peak but moderate swelling below 550 C, suggesting a beneficial effect of Mo (this being the main compositional difference between the PE16 matrix alloy and the modified Incoloy DS alloys)
Radiation Effects in Nickel-Based Alloys
4.0
at lower temperatures. However, swelling in the PE16 matrix alloy remained significantly higher at all temperatures than in STA Nimonic PE16 (containing 0.15% Si), indicating a significant benefit of the g0 forming elements Al and Ti.
3.0
4.04.2.2
6.0 68 dpa
116 dpa
Swelling (%)
5.0
2.0 1.0 0.0 ST
A1
A2
OA
Figure 3 Effect of heat treatment on void swelling of Nimonic PE16 irradiated in Experimental Breeder Reactor-II at 538 C. Adapted from Bates, J. F.; Powell, R. W. J. Nucl. Mater.1981, 102, 200–213; Garner, F. A.; Gelles, D. S. In Effects of Radiation on Materials: 14th International Symposium; Packan, N. H., Stoller, R. E., Kumar, A. S., Eds.; American Society for Testing and Materials: Philadelphia, PA, 1990; Vol. II, pp 673–683, ASTM STP 1046.
12.0 15
8.0 10 6.0
4.0
5
Fluence (1026 n m–2, E > 0.1 MeV)
10.0
Swelling (%)
129
2.0
0.0 300
350
400
450 500 550 Temperature (°C)
PE16 STA PE16 Matrix DS (0.19 Si)
600
0 650
DS (0.56 Si) DS (1.05 Si) DS (2.05 Si)
Figure 4 Void swelling data derived from density measurements for Nimonic PE16, a PE16 matrix alloy, and modified Incoloy DS alloys, irradiated in the UK-1 rig in Experimental Breeder Reactor-II. Unpublished data from Boothby, R. M.; Cattle, G. C. Void Swelling in EBR-2 Irradiated Nimonic PE16 and Incoloy DS; FPSG/P(90)10, with permission from AEA Technology Plc.
Void-Swelling Models
Point defects created by atomic displacements are lost either through mutual recombination or by migration to sinks. Void swelling requires a mobile population of excess vacancies and can only occur over a limited temperature range, typically 350– 700 C in neutron-irradiated steels and nickel-based alloys. Rapid diffusion at higher temperatures reduces the concentration of radiation-induced vacancies to near thermal equilibrium levels. Recombination dominates at lower temperatures, where reduced vacancy mobility prevents the formation of voids as the necessary counter-migration of matrix atoms cannot occur. In the swelling regime, an increased bias for interstitials over vacancies at dislocation sinks gives rise to the surplus vacancies which agglomerate to form voids. The flux of point defects to sinks, including void surfaces, dislocations, and grain boundaries, results in the segregation of particular solute atoms at the sinks and the depletion of others. In austenitic steels and nickel-based alloys, it is generally found that nickel segregates at the point defect sinks. This is generally attributed to the inverse Kirkendall effect described by Marwick,29 whereby faster diffusing solutes such as Cr move in the opposite direction to the vacancy flux and are depleted at the sink, and slower diffusing solutes such as Ni are enriched. One of the earliest observations of nickel segregation at void surfaces due to the inverse Kirkendall effect was made by Marwick et al.30 in an alloy with a composition corresponding to that of the matrix phase in Nimonic PE16. (For more detailed discussions on radiationinduced segregation effects, see the reviews of Wiedersich and Lam,31 and Rehn and Okamoto.32) Venker and Ehrlich33 recognized that differences in the partial diffusion coefficients of alloy constituents might account for the effects of composition on swelling. Any effect of this kind would generally be expected to be more significant the larger are the differences between the partial diffusion coefficients of the alloy components. Garner and Wolfer34 examined Venker and Ehrlich’s conjecture and concluded that the addition of even small amounts of a fastdiffusing solute such as silicon to austenitic alloys would greatly increase the effective vacancy diffusion
130
Radiation Effects in Nickel-Based Alloys
coefficient (i.e., would enhance the diffusion rate for all matrix elements). The overall effect is analogous to an increase in temperature – resulting in an effective decrease in the vacancy supersaturation and hence a reduction in the void nucleation rate. This mechanism is generally accepted as the explanation for the beneficial effect of silicon in reducing swelling in austenitic steels and nickel-based alloys. Although this relies on the diffusion of silicon via vacancy exchange, silicon is also generally observed to segregate to point defect sinks and since it is an undersized solute, this is believed to occur by the migration of interstitial–solute complexes. There is, however, no reason to suppose that both diffusion mechanisms cannot operate simultaneously. Garner and Wolfer34 originally considered that since nickel diffuses relatively slowly in austenitic alloys, an increase in nickel content would have the opposite effect to silicon. However, a later assessment made by Esmailzadeh and Kumar,35 based on diffusion data reported by Rothman et al.,36 indicated that the void nucleation rate in Fe–15Cr–Ni alloys would decrease with an increase in nickel content from 20 to 45%. This result is obtained because, although nickel remains the slowest diffusing species, the effective vacancy diffusion coefficient of the system is calculated to increase at the higher nickel content. Esmailzadeh and Kumar’s calculations also confirmed the beneficial effect of silicon, with the addition of 1% Si predicted to be as effective in suppressing void nucleation as increasing the nickel content from 20 to 45%. Effects at nickel contents above 45% could not be examined due to a lack of appropriate diffusion data. As well as affecting the nucleation of voids, differences in the diffusion rates of the various solutes might also be expected to influence void growth. Simplistically, this can be thought of as being partly due to the segregation of slower diffusing solutes reducing the rate of vacancy migration in the vicinity of the voids. However, a further consequence of such nonequilibrium solute segregation was identified by Marwick,29 who realized that it would give rise to an additional vacancy flux which would oppose the radiation-induced flux to the sink. As discussed by Marwick, this additional flux (the Kirkendall flux) may itself be an important factor in limiting void growth, since it will reduce the probability of vacancy annihilation at sinks and increase the likelihood of point defect recombination. The effect of nickel content on void swelling was considered further in a model developed by Wolfer
and coworkers.37,38 The model examined the compositional dependence of the void bias and focused on the effects of nickel segregation at void surfaces. Wolfer’s model indicated that the compositional gradients produced by radiation-induced segregation give rise to additional drift forces which affect the point defect fluxes and thereby modify the bias terms. These additional drift forces arise from the effects of composition on point defect formation and migration energies, on the lattice parameter and the elastic moduli, and from the Kirkendall flux. Wolfer’s calculations for binary Fe–Ni alloys indicated that the effect of the Kirkendall flux is small for interstitials but significant for vacancies. Nevertheless, it was considered that the overall effect of compositional gradients on the bias terms is likely to be greater for interstitials than for vacancies due to other factors, particularly the effect of variations in the elastic moduli. As noted by Garner and Wolfer,39 an increase in the shear modulus in the segregated regions around voids would reduce the bias for interstitials and therefore help to stabilize voids. It is difficult to predict the significance of this effect in complex alloys, however, since depletion of Cr in the segregated region will tend to reduce the shear modulus, whereas enrichment of Ni in high-Ni alloys will tend to increase it.38 A more significant result of the model with regard to the effect of nickel on swelling is that there is a reversal in the sign of the Kirkendall force for vacancies in Fe–Ni alloys at 35% Ni. Below this level, vacancies are predicted to be attracted into regions of higher Ni concentration, but above it, the opposite occurs. Wolfer et al. considered that this effect may account for the dependence of swelling on Ni content in austenitic alloys containing less than 35% Ni. A generalized description of the swelling behavior of austenitic alloys, which was consistent with the model developed by Wolfer et al., was put forward by Garner40 (see also Chapter 4.02, Radiation Damage in Austenitic Steels). Garner’s ideas were largely based on the results of the EBR-II irradiation experiments and the earlier ion bombardment work of Johnston et al., both of which showed a strong dependence of swelling on nickel content. It was considered that swelling was characterized by a transient period followed by a regime in which the swelling rate became constant. In neutron-irradiated alloys, the swelling rate in the posttransient regime was generally found to be 1% per dpa. In swellingresistant alloys, however, it was argued that such high swelling rates might not be observed owing to extended transient periods. The duration of the
Radiation Effects in Nickel-Based Alloys
transient regime was shown to be dependent on alloy composition and could extend for many tens of dpa in low-swelling materials. The duration of the transient regime was implicitly linked to the completion of void nucleation but, at the time these ideas were put forward, relatively few measurements of void concentrations were available, as swelling data were mainly derived from dimensional or density changes. Factors that were proposed to account for the influence of nickel on the void nucleation rate included the effect on vacancy diffusivity described by Esmailzadeh and Kumar35; a possible correlation with the development of fine scale compositional fluctuations by a spinodal-like decomposition process (observed by Dodd et al.41 in ion-irradiated ternary Fe–Cr–Ni alloys); and an effect of nickel on the minimum critical radius for the formation of stable voids.42 Voids are unstable below a critical size, and will generally shrink unless stabilized by gas atoms; the minimum stable void radius is dependent on a number of factors, including temperature and defect bias, and Coghlan and Garner suggested that the compositional dependence of the vacancy diffusivity would also affect this critical size. In other words, it was considered that the transition from gas bubble to void would require a larger bubble size in highnickel alloys, particularly at relatively high temperatures in the swelling regime where void nucleation becomes increasingly difficult. Hoyt and Garner43 subsequently argued that the minimum critical void radius concept might account for the minimum in swelling found at the intermediate nickel contents, provided that a compositional-dependent bias factor for dislocations was also incorporated into the model. The compositional dependence of the bias factor arises from solute segregation, which reduces the strain energy of dislocations and decreases the ratio of the bias for interstitials compared to vacancies. It is of interest that early evidence for the operation of the bubble to void transition was obtained by Mazey and Nelson,44 who implanted Nimonic PE16 (STA condition) and a PE16 matrix alloy (ST condition) with 1000 appm He to produce a high density of gas bubbles before subsequent irradiation with 46.5 MeV Ni6+ ions. The PE16 matrix alloy used in this particular experiment was a low Si variant (<0.02 wt%) which was known to exhibit relatively high swelling. The mean bubble size following helium implantation at 625 C was higher by a factor of about two in the matrix alloy (11 nm diameter) than in the commercial PE16 alloy (5 nm diameter). Examination of the alloys following subsequent
131
irradiation also at 625 C revealed high swelling (12% at 60 dpa) with a uniform distribution of large voids but no remaining helium bubbles in the matrix alloy, and low swelling (1% at 60 dpa) with a bimodal distribution of bubbles plus voids in the standard PE16 alloy (see Figure 5). These results were interpreted as providing evidence for the concept of a critical stable void size, with only a small fraction of bubbles in the commercial PE16 alloy, but all of the bubbles in the matrix alloy, being sufficiently large to grow as voids. Although not specifically discussed by Mazey and Nelson, the compositional differences between the two alloys suggest that the presence of Si and/or the g0 forming solutes Al plus Ti may help to reduce void nucleation in PE16. The belief advanced by Garner,40 that sluggish void nucleation generally accounted for low swelling in nickel-based alloys, persisted for some time. However, data reported by Muroga et al.45,46 largely overturned this view. Muroga et al. carried out microstructural examinations of a series of EBR-II-irradiated Fe–15Cr–Ni ternary alloys with Ni contents ranging from 15 to 75 wt%, and of archived samples of similar alloys from the heavy-ion bombardment experiments of Johnston et al.12 Examination of alloys irradiated in EBR-II at 510 C showed that the saturation void concentration was dependent on nickel content and was minimized at 35–45% Ni, but revealed that there was no increase in void numbers in any of the materials above a fluence of 2.6 1026 n m2 (E > 0.1 MeV) (see Figure 6). Alloys containing 19% and 30% Ni exhibited high swelling rates at higher fluences, but swelling remained low in higher nickel alloys. Similar effects were found in the ion-bombarded samples, where, for example, it was shown that there was no significant change in the void concentration in Fe–15Cr–45Ni at doses above 50 dpa in irradiations at 675 C, yet a marked increase in swelling rate occurred above 120 dpa. Thus, contrary to earlier ideas, these investigations clearly demonstrated that the onset of a high swelling rate was not related to the cessation of void nucleation. It follows that the transition to a high rate of swelling must be due to an increase in the growth rate of existing voids. Muroga et al.45,46 observed that the total dislocation density in the irradiated Fe–15Cr–Ni alloys was only weakly dependent on nickel content. This suggested that at the intermediate nickel levels, where the void concentration was low, dislocations were weak sinks (for both vacancies and interstitials) relative to voids. In addition, it was observed that
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Radiation Effects in Nickel-Based Alloys
20
400 300
16
0.2 dpa
Bubbles
0.2 dpa
12 200
30 dpa
300 200
15 10
Bubbles Voids
100
5
0
0 Bubbles
300
60 dpa
15
Number of cavities in interval Nc ⫻ 10–19 m–3
0
Number of voids in interval Nv ⫻ 10–19 m–3
Number of bubbles in interval Nb ⫻ 10–19 m–3
8 100
Voids
4 0 16
30 dpa
12 8 4 0 16
60 dpa
12
200
10
100
5
4
0
0
8
0 0 (a)
5
10
15
20
25
30
35
40
45
50
55
0
60
Cavity diameter d nm
(b)
10 20 30 40 50 60 70 80 90 100 110 120 130 Cavity diameter, d (nm)
Figure 5 Histograms showing size distributions of bubbles/voids in (a) solution treated and aged Nimonic PE16 and (b) solution treated PE16 matrix alloy, irradiated with Ni6þ ions at 625 C to damage levels of 30 and 60 dpa following implantation with 1000 appm. He (producing 0.2 dpa) at the same temperature. Reproduced from Mazey, D. J.; Nelson, R. S. J. Nucl. Mater.1979, 85–86, 671–675.
dislocation loops persisted to higher doses at the intermediate nickel contents, indicating a lower growth rate for the loops – again implying an effect of nickel on dislocation sink strength. Based on these observations, Muroga et al. suggested that a reduced dislocation bias for interstitials at the intermediate nickel contents might explain the influence of nickel on the early stages of void development. An additional factor was required to account for the eventual transition to a high swelling rate. Microchemical data presented by Muroga et al.46 suggested that this transition was related to the depletion of nickel in the matrix owing to its enrichment at void surfaces. A complete description which incorporates all of the composition-dependent factors which affect the nucleation and growth of voids is lacking. However, there is a general consensus that the major influence of alloy composition arises through its effects on the effective vacancy diffusivity and on segregation arising from the inverse Kirkendall effect. A correlation between the magnitude of void swelling and radiation-induced segregation was shown for Fe–Cr–Ni
ternary alloys by Allen et al.48 The compositional dependence of radiation-induced segregation was determined using a model based on the earlier work of Marwick,29 which incorporates both the vacancy flux to the voids and the back-diffusion of vacancies due to the solute gradients set up by the inverse Kirkendall effect. Vacancy diffusivities for various alloy compositions were determined by the measurements of grain boundary segregation in protonirradiated samples. Swelling data for ion and neutron-irradiated alloys were then compared with the expected swelling propensity defined by the ratio of the forward to back diffusion terms calculated at the appropriate irradiation temperature. The materials for which vacancy diffusivity data were determined included Fe-based alloys containing 16–24% Cr and 9–24% Ni, and Ni-based alloys containing 18% Cr and either zero or 9% Fe. This work did not specifically examine 40–50% Ni alloys corresponding to the highest swelling resistance, though the results indicated that swelling generally decreased with increasing levels of nickel enrichment and chromium depletion at void surfaces.
Radiation Effects in Nickel-Based Alloys
133
4.0 Immersion density Garner and Kumar47
19Ni
3.5
DFR ~ 17 dpa DFR ~ 80 dpa
3.0
10
35Ni
75Ni 45Ni
Swelling (%)
20
0 Void density (m−3)
30Ni
1021
2
4 6 8 Fluence (1026 n m-2)
1.5
0.0 300
75Ni 30Ni 35Ni 45Ni
0
2.0
0.5
19Ni
1020
2.5
1.0
Fe–15Cr–XNi 510 ⬚C
350
400
450 500 550 Temperature (°C)
600
650
600
650
600
650
1022
10
12
Figure 6 Fluence dependence of swelling and void density of Fe–15Cr–Ni alloys irradiated in Experimental Breeder Reactor-II at 510 C. Swelling data obtained by immersion density measurements by Garner and Kumar47 are also shown. Reproduced from Muroga, T.; Garner, F. A.; Ohnuki, S. J. Nucl. Mater.1991, 179–181, 546–549.
Void concentration (m−3)
Swelling (%)
30
DFR ~ 17 dpa DFR ~ 80 dpa 1021
1020
1019
4.04.2.3 Swelling Behavior of NeutronIrradiated Nimonic PE16
1018 300
Brown et al.49 compared the swelling behavior of STA Nimonic PE16 and two cold-worked austenitic steels (M316 and Nb-stabilized FV548) which were irradiated in DFR as fuel pin cladding. Two PE16 clad pins were examined, which were irradiated to burn-ups of 6.1% and 21.6% of heavy atoms, corresponding to peak damage levels of about 17 and 80 dpa, respectively. Void concentrations and swelling were lower in PE16 than in the austenitic steels. Swelling data, void concentrations, and void diameters for the two PE16 pins examined by Brown et al. are shown in Figure 7. Note that Brown et al.49 only showed trend lines for void concentration and void size in the less highly irradiated pin and compared the swelling tendencies of the two pins; the individual data points were not plotted and those shown in Figure 7 are previously unpublished data obtained by Sharpe. Brown et al. stated that the void concentration in PE16 decreased with increasing irradiation temperature but did not alter greatly with an increasing dose above 17 dpa. It should be noted, however, that swelling measurements for the higher burn-up pin were restricted to temperatures
100 90 Void diameter (nm)
80
350
400
450 500 550 Temperature (°C)
DFR ~ 17 dpa DFR ~ 80 dpa
70 60 50 40 30 20 10 0 300
350
400 450 500 550 Temperature (°C)
Figure 7 Swelling data, void concentrations, and void diameters for Nimonic PE16 Dounreay Fast Reactor fuel pin cladding. Unpublished data from Sharpe, R. M. Void Swelling in Fast Reactor Irradiated Commercial High Nickel Alloy; DFMC/P(82)27, with permission from AEA Technology Plc.
below 525 C, so that a direct comparison of void concentrations in the two pins cannot be made at higher temperatures. Although there were fewer voids in PE16 than in the two steels, the voids
10.0
Swelling (%)
8.0
STA DAA766 OA DAA766 STA Z260D STA Z184
Fluence
15
10
6.0 4.0
5 2.0 0.0 300
350
400
450
500
550
600
0 650
600
650
600
650
Temperature (°C) 1022 Void concentration (m–3)
appeared to be homogeneously distributed and to have developed during the early stages of irradiation; once nucleated, the growth rate of voids in PE16 remained low. These observations are clearly contrary to early models which suggested that low swelling rates result from incomplete void nucleation and extended transient regimes. Rather, in agreement with the more recent observations of Muroga et al.,45,46 it appears that the swelling resistance of PE16 is due to a combination of a comparatively low saturation void concentration, which is reached at a relatively low displacement dose, and a low void growth rate. There does not appear to be any evidence of an accelerated swelling rate in PE16 once void nucleation is complete. Additional data on void concentrations in neutronirradiated PE16 are available from Cawthorne et al.,8 Sklad et al.,50 and Boothby.28 The results presented by Cawthorne et al. for PE16 fuel pin cladding irradiated in DFR to a peak fluence of 5.6 1026 n m2 (28 dpa) differ from those shown in Figure 7 in that, although void number densities are similar for irradiations at 380–520 C, void concentrations are about an order of magnitude higher at 350 C and 600–630 C. Such discrepancies might arise from uncertainty and/or variability in irradiation temperatures. Another possibility is that void nucleation was incomplete at the higher irradiation temperatures in the lower burn-up pin examined by Brown et al. Data from Sklad et al. show an increase in void numbers in unstressed PE16 specimens irradiated in EBR-II at 500 C from an average (for two differently heat treated conditions) of about 4 1019 to 1.2 1020 m3 with increasing fluence from 1.2 1026 to 4.0 1026 n m2 (E > 0.1 MeV), that is, from 6 to 20 dpa. In this case, the void concentration and overall swelling of 0.2% at 20 dpa remain below the levels shown in Figure 7 for the DFR-irradiated pin at 17 dpa; this may reflect the effect of stress on swelling for fuel pin cladding. Void swelling data determined from Transmission electron microscope (TEM) examinations of three heats of PE16 which were irradiated in the UK-1 rig in EBR-II are shown in Figure 8, which includes previously unpublished results for the low boron (4 ppm) heat Z184 as well as data for heats DAA766 and Z260D (with 18 and 70 ppm boron, respectively) which were reported by Boothby.28 Data are shown for all three heats in the STA condition (ST 1020 C and aged 4 h at 750 C) and for DAA766 in the OA condition (a multistage heat treatment that included aging at 900 C, slow cooling to 750 C, and then aging for 16 h at that temperature, resulting in the
Fluence (1026 n m-2, E > 0.1 MeV)
Radiation Effects in Nickel-Based Alloys
1021
STA DAA766 OA DAA766 STA Z260D STA Z184
1020
1019
1018 300
350
400 450 500 550 Temperature (°C)
100 90 Void diameter (nm)
134
80 70
STA DAA766 OA DAA766 STA Z260D STA Z184
60 50 40 30 20 10 0 300
350
400 450 500 550 Temperature (°C)
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
Radiation Effects in Nickel-Based Alloys
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
135
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
136
Radiation Effects in Nickel-Based Alloys
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
Radiation Effects in Nickel-Based Alloys
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
5.0
Displacement rate (10–7 dpa s–1) 10.0
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
137
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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Radiation Effects in Nickel-Based Alloys
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
Radiation Effects in Nickel-Based Alloys
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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Radiation Effects in Nickel-Based Alloys
below 600 C with an activation energy that (in agreement with theoretical predictions for a process governed by point defect recombination) was 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
Radiation Effects in Nickel-Based Alloys
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
1200
Yield at Ti
18 16 14 12
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10 600
8 6
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Total elongation (%)
Elong. at Ti
1000 Yield stress (MPa)
141
4 200 2 0 400
450
500
550
600
650
700
750
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)
1000
14 12
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16
4 200 2 0 400
450
500
550 600 650 Temperature (°C)
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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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Radiation Effects in Nickel-Based Alloys
precipitate coarsening resulted in lower tensile strength. Elongations to failure for tests carried out at the irradiation temperature were between 1.5% and 3% up to 625 C, compared to >8% in 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
Total elongation (%)
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.
Radiation Effects in Nickel-Based Alloys
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
144
Radiation Effects in Nickel-Based Alloys
by Yang et al.4 The alloys were all irradiated in a ST condition; additionally, D25 was irradiated in an aged (24 h at 700 C) condition (STA), and D66 in a 30% cold-worked plus aged (11 h at 800 C plus 2 h at 700 C) condition (CWA). Specimens were irradiated at 450–735 C to a fast neutron fluence of 4 1026 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
Radiation Effects in Nickel-Based Alloys
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 (<2%) for irradiation temperatures up to 550 C and then increased sharply. In tests carried out at the irradiation temperatures, however, premature failures occurred above about 500 C, with tensile strengths reduced to 300 MPa and ductilities close to zero. All of the materials examined by Cauvin and Le Naour et al. showed similar properties, with no systematic influence of composition or heat treatment. Plitz et al.57 listed three fuel pin failures in Inconel 706 cladding in the Phenix reactor, with a further 12 failures in austenitic steel cladding; two of the Inconel 706 failures were associated with long periods of low power operation followed by a rise to full power conditions, resulting in mechanical interaction between the fuel and the low ductility cladding material. 4.04.5.2
Helium Implantation Experiments
Relatively few experiments have used helium implantation to investigate the embrittlement of nickel-based alloys, although this technique has been more widely used for austenitic steels. However, some data for high-Ni alloys have been published by Shiraishi et al.95 and Boothby.96 Shiraishi et al.95 compared the effects of helium injection and neutron irradiation on the tensile properties of developmental g0 -Ni3(Ti,Al) (alloy 7817) and g00 -Ni3Nb (alloy 7818) precipitation-hardened 40Ni–15Cr alloys. A number of alloy conditions, including ST, aged, and cold worked, were tested. Cyclotron injections of helium were made at 650 C to levels of 5 or 10 appm. Neutron irradiations were made at the same temperature to a fast fluence (E > 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
145
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 <10% and as low as 1–2% in the highest strength conditions). The g00 -hardened alloy 7818 showed lower tensile strength (typically 500–600 MPa) but maintained good ductility, with total elongation values always exceeding 10% and generally being above 20% in STA conditions. Boothby96 examined the effects of helium and/or lithium injection on the tensile properties of STA (ST 1050 C, aged 8 h at 700 C) Nimonic PE16. The effects of lithium were examined since this element is also produced from the 10B(n,a)7Li reaction in neutron-irradiated alloys. Helium and lithium were implanted at ambient temperature, either singly or in combination, to levels of 10 appm each. Samples were tested in the as-implanted condition or following an additional aging treatment of 72 h at 750 C. Tensile tests at a strain rate of 3 105 s1 on asimplanted samples showed no effect of helium or lithium on ductility at 200–550 C. However, at 650 C, the as-implanted samples were all embrittled to a similar extent, with the total elongations generally reduced to about half of the unimplanted levels regardless of whether He, Li, or (He þ Li) ions were injected. Postimplant aging of samples containing He or (He þ Li) resulted in further ductility loss in tests at 650 C, with significant embrittlement also evident at 550 C though not at 450 or 200 C. Postimplant aging of samples containing only lithium, however, resulted in some recovery in ductility compared to the as-implanted condition at 650 C but some ductility loss at 550 C. Thus, although it was clear that lithium had a detrimental effect on ductility, it did not appear to exacerbate the effects of helium. A mechanism for lithium embrittlement was not identified, though ductility loss in both Li and He implanted samples was associated with an increased propensity for intergranular fracture.
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Radiation Effects in Nickel-Based Alloys
Some additional, previously unpublished, data from a helium injection experiment conducted by Boothby and Cattle are given in Table 2. In this experiment, helium was implanted at ambient temperature to levels of 2, 10, and 50 appm into Nimonic PE16 that had been given a two-stage aging treatment (ST 1050 C, aged 4 h at 800 C plus 16 h 750 C). As before, helium-doped samples were either retained in the as-implanted condition or given an additional aging treatment to coarsen the dispersion of gas bubbles. Tensile tests were carried out at 650 C at strain rates of 3 105 and 3 106 s1. The results show a significant loss of ductility even at 2 appm helium. In tests carried out at the higher strain rate, the total elongation values decreased progressively with increasing helium content and were further reduced by postimplant aging. The ductility of as-implanted samples was generally lower but less sensitive to helium concentration in tests at a strain rate of 3 106 than at 3 105 s1. However, there was little effect of the strain rate on the ductility of the postimplant aged samples. Figure 15 illustrates grain boundary structures in a tensile-tested PE16 sample which had been aged subsequent to helium injection. Failure in this case appeared to occur by the growth and coalescence of cavities which were nucleated at grain boundary gas bubbles.97 The nucleation of unstable cavities at grain boundary helium bubbles requires the application of a critical tensile stress, which is an inverse function of the bubble radius, normal to the boundary. Cavity growth then occurs via the stress-induced absorption of vacancies. In the as-implanted condition, however, Table 2
Tensile properties of helium-implanted Nimonic PE16 at 650 C
Strain rate (s1) 5
3 10
3 106
the helium dispersion was too fine to enable grain boundary cavities to nucleate during tensile testing. Reduced ductility in the as-implanted samples appeared to be associated with grain boundary wedge cracking, where, as discussed by van der Schaaf and Marshall98 in relation to helium embrittlement of type 316 steel, the role of helium may be simply to decrease the effective surface energy for fracture. Although it is evident from simulation experiments that helium alone can largely account for irradiation embrittlement, it is more difficult to assess the significance of other radiation-induced effects such as matrix hardening and grain boundary segregation and/or precipitation. One experiment which examined the effect of the radiation-induced precipitation of the Ni3Si g0 phase on the ductility of a binary Ni-8 at.% Si alloy was described by Packan et al.99 In this experiment, thin foil tensile specimens were bombarded with either protons or a-particles to damage levels of 0.1– 0.3 dpa at 750 K; irradiation with a-particles resulted in the introduction of high helium concentrations of about 750 appm per 0.1 dpa. Proton and a-particle irradiations both resulted in the formation of g0 layers about 20–30 nm thick at the grain boundaries, but the material remained relatively ductile, exhibiting transgranular failures, in tensile tests carried out at a strain rate of 3 104 s1 at room temperature and, for the proton-irradiated case only, 720 K. Unfortunately, no tests were carried out at higher temperatures and samples which were irradiated with a-particles only at 750 K were not tested except at room temperature. However, low ductility intergranular failure was
Helium (appm)
0.2% PS (MPa)
UTS (MPa)
Uniform elongation (%)
Total elongation (%)
0 2 10 50 2a 10a 50a 0 2 10 50 2a 10a 50a
487 495 432 505 445 443 403 434 473 430 404 425 404 430
575 603 538 569 500 499 483 491 491 479 458 438 431 439
3.9 5.2 4.9 3.7 4.2 2.8 2.2 2.3 0.8 1.2 1.1 0.6 2.3 0.4
35.4 19.3 11.7 7.1 5.4 4.4 2.4 31.8 6.8 7.2 6.0 5.8 4.3 1.3
a Postimplant aged 72 h at 750 C. Unpublished data from Boothby, R. M.; Cattle, G. C. Development of g0 -hardened 25Ni-based Alloys for Fast Reactor Core Applications; FPSG/P(91)9, with permission from AEA Technology PLC.
Radiation Effects in Nickel-Based Alloys
100 nm
147
200 nm
Figure 15 Transmission electron micrographs illustrating (left) bubble dispersion on a grain boundary parallel to the tensile axis and (right) cavitation on a boundary approximately normal to the tensile axis in Nimonic PE16 (implanted with 10 appm helium at ambient temperature, and subsequently aged for 60 h at 750 C prior to tensile testing at 650 C). Reproduced from Boothby, R. M. J. Nucl. Mater. 1990, 171, 215–222.
induced in a test carried out at 720 K in a sample which was preimplanted with 1000 appm He at 970 K, then irradiated to 0.3 dpa, introducing an additional 2300 appm He at 750 K. Preimplantation of helium at 970 K produced grain boundary bubbles which were 10–20 nm in diameter, compared to 1.5–2 nm in material that was only irradiated with a-particles at 750 K. The results of this experiment therefore indicated that the radiation-induced precipitation of g0 at grain boundaries did not give rise to embrittlement unless helium was also implanted into the specimens.
4.04.6 Concluding Remarks The effects of irradiation on the microstructure and mechanical properties of nickel-based alloys are complex and, although the main factors affecting their behavior have been identified, a full understanding of radiation-induced effects remains elusive. This is particularly true of the precipitation-hardened alloys, typified by Nimonic PE16 and Inconel 706, where the role of the hardening phases – which confer high thermal creep strength, but are redistributed during irradiation and may possibly influence swelling behavior and contribute to intergranular embrittlement – is unclear. The radiation-induced effects considered in this chapter – void swelling, irradiation creep, the evolution of precipitate and dislocation structures, and irradiation embrittlement – are interrelated in several ways, but particularly through the effects of point defect fluxes and the consequent redistribution of solute atoms. The beneficial effect of nickel on the swelling resistance of austenitic alloys is well known, but a
clear explanation for the minimum in swelling that is found in alloys containing about 40–45% Ni has not been forthcoming. There is general agreement that the major influence of alloy composition on swelling arises through its effects on the effective vacancy diffusivity and on segregation via the inverse Kirkendall effect. However, on what appears to be the mistaken assumption that the swelling resistance of nickel-based alloys derives from an extended void nucleation period, swelling models have largely focused on factors affecting the nucleation rather than the growth of voids. Data for neutron-irradiated Nimonic PE16, for example, indicate that its swelling resistance is due to a combination of a comparatively low saturation void concentration, which is reached at a relatively low displacement dose, and a low void growth rate. The minimum critical void radius concept appears to offer the most plausible explanation for the minimum in swelling found at intermediate nickel contents, although experimental data comparing the behavior of PE16 and a nonprecipitation hardenable alloy with a similar matrix composition indicate that, in addition to the Ni content of the alloy, the presence of Si and/or the g0 forming solutes Al plus Ti may also be important. The dependence of the void growth rate on Ni may be related to the effects of radiation-induced segregation on the bias terms for the absorption of point defects at sinks, though again there is evidence that minor solutes, including Si, B, and Mo, as well as the g0 -forming elements, have a beneficial effect on the overall swelling behavior of nickel-based alloys. The irradiation creep behavior of nickel-based alloys generally appears to be similar to that of austenitic steels, though the higher thermal creep
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strength of precipitation-hardened alloys is an advantage at higher operating temperatures. The main drawback of precipitation-hardened nickelbased alloys for reactor core applications is perceived to be a high susceptibility to irradiation embrittlement. Although it has been suggested that a combination of matrix hardening and of grain boundary weakening due to the formation of brittle intergranular layers of g0 (e.g., in the case of Nimonic PE16) or Z phase (in the case of Inconel 706) is responsible for the irradiation embrittlement of these alloys, there is strong evidence, at least for g0 hardened materials, that helium is the primary cause of low ductility failures. Experimental data have shown that the implantation or generation of relatively small amounts of helium can give rise to low tensile ductility, with intergranular failures initiated by either the growth and linkage of cavities or by wedge cracking depending on test conditions and helium distribution, under conditions where grain boundary g0 layers are not formed. However, irrespective of the details of the embrittlement mechanism, it is evident that this aspect of radiation damage does not preclude the in-core application of nickel-based alloys, as has clearly been demonstrated by the successful use of PE16 for fuel element cladding irradiated to high burn-ups in PFR. The long-term integrity of PE16 cladding is attributable to the dimensional stability of the alloy, arising from a combination of good swelling resistance and high creep strength, and relatively low operating stresses which allay irradiation embrittlement concerns.
4. 5. 6.
7.
8.
9.
10.
11. 12. 13.
14.
15.
Acknowledgments The inclusion of previously unpublished information from the UKAEA fast reactor program (in particular, data obtained by RM Sharpe and by RM Boothby and GC Cattle) is with the kind permission of AEA Technology PLC.
16. 17.
18.
References 1. 2. 3.
Rowcliffe, A. F.; Mansur, L. K.; Hoelzer, D. T.; Nanstad, R. K. J. Nucl. Mater. 2009, 392, 341–352. Angeliu, T. M.; Ward, J. T.; Witter, J. K. J. Nucl. Mater. 2007, 366, 223–237. Benes, O.; Cabet, C.; Delpech, S.; et al. Review report on liquid salts for various applications. ALISIA (Assessment of Liquid Salts for Innovative Applications) Deliverable D-50 (version V4); 2009, Cordis website: http://cordis.europa. eu/home_en.html.
19. 20. 21. 22.
Yang, W. J. S.; Gelles, D.; Straalsund, J. L.; Bajaj, R. J. Nucl. Mater. 1985, 132, 249–265. Toloczko, M. B.; Garner, F. A.; Standring, J.; Munro, B.; Adaway, S. J. Nucl. Mater. 1998, 258–263, 1606–1612. Bramman, J. I.; Etherington, E. W.; Nelson, R. S.; Norgett, M. J. Radiation damage units for fast reactor steels. Proceedings of the British Nuclear Energy Society Conference: Irradiation Embrittlement and Creep, London, UK, Nov 9–10 1972; pp 27–31. Bramman, J. I.; Cawthorne, C.; Fulton, E. J.; Sinclair, W. D. J. In Proceedings of the British Nuclear Energy Society Conference: Voids Formed by Irradiation of Reactor Materials, Reading, UK, Mar 24–25, 1971; Pugh, S. F., Loretto, M. H., Norris, D. I. R., Eds.; 1971; pp 27–33. Cawthorne, C.; Fulton, E. J.; Bramman, J. I.; Linekar, G. A. B.; Sharpe, R. M. In Proceedings of the British Nuclear Energy Society Conference: Voids Formed by Irradiation of Reactor Materials, Reading, UK, Mar 24–25, 1971; Pugh, S. F., Loretto, M. H., Norris, D. I. R., Eds.; 1971; pp 35–43. Hudson, J. A.; Mazey, D. J.; Nelson, R. S. In Proceedings of the British Nuclear Energy Society Conference: Voids Formed by Irradiation of Reactor Materials, Reading, UK, Mar 24–25, 1971; Pugh, S. F., Loretto, M. H., Norris, D. I. R., Eds.; 1971; pp 213–221. Bullough, R.; Perrin, R. C. In Proceedings of the British Nuclear Energy Society Conference: Voids Formed by Irradiation of Reactor Materials, Reading, UK, Mar 24–25, 1971; Pugh, S. F., Loretto, M. H., Norris, D. I. R., Eds.; 1971; pp 79–107. Williams, K. R.; Fisher, S. B. Radiat. Eff. Defects Solids 1972, 15, 243–250. Johnston, W. G.; Rosolowski, J. H.; Turkalo, A. M.; Lauritzen, T. J. Nucl. Mater. 1974, 54, 24–40. Bates, J. F.; Johnston, W. G. In Radiation Effects in Breeder Reactor Structural Materials, Scottsdale, AZ, June 19–23, 1977; Bleiberg, M. L., Bennett, J. W., Eds.; Metallurgical Society of AIME: New York, 1977; pp 625–644. Makin, M. J.; Hudson, J. A.; Mazey, D. J.; Nelson, R. S.; Walters, G. P.; Williams, T. M. In Radiation Effects in Breeder Reactor Structural Materials, Scottsdale, AZ, June 19–23, 1977; Bleiberg, M. L., Bennett, J. W., Eds.; Metallurgical Society of AIME: New York, 1977; pp 645–665. Bajaj, R.; Diamond, S.; Chickering, R. W.; Bleiberg, M. L. 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 541–561, ASTM STP 725. Harries, D. R. In The Physics of Irradiation Produced Voids, Harwell, UK, Sept 9–11, 1974; Nelson, R. S., Ed.; 1974; pp 287–298, UKAEA report AERE-R7934. Watkin, J. S. Irradiation Effects on the Microstructure and Properties of Metals; American Society for Testing and Materials: Philadelphia, PA, 1976; pp 270–283, ASTM STP 611. Sims, C. T. In Superalloys II; Sims, C. T., Stoloff, N. S., Hagel, W. C., Eds.; John Wiley & Sons: New York, 1987; pp 217–240. Bates, J. F.; Powell, R. W. J. Nucl. Mater. 1981, 102, 200–213. Powell, R. W.; Peterson, D. T.; Zimmerschied, M. K.; Bates, J. F. J. Nucl. Mater. 1981, 103–104, 969–974. Gelles, D. S. J. Nucl. Mater. 1984, 122–123, 207–213. Garner, F. A.; Gelles, D. S. In Effects of Radiation on Materials: 14th International Symposium; Packan, N. H, Stoller, R. E., Kumar, A. S., Eds.; American Society for Testing and Materials: Philadelphia, PA, 1990; Vol. II, pp 673–683, ASTM STP 1046.
Radiation Effects in Nickel-Based Alloys 23. Garner, F. A.; Brager, H. R. In Effects of Radiation on Materials: 12th International Symposium; Garner, F. A., Perrin, J. S., Eds.; American Society for Testing and Materials: Philadelphia, PA, 1985; Vol. 1, pp 187–201, ASTM STP 870. 24. Mazey, D. J.; Hanks, W. In Materials for Nuclear Reactor Core Applications, Bristol, UK, Oct 27–29, British Nuclear Energy Society: London, 1987; Vol. 1, pp 129–135. 25. Mazey, D. J.; Hanks, W.; Titchmarsh, J. M. J. Nucl. Mater. 1988, 160, 153–165. 26. Lee, E. H.; Mansur, L. K. Metall. Trans. 1992, 23A, 1977–1986. 27. Pechenkin, V. A.; Epov, G. A. J. Nucl. Mater. 1996, 233–237, 1009–1015. 28. Boothby, R. M. J. Nucl. Mater. 1996, 230, 148–157. 29. Marwick, A. D. J. Phys. F: Metal Phys. 1978, 8, 1849–1861. 30. Marwick, A. D.; Kennedy, W. A. D.; Mazey, D. J.; Hudson, J. A. Scripta Metall. 1978, 12, 1015–1020. 31. Wiedersich, H.; Lam, N. Q. In Phase Transformations During Irradiation; Nolfi, F. V., Ed.; Applied Science Publishers: London, 1983; pp 247–290. 32. Rehn, L. E.; Okamoto, P. R. In Phase Transformations During Irradiation; Nolfi, F. V., Ed.; Applied Science Publishers: London, 1983; pp 1–46. 33. Venker, H.; Ehrlich, K. J. Nucl. Mater. 1976, 60, 347–349. 34. Garner, F. A.; Wolfer, W. G. J. Nucl. Mater. 1981, 102, 143–150. 35. Esmailzadeh, B.; Kumar, A. S. In Effects of Radiation on Materials: 12th International Symposium; Garner, F. A., Perrin, J. S., Eds.; American Society for Testing and Materials: Philadelphia, PA, 1985; Vol. 1, pp 468–480, ASTM STP 870. 36. Rothman, S. J.; Nowicki, L. J.; Murch, G. E. J. Phys. F: Metal Phys. 1980, 10, 383–398. 37. Wolfer, W. G. J. Nucl. Mater. 1982, 114, 292–304. 38. Wolfer, W. G.; Garner, F. A.; Thomas, L. E. In Effects of Radiation on Materials: 11th International Symposium; Brager, H. R., Perrin, J. S., Eds.; American Society for Testing and Materials: Philadelphia, PA, 1982; pp 1023–1041, ASTM STP 782. 39. Garner, F. A.; Wolfer, W. G. J. Nucl. Mater. 1984, 122–123, 201–206. 40. Garner, F. A. J. Nucl. Mater. 1984, 122–123, 459–471. 41. Dodd, R. A.; Garner, F. A.; Kai, J. J.; Lauritzen, T.; Johnston, W. G. In Radiation-Induced Changes in Microstructure: 13th International Symposium (Part 1); Garner, F. A., Packan, N. H., Kumar, A. S., Eds.; American Society for Testing and Materials: Philadelphia, PA, 1987; pp 788–804, ASTM STP 955. 42. Coghlan, W. A.; Garner, F. A. In Radiation-Induced Changes in Microstructure: 13th International Symposium (Part 1); Garner, F. A., Packan, N. H, Kumar, A. S., Eds.; American Society for Testing and Materials: Philadelphia, PA, 1987; pp 315–329, ASTM STP 955. 43. Hoyt, J. J.; Garner, F. A. J. Nucl. Mater. 1991, 179–181, 1096–1099. 44. Mazey, D. J.; Nelson, R. S. J. Nucl. Mater. 1979, 85–86, 671–675. 45. Muroga, T.; Garner, F. A.; Ohnuki, S. J. Nucl. Mater. 1991, 179–181, 546–549. 46. Muroga, T.; Garner, F. A.; McCarthy, J. M.; Yoshida, N. In Effects of Radiation on Materials: 15th International Symposium; Stoller, R. E., Kumar, A. S., Gelles, D. S., Eds.; American Society for Testing and Materials: Philadelphia, PA, 1992; pp 1015–1033, ASTM STP 1125. 47. Garner, F. A.; Kumar, A. S. In Radiation-Induced Changes in Microstructure: 13th International Symposium (Part 1); Garner, F. A., Packan, N. H., Kumar, A. S., Eds.; American
48.
49.
50.
51. 52.
53.
54. 55. 56. 57. 58.
59. 60. 61.
62.
63.
64. 65. 66. 67.
149
Society for Testing and Materials: Philadelphia, PA, 1987; pp 289–314, ASTM STP 955. Allen, T. R.; Busby, J. T.; Gan, J.; Kenik, E. A.; Was, G. S. In Effects of Radiation on Materials: 19th International Symposium; Hamilton, M. L., Kumar, A. S., Rosinski, S. T., Grossbeck, M. L., Eds.; American Society for Testing and Materials: West Conshohocken, PA, 2000; pp 739–755, ASTM STP 1366. Brown, C.; Sharpe, R. M.; Fulton, E. J.; Cawthorne, C. In Dimensional Stability and Mechanical Behaviour of Irradiated Metals and Alloys, Brighton, UK, Apr 11–13, 1983; British Nuclear Energy Society: London, 1983; Vol. 1, pp 63–67. Sklad, P. S.; Clausing, R. E.; Bloom, R. E. Irradiation Effects on the Microstructure and Properties of Metals; American Society for Testing and Materials: Philadelphia, PA, 1976; pp 139–155, ASTM STP 611. Dumbill, S.; Boothby, R. M.; Williams, T. M. Mater. Sci. Technol. 1991, 7, 385–390. Brown, C.; Linekar, G. A. B. Materials for Nuclear Reactor Core Applications, Bristol, UK, Oct 27–29, 1987; British Nuclear Energy Society: London, 1987; Vol. 1, pp 219–222. Batey, W.; Linekar, G. A. B. Proceedings of International Conference on Materials for Nuclear Reactor Core Applications, Bristol, UK, Oct 27–29, 1987; British Nuclear Energy Society: London, 1987; Vol. 1, pp 319–322. Cole, G. C. Fast Reactor Core and Fuel Structural Behaviour, Inverness, UK, June 4–6, 1990; British Nuclear Energy Society: London, 1990; pp 25–31. IAEA Status of Liquid Metal Cooled Fast Reactor Technology; International Atomic Energy Agency: Vienna, 1999; pp 273–316, IAEA Tec Doc 1083. Allbeson, K. F.; Brown, C.; Gillespie, J. Nucl. Eng. 1990, 31, 87–89. Plitz, H.; Crittenden, G. C.; Languille, A. J. Nucl. Mater. 1993, 204, 238–243. Naganuma, M.; Koyama, S.; Asaga, T.; et al. In Proceedings of International Symposium on MOX Fuel Cycle Technologies for Medium and Long Term Deployment, Vienna, May 17–21 1999; pp 311–321, IAEA-CSP-3/P (2000). Harries, D. R. J. Nucl. Mater. 1977, 65, 157–173. Ehrlich, K. J. Nucl. Mater. 1981, 100, 149–166. Garner, F. A. In Nuclear Materials; Frost, B. R. T., Ed.; Materials Science and Technology: A Comprehensive Treatment; Cahn, R. W., Haasen, P., Kramer, E. J., Eds.; VCH: New York, 1994; Vol. 10A, pp 419–543. Mosedale, D.; Harries, D. R.; Hudson, J. A.; Lewthwaite, G. W.; McElroy, R. J. In Radiation Effects in Breeder Reactor Structural Materials, Scottsdale, AZ, June 19–23, 1977; Bleiberg, M. L., Bennett, J. W., Eds.; Metallurgical Society of AIME: New York, 1977; pp 209–228. 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. Garner, F. A.; Toloczko, M. B. J. Nucl. Mater. 1997, 251, 252–261. Garner, F. A.; Toloczko, M. B.; Grossbeck, M. L. J. Nucl. Mater. 1998, 258–263, 1718–1724. Paxton, M. M.; Chin, B. A.; Gilbert, E. R.; Nygren, R. E. J. Nucl. Mater. 1979, 80, 144–151.
150
Radiation Effects in Nickel-Based Alloys
68. Gilbert, E. R.; Chin, B. A. In Effects of Radiation on Materials: 12th International Symposium; Garner, F. A., Perrin, J. S., Eds.; American Society for Testing and Materials: Philadelphia, PA, 1985; Vol. 1, pp 52–60, ASTM STP 870. 69. Thomas, L. E. Phase Stability During Irradiation; American Institute of Mining, Metallurgical and Petroleum Engineers: Warrendale, PA, 1981; pp 237–255. 70. Gelles, D. S. J. Nucl. Mater. 1993, 205, 146–161. 71. Porollo, S. I.; Dvoriashin, A. M.; Vorobjev, A. N.; et al. J. Nucl. Mater. 2000, 283–287, 239–243. 72. Gelles, D. S.; Garner, F. A.; Brager, H. R. In Effects of Radiation on Materials: 10th International Symposium; Kramer, D., Brager, H. R., Perrin, J. S., Eds.; American Society for Testing and Materials: Philadelphia, PA, 1981; pp 735–753, ASTM STP 725. 73. Gelles, D. S. In Effects of Radiation on Materials: 12th International Symposium; Garner, F. A., Perrin, J. S., Eds.; American Society for Testing and Materials: Philadelphia, PA, 1985; Vol. 1, pp 98–112, ASTM STP 870. 74. Garner, F. A.; Gelles, D. S. J. Nucl. Mater. 1988, 159, 286–309. 75. Garner, F. A.; Wolfer, W. G.; Brager, H. R. In Effects of Radiation on Structural Materials; Sprague, J. A., Kramer, D., Eds.; American Society for Testing and Materials: Philadelphia, PA, 1979; pp 160–183, ASTM STP 683. 76. Nelson, R. S.; Hudson, J. A.; Mazey, D. J. J. Nucl. Mater. 1972, 44, 318–330. 77. Baron, M.; Chang, A.; Bleiberg, M. L. In Radiation Effects in Breeder Reactor Structural Materials, Scottsdale, AZ, June 19–23, 1977; Bleiberg, M. L., Bennett, J. W., Eds.; Metallurgical Society of AIME: New York, 1977; pp 395–404. 78. Chang, A. L.; Baron, M. J. Nucl. Mater. 1979, 83, 214–222. 79. Gelles, D. S. In Effects of Radiation on Structural Materials; Sprague, J. A., Kramer, D., Eds.; American Society for Testing and Materials: Philadelphia, PA, 1979; pp 194–206, ASTM STP 683. 80. Gelles, D. S. In Effects of Radiation on Materials: 10th International Symposium; Kramer, D., Brager, H. R., Perrin, J. S., Eds.; American Society for Testing and Materials: Philadelphia, PA, 1981; pp 562–582, ASTM STP 725. 81. Yang, W. J. S. J. Nucl. Mater. 1982, 108–109, 339–346. 82. Broomfield, G. H.; Harries, D. R.; Roberts, A. C. J. Iron Steel Inst. 1965, 203, 502–506. 83. Broomfield, G. H. Irradiation Effects in Structural Alloys for Thermal and Fast Reactors; American Society for Testing and Materials: Philadelphia, PA, 1969; pp 38–66, ASTM STP 457. 84. Hamm, C. D.; Miller, D. A. In Creep and Fracture of Engineering Materials and Structures; Wilshire, B., Owen, D. R. J., Eds.; Swansea, UK, Apr 1–6, 1984; Pineridge Press: Swansea, 1984; Vol. 1, pp 359–370.
85.
86. 87.
88.
89.
90.
91.
92.
93.
94.
95. 96. 97. 98.
99.
McElroy, W. N.; Farrar, H. In Radiation-Induced Voids in Metals, Albany, NY, June 9–11, 1971; Corbett, J. W., Ianniello, L. C., Eds.; USAEC Technical Information Center, Oat Ridge, TN, 1971; pp 187–229. Rowcliffe, A. F.; Horak, J. A. Am. Nucl. Soc. Trans. 1981, 38, 266–267. 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. Fisher, S. B.; Callen, V. M.; Rose, P. K. In Effects of Radiation on Materials: 15th International Symposium; Stoller, R. E., Kumar, A. S., Gelles, D. S, Eds.; American Society for Testing and Materials: Philadelphia, PA, 1992; pp 667–688, ASTM STP 1125. Boothby, R. M.; Harries, D. R. In Proceedings of International Conference on Mechanical Behaviour and Nuclear Applications of Stainless Steel at Elevated Temperatures, Varese, Italy, May 20–22, 1981; The Metals Society: London, 1982; pp 157–164. Vaidyanathan, S.; Lauritzen, T.; Bell, W. L. In Effects of Radiation on Materials: 11th International Symposium; Brager, H. R., Perrin, J. S., Eds.; American Society for Testing and Materials: Philadelphia, PA, 1982; pp 619–635, ASTM STP 782. Huang, F. H.; Fish, R. L. In Effects of Radiation on Materials: 12th International Symposium; Garner, F. A., Perrin, J. S., Eds.; American Society for Testing and Materials: Philadelphia, PA, 1985; Vol. 2, pp 720–731, ASTM STP 870. Yang, W. J. S.; Makenas, B. J. In Effects of Radiation on Materials: 12th International Symposium; Garner, F. A., Perrin, J. S., Eds.; American Society for Testing and Materials: Philadelphia, PA, 1985; Vol. 1, pp 127–138, ASTM STP 870. Cauvin, R.; Schauff, R.; Rabouille, O. Materials for Nuclear Reactor Core Applications, Bristol, UK, Oct 27–29, 1987; British Nuclear Energy Society: London, 1987; Vol. 1, pp 187–190. Le Naour, F.; Hugon, M. P.; Grosjean, P.; Maillard, A.; Seran, J. L. In Materials for Nuclear Reactor Core Applications, Bristol, UK, Oct 27–29, 1987; British Nuclear Energy Society: London, 1987; Vol. 1, pp 211–217. Shiraishi, H.; Yamamoto, N.; Hasegawa, A. J. Nucl. Mater. 1989, 169, 198–205. Boothby, R. M. J. Nucl. Mater. 1992, 186, 209–211. Boothby, R. M. J. Nucl. Mater. 1990, 171, 215–222. Van der Schaaf, B.; Marshall, P. In Dimensional Stability and Mechanical Behaviour of Irradiated Metals and Alloys, Brighton, UK, Apr 11–13, 1983; British Nuclear Energy Society: London, 1983; Vol. 1, pp 143–148. Packan, N. H.; Schroeder, H.; Kesternich, W. J. Nucl. Mater. 1986, 141–143, 553–558.
4.05
Radiation Damage of Reactor Pressure Vessel Steels
C. English and J. Hyde National Nuclear Laboratory, Harwell, Oxfordshire, UK
ß 2012 Elsevier Ltd. All rights reserved.
4.05.1
Introduction
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4.05.2 4.05.3 4.05.3.1 4.05.4
Pressure Vessel Steels Effect of Neutron Irradiation on Bulk Properties Summary Development of Mechanistic Insight of Factors Controlling Current Plant Lifetimes Introduction Radiation Damage Mechanisms Mechanistic Framework Cluster Development Under Irradiation: Cu Introduction CEC characterization Development with increasing fluence Development of Matrix Defects Introduction Nature Development with flux and fluence and irradiation temperature Effect of alloying MD and hardening Effect of Radiation Damage on Hardening Segregation to Grain Boundaries Summary Development of Mechanistically Based DDRs Introduction DDRs for CMn Steels US Mechanistically Guided DDRs Japanese Embrittlement Correlations Summary Current Issues in the Development of DDRs
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4.05.4.1 4.05.4.2 4.05.4.3 4.05.4.4 4.05.4.4.1 4.05.4.4.2 4.05.4.4.3 4.05.4.5 4.05.4.5.1 4.05.4.5.2 4.05.4.5.3 4.05.4.5.4 4.05.4.5.5 4.05.4.6 4.05.4.7 4.05.4.8 4.05.5 4.05.5.1 4.05.5.2 4.05.5.3 4.05.5.4 4.05.5.5 4.05.6 References
Abbreviations 3DAP AES ASME ASTM BWR CDB CEC CRP Cumatrix
Three-dimensional atom probe Auger electron spectroscopy American Society of Mechanical Engineers American Society for Testing and Materials Boiling water reactor Coherent Doppler broadening Cu-enriched cluster Cu-rich cluster Level of Cu in matrix
DBTT DDR DIDO dpa DS DHv EFTEM EOL EONY
160 160 160 161 162 162 162 163 166 166 166 167 168 168 169 169 170 170 170 171 172 174 175 176 177
Ductile-to-brittle transition temperature Dose–damage relationships DIDO was a MTR in the UK Displacements per atom Change in PALA ‘S’ parameter Change in Vickers hardness Energy filtered transmission electron microscopy End-of-life Eason, Odette, Nanstad, Yamamoto
151
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EPRI EWO FEGSTEM/ EDX FIM FIS HERALD HSSI IAEA IS IVAR
JEAC KMC KTA LEAP LWR MD MMA MS MTR ORNL OSIRIS PA PALA PAS-OEMS
PA-t PISA PLUTO PWHT PWR f ft RPV SANS SAW SIA SMA SMD
Electric Power Research Institute (EPRI) Eason Wright and Odette Field emission gun scanning transmission electron microscopy/ energy dispersive X-ray Fragilisation par Irradiation Moyenne Fragilisation par Irradiation Superieure HERALD was a MTR in the UK Heavy Section Steel Irradiation International Atomic Energy Authority Interstitial solute Irradiation Variable Program database on irradiation-induced changes in reactor Pressure Vessel Steels Japan Electric Association Code Kinetic Monte Carlo (Models) German Nuclear Standards Commission Local electrode atom probe Light water reactor Matrix damage Manual metal-arc welds Molecular statics Materials Test Reactor Oak Ridge National Laboratory OSIRIS was a MTR in France Positron annihilation Positron annihilation lineshape analysis Positron annihilation spectroscopyorbital angular momentum distribution Positron Annihilation Lifetime Phosphorus in Steel Ageing 5th Framework Programme of the EEC PLUTO was a MTR in the UK Postweld heat treatment Pressurized water reactor Flux or dose rate Fluence or dose Reactor pressure vessel Small-angle neutron scattering Submerged-arc weldment Self-interstitial atom Submerged-arc welds Stable matrix defect
SOL SRM TAGSI TEM UCRR UMD USE USNRC VVER
Start-of-life Standard Reference Monitor Technical Advisory Group on Structural Integrity Transmission electron microscope Union Carbide Research Reactor Unstable matrix defect Upper shelf energy US Nuclear Regulatory Commission From the Russian. Translates as Water-Water Energetic Reactor
4.05.1 Introduction The ferritic steel reactor pressure vessel (RPV) of a light water reactor (LWR) is unique in terms of nuclear plant safety. This is because the RPV is a pressure boundary component whose catastrophic failure by brittle fracture could lead to severe core damage and, potentially, to the widespread release of radioactivity. In addition to the thermal, mechanical, and chemical degradation processes common to all primary circuit components, the ferritic steel RPV, because of its close proximity to the reactor core, also undergoes changes in mechanical properties because of radiation damage from the flux of fast neutrons arising from nuclear reactions in the fuel. The integrity of the RPV has to be assured throughout the life of the plant. Irradiation results in hardening and embrittlement processes, the most important effect of which is the rise in the ductile–brittle transition temperature (DBTT) and the decrease in the fracture toughness of the RPV. This can be an important issue in the region of the beltline which experiences the highest neutron fluence. These changes in mechanical properties make the RPV more susceptible to brittle fracture as the plant ages. The primary purpose of this chapter is to demonstrate our current understanding of such radiation damage effects in ferritic pressure vessel steels. Low alloy ferritic steel pressure vessels are employed in all western LWRs (both boiling water reactors (BWRs) and pressurized water reactors (PWRs)), and also in VVER (LWR) reactors in Russia and a number of European countries (see, e.g., Steele and Sterne1). In the past, ferritic pressure vessel steels were also employed in gas-cooled Magnox reactors in the United Kingdom2 (Magnox reactors with steel vessels are now undergoing decommissioning3).
Radiation Damage of Reactor Pressure Vessel Steels
Demonstrating the safe operation of such a plant has led to extensive international research over the last 40–50 years on the aging effects in ferritic steels. The need for such scientific understanding has been raised at the highest level. For example, Sir Alan Cottrell, then UK Government Chief Scientist, in a memorandum concerning the integrity of LWR pressure vessels, dated 22 January 1974, to the UK Parliament Select Committee on Science and Industry stated ‘‘The possible gradual growth of small cracks in highly stressed regions, by ageing and corrosion effects during service needs further scientific investigation.’’4 The discussion here establishes that in the field of radiation damage, it is the direct application of fundamental research to operating reactors that is significant. This chapter demonstrates that developments in the understanding of the damage mechanisms have enabled an improved description of the in-service properties of RPVs of operating reactors. Independent peer review has been central to the process and particularly important from a regulatory perspective. Frequently, it is not simply the research community directly involved that has to assess any improved description of the degradation processes; for example, the safety authorities within the utility, operating the reactor (or fleet), or the national regulator will become involved in the process. Most importantly it has not always been possible to predict the end of life (EOL) vessel properties from the data obtained from materials irradiated as part of vessel surveillance programs. The vessel surveillance programs for commercial nuclear reactors are intended to monitor the irradiation-induced changes in mechanical properties of life-limiting structural materials subjected to significant neutron fluence. Thus, they are designed to provide advance information concerning the state of degradation in the mechanical properties of key structural components. However, because of the inevitable differences in neutron dose rate between the vessel wall and such surveillance samples (typically a factor of 5 or more), the scarcity of such data at the start of plant operation, and complex and unexpected embrittlement dependencies on steel composition, it has become necessary to develop dose–damage relationships (DDRs) on the basis of mechanistic understanding that predict the embrittlement dependence on material and irradiation variables. The vast majority of investigations on the aging effects in ferritic steels over the last 40–50 years have been concerned with the effects of neutron irradiation
153
over quite a narrow set of irradiation conditions, for example, irradiation temperatures of 250–300 C (although there was interest in irradiation temperatures as low as 160 C) and irradiation doses typical of in-service exposures (<0.1 dpa). Early studies focused, primarily, on the effects of irradiation on mechanical properties, while in the last 20 years such studies have been combined with rigorous investigations of the effects of materials and irradiation variables on the microstructure developed under irradiation. It is this combination that has allowed significant insight into the mechanisms that determine the effect of radiation damage on the bulk properties of ferritic RPV steels. As will be shown, such studies have encompassed simple model alloys, steels with a controlled composition, and commercial steels employed in real vessels. A major outcome in recent years is the development of mechanistically based or guided DDRs that are employed, frequently in a regulatory framework, to predict the behavior of reactor vessels. (Some workers refer to DDRs as ‘embrittlement correlations’.) In formulating such DDRs, it has been necessary to encompass not only the irradiation variables such as flux, fluence, and irradiation temperature but also material factors, such as composition, heat treatment, and product form. In reviewing such a long-standing field, it is necessary to provide a focus to enable presentation of a manageable set of data. As stated earlier, the chapter focuses on the mechanisms that control RPV embrittlement, how such understanding has been incorporated into mechanistically based DDRs, and the limitations or current research issues associated with their development. These mechanistically based DDRs have been developed primarily for LWRs and Magnox reactors located in the West and this chapter gives prominence to such studies; that is, the chapter does not provide an extensive review of studies on the embrittlement of the RPVs of VVER reactors (see Nikolaev et al.,5 Kryukov et al.,6 Shtrombakh and Nikolaev,7 and Brumovsky8 for such a review). Its organization reflects the overall focus; first, the chapter provides a description of RPV steels (Section 4.05.2) and then provides an overview of the effects of radiation damage on their mechanical properties (Section 4.05.3). Section 4.05.4 describes the status of our mechanistic understanding of embrittlement caused by radiation damage. This understanding has proved pivotal to the development of the DDRs described in Section 4.05.5. Outstanding issues, particularly those arising through plant life extension, are outlined in Section 4.05.6.
154
Radiation Damage of Reactor Pressure Vessel Steels
4.05.2 Pressure Vessel Steels In reviewing radiation damage effects in ferritic steels, it is important to recognize that a range of ferritic steels have been employed in commercial reactors. Such steels were necessarily employed in thick sections and the fabrication of the vessels involved welding of preformed plates or forgings (for a description of typical vessels, see Steele and Sterne1). Frequently, in early designs, the welds were located opposite the center of the reactor core and received the highest neutron dose. Commercially available ferritic steels were employed in the construction of the first reactors (both for Magnox and LWR designs). For LWRs, which have to contain higher pressure than the pressure vessels of Magnox reactors, the desirability of using steels of high toughness, adequate strength, and weldability in thick sections, combined with good service experience, has narrowed the choice to a number of low alloy steels; for example, those containing manganese, nickel, and Table 1
molybdenum. Indeed, the ASME specification for quenched and tempered vacuum-treated carbon and alloy steel forgings for pressure vessels permits only SA-508 Class 2 and Class 3 compositions. Early vessels were constructed from A302 and A302B plates. Vessels constructed in Russia (so-called VVER reactors) are fabricated from high-strength CrMoV steels. The steels used in vessel construction of Magnox pressure vessels generally consisted of simple C–Mn ferritic plates, either Si-killed or Al-grain refined, in the normalized and stress-relieved condition. Table 12 shows the average, or range of, chemical compositions for the plates, manual metal-arc (MMA) welds, submerged-arc (SMA) welds, and forgings, respectively, as derived from the populations of vessel cut-outs ex-construction, surveillance samples, and contemporary reproduction vessel materials. Compositions of commercial MnMoNi steels used for modern LWR pressure vessels are given in Table 2.9–12 A detailed description of the development of modern pressure vessel steels can be found in
Chemical composition of Magnox vessel materials (wt%)
Material
C
Mn
Si
S
P
Cu
Plate MMA SMA Forgings
0.09–0.17 0.086 0.088 0.18
1.04–1.32 0.91 1.49 1.30
0.10–0.60 0.42 0.52 0.36
0.02–0.04 0.022 0.037 0.024
0.01–0.04 0.025 0.031 0.024
0.03–0.15 0.08 0.23 0.10
Ni and Cr each <0.1 wt%, Mo, Sn, As, and Ti each <0.05%. Source: Jones, R. B.; Bolton, C. J. Neutron Radiation Embrittlement Studies in Support of Continued Operation and Validation by Sampling of Magnox Reactor Steel Pressure Vessels and Components, Twenty Fourth Water Reactor Safety Information Meeting, Bethesda, 1996; Vol. 2, pp 25–48.
Table 2
Typical compositions of modern LWR pressure vessels
Material specification
Country of origin
Number of analyses
C
Mn
Mi
Mo
Si
Cr
P
S
Cu
Al
S533B Cl 1 S533B Cl 1
Japana United States Japana
175 13
0.201 0.218
1.371 1.367
0.616 0.547
0.520 0.547
0.243 0.236
0.136 0.074
0.007 0.009
0.006 0.014
0.049 0.117
0.025
166
0.200
1.398
0.753
0.505
0.243
0.097
0.007
0.007
0.054
0.028
Franceb
125
0.161
1.338
0.722
0.503
0.235
0.295
0.010
0.010
0.065
Japana United States
64 5
0.206 0.238
0.803 0.682
0.844 0.600
0.585 0.595
0.231 0.284
0.374 0.374
0.006 0.006
0.006 0.011
0.048 0.040
SA508 Cl 3 20MnMoNi 55 forging grade SA508 Cl 3 (etc.) SA508 Cl2 SA508 Cl 2 a
Additional analyses of Japanese steels shows that As 0.007–0.009, Sn 0.008, Sb 0.002, Co 0.009–0.010 wt%. Analyses of French SA508 Cl 3 showed As 0.016, Sn 0.011, Co 0.017 wt%.
b
0.029
Radiation Damage of Reactor Pressure Vessel Steels
Table 3
155
Chemical composition of surveillance materials for Sizewell ‘B’ PWR
Wt%
C
Si
Mn
P
S
Ni
Cr
Mo
Al
V
Cu
A508 Class III forging Submerged-arc weld
0.16 0.07
0.27 0.38
1.42 1.53
0.005 0.008
0.004 0.009
0.74 0.62
0.13 0.02
0.51 0.58
0.022 0.004
<0.01 <0.01
0.05 0.02
Source: Priest, R. H.; Belcher, W. P. A.; Mendes, C. M.; Neale, B. K. Int. J. Press. Vessels Piping 2000, 77, 621–628.
4.05.3 Effect of Neutron Irradiation on Bulk Properties Neutron irradiation of RPV steels results in an increase in the yield stress (sY) and an upward shift in the DBTT. The increase in the DBTT is generally determined from the shift in the Charpy impact curve at a specific energy level, typically 41 J (DT41 J). Historically, most data on effects of irradiation on the DBTT arose from irradiation and testing of Charpy specimens, although frequently increases in yield stress were also reported. The specimen geometries in both cases were convenient for irradiations in the restricted space available for either surveillance capsules1 or rigs inserted into the cores of material test reactors (MTRs).15 More recently, with the advent of the Master Curve technique,16 there has been greater interest in acquiring
300 250
Charpy energy (J)
Druce and Edwards.13 It is to be noted that commercial steels are complex with many elements present. Although the concentrations of additional elements are low, all have the potential to influence the radiation damage response of the steel. The deleterious effects of residual impurities such as Cu and P on the in-service degradation in properties (see Section 4.05.3) have been increasingly recognized. Consequently, over time, the levels of Cu and P in ferritic pressure vessel steels have been reduced and in modern steels, such as the forgings and welds employed in fabricating the vessel for the PWR at Sizewell in the United Kingdom, the levels of such impurities are well-controlled (see Table 3). Finally, it is important to recognize that these steels have a complex microstructure at start of life (SOL), and that there may be significant spatial variation in the local microstructure.14 For example, the local dislocation density or size and number of second-phase precipitates may vary significantly both within a given material and also between plates, welds, or forgings produced to the same nominal specification.
JRQ
200
Unirradiated 150
Irradiated 290 ⬚C 27 mdpa
100 50 0 -100
15.7 mdpa -50
0
50
100
150
200
250
300
Temperature (⬚C) Figure 1 Effect of irradiation on the Charpy-V USE of the JRQ (A508 cl.3 forging) irradiated in a materials test reactor to doses of 15.7 and 27 mdpa (1 mdpa ¼ 0.001 dpa).
data directly on fracture toughness (see, e.g., Viehrig and Lucon17) (see also Chapter 4.14, Fracture Toughness Master Curve of bcc Steels). Correlations have been established between the different measures; for example, see Sokolov and Nanstad18 for experimental data on the relationship between Charpy and static fracture toughness shifts and between Charpy and yield strength shifts. Williams et al.19 have published hardness/Charpy shift correlations for A533B plate and weld. From the early test programs in the 1960s to the present day, great reliance has been placed on measuring the shift in the Charpy transition from lower energy to the upper shelf energy (USE) as a means of determining the effect of irradiation damage on bulk mechanical properties. Figure 1 shows an example for an RPV forging; it can be seen that not only is there a shift in the transition temperature (usually measured at the 41 J level), but there is also a drop in the upper shelf. Note that the decrease in the USE is accompanied by a decrease in the slope of the Charpy curve in the transition region; this can create issues with estimating the Charpy shift if the upper shelf level approaches the indexing level. The greatest attention has been given to the shift in the transition region.
156 Table 4
Radiation Damage of Reactor Pressure Vessel Steels Parameter range of interest for different reactor types
Reactor type
Temperature range
Dose range (n cm2) > 1 MeV a
Dose range (mdpa)
Dose rate (n cm2) E > 1 MeV s1
Magnox PWR BWRs
160–>390 C 270 C–296 C Predominantly 270 C
1016–2 1018 fast n cm2 (Ni doses) 6–8 1019 <1–2 1018
0.02–4 60 <3
1 1011 1 1010
a In the LWR, community fluence (or equivalently dose) is generally expressed in terms of neutron exposure (e.g., n cm2, E > 1 MeV); this reflects the fact that the neutron spectrum does not vary significantly from location to location or plant to plant. In contrast, in the United Kingdom a key feature of Magnox reactors is that the neutron spectrum varies significantly with location, and displacements per atom, dpa, is a more appropriate measure.
The change in mechanical properties is required as a function of irradiation fluence (dose), dose rate (flux), irradiation temperature, steel type and composition, product form (plate, forging, or weld), and material heat treatment. Insight into the ranges of the irradiation variables of interest can be seen from Table 4. At the design stage of the early reactors, an allowance for the potential irradiation-induced embrittlement was included. The magnitude of this allowance was based on the judgment of potential irradiation effects. At the time of the design of Magnox reactors, the shift allowance for SMA welds was 40 C.20 Such allowances were frequently exceeded once data relevant to plant conditions became available. The intent of this section is not to provide an exhaustive description of the exact mechanical property response of individual steels to specific irradiation conditions; rather the intent is to identify the main parameters that control the radiation response of ferritic pressure vessel steels and to give an indication of the potential changes in mechanical properties that may occur through typical in-service conditions and lifetimes. The fluence dependence of RPV embrittlement has been studied since experimental programs were initiated in the 1950s. Most data available refer to the case in which embrittlement is determined by irradiation-induced hardening (rather than nonhardening embrittlement – see Section 4.05.4) at temperatures greater than 150 C and less than 300 C. Indeed, the available data are dominated by experiments that investigated the embrittlement of MnMoNi steels at an irradiation temperature of 270–295 C. It should also be recognized that after over 50 years of experimentation, a significant quantity of data has been obtained. For example, the US surveillance database arising from BWR and PWR reactors now encompasses 800 individual data points on the embrittlement observed from various steels irradiated at a range of fluences (and a
restricted range of irradiation temperatures).21,22 Data on the results of French surveillance programs have also been published.23 In parallel, wellcontrolled experiments in MTRs, frequently making use of steels with well-controlled compositional variations, have been performed. These have provided incredibly valuable information that has helped develop an understanding of the radiation damage processes in RPV steels. Indeed, early irradiation programs by Odette and coworkers,24 and Hawthorne and coworkers,25 focused on a restricted fluence range but on a number of steels with well-controlled compositional variation. Other notable programs were the irradiations performed as part of the Heavy-Section Steel Irradiation (HSSI) program at Oak Ridge (see, e.g., Taboda et al.26 and Nanstad and Bergen27) and a number of IAEA coordinated programs (see, e.g., International Atomic Energy Agency28). The most recent program by Odette and coworkers at the Ford Nuclear Reactor, University of Michigan, focuses on a range of fluxes, fluences, and material compositions. The resulting IVAR irradiation database encompasses 57 alloys that were irradiated at three different fluxes and three different temperatures over a range of fluences, giving in total 1537 different alloy/irradiation conditions. Irradiations were at 270, 290, and 310 C, fluxes between 5 1010 and 1 1012 n cm2 s1, and fluences between 0.004 1019 and 4 1019 n cm2, E > 1 MeV (see Heatherly et al.15 and Eason et al.29). The irradiation temperature of the irradiation rig was restrained to <5 C and the uncertainty in fluence was 7%. (It should be noted that there are welldefined techniques for establishing the neutron spectra in specific locations in reactors and for evaluating the resultant damage in the irradiated material; see Heatherly et al.15 and ASTM Standard 185 on ‘Practice for Conducting Surveillance Tests for Light-Water Cooled Nuclear Power Reactor Vessels’ and related ASTM Standards for a discussion of these techniques.) It is to be noted that typical MTR dose rates
157
Radiation Damage of Reactor Pressure Vessel Steels
are 1 1012–1 1013 n cm2, E > 1 MeV s1, that is, one or two orders of magnitude higher than the PWR surveillance dose rates given in Table 4. Overall, it was found in all RPV steels of interest that embrittlement increases with increasing fluence, but the rate of embrittlement may decrease (with increasing fluence). Furthermore, embrittlement does not saturate in the fluence range of interest to power reactor applications. These trends are illustrated in Figure 2 for Magnox CMn steel SMA weld transition shifts and for a MnMoNi plate HSST-02. Irradiations in the 1960s demonstrated that composition was a major factor controlling the response of the ferritic low alloy steels employed in operating reactors. By the mid-1960s, it was thought likely that residual elements in steels could be responsible for much of the observed scatter in the irradiation embrittlement response.31,32 The work suggested that reducing the residual element content of A302-B steel would markedly improve the resistance
Transition shift (⬚C)
150 DTMeasured ¯ DTConverted ¯
100
Peak Cu precipitation
DTTotal 50 DTMatrix 0
(a)
0
5
Ö(dpa ⫻ 105)
10
15
Measured Δ T30 ft-lbs (⬚C)
120 100
60
Table 5 Comparison of shift in 30 ft-lbs (41 J) transition temperature (DT30 ft-lb) due to irradiation at 288 C for experimental and commercial weld deposits and the A543 reference plate studied by Potapovs and Hawthorne33
40
Material
80
20
Measured ΔT30 ft-lbs (⬚C)
0 0 (b)
to irradiation embrittlement at typical service temperatures of 550 oF (288 oC). In a pioneering set of studies, Hawthorne and coworkers undertook studies that explored the effect of material composition on irradiation embrittlement in a systematic manner. Studies were undertaken on materials irradiated at a controlled temperature of 288 C in the Union Carbide Research Reactor (UCRR) and in the light water-cooled and moderated test reactor, UBR, at the Buffalo Materials Research Center.25 A series of small (150 kg) laboratory melts were produced to the nominal plate steel specifications using pure elements. These were then split, generally into three blocks, two of which were remelted and selected residual element additions added, while the third was kept to provide a low residual element reference. Each steel was also compared to material obtained from normal commercial production. Potapovs and Hawthorne33 demonstrated that additions of Cu, and Cu and Ni, to a laboratory melt containing low level of residuals greatly increased the observed embrittlement (see Table 5). This must be regarded as a landmark paper in the understanding of the factors that control radiation damage in RPV steels. Note that it was 15 years before the underlying mechanisms were elucidated (see Section 4.05.4). The effect of different Cu and Ni levels in steels irradiated as part of US surveillance programs is illustrated in Figure 3. The effect of increased levels of Cu in steels of the same Ni level and the effect of increased levels of Ni at constant Cu level is clear (data taken from Eason et al.33).
1 ⫻ 1019 2 ⫻ 1019 3 ⫻ 1019 4 ⫻ 1019 Fluence (n cm−2, E > 1 MeV)
5 ⫻ 1019
Figure 2 (a) Magnox submerged-arc weld transition shift data. Reprinted with permission from Buswell, J. T.; Jones, R. B. In Effects of Radiation on Materials, 16th International Symposium, ASTM STP 1175; Kumar, A. S., Gelles, D. S., Nanstad, R. K., Eds.; American Society for Testing and Materials: Philadelphia, PA, 1994; pp 424–443. Copyright ASTM International and (b) HSST-02 reference plate irradiated in surveillance schemes in a number of US LWRs.22
Expt. weld 1934 Expt. weld 1938 Expt. weld 1938 Expt. weld 1948 Commercial Filler (0.24Cu)
Composition (wt%)
Fluence (1019 n cm2) [E > 1 MeV]
DT41 J (DT30 ft-lb ) C
Cu
Ni
0.06
0.77
3.5
53
0.07
1.62
3.5
111
0.07
1.62
3.5
200
0.03
1.56
3.5
110
0.24
1.58
3.5
415
158
Radiation Damage of Reactor Pressure Vessel Steels
140 120
Melt 67*
Melt 68*
<0.01% Cu 0.70% Ni
0.30% Cu 0.70% Ni (D⬚C)
(D⬚F) US weld 0.25 wt% Cu 0.54 wt% Ni US forgings 0.06 wt% Cu 0.7 wt% Ni
80 60 40 20 0 0
1⫻1019
3⫻1019
2⫻1019
−2
(a)
Fluence n cm
4⫻1019
0.015 0.025 0.025 + 0.021% Sn (%P) 150 0.003
300
5⫻1019
(E > 1 MeV)
Cv41J (30 ft-lb) transition temperature increase
Charpy shift ΔT41J (⬚C)
100
*Irradiated 288 ⬚C (550F) ~2.5 ´ 1019 n cm–2 > 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. <0.004% Sn) to the high phosphorus alloy did not affect embrittlement. They further established that an arsenic content of 0.035% does not have an observable effect on the radiation sensitivity of plates containing 0.020% P and <0.18% Cu. Alloying with 0.50% chromium also did not alter the radiation resistance of plates having a copper content of 0.30% Cu. In the United Kingdom, there was considerable interest in the effect of irradiation temperature because of the wide variation of temperatures in Magnox pressure vessels. Barton et al.34 reported the irradiation temperature, the dose rate, and the dependencies of the yield stress increase for CMn steels following irradiation in the PLUTO MTR. The work focused on EN2, a Si-killed mild steel and an Al-killed grain-size-controlled mild steel. The Cu contents of the steels were 0.14, 0.18, and
Radiation Damage of Reactor Pressure Vessel Steels
FT ðDsÞ ¼ 1:869 4:57 103 T
½1
This parameter has been important for the correlation of data from similar steels irradiated at different irradiation temperatures.2 Understanding the effect of flux, or dose rate, on radiation damage of ferritic steels has proved particularly important in the formulation of mechanistically derived DDRs. There was particularly strong interest in the United Kingdom in understanding the effects of flux on bulk properties as Magnox RPVs operated within a range of fluxes.2 A number of experimental investigations have examined the dose rate dependence of hardening.37 In the absence of precipitation effects, no influence of dose rate on irradiation hardening has been detected. Data obtained from over five orders of magnitude change in dose rate for C–Mn plate steels at an irradiation temperature of 200 C, typical of Magnox applications,37 demonstrated no dependence on dose rate. However, there is agreement that there is a strong effect of flux on the embrittlement of Cu-containing steels.38 In these steels, it was found that at doses before the ‘saturation’ of embrittlement (see Section 4.05.4) the rate of embrittlement with fluence increased with decreasing dose rate. Williams et al. studied the effect of dose rate on the embrittlement in low Ni welds at preplateau doses19 (1 1019 n cm2, E > 1 MeV is approximately 1.5 102 dpa). They reported the irradiation-induced shift in the 41 J transition temperature of a number of Mn–Mo–Ni SMA welds after irradiation in MTRs at dose rates between 6 1010 and 2 108 dpa s1 and doses of less than 30 mdpa. It was observed that for the welds SD and SL, in which the Cu levels are low (<0.15 wt%), there is no apparent effect of dose rate (Figure 5). At higher copper levels (0.56 wt% for weld SH, 0.36% for SG, and 0.24% for SF), there was a marked effect of dose rate at low
100 Weld SF Weld SH
Shift (⬚C)
0.13%, respectively, by weight. Irradiation temperatures were reported to have been controlled to 1 C34 and the fast neutron doses were restricted to a maximum of 2.5 1017 n cm2 (as measured by Ni monitors). The results between 100 and 350 C exhibited a simple linear dependence of yield strength increase as a function of irradiation temperature. Jones and Williams35 carried out an analysis of the Barton data34 and another dataset on the irradiation temperature dependence of similar steels from Grounes,36 and pointed out that the combined data form a homogeneous data set with relatively little variability. The least squares regression is
159
100
Welds SD and SL
Weld SG 0.0
0.010
0.0 Dose (dpa)
Herald Herald VT4 (low flux)
0.010 Dido Osiris
Figure 5 Comparison between results obtained at different dose rates and in different irradiations of quenched and tempered low alloy welds (DIDO, HERALD, and OSIRIS are all MTRs in which samples were irradiated). Reproduced from Williams, T. J.; Ellis, D.; Swan, D. I.; et al. In Proceedings of the 2nd International Symposium on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, Sept 1985; ANS: Monterey, CA, 1986.
doses (<0.01 dpa), where low dose rate produced a significantly higher shift (Figure 5). At higher doses, no such effect was observed. In studies where both Charpy specimens and tensile samples were irradiated, it was found that radiation damage in RPV steels generally caused an increase in both Charpy shift and yield strength. This is significant, as irradiation can also lower the fracture strength of materials (see Section 4.05.4) and cause nonhardening embrittlement. As will be shown, none of the DDRs developed for MnMoNi steels have found it necessary to account for the embrittlement due to nonhardening embrittlement. It is only in the case of CMn steels that evidence has been found for nonhardening embrittlement in irradiated SMA welds.2 There have been investigations of nonhardening embrittlement in experiments mounted in MTRs; for example, McElroy et al.39 and Nanstad et al.40 investigated nonhardening embrittlement of simulated coarse-grained heat-affected zones (CGHAZ). A review of intergranular embrittlement in RPV steels can be found in English et al.41 4.05.3.1
Summary
In contrast to the assumptions made at the design stage, it was found that radiation damage in ferritic RPV steels could cause a significant change in bulk properties, primarily an increase in DBTT (as manifested by an increase in Charpy 41 J level) and an increase in yield strength. Evidence of radiation
160
Radiation Damage of Reactor Pressure Vessel Steels
damage causing nonhardening embrittlement was also found in CMn steels. There are proven mechanical test techniques for determining the change in bulk properties and large irradiation programs have been performed. These have made use of both surveillance programs and MTRs. The level of the measured embrittlement depends on the fluence, flux, and irradiation temperature. The most important discovery was the sensitivity to steel composition, in particular a strong dependence of embrittlement on the levels of Cu, Ni, and P.
4.05.4 Development of Mechanistic Insight of Factors Controlling Current Plant Lifetimes 4.05.4.1
Introduction
The previous section included a description of how the effect of radiation damage on the bulk properties of ferritic steels was established from the 1950s and how critical insight into the important role of Cu arose in the late 1960s and early 1970s. Equivalent advances in mechanistic understanding did not occur for another 10 or 15 years. The improved mechanistic understanding in the 1980s had its origin in the identification of the controlling variables that emerged from the experiments discussed in Section 4.05.3. This stimulated considerable interest on the possible role of elements such as Cu in the embrittlement process. The other, possibly more important, reason was that in the 1970s there were only a few microstructural techniques available for characterizing irradiation damage. Transmission electron microscopy (TEM), the dominant technique, could not resolve the irradiation-induced damage in steels that resulted in a significant change in mechanical properties.14 However, in the mid-1980s there was an explosion of information resulting from the application of a range of different and improved microstructural techniques. These techniques have now been applied to PWR, BWR, and Magnox steels irradiated in surveillance locations of power reactors and to representative materials irradiated in materials testing reactors. A more recent advance that is highly relevant to developing detailed mechanistic insight is the advent of ‘multiscale’ modeling. Here, the power of modern computing tools is such that microstructural development (and the resultant change in mechanical properties) can be modeled across the
various length and timescales involved in RPV embrittlement. Such models are subject to intense R&D and the current capability can be seen in Wirth et al.,42 Soneda et al.,43 Becquart,44 and Domain et al.45 However, these models have not, as yet, made a direct impact on the development of DDRs and so the current state of multiscale modeling is not reviewed here. 4.05.4.2
Radiation Damage Mechanisms
Microstructural development of ferritic steels during service is driven by the interaction of neutrons and metal atoms. Collisions between the incident neutron and constituent atoms result in momentum transfer to the lattice atoms. If the transfer is above 40 eV, atoms can be permanently displaced from their lattice sites resulting in vacancy–interstitial pairs. Indeed, lattice atoms with tens of kiloelectron volts may be created and a branching, tree-like distribution of displaced atoms formed, termed a displacement cascade. Vacancy and interstitial clustering may occur within the cascade. Vacancies and interstitials escaping from the cascade give rise to concentrations of vacancy and interstitial point defects throughout the material. The fate of the point defects formed in the irradiation depends most sensitively on irradiation dose rate and temperature, and also material factors such as composition. This chapter focuses on the microstructural development of RPV steels. These steels experience a relatively low dose rate (and thus lifetime dose) at a relatively low temperature. The microstructure developed under these conditions is very different to that developed at high doses and high temperatures in the operating regime typical of fusion or fast reactors. The critical features are that at the temperature range of most operating pressure vessels both vacancy and interstitial point defects are mobile. Freely migrating defects (vacancies or interstitials) and mobile interstitial clusters escaping from the initial damage event may interact with point defect sinks, such as preexisting dislocations, recombine with each other, either directly or at solute traps, or cluster to form vacancy or interstitial clusters. The end result of the interactions described above is a microstructure comprising of a high density of small clusters in the matrix. These clusters may be point defect solute complexes, and, depending on steel composition, solute clusters formed from radiation-enhanced precipitation. At the dose, dose rate range, and temperature of interest to operating
Radiation Damage of Reactor Pressure Vessel Steels
Mechanistic Framework
The effect of radiation damage on ferritic RPV steels46–50 has for some time been considered in terms of The formation of matrix damage (MD), that is, defect clusters and dislocation loops. It is well established that in low copper steels the shift in impact or yield strength properties depends on √dose. The irradiation enhanced formation of copperenriched clusters (CECs). (CEC are also referred to as CRPs (Cu-rich precipitates) as they were originally assumed to be strictly Cu-rich rather than simply Cu-enriched.) It has been demonstrated that, in many low-to-medium Ni steels and alloys, the yield strength change due to copper precipitation rises to a plateau value that is then unchanged by subsequent irradiation. The irradiation induced/enhanced grain boundary segregation of embrittling elements such as P. The first two mechanisms contribute to embrittlement by increasing the steels’ hardness as illustrated in Figure 6(a). The third mechanism induces embrittlement without hardening. The latter mechanism is not necessarily found in all RPV steels under operating conditions. Indeed, for MnMoNi steels irradiated in surveillance schemes in Western LWRs, the observed embrittlement is associated with the first two mechanisms; that is, the total shift in the ductile to brittle transition, as measured at the Charpy 41 J level, is DT41J ¼ DT41JðCRPÞ þ DT41JðMDÞ
½2
where CRP ¼ Cu-rich precipitate and MD ¼ matrix damage, or equivalently the increases in yield strength, Dsy, is given by Dsy ¼ DsyðCRPÞ þ DsyðMDÞ
½3
The first two mechanisms serve to harden the material and increase the yield strength sy, while the third mechanism causes a drop in the fracture strength, sF. The effects of these changes on the fracture behavior are illustrated in Figure 6(b), where the temperature dependence of the yield stress, sy, and the fracture stress, sF, are plotted. It can be seen that the effect of
Total (high Cu)
Increase in yield strength or Charpy temperature
4.05.4.3
Total (low Cu) MD CEC for high Cu
CEC for low Cu (ft)1/2
(a) Yield stress (s y) or fracture stress (s F)
reactors, the clusters formed are usually thermally stable. Lastly, segregation of solutes to grain boundaries or other sinks may occur.
(b)
161
sF Embrittlement s F-
ΔTT1
ΔTT2
Hardening
s y+ sy
ΔTT3
Temperature
Figure 6 (a) Schematic showing the dose dependence of matrix damage and copper clustering, and (b) variation of yield stress, sy, and fracture stress, sF, with temperature.
irradiation in causing hardening or a change in sF is to cause a change in the transition temperature. In the figure, DTT1 is caused by an increase in sy, while segregation of P to grain boundaries can lower the fracture stress and result in a shift DTT2. If both mechanisms are operative, then a combined shift of DTT3 occurs. It is important to note that Ni and Mn are known to strongly influence hardening in steels containing low levels of Cu and also CEC hardening in Cu-containing steels. In Cu-containing steels, saturation of the cluster hardening has been demonstrated in steels containing up to 1 wt% Ni. At steel Ni levels above 1.5 wt% (and with Mn 1.2–1.7 wt%), cluster hardening has not been observed to saturate. The precise Cu, Ni, and Mn levels at which the plateau is suppressed have not been fully characterized, and are the subject of current research. Similarly, the exact influence of Ni and Mn on the embrittlement of low Cu steels has not been fully established and is again a subject of ongoing research. (The term standard MnMoNi steels is used to refer to steels with typical Mn levels (<1.5 wt%) and Ni levels 1–1.2 wt%. The limit on the level of Ni is the subject of ongoing debate
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Radiation Damage of Reactor Pressure Vessel Steels
with some workers preferring a limit of 1.0 wt%, with others promoting a higher limit.) Indeed, as will be described in Section 4.05.6, there is concern about whether at high doses (typical of those achieved in plant with extended lives) there may be deviations from the simple framework established above. 4.05.4.4 Cluster Development Under Irradiation: Cu 4.05.4.4.1 Introduction
The purpose of this subsection is to establish the level of understanding regarding the irradiation-induced formation of CECs, and how this understanding supports the mechanistic framework outlined above. Insight has been developed by characterizing populations of irradiation-induced CECs in Cu-containing steels and, in particular, the dependence of CEC structure, composition, size, and number density on material and irradiation parameters (e.g., steel composition, fluence, flux, etc.). Measurements have also been made on the level of Cu in the matrix (Cumatrix) which is not associated with any CEC or precipitate. This is an important measure as, clearly, Cumatrix will decrease during irradiation as CECs are formed, and at SOL it may be lower than at the bulk level if Cu is precipitated during the final heat treatment. Several techniques have been successfully employed to characterize the CECs formed during neutron irradiation of Cu-containing pressure vessel Table 6
steels (Table 6). TEM was first used to observe CECs, but the cluster sizes were close to the resolution limit for most TEMs, so atom probe (3DAP) and small angle neutron scattering (SANS) became the most commonly used methods to characterize CECs. However, these are not the only techniques; for example, an important development has been the advent of experiments employing a positron annihilation technique, coincidence Doppler broadening (CDB) spectra of positrons annihilating in aged or irradiated Cu-containing alloys or steels. CDB provides a means of identifying the elements around the annihilation site.52 It has become the standard practice to characterize the same as-irradiated specimens with a number of techniques.53,54 Part of the logic behind this is that all the techniques have limitations, either in terms of volume analyzed or complexity of interpreting experimental data, and that combining techniques provides better information. These microstructural techniques have been discussed in greater detail in English and Hyde.55,56 4.05.4.4.2 CEC characterization
Evidence from various independent studies using techniques such as FEGSTEM or 3DAP has confirmed the small solute clusters to be multi-alloyed with minor constituents of the steel. An example of the composition of a cluster formed in an irradiated low Ni A533B weld metal is shown in Figure 7. The results of the 3DAP study clearly show that the CECs are alloyed
Techniques for the characterization of irradiation-induced solute clusters
Property
Existing techniques
Comment
Shape and size of solute clusters
3DAP, SANS, TEM/FEGSTEM
Composition of solute clusters
3DAP, SANS, FEGSTEM/EDX
Number density (Nd) and volume fraction of solute clusters Level of matrix Cu
3DAP, SANS, FEGSTEM/ EFTEM 3DAP, FEGSTEM/EDX
Analysis of the irradiation-induced scattering data from SANS provides a feature size distribution (for clusters 1–10 nm diameter) averaged over a volume approximately 50 mm3. 3DAP (LEAP) provides a direct estimate of the size of clusters formed in a region 60 60 100 nm3 3DAP yields a direct measure of the composition of each cluster observed. SANS measurements of irradiation-induced scattering in a strong magnetic field provide indirect evidence of cluster composition. Techniques are starting to be developed for measuring the level of Fe in Cu-enriched precipitates using advanced TEM techniques. At present, uncertainties on composition of small (2 nm diameter) are large FEGSTEM and 3DAP can directly measure the number density of visible clusters in the analyzed volume. In practice, Nd estimates from all techniques are subject to considerable uncertainty 3DAP provides a direct measure of the level of Cu in the matrix that is not in clusters. In the FEGSTEM techniques have been developed for measuring the level of matrix Cu from areas where clusters are not observed51
Radiation Damage of Reactor Pressure Vessel Steels
163
those measured. Furthermore, he found that the concentration of Fe in the precipitate phase is a function of aging temperature with less Fe at higher aging temperatures. Consensus has not yet been reached on the precise Fe content in irradiation-induced clusters in ferritic steels and is the subject of ongoing research. 4.05.4.4.3 Development with increasing fluence
Silicon Manganese Nickel Copper Figure 7 Example of single irradiation-induced cluster approximately 3 nm in diameter.
with Mn, Ni, and Si; there is also evidence of an association with P near the cluster–matrix interface (not shown). Similar cluster compositions have been found by other workers using other atom probes and FEGSTEM (which is not sensitive to Fe in clusters), and there is good agreement between the techniques. The atom probe data indicate that the clusters are dilute in that the majority element is Fe.57–59 In contrast, analysis of SANS data is often undertaken assuming that irradiation-induced clusters are nonmagnetic (implicitly Fe-free). The assumption of low levels of Fe in the Cu clusters was supported by thermodynamic calculations that predicted low levels of Fe in such precipitates.47 Furthermore, PAS-OEMS on as-irradiated (290 C) Fe–0.9Cu or Fe–0.9Cu–1Mn indicate that the clusters in this condition are not magnetic and therefore very unlikely to contain Fe,60 although it should be noted that clusters in model alloys may be different from those found in RPV steels. Carter demonstrated that the scattering observed in SANS experiments is not inconsistent with the presence of magnetic clusters.57 More recently, Morley et al.61 attempted to characterize the extent of trajectory aberrations in the atom probe that might give rise to an incorrect estimate of the true Fe content in small clusters in thermally aged (330–405 C) RPV steels. He concluded that the clusters do contain Fe, but the levels are lower than
The progressive precipitation of the Cu present in solution at the SOL as a result of irradiation has been recognized since the 1980s. There are a large number of studies that show that a high density of small CECs are formed under irradiation, and that the number density, size, and volume fraction are strongly dependent on the irradiation fluence, flux, and the material composition. For MnMoNi steels and Fe–Cu alloys, the number density and volume fraction increase rapidly with increasing fluence, and then there is an appearance of saturation, that is, a pattern of behavior that mirrors the shape of the curve in Figure 6(a). This is illustrated in Figure 8 from the work of Odette et al. (see data presented in Eason et al.29 where the volume fraction, fp, and radius have been derived from SANS data). It can be seen that the radius increases with fluence in this example. Auger et al.62 had found a similar pattern of behavior in SANS and AP data from ten steels and two Fe–Cu alloys with 0.2 wt% Cu. Saturation occurred at a similar fluence to that of Odette et al., that is, 1 1019 n cm2 (for irradiation temperatures close to 290 C). The rate at which the volume fraction increases with fluence is also strongly dependent on the irradiation flux and the composition of other elements such as Ni. Figure 9 shows the SANS measurements of Williams and Phythian63 on MnMoNi SAWs. It shows the effect of dose rate, Cu, and Ni on the development of CEC volume fractions with dose. Decreasing the Cu decreases both the volume fraction and CEC size. In the high Cu welds (seen most clearly at high dose rate), increasing the Ni clearly increases the total volume fraction of CECs formed at all doses. At the same time, the mean CEC diameter is somewhat decreased in the higher Ni weld (thus the volume fraction increase is associated with a large increase in cluster number density). It is also evident that, while the precipitated volume fraction appears to be saturating at the highest dose in the low Ni welds, there is no sign that saturation is close in the high Ni welds. A number of authors have found similar results.29,64,65
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Radiation Damage of Reactor Pressure Vessel Steels
2
0.6
fp (%)
(nm)
0.4
0.2
0.0 0.01
0.1 1 ft (1023 n m−2)
1
0 0.01
10
0.1
1
10
ft (1023 n m−2)
6
Volume fraction (⫻10–3)
Volume fraction (⫻10–3)
Figure 8 SANS data on volume fraction, fp, and, radius, rp, for 0.4 wt% Cu, 1.25 wt% Ni split melt model steel alloys (LD) irradiated at three flux levels between 0.6 and 10 1015 n m2 s1 in IVAR at 290 C, plotted as a function of fluence, ft. Reproduced from Eason, E. D.; Odette, G. R.; Nanstad, R. K.; Yamamoto, T.; EricksonKirk, M. T. A Physically Based Correlation of Irradiation-Induced Transition Temperature Shifts For RPV Steels; Oak Ridge Report ORNL/TM-2006/530, 2007.
5 4 3 2 1 0 0
(a)
10
20 30 Dose (mdpa)
6
4
2
0 0
40 (b)
10
20 30 Dose (mdpa)
High Cu, High dose rate
High Cu, Medium dose rate
High Cu, Low dose rate
Low Cu, High dose rate
40
Figure 9 Effect of bulk composition and dose rate on CEC size and volume fraction in MnMoNi SAWs. In (a) data are from low Ni welds, high Cu 0.08Ni, 0.55Cu; low Ni, low Cu 0.08Ni, 0.15Cu; and in (b) data are from high Ni welds, high Cu 1.65Ni, 0.55Cu; high Ni, low Cu 1.5Ni, 0.05Cu (all wt%). In both figures could be seen high flux 5.5 109, medium flux 6.3 1010, and low flux 9 1011 (all dpa s1). Reprinted with permission from Williams, T. J.; Phythian, W. J. In Effects of Radiation on Materials, 17th International Symposium, ASTM STP 1270; Gelles, D. S., Nanstad, R. K., Kumar, A. S., Little, E. A., Eds.; American Society for Testing and Materials: Philadelphia, 1996; p 191. Copyright ASTM International.
It is to be noted that Soneda65 found that when comparing Cu clusters in low-fluence-irradiated steels formed at a flux of 109 and 1010 n cm2 s1 E > 1 MeV, there was significant effect of dose rate on cluster size rather than number density. The effect of Ni was established early on,66 but it has only just been established that Mn also has an important effect. Odette et al. demonstrated from SANS studies that at constant Cu and Ni, increasing Mn decreased the size of the clusters but increased their number density as illustrated for 0.4 wt% Cu, 0.8 wt% Ni, and 3.4 1019 E > 1 MeV in Figure 10. In higher Ni steels, Burke et al.67 have also demonstrated that removing Mn from a steel significantly lowers the resultant embrittlement and the level of observable solute-related damage.
It was thought for many years (e.g., Jones and Bolton2) that CEC-related hardening would reach a maximum level once all the Cu had precipitated, and remain at this level as the CEC size remained constant – probably as a result of a balance between cluster nucleation and growth, and cluster destruction in cascades, that is, overaging did not occur. Jones and Bolton2 reported measurements of Cu cluster diameter using SANS on unirradiated and irradiated C–Mn SMA welds. It was shown that under surveillance conditions, and at temperatures below about 300 C, Cu clusters grew to about 2 nm in diameter. Even after subsequent accelerated irradiations (of the surveillance samples irradiated to the lowest doses) to doses between 200 105 and 1200 105 dpa, the mean precipitate diameter was still 2 nm.
2.0
6.0
1.8
5.0 NP (1023 m−3)
(nm)
Radiation Damage of Reactor Pressure Vessel Steels
1.6 1.4 1.2 1.0 0.0
0.5
1.0 Mn (%)
1.5
4.0 3.0 290 ⬚C 3.4 ⫻ 1023 n m–2
2.0
0.4 wt% Cu 0.8 wt% Ni 2.0
165
1.0 0.0
0.5
1.0 Mn (%)
1.5
2.0
3 Increasing Cu and Ni CEC volume fraction
Number density (⫻1017 cm−3)
Radius of gyration (nm)
Figure 10 SANS data on cluster radius, rp, number density, Np, and volume fraction, fp for 0.4 wt% Cu split melt model steels irradiated at high IVAR flux at 290 C. Effects on Mn variations in alloys with 0.8 wt% Ni. Reproduced from Eason, E. D.; Odette, G. R.; Nanstad, R. K.; Yamamoto, T.; EricksonKirk, M. T. A Physically Based Correlation of Irradiation-Induced Transition Temperature Shifts For RPV Steels; Oak Ridge Report ORNL/TM-2006/530, 2007.
2 1 0 5 4 3 2
Decreasing flux Ni
(ft)1/2
1 0
May decrease at high fluence
0
2 4 Dose (⫻1019 n cm−2) Doel-1
6
Figure 12 Schematic of the effect of selected material and irradiation variables on CEC volume fraction.
Doel-2
Figure 11 Radius of gyration and number density of CRPs for the RPV surveillance test specimens of Doel-1 (△) and Doel-2 (). Reproduced from Toyama, T.; Nagai, Y.; Tang, Z.; et al. Acta Mater. 2007, 55, 6852–6860.
More recent observations at the onset of the plateau in CEC formation in a number of commercial steels have shown that only around half of the available Cu had precipitated (Auger et al.)62 This suggests that particle coarsening and overaging could occur in irradiated RPV steels as well as in thermally aged steels, once the precipitation level was high. Particle coarsening has been reported in MnMoNi surveillance material from the Doel-1 and Doel-2 reactors,68,69 as shown in Figure 11. This particle coarsening should result in overaging (i.e., softening) after the hardness reaches a maximum value. Figure 12 illustrates schematically the influence of selected material and irradiation variables on the
CEC volume fraction (nominally for a Cu-containing MnMoNi steel irradiated at 290 C). The Cu level in the matrix at the SOL is an important parameter as it is the matrix Cu that is available for precipitation of CECs. This has been determined by either thermodynamic modeling or by direct measurement. For example, modeling calculations were performed by Buswell and Jones70 to determine matrix Cu levels in Magnox SMA welds with bulk Cu contents between 0.13 and 0.31 wt%. They found that the precipitation of Cu during the final weld stress-relief heat treatment (at 590–600 C), and also during the subsequent extremely slow cooling (5 C h1) of the RPV before reactor operation, reduced the maximum Cu available to precipitate during irradiation to no more than 0.15–0.20 wt%.70 Precipitation during the final weld stress relief also occurs in US steels. Indeed, a consensus has emerged that there is an upper limit to SOL-dissolved Cu, which is dependent on heat treatment. McElroy and
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Radiation Damage of Reactor Pressure Vessel Steels
Lowe have shown that even differences of 20–30 C in the heat treatment temperature can markedly affect the dissolved Cu content,71 as can the slow cooling of real structures from the stress-relief temperature. 4.05.4.5
Development of Matrix Defects
4.05.4.5.1 Introduction
Historically, the hardening observed in low Cu steels has been considered to arise from matrix defects. MD arises from the clustering of irradiation-induced point defects to form either vacancy or interstitial clusters, and/or solute–defect complexes (see, e.g., Odette and Lucas53). In addition, various solutes may diffuse to these clusters giving rise to complex defect–solute configurations. A critical factor in attempting to develop insight is that there is no one technique that allows a direct characterization of matrix defects. Insight has been obtained from indirect studies using techniques such as positron annihilation, where the nature of matrix defects can be inferred only after data analysis or modeling. Within these constraints, there are two aspects which should be discussed. First, the nature of MD, in particular whether it is vacancy or interstitial clusters that give rise to the observed hardening, and, second, the evolution of MD clusters at different fluxes and increasing fluence. The discussion also needs to include an assessment of whether MD is dependent on the presence of other solutes and whether MD can be treated independently of CEC formation. It should be noted that in generating insight into MD, studies of model alloys (e.g., Fe–Cu alloys) have been particularly important, including ion and electron irradiation studies. This is in contrast to the study of Cu precipitation where the majority of information has arisen from studies of Cu-containing neutron-irradiated steels. 4.05.4.5.2 Nature
Experimental information on the nature of matrix defects necessarily splits into evidence for clusters formed by vacancies and interstitial point defects. Significant insight into potential vacancy clusters has come from positron studies of irradiated and model steels. Positrons are well established as a tool for probing vacancy-type defects in solids, where the defect concentration is typically 1018 cm3. (Positrons are attracted to regions of the lattice which are more ‘open’ than average. The most obvious such regions are vacancies and vacancy clusters (the larger clusters being stronger positron traps). Less
obvious open regions include the tensile parts of a dislocation strain field (even around an interstitial loop) or incoherent particle–matrix interfaces. The positron annihilation techniques have included both positron lifetime (PA-t) and lineshape (PALA) analyses or, more recently, CDB. The latter is an important recent development. CDB (also known as positron annihilation spectroscopy-orbital angular momentum distribution, PAS-OEMS) measures the energy spectrum of the gamma rays produced at the positron annihilation sites, and thus the momenta of the electrons at those sites. The energy (or momentum) spectra characterize the elements surrounding the positron traps.72 Overall, the experimental data provide strong evidence for open-volume clusters that are sensitive to positrons, and for PA being a useful technique for studying MD (see, e.g., Buswell and Highton,73 Dai et al.,74 Carter et al.54). In model alloys, a number of authors report evidence for large vacancy clusters or microvoids after irradiation. Most interestingly, although the positron lifetime studies have indicated the presence of vacancy clusters in model alloys, such studies have not identified vacancy clusters in commercial steels (e.g., Dai et al.74). The study by Valo et al.75 is most convincing in this aspect, as the authors investigated both model alloys and RPV steels. Evidence for interstitial clusters has come primarily from studies of model alloys or steels irradiated to very high doses in excess of that of interest for most power reactor applications. Such studies have normally employed TEM techniques that are sensitive to the strain field associated with small dislocation loops. The imaging of such features is difficult in ferritic materials because of the need to correct for the image distortion caused by the ferromagnetic behavior of the samples, and because of the contrast arising from surface oxide on the thinned specimens. Critically, the resolution for small dislocation loops is 2–3 nm in even well-prepared samples imaged in higher voltage microscopes. Krishnamoorthy and Ebrahimi76 and Hoelzer and Ebrahimi77 reported the formation of visible interstitial loops in Fe irradiated to 4.63 1023 n m2; E > 1 MeV at 280 C; the loops increased in size and decreased in number density after annealing at 500 C. There have been fewer studies on irradiated steels, but in MnNiMo steels little evidence for dislocation loops has been reported. Soneda and coworkers have undertaken weak-beam TEM observations of RPV surveillance steels containing 0.06 and 0.12 wt% Cu irradiated to 4 1019 n cm2.65
Radiation Damage of Reactor Pressure Vessel Steels
(b)
(a)
g 20 nm
g
Figure 13 Contrast analysis on dislocation loops: weak-beam images of the same area imaged (a) with diffraction vector g ¼ (011) and (b) g ¼ (200) close to the [011] pole, in the foil at 300 nm depth. Reproduced from Fujii, K.; Fukuya, K. J. Nucl. Mater. 2005, 336, 323–330.
Soneda reported the formation of interstitial dislocation loops, whose diameter and number density are 2–3 nm and of the order of 1021 m3, respectively. At very high doses (and dose rates) a uniform density of loops has been observed. Fuji and Fukura78 undertook a weak-beam TEM study for MD in A533B RPV steel produced by 3 MeV Ni2þ ion irradiation to a dose of 1 dpa at 290 C. The MD was found to consist of small dislocation loops. The observed and analyzed dislocation loops have Burgers vectors b ¼ a <100> (Figure 13). The dislocation loops have a mean image size d ¼ 2.5 nm and the number density is about 1 1022 m3. Most of the loops are stable after thermal annealing at 400 C for 30 min. This indirect evidence suggests that their nature is interstitial. Kocik et al.79 examined the radiation damage microstructures in Cr–Mo–V surveillance base metal and weld containing (0.06–0.07 wt% Cu and 0.012–0.014 wt% P) irradiated up to 6 1024 n m2 (E > 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.
167
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.
Radiation Damage of Reactor Pressure Vessel Steels
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
Radiation Damage of Reactor Pressure Vessel Steels
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 <0.02 wt% P.’’101 4.05.4.8
Summary
The formation of matrix defects has been established as an important contribution to RPV embrittlement. The studies of low Cu steels have established that they are produced continuously during irradiation but the nature of these defects is still uncertain. The experimental data provide strong evidence for clusters that are sensitive to positrons; and for PA being a useful technique for studying MD. Although microvoids have been identified in model alloys from positron lifetime studies, none has been identified in commercial steels from such studies. It is important to note that there is increasing evidence that vacancy
clusters are not responsible for the observed hardening; it has been inferred that interstitial clusters are responsible for the observed hardening in RPV steels. Postirradiation annealing has been shown to be a powerful means of investigating the nature of the matrix defects further. A major development has been characterizing MD as being due to two components; first, SMD and second, at high fluxes, UMD. UMD are matrix defects which, although thermally unstable at the irradiation temperature, are frozen into the microstructure during the cooldown after irradiation.
4.05.5 Development of Mechanistically Based DDRs 4.05.5.1
Introduction
DDRs are equations that describe the changes in mechanical properties (yield stress, Charpy transition temperature, upper shelf toughness, hardness, etc.) as a function of neutron dose. The purpose of this section is to set out the mechanistic DDRs that have been developed to describe the embrittlement of RPV steels incorporated in the different reactor classes worldwide. However, it is important to note that the earliest DDRs were purely statistical in nature. The formulas were derived in different countries from the analysis of (different) databases. The first prediction formula in the United States that included the effect of residuals (‘residuals’ refers to levels of solutes such as Cu) was published in 1975. It was then followed by the Regulatory Guide 1.99 revision 1.106 This Guide has been particularly influential and presents some important features, such as the existence of thresholds in chemical composition, and an explicit dependence on Cu and Ni. It was followed in 1988 by USNRC Regulatory Guide 1.99 Revision 2.107 Petroquin108 has reviewed and compared the formulas employed for the prediction of irradiation embrittlement of reactor vessel materials, including the empirical DDRs developed in France (FIS and FIM formulae109), Germany (KTA110,111), and Japan ( Japan Electric Association Code ( JEAC) 4201 1991112,113). We focus here on DDRs (or embrittlement correlations) that have been developed with the form of the equations reflecting mechanistic understanding of the development of radiation damage in RPV steels, while the exact parameterization has been undertaken through fitting the models to large mechanical property databases arising from the testing of irradiated surveillance samples. The major driving force for such developments has been the
Radiation Damage of Reactor Pressure Vessel Steels
need to take advantage of the greatly improved understanding of embrittlement mechanisms in DDRs that enable interpolation or extrapolation with improved confidence to a parameter space poorly covered by a given (national) surveillance database. A common feature of such DDRs is that they follow the same mechanistic framework (described in Section 4.05.4), but give different weights to the parametric dependencies of radiation damage that have been described in the previous sections. Properly describing the effect of flux on the embrittlement of both low Cu and Cu-containing steels has been subject to extensive debate (see Section 4.05.6). A further common feature is that the DDRs have been refined as new surveillance data have become available, frequently with changes in the form of the equations in order to accommodate more sophisticated mechanistic understanding. It is important to note that they have been developed to describe embrittlement under relatively low dose rate conditions that apply to specific steel types, that is, CMn or MnMoNi steels. Two classes of steels have been described by such DDRs, namely, CMn RPV steels employed in Magnox reactors and MnMoNi steels used in Western LWRs (primarily in the United States and Japan). The mechanistically based DDRs for CMn steels were developed in the 1980s while it was not before 1998 that the first such DDR was published describing embrittlement in US RPV steels. 4.05.5.2
DDRs for CMn Steels
DDRs used to predict the embrittlement of the C–Mn steels used in the UK Magnox RPV had been mechanistically based from the 1980s.2 These predict radiation-induced changes in yield stress (hardening) or embrittlement (Charpy impact energy transition temperature or fracture toughness transition temperature) as a function of radiation dose and temperature. The approach adopted has been set out in Jones and Bolton2 and Wootton et al.3,114 The advantage of this approach was that the derived relationships could be used with confidence when limited extrapolation was required into regions of neutron dose, dose rate, or irradiation temperatures that were not specifically included in the surveillance database. It is important to note that it was the advances in understanding that enabled the adoption of a mechanistic approach (rather than adopting an empirical approach which had been followed in all other embrittlement correlations of this time). More specifically,
171
the seminal work of Fisher and coworkers in the early 1980s50 assumed that changes in yield stress arose from the combined effects of irradiation damage clusters and copper precipitates. Subsequently,2,3,114 a two-term relationship was finally adopted2 to model both hardening (Ds) and embrittlement (DT40 J) and had the following form: 9 DT40J > = 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
172
Radiation Damage of Reactor Pressure Vessel Steels
showing minimal effects. DDRs had to be developed that accommodated a nonhardening embrittlement mechanism. In addition, it was established that thermal neutrons could make a significant contribution to the irradiation damage in side-core locations, and that they were not conservatively covered by the DDRs.115,116 This conclusion was reached from an analysis of 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 <1 keV), respectively; the constant k is the thermal neutron effectiveness factor for the material. The thermal neutron effectiveness factor was found to be material dependent,114 and separate values were estimated for the different RPV materials. It should also be noted that a large-scale sampling and testing program of SMA weld metal removed from a decommissioned RPV validated the assessment process114 detailed above. Wotton et al.114 note that as a result of successfully addressing these and other challenges when the last two steel pressure vessel stations closed in December 2006, they had achieved lifetimes of nearly 40 years. To quote the authors,114 ‘‘This radical approach was subjected to rigorous peer review and its acceptance by the UK Nuclear Installations Inspectorate (NII) regulator was a major achievement.’’ Part of the peer review process may be illustrated by Knott and coworkers117,118 which detail the result of independent peer review by the UK Technical Advisory Group on Structural Integrity (TAGSI) of, first, the principles underlying the assessment of mechanical properties of irradiated ferritic RPV steels, and, second, the effects of gamma irradiation dose on the properties of C–Mn steels used in RPVs.
4.05.5.3
US Mechanistically Guided DDRs
In the late 1990s, mechanistically guided correlations were developed to describe the embrittlement of RPV materials employed in the United States.119,120 In common with other mechanistic DDRs, the form of these correlations is determined by current mechanistic understanding, but, in this case, the coefficients employed in the various terms in the correlation are determined by fitting to the extensive mechanical property data on the embrittlement of RPV materials acquired in vessel surveillance programs (Charpy V-notch shift at the 41 J transition temperature, DT41 J). These correlations are developed to describe a specific fleet of reactors, namely the embrittlement of RPV materials in US boiling water and pressurized water reactors (BWRs and PWRs). The initial mechanistically based or guided correlation models were presented in NUREG/CR-6551, published in November 1998.119 The models discussed incorporated material chemical composition and various exposure variables to enable predictions of TTS and USE changes. Another embrittlement shift model was developed at about the same time on the same mid-2000 database under the auspices of the Electric Power Research Institute (EPRI) and the American Society for Testing and Materials (ASTM) E10.02 subcommittee (the E900 model120), published as E900-02 in 2002. This was a simplified form but did not have a strong dependence on flux.121 The embrittlement shift model in NUREG/CR6551 was updated in July 2000 with additional surveillance data collected since the earlier work; this is referred to in this report as the Draft 2000 model.122 Significantly, motivation for a new modeling effort came from the fact that 62 additional low flux BWR shifts became available in 2003 (described below). These data were significantly underpredicted by the previous shift models,119,120 so it was necessary to investigate the cause of the underprediction. Additional pressurized water reactor (PWR) data from surveillance reports (about 140 shifts) were also added to the database in 2003 and 2004. Finally, the reliability of the database was improved when all old and new surveillance data were reviewed for completeness, duplicates, and discrepancies, during the summer and fall of 2004, in cooperation with the ASTM Subcommittee E10.02 on Radiation Effects in Structural Materials.123 The DDR that has been incorporated into the latest USNRC Regulatory Guide on screening limits for pressurized thermal shock124 was produced by
Radiation Damage of Reactor Pressure Vessel Steels
DT ¼ MF þ CRP ðin FÞ
½7
where
pffiffiffiffiffiffiffiffiffiffiffi MF ¼ Að1 0:001718Ti Þ 1 þ 6:13PMn2:47 ðfte Þ
½8
and A ¼ 1.140 107 for forgings, 1.561 107 for plates, and 1.417 107 for welds Ti ¼ irradiation temperature ( F); P ¼ bulk P (wt%); Mn ¼ bulk Mn (wt%)
9 8 10 2 1 > > = < 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
Radiation Damage of Reactor Pressure Vessel Steels
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
Radiation Damage of Reactor Pressure Vessel Steels
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.
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Radiation Damage of Reactor Pressure Vessel Steels
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modifications in the form, or the values of, the fitting parameters. The major topics are the following:
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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 <60 mdpa can be understood in terms of the framework of the understanding developed for Cu-containing MnMoNi steels with <1.2 wt% Ni, that is, hardening from MD and solute-enriched clusters (see, e.g., the work of Williams et al.129). A notable difference is that there is little evidence for a plateau in hardening from CECs as the fluence increases; rather a continuous increase with fluence is observed (for fluences up to 60 mdpa). It is possible, however, that an additional high fluence embrittlement mechanism may operate in Ni and Mn-containing steels. Specifically, it has been suggested that at long times and or at high
Radiation Damage of Reactor Pressure Vessel Steels
doses, Mn, Ni, and Si could form a new phase in RPV steel.47,130 This late-blooming phase (LBP) would produce an additional increment of hardening at high fluences, that is, late-onset embrittlement or anomalous hardening at high doses. If this is the case, then the DDRs described in the previous section may become nonconservative. Recently, there have been an increasing number of observations of MnNiSi clusters in irradiated low Cu steels (see, e.g., Auger et al.131 and Soneda et al.127), and there is intensive research aimed at establishing whether NiMnSi clusters represent segregation to small microstructural features (thereby lowering interfacial or strain energies) or represent precipitates of a distinct Ni–Si–Mn-enriched phase that is thermodynamically favored at RPV operating temperatures and RPV steel compositions. It is the latter possibility that gives rise to the possibility of a late-onset embrittlement.
References 1. 2.
3. 4. 5. 6. 7. 8.
9. 10.
11.
12.
Steele, L. E.; Sterne, R. H. Nucl. Eng. Des. 1969, 10, 259–307. Jones, R. B.; Bolton, C. J. Neutron Radiation Embrittlement Studies in Support of Continued Operation and Validation by Sampling of Magnox Reactor Steel Pressure Vessels and Components, Twenty Fourth Water Reactor Safety Information Meeting, Bethesda, 1996; Vol. 2, pp 25–48. Wootton, M. R.; Moskovic, R.; Bolton, C. J.; Flewit, P. E. J. Energy Mater.: Mater. Sci. Eng. Energy Syst. 2008, 3(1), 45–56. Quoted in ‘‘An Assessment of the Integrity of PWR Pressure Vessels’’, Report by a study Group under Marshall, W., Dr., C.B.E., F.R.S (UKAEA 1976). Nikolaev, Yu. A.; Nikolaeva, A. V.; Shtrombakh, Y. I. Int. J. Press. Vessels Piping 2002, 79(8–10), 619–636. Kryukov, A. M.; Nikolaev, Yu. A.; Planman, T.; Platonov, P. A. Nucl. Eng. Des. 1997, 173(1–3), 333–339. Shtrombakh, Ya. I.; Nikolaev, Yu. A. J. ASTM Int. 2007, 4(5). Paper ID: JAI100276; DOI: 10.1520/JAI100276. Brumovsky, M. Surveillance Specimen Programs for WWER Type Reactors; American Society of Mechanical Engineers, Pressure Vessels and Piping Division (Publication) PVP, 2005; Vol. 7, pp 23–31; PVP2005-71475. Cadiou, L.; Comon, J.; Dunand-Roux, L.; Vassal, J.; Houssin, E. Materiaux et Techniques 1977, 6, 341. Recent Test Result of Chemical Composition and Fracture Toughness, SA508 Class 2 an Class 3 and SA-533B Grade B Class 1; Japan Steel Works Internal Report, Nov 1977. Onodera, S.; Fukioka, K.; Tsukada, H.; Suzuki, K. Advantages in application of integrated flange forgings for reactor vessels. In Presented at 3rd MPA Seminar, Stuttgart, Sept 15, 1977. Nuclear Pressure Vessel Steel Data Base; EPRI NP9343 Proj. 886-1; December 1978.
13. 14.
15.
16.
17. 18.
19.
20. 21.
22.
23. 24.
25.
26.
27.
28.
29.
177
Druce, S. G.; Edwards, B. C. Nucl. Energy 1980, 19(5), 347–360. Buswell, J. T. Examination of the Materials by Electron Microscopy, Annex 9 to Analysis of the Behaviour of Advanced Reactor Pressure Vessel Steels Under Neutron Irradiation: The UK Programme (Report from the UK for the IAEA Coordinated Research Program on the analysis of the behaviour of Advanced Reactor Pressure Vessel Steels under Neutron Irradiation); UKAEA, April 1983; pp 281–315. Heatherly, D. W.; Thoms, K. R.; Hurst, M. T.; Giles, G. E. Heavy Section Steel Irradiation Program’s Reusable Irradiation Facilities; Oak Ridge National Report ORNL/ TM-2002/77, April 2005. McCabe, D. E.; Merkle, J. G.; Wallin, K. In Fatique and Fracture Mechanics, 30th Volume, ASTM 1360; Paris, P. C., Jerina, K. L., Eds.; American Society for Testing Materials: West Conshohocken, PA, 2000; pp 21–33. Viehrig, H.; Lucon, E. In Proceedings of PVP2007 2007 ASME Pressure Vessels and Piping Division Conference, San Antonio, TX, USA, July 22–26, 2007; ASME: USA. Sokolov, M. A.; Nanstad, R. K. Comparison of irradiation-induced shifts of KJc and Charpy impact toughness for reactor pressure vessel steels. . In In ; Effects of Radiation on Materials, 18th International Symposium, ASTM STP 1325; Nanstad, R. K., Hamilton, M. L., Garner, F. A., Kumar, A. S., Eds.; American Society for Testing and Materials: West Conshohocken, PA, 1999; p 167. Williams, T. J.; Ellis, D.; Swan, D. I.; et al. In Proceedings of 2nd International Symposium on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, Sept 1985; ANS: Monterey, CA, 1986. Robertson, T. S. J. Iron Steel Inst. 1953, 175, 361–374. Papers in Radiation Embrittlement and surveillance of Nuclear Reactor Pressure Vessels: An International Review, 3rd Volume, ASTM STP 1011; Steele, L. E., Ed.; American Society for Testing and Materials: Philadelphia, 1996. Eason, E. D.; Wright, J. E.; Odette, G. R. Improved Embrittlement Correlations for Reactor Pressure Vessel Steels; NUREG/CR-6551; US Nuclear Regulatory Commission: Washington, DC, 1998. Brillaud, C.; Hedin, F. Effects of Irradiation on Materials, 15th ASTM Syrup.; Nashville, TN, 1990. Odette, G. R; Lucas, G. E. An investigation of Kinetic aspects of Irradiation Embrittlement of Light Water Reactor Pressure Vessel Steels; EPRI Final Report RP-1021-7, RP-2455-11. Hawthorne, J. R. An Exploratory Study of Element Interactions and Composition Dependencies in Radiation Sensitivity Development; Final Report Prepared NUREG/ CR-5357 – MEA-2341; 1989. Taboda, A.; Randall, P. N.; Serpan, C. Z. In Radiation Embrittlement of Nuclear Reactor Pressure Vessel Steels: An International Review, 3rd Volume, ASTM STP 1011; Steele, L. E., Ed.; American Society for Testing and Materials: Philadelphia, 1989; pp 27–38. Nanstad, R. K.; Bergen, R. G. Irradiation Effects on Charpy and Tensile Properties of Low Upper-shelf Welds; HSSI Series 2 and 3 USNRC Report NUREG/CR5859, August 1992. Analysis of the behaviour of Advanced Pressure Vessel Steels under Neutron Irradiation; Technical Report No. 265; International Atomic Energy Agency: Vienna, 1986. Eason, E. D.; Odette, G. R.; Nanstad, R. K.; Yamamoto, T.; EricksonKirk, M. T. A Physically Based Correlation of Irradiation-Induced Transition
178
30.
31.
32.
33. 34. 35.
36.
37.
38.
39.
40.
41.
42.
43. 44.
Radiation Damage of Reactor Pressure Vessel Steels Temperature Shifts for RPV Steels; Oak Ridge Report ORNL/TM-2006/530, 2007. Odette, G. R. O; Yamamoto, T. In Proceedings of the International Symposium on Research for Aging Management of Light water Reactors, Japan, Oct 22–23, 2007; pp 279–306, ISBN 978-4-902-71-02-3 C3053. Hawthorne, J. R.; Steele, L. E. In Effects of Radiation on Structural Metals, ASTM STP 426; American Society for Testing and Materials: Philadelphia, 1967; pp 534–572. 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. Potapovs, U.; Hawthorne, J. R. Nucl. Appl. 1969, 6(1), 27–46; and NRL Report 6803, Nov 22, 1968. Barton, P. J.; Harries, D. R.; Mogford, I. L. J. Iron Steel Inst. 1965, 203, 507–511. Jones, R. B.; Williams, T. J. In Effects of Irradiation on Materials, 17th International Symposium, ASTM STP 1270; Gelles, D. S., Nanstead, R. K., Kumar, A. V., Little, E. A., Eds.; American Society for Testing and Materials: West Conshohocken, PA, 1996; p 569. Grounes, M. Effects of Radiation on Structural Metals, ASTM STP 426; American Society for Testing and Materials: Philadelphia, 1967; pp 534–572; early considerations summarised in Steels for Reactor Pressure Circuits, Report of a Symposium held in London, Nov 30 to Dec 2, 1960; Iron and Steel Institute: London, 1961; pp 422–426. Jones, R. B.; Buswell, J. T. In Effects of Radiation on Materials, 17th International Symposium, ASTM STP 1270, Sun Valley, Idaho, June 20–23, 1994; Gelles, D. S., Nanstad, R. K., Kumar, A. S., Little, E. A., Eds.; American Society for Testing and Materials: Philadelphia, 1996. Carter, R. G.; Soneda, N. Proceedings of EPRI-CRIEPI Workshop on Dose Rate Effects in RPV Materials; EPRI: Palo Alto, CA, 2002, 1006981 and CRIEPI, Tokyo, Japan, T980203, July 2002. McElroy, R. J.; English, C. A.; Foreman, A. J. E.; et al. In Effects of Radiation on Materials, 18th International Symposium, ASTM STP 1325; Nanstad, R. K., Hamilton, M. L., Garner, F. A., Kumar, A. S., Eds.; American Society for Testing and Materials: West Conshohocken, PA, 1999; pp 296–316. Nanstad, R. K.; McCabe, D. E.; Sokolov, M. A.; English, C. A.; Ortner, S. R. In Effects of Radiation on Materials, 20th International Symposium, ASTM STP 1405; Rosinski, S. T., Grossbeck, M. L., Allen, T. R., Kumar, A. S., Eds.; American Society for Testing and Materials: West Conshohocken, PA, 2001. English, C. A.; Ortner, S. R.; Gage, G.; Server, W. L.; Rosinkski, S. T. In Effects of Radiation on Materials, 20th International Symposium, ASTM STP 1405; Rosinski, S. T., Grossbeck, M. L., Allen, T. R., Kumar, A. S., Eds.; American Society for Testing and Materials: West Conshohocken, PA, 2001; pp 151–173. Wirth, B. D.; Odette, G. R.; Stoller, R. E. In Advances in Materials Theory and Modeling – Bridging Over Multiple-Length and Time Scales; Bulatov, V., Cleri, F., Colombo, L., Lewis, L., Mousseau, N., Eds.; Materials Research Society: Pittsburgh, PA, 2001; Vol. 677, pp 5.2.1–5.2.6. Soneda, N.; Ishino, S.; Takahashi, A.; Dohi, K. J. Nucl. Mater. 2003, 323, 169–180. Becquart, C. S. Nucl. Instrum. Methods Phys. Res. B 2004, 228(1–4), 111–121.
45.
Domain, C.; Becquart, C. S.; Malerba, L. J. Nucl. Mater. 2004, 335, 121–145. 46. Phythian, W. J.; English, C. A. J. Nucl. Mater. 1993, 205, 162. 47. Odette, G. R. In Microstructure of Irradiated Materials MRS Symposium Proceedings Symp. held Boston, USA, Nov 1994; Robertson, I. M., et al. Ed.; Materials Research Society: Pittsburgh, 1995; Vol. 373, 137–148. 48. Knott, J. F.; English, C. A. Int. J. Press. Vessels Piping 1999, 76, 891–908. 49. Bolton, C. J.; Buswell, J.; Jones, R.; Moskovic, R.; Priest, R. The modelling of irradiation embrittlement in submerged-arc welds. In Effects of Radiation on Materials, 17th International Symposium, ASTM STP 1270; Gelles, D. S., Nanstad, R. K., Kumar, A. S., Little, E. A., Eds.; American Society for Testing and Materials: Philadelphia, 1996; p 103. 50. Fisher, S. B.; Harbottle, J. R.; Aldridge, N. Radiation Hardening in Magnox Pressure Vessel Steels; Philosophical Transactions Royal Society: London, 1985; Vol. A315; pp 301–332. 51. Dumbill, S.; Ortner, S. R. In IAEA/LMNPP Specialists Meeting on Irradiation Embrittlement and Mitigation Gloucester, England, UK, May 14th to 17th, 2001. 52. Nagai, Y.; Hasegawa, M.; Tang, Z.; et al. Phys. Rev. B 2000, 61, 6574–6578. 53. Odette, G. R.; Lucas, G. E. Radiat. Eff. Defects Solids 1998, 144, 189–231. 54. Carter, R. G.; Soneda, N.; Dohi, K.; et al. In Proceedings of International Symposium Fontevraud IV, Contribution of Materials Investigation to the Resolution of Problems Encountered in Pressured Water Reactors; Sept 14–18; SFEN: Fontevraud, 1998; pp 89–100. 55. English, C. A.; Hyde, J. M. In Comprehensive Structural Integrity; Milne, I., Ritchie, R. O., Karihaloo, B., Eds.; Elsevier: Amsterdam, 2003; Vol. 6, 351–398, doi: 10.1016/B0-08-043749-4/06110-3. 56. English, C. A.; Hyde, J. M. In Recent progress in the understanding of RPV embrittlement, In Proceedings of the International Symposium on Research for Aging Management of Light Water Reactors, Oct 22–23, 2007; Institute of Nuclear Safety System, Incorporated: Fukui City, Japan. 57. Carter, R. G.; Soneda, N.; Dohi, K.; Hyde, J. M.; English, C. A.; Server, W. J. Nucl. Mater. 2001, 298(3), 211–224. 58. Pareige, P.; Stoller, R. E.; Russell, K. F.; Miller, M. K. J. Nucl. Mater. 1997, 249, 165. 59. Miller, M. K.; Burke, M. G. In Effects of Radiation on Materials, 14th International Symposium, ASTM STP 1046; ASTM: Philadelphia, PA, 1990; Vol. 2, p 107. 60. Asoka-Kumar, P.; Wirth, B. D.; Sterne, P. A.; Howell, R. H.; Odette, G. R. Philos. Mag. Lett. 2002, 82(11), 609–615. 61. Morley, A.; Sha, G.; Hirosawa, S.; Cerezo, A.; Smith, G. D. W. Ultramicroscopy 2009, 109, 535–540. 62. Auger, P.; Pareige, P.; Welzel, S.; Van Duysen, J. C. J. Nucl. Mater. 2000, 280, 333–344. 63. Williams, T. J.; Phythian, W. J. In Effects of Radiation on Materials, 17th International Symposium, ASTM STP 1270; Gelles, D. S., Nanstad, R. K., Kumar, A. S., Little, E. A., Eds.; American Society for Testing and Materials: Philadelphia, 1996; p 191. 64. Odette, G. R.; Yamamoto, T.; Klingensmith, R. D. Phil. Mag. 2005, 85, 779. 65. 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.
Radiation Damage of Reactor Pressure Vessel Steels 66.
Odette, G. R.; Lucas, G. E. In Effects of Irradiation on Materials, 14th Symposium, ASTM-STP-1046; American Society for Testing and Materials: Philadelphia, 1989; pp 323–347. 67. Burke, M. G.; Stofanak, R. J.; Hyde, J. M.; English, C. A.; Server, W. L. In Effects of Radiation on Materials, ASTM STP 1447; Grossbeck, M. L., Lott, R. G., Eds.; American Society for Testing and Materials: West Conshohocken, PA, 2002; pp 194–207. 68. Hasegawa, M.; Nagai, Y.; Toyama, T.; et al. In Proceedings of the International Symposium on Research for Ageing Management of Light Water Reactors, held Fukui City, Japan October 2007; Eyre, B. L., Kimure, I., Eds.; INSS: Japan, 2008; pp 327–344. 69. Toyama, T.; Nagai, Y.; Tang, Z.; et al. Acta Mater. 2007, 55, 6852–6860. 70. Buswell, J. T.; Jones, R. B. In Effects of Radiation on Materials, 16th International Symposium, ASTM STP 1175; Kumar, A. S., Gelles, D. S., Nanstead, R. K., Little, E. A., Eds.; ASTM: USA, 1993; pp 424–443. 71. McElroy, R. J.; Lowe, A. L. In Effects of Radiation on Materials, 17th International Symposium, ASTM 1270; Gelles, D. S., Nanstead, R. K., Kumar, A. S., Little, E. A., Eds.; ASTM: USA, 1996; p 68. 72. Asoka-Kumar, P.; Alatalo, M.; Gosh, V. J.; Kruseman, A. C.; Nielsen, B.; Lynn, K. G. Phys. Rev. Lett. 1996, 77, 2097. 73. Buswell, J. T.; Highton, J. P. A Positron Annihilation Examination of Irradiated Pressure Vessel Steels; CEGB Report, TPRD/B/0787/R86; Berkeley Nuclear Laboratories, 1986. 74. Dai, G. H.; Moser, P.; Van Duysen, J. C. Study of the Nature of the Defects in a Highly Neutron Irradiated RPV Steel; ICPA-9; Hungary, Aug 25–30, 1991. 75. Valo, M.; Krause, R.; Saarinen, K.; Hautojarvi, P.; Hawthorne, J. R. In Effects of Radiation on Materials, 15th International Symposium, ASTM 1125; Stoller, R. E., Kumar, A. S., Gelles, D. S., Eds.; ASTM: Philadelphia, 1992; p 172. 76. Krishnamoorthy, V.; Ebrahimi, F. In Effect of Alloying Elements and Annealing on Irradiation-Induced Defects in Iron Alloys, Materials Research Society Proceedings, 1989; Vol. 138, pp 99–104. 77. Hoelzer, D. T.; Ebrahimi, F. In Iron Alloys, Mat. Res. Soc. Symp. Proc.; Materials Research Society: Pittsburgh, 1995; Vol. 373, p 57. 78. Fujii, K.; Fukuya, K. J. Nucl. Mater. 2005, 336, 323–330. 79. Kocik, J.; Keilov, E.; Cizek, J.; Prochazka, I. J. Nucl. Mater. 2002, 303, 52–64. 80. Miller, M. K.; Russell, K. F.; Kocik, J.; Keilova, E. J. Nucl. Mater. 2000, 282, 83–88. 81. Makin, M. J.; Whapham, A. D.; Minter, F. J. Phil. Mag. 1962, 7, 285. 82. Makin, M. J.; Minter, F. J. Acta Metall. 1960, 8, 691. 83. Jones, R. B.; Williams, T. J. In Effects of Radiation on Materials, 17th International Symposium, ASTM 1270; Gelles, D. S., Nanstead, R. K., Kumar, A. S., Little, E. A., Eds.; ASTM: Philadelphia, PA, 1996; p 569. 84. Kampmann, R.; Frisius, F.; Hackbarth, H.; Beavan, P. A.; Wagner, R.; Hawthorne, J. R. In Proceedings of 5th International Symposium on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, held Monterey, CA, 1991; American Nuclear Society: La Grange Park, IL, 1992; p 679, Order No 700176. 85. Mader, E.; Lucas, G. E.; Odette, G. R. In Effects of Radiation on Materials, 15th International Symposium, ASTM 1125; Stoller, R. E., Kumar, A. S., Gelles, D. S., Eds.; ASTM: Philadelphia, PA, 1992; p 151.
86.
87. 88.
89. 90. 91. 92. 93.
94. 95.
96. 97. 98. 99.
100. 101.
102.
103.
104.
179
Soneda, N.; Dohi, K.; Ishino, S.; Takahashi, A. Computer Simulation of the Effect of Neutron Irradiation on the Microstructural Evolution in bcc-Fe; CRIEPI T03076, 2004. Elliot, R. P. Constitution of Binary Alloys, First Supplement; McGraw-Hill: New York, 1965; p 424. Jones, R. B.; Buswell, J. T. In Proceedings of the 3rd International Symposium on Environmental Degradation of Reactor Materials – Water Reactors; Metallurgical Society Inc.: Warrendale, PA, USA, 1998; pp 111–120. Nagai, Y.; Tang, Z.; Hasegawa, M.; Kanai, T.; Saneyasu, M. Phys. Rev. B. 2001, 63, 134110, 10.1103/PhysRevB.63.134110. Nagai, Y.; Takadate, K.; Tang, Z.; et al. Phys. Rev. B 2003, 67, 224202. Miller, M. K.; Nanstad, R. K.; Sokolov, M. A.; Russell, K. F. J. Nucl. Mater. 2006, 351, 187. Little, E. A.; Harries, D. R. Metal Sci. J. 1970, 4, 188. Ortner, S. Microstructural Characterization of Reactor Pressure Vessel Steels: Post-irradiation Annealing Experiments, Joint EPRI-CRIEPI RPV Embrittlement Studies (1999–2004). EPRI: Palo Alto, CA, 2004, 1003531, and CRIEPI, Tokyo, Japan: Q980401. Russell, K. C.; Brown, L. M. Acta Metall. 1972, 20, 969–974. Phythian, W. J.; Foreman, A. J. E.; English, C. A.; et al. In Effects of Radiation on Materials, 15th International Symposium, ASTM STP 1125, Nashville, 1990; Stoller, R. E., Kumar, A. S., Gelles, D. S., Eds.; American Society for Testing and Materials: Philadelphia, 1992; pp 131–150. Foreman, A. J. E; Makin, M. J. Phil. Mag. 1966, 14, 911. English, C. A.; Phythian, W. J.; McElroy, R. J. Materials Research Society Symposium Proceedings; 1996; Vol. 439, pp 471. Odette, G. R.; Lucas, G. E. Radiat. Eff. Defects Solids 1998, 144, 189. Wirth, B. D.; Odette, G. R.; Pavenich, W. A.; Lucas, G. E.; Spooner, S. E. In Effects of Radiation on Materials, 18th International Symposium, ASTM STP 1325; Nanstad, R. K., Hamilton, M. L., Garner, F. A., Kumar, A. S., Eds.; American Society for Testing and Materials: Philadelphia, 1999; pp 102–124. Bacon, D. J.; Osetsky, Y. N. J. Nucl. Mater. 2004, 329–333, 1233. Final Technical Report Phosphorus Influence on Steels Ageing (PISA) Contract No: FIKS-CT-2000-00080. Work performed as part of the European Atomic Energy Community’s R&T Specific Programme Nuclear Energy, key action Nuclear Fission Safety, 1998–2002. http://cordis.europa.eu/documents/documentlibrary/ 87410651EN6.pdf. Druce, S. G.; English, C. A.; Foreman, A. J. E.; et al. The modelling of irradiation-enhanced phosphorus segregation in neutron irradiated reactor pressure vessel submerged-arc welds. . In In Effects of Radiation on Materials, 17th International Symposium, ASTM STP 1270; American Society for Testing and Materials: West Conshocken, PA, 1996; p 119. Miller, M. K.; Burke, M. G. In Effects of Radiation on Materials, 16th International Symposium, ASTM STP 1175; Kumar, A. S., Gelles, D. S., Nanstad, R. K., Little, E. A., Eds.; American Society for Testing and Materials: West Conshocken, PA, 1994; p 492. Nikolaeva, A. V.; Kevorkyan, Y. R.; Nikolaev, Y. A. In Effects of Radiation on Materials, 19th International Symposium, ASTM STP 1366; Hamilton, M. L., Kumar, A. S., Rosinski, S. T., Grossbeck, M. L., Eds.; American Society
180
105. 106.
107.
108. 109.
110. 111.
112.
113. 114.
115.
116.
117.
Radiation Damage of Reactor Pressure Vessel Steels for Testing and Materials: West Conshohocken, PA, USA, 1999; pp 399–411. Faulkner, R. G.; Jones, R. B.; Lu, Z.; Flewitt, P. E. J. Phil. Mag. 2005, 19(85), 2065–2099. Effects of residual elements on predicted radiation damage to reactors vessels materials, Regulatory Guide 1.99, Revision 1; US Nuclear regulatory Commission: Washington, DC; April 1977. Radiation embrittlement of reactor vessel materials, Regulatory Guide 1.99, Revision 2; Office of Nuclear Regulatory Research, US Nuclear Regulatory Commission: Washington, DC, May 1988. Petrequin, P. A Review of Formulas for Predicting Irradiation Embrittlement of Reactor Vessel Materials; AMES Report No. 6. EUR 16455, Nov 1996. Brillaud, C.; Hedin, F.; Houssin, B. In Effects of Radiation on Materials, 13th International Symposium, ASTM STP 956; Stoller, R. E., Garner, F. A., Henager, C. H., Igata, N., Eds.; American Society for Testing and Materials: Philadelphia, 1987; pp 420–447. Surveillance of the Irradiation Behaviour of Reactor Pressure Vessel Materials for LWR Facilities, Nuclear Safety Standard Commission, KTA 3203, GRS, Koln, Issue 3/84. Monitoring the Radiation Embrittlement of Materials of the Reactors Pressure Vessel of Light Water Reactors, Nuclear Safety Standard Commission, KTA 3203, GRS, Koln, Issue 6/01. Tomimatsu, M.; Urabe, Y.; Sanoh, J.; Iida, M.; Nakamura, T.; Tamura, A. Evaluation of RPV steel surveillance program in Japanese PWR: Radiation embrittlement prediction. In Proceedings of the 3rd International Symposium on Contribution of Materials Investigation to the Resolution of Problems Encountered in Pressurized Water Reactors, Fontevraud, France, Sept 1994; French Nuclear Energy Society; Vol. 1, p 627. Japan Electric Association. Method of Surveillance Tests for Structural Materials of Nuclear Reactors; JEAC4201-1991, 1991. Wootton, M.; Moskovic, R.; Bolton, C.; Flewitt, P. Magnox Steel Reactor Pressure Vessel Monitoring Schemes – An Overview January 2010JAI (January) 2010, 7(1); DOI: 10.1520/JAI101948. Jones, R. B.; Edens, D. J.; Wootton, M. R. In Effects of Radiation on Materials, 19th International Symposium, Seattle, June 1998, ASTM STP 1366; American Society for Testing and Materials: West Conshocken, PA, 2000; 366–382. Jones, R. B.; Cowan, J. R.; Corcoran, R. C.; Walmesly, J. R. In Effects of Radiation on Materials, 19th International Symposium, Seattle, June 1998, ASTM STP 1366; American Society for Testing and Materials: West Conshocken, PA, 2000; pp 473–491. Knott, J. F.; English, C. A.; Weaver, D. R.; Lidbury, D. P. G. Int. J. Press. Vessels Piping 2005, 82(12), 929–940.
118. 119. 120. 121.
122.
123.
124.
125.
126.
127.
128.
129. 130. 131.
Knott, J. F.; English, C. A. Int. J. Press. Vessels Piping 1999, 76, 891–908. Eason, E. D.; Wright, J. E.; Odette, G. R. Improved Embrittlement Correlations for Reactor Pressure Vessel Steels; NUREG/CR-6551, Nov 1998. ASTM E 900. Annual Book of ASTM Standards; American Society for Testing and Materials: West Conshohocken, PA, 2002; Vol. 12.02. Server, W.; English, C.; Naiman, D.; Rosinski, S. Charpy Embrittlement Correlations – Status of Combined Mechanistic and Statistical Bases for U.S. Pressure Vessel Steels (MRP-45); EPRI 1000705: Palo Alto, CA, 2001. Kirk, M.; Santos, C.; Eason, E.; Wright, J.; Odette, G. R. In Transactions of the 17th International Conference on Structural Mechanics in Reactor Technology (SMiRT 17), Prague, Czech Republic, Aug 17–22, 2003; International Association for Structural Mechanics in Reactor Technology: Raleigh, NC, 2003. Eason, E.; Odette, G. R.; Nanstad, R. K.; Yamamoto, T. A Physically Based Correlation of Irradiation-Induced Transition Temperature Shifts for RPV Steels; Oak Ridge Report, ORNL Report ORNL/TM-2006/530, Nov 2007. EricksonKirk, M. T.; Dickson, T. L. Recommended Screening Limits for Pressurized Thermal Shock (PTS); US Nuclear Regulatory Commission Report NUREG-1874; Office of Nuclear Regulatory Research, US Nuclear Regulatory Commission: Washington, DC, 2007. Carter, R. G.; Soneda, N.; Dohi, K.; et al. In Effect of dose rate on irradiation damage in high Cu weld, In Proceedings of the Workshop on Dose Rate Effects in Reactor Pressure Vessel Materials, held Squaw Creek, CA, USA, Nov 12–14, 2001; EPRI: Palo Alto, CA. Hiranumu, N.; Soneda, N.; Dohi, K.; Ishino, S.; Dohi, N.; Ohata, H. Mechanistic Modeling of Transition Temperature Shift of Japanese RPV Materials, presented at the 30th MPA-Seminar in conjunction with the 9th German-Japanese Seminar, Stuttgart, Germany, Oct, 2004. Soneda, N.; Dohi, K.; Nomoto, A.; Nishida, K.; Ishino, S. In Proceedings of the International Symposium on Research for Ageing Management of Light Water Reactors, held Fukui City, Japan October 2007; Eyre, B. L., Kimura, I., Eds.; INSS: Japan, 2008; pp 355–370. Carter, R.; Server, W.; English, C. Predictions of neutron irradiation embrittlement in BWR vessel steels; Proceedings of Fontevraud 6; French Nuclear Energy Society: France, 2006. Williams, T. J.; Ellis, D.; English, C. A.; Hyde, J. M. Int. J. Press. Vessels Piping 2002, 79, 649–660. Liu, C. L.; Odette, G. R.; Wirth, B. D.; Lucas, G. E. Mater. Sci. Eng. A 1997, 238, 202–209. Auger, P.; Pareige, P.; Welzel, S.; Van Duysen, J. C. J. Nucl. Mater. 2000, 280, 331–344.
4.06
Radiation Effects in Refractory Metals and Alloys
K. J. Leonard Oak Ridge National Laboratory, Oak Ridge, TN, USA
Published by Elsevier Ltd.
4.06.1
Introduction
181
4.06.2 4.06.2.1 4.06.2.2 4.06.2.3 4.06.3 4.06.3.1 4.06.3.2 4.06.3.3 4.06.4 4.06.4.1 4.06.4.2
Niobium and Nb-Base Alloys Introduction and History of Nb and Nb Alloys Radiation-Induced Swelling of Nb and Nb-Base Alloys Mechanical Properties of Irradiated Nb and Nb Alloys Tantalum and Ta-Base Alloys Introduction and History of Ta and Ta Alloys Irradiation-Induced Swelling of Ta and Ta-Base Alloys Mechanical Properties of Irradiated Ta and Ta-Base Alloys Molybdenum and Mo-Base Alloys Introduction and History of Mo and Mo Alloys Irradiation-Induced Swelling and Physical Property Changes in Mo and Mo-Base Alloys Mechanical Properties of Irradiated Mo and Mo Alloys Tungsten and W-Base Alloys Introduction and Irradiated Properties Database for W and W Alloys Irradiation-Induced Swelling and Physical Property Changes in W and W Alloys Irradiated Mechanical Properties of W and W Alloys Outlook
182 182 183 185 188 188 189 190 194 194
4.06.4.3 4.06.5 4.06.5.1 4.06.5.2 4.06.5.3 4.06.6 References
Abbreviations
Symbols
bcc C-103 Cb-752 DBTT
D n T Tirr Tm a DV/V w
Body-centered cubic Nb–10Hf–1Ti alloy Nb–10W–2.5Zr alloy Ductile–brittle transition temperature FS-85 Nb–10W–28Ta–1Zr alloy HP High purity JIMO Jupiter icy moons orbiter LCAC Low-carbon arc cast NERVA Nuclear experiment for rocket vehicle applications ODS Oxide dispersion strengthened RIS Radiation-induced segregation SNAP Systems nuclear auxiliary power T-111 Ta–8W–2Hf alloy TZM Mo–0.5Ti–0.1Zr alloy UTS Ultimate tensile strengths UWMAK-III University of Wisconsin Madison fusion reactor concept
194 197 206 206 206 207 209 211
Thermoelectric power Neutron particle Temperature Irradiation temperature Melting temperature Alpha particle Volume fraction swelling Fluence
4.06.1 Introduction Refractory metals and alloys offer attractive and promising high-temperature properties, including high-temperature strength, good thermal conductivity, and compatibility with most liquid metal coolants, many of which are suitable for applications in nuclear environments. Though many of the refractory alloys have been known for more than 60 years, there are significant gaps in the materials property database for both unirradiated and irradiated 181
182
Radiation Effects in Refractory Metals and Alloys
conditions. In addition, significant issues related to low-temperature irradiated mechanical property degradation at even low neutron fluences restrict the use of refractory metals. Protection from oxidizing environments also restricts their use, unless suitable protection or a liquid metal coolants is used. Much of the early research on refractory metal alloys was centered on applications in aerospace as well as cladding and structural materials for fission reactors, with particular emphasis on space reactor applications. Reviews concerning the history of these programs and the development of many of the alloys whose irradiated properties are discussed in this chapter can be found elsewhere.1–5 Due to cancellations and reintroduction of new mission criteria for these space reactor programs, the materials database shows similar waves in the gains of intellectual knowledge regarding refractory alloy and irradiated property behavior. Unfortunately, as seen in the subsequent sections of this chapter, much of the irradiated property database for refractory metals consists of scoping examinations that show little overlap in either material type, metallurgical conditions (i.e., grain size, impurity concentrations, thermomechanical treatments), radiation conditions (i.e., spectra, dose and temperature), or postirradiation test conditions or methods. The irradiation behavior of body-centered cubic (bcc) materials is known. Irradiation-induced swelling because of void formation in the material lattice is typical for temperatures between 0.3 and 0.6 Tm, where Tm is the melting temperature. Maximum swelling in refractory metals is <10% for displacement damage levels up to 50 dpa (displacements per atom), but typical values for fission reactor applications are <4%. Alloy additions can further reduce the sensitivity to swelling, for example, rhenium additions to molybdenum or tungsten. These levels of swelling are manageable through the appropriate engineering design of components. The generation of dislocation loops and point defects provide significant irradiation-induced strengthening or hardening of refractory metals and alloys. This in turn creates reductions in the ductility and fracture toughness of the material. This is most pronounced at temperatures <0.3 Tm, where defect mobility is reduced. The increase in the yield strength of the material because of the irradiationinduced defects can exceed the fracture strength of the material, leading to brittle behavior. These degradations in material property can begin to occur at neutron fluences as low as 1 1020 n cm2,
or 0.03 dpa3 and increase in severity with dose. As irradiation temperatures increase, dislocation loop and void sizes increase, whereas their densities are reduced, providing improvements in ductility, though at a reduced strength of the material. At high enough temperatures, recovery of properties to levels close to that of the unirradiated values is possible, though changes in material properties may be further influenced by microstructural changes such as segregation or precipitate formation of solute and transmuted species or recrystallization, which can lead to further deterioration of properties. Detailed information on the effects of radiation on materials is presented in Chapter 1.03, Radiation-Induced Effects on Microstructure, and in Chapter 1.04, Effect of Radiation on Strength and Ductility of Metals and Alloys. In general, the use of refractory alloys in radiation environments is not recommended at temperatures <0.3 Tm. However, new research work, particularly on molybdenum and its alloys, has shown that control over interstitial element contamination levels, grain size, and morphology, as well as the introduction of oxide dispersion strengthening, can lead to improvements in the lowtemperature irradiation behavior. This is discussed in detail in this chapter. The following sections of this chapter deal with the irradiated properties database of niobium, tantalum, molybdenum, and tungsten, as well as their alloys. While vanadium may sometimes be considered a refractory metal, its melting temperature is considerably lower than that of the other materials mentioned. However, its radiation effects database is considerable and well advanced relative to some refractory metals and it is therefore discussed separately in Chapter 4.12, Vanadium for Nuclear Systems. The irradiated properties database for refractory alloys is particularly thin, especially involving fracture toughness properties, irradiation creep effects, and combined radiation effects with high thermomechanical loads such as those experienced in plasma facing components or spallation target materials. Where needed, a comparison of the unirradiated and irradiated properties of a material is given.
4.06.2 Niobium and Nb-Base Alloys 4.06.2.1 Introduction and History of Nb and Nb Alloys The push for higher operating temperatures in turbine engines, as well as in reactor designs for both terrestrial and space applications, has frequently
Radiation Effects in Refractory Metals and Alloys
governed the periodic scientific examinations of refractory alloys. The historical examination of Nb and its alloys is typical of this, with early studies of the irradiation properties exploring the potential uses of these alloys in fusion energy and fission type space reactor. While these alloys have favorable properties, such as elevated temperature capability and compatibility with liquid alkali metals for energy applications, and attractive physical properties such as thermal conductivity, much of the work on Nb and Nb-base alloys has examined the nonirradiated properties. It is worth putting into perspective the relatively small commercial market for niobium-base alloys. Approximately, 75% of all niobium metal is used as minor alloying additions in steel, and only 1–2% is produced in the form of niobium-base alloys. The total market for niobium-base alloys in the mid-1990s was <105 kg year1 (100 metric tonnes year1).6 By comparison, 60 tons of niobium was used in 1961 for the SNAP-50 reactor program alone, and substantial additional quantities were used for other research projects such as the NERVA (Nuclear Experiment for Rocket Vehicle Applications) program.7 Throughout the history of the various space reactor programs, dozens of alloys were examined, with several brought to near-commercial production. However, today only the Nb–1Zr and C-103 (Nb–10Hf–1Ti) alloys remain commercially available for use in the sodium vapor lamp and rocket or turbine engine exhaust nozzles. Nb–1Zr has historically been considered the only niobium-base alloy with a sufficiently mature database (mechanical properties including thermal creep, chemical compatibility, fabrication, and welding knowledge) to be considered a near-term candidate for radiation environments.7–11 Developed for high ductility and good weld characteristics, the alloy shows less-than-desirable thermal creep strength at elevated temperatures compared to other refractory alloys. Though the C-103 alloy has greater short-term elevated temperature strength than that of Nb–1Zr, its long-term properties show no improvement over Nb–1Zr.12 Nonetheless, Nb–1Zr is the only Nb-base alloy with a significant radiation effects database. Despite the periodic programmatic interest in the use of Nb and Nb-base alloys, no clear fundamental study of the irradiated properties for a specific application has been performed or completed. Much of the data available on the irradiated properties is scattered and easily spans a time frame of several decades, which can lead to misinterpretations of results on the basis of either the limited scientific knowledge
183
of the time, lack of understanding of the sensitivity of properties on impurity concentrations, or aging effects. Radiation effects data are limited to the examination of swelling and tensile properties, with no information regarding fracture toughness or irradiation creep performance. The following sections deal with radiation effects on the properties of Nb and Nb–1Zr specifically. While some initial scoping examinations have been performed on other Nb-base alloys, these are relatively inconsequential and based on the less-thandesirable ductility, thermal stability, or welding capabilities of these alloys. 4.06.2.2 Radiation-Induced Swelling of Nb and Nb-Base Alloys Like all group V transition metals, the affinity of Nb for, and its ability to dissolve, high concentrations of interstitial atoms such as hydrogen, oxygen, nitrogen, and to a lesser extent, carbon can strongly influence the properties of the metal through defect–impurity interactions. Hydrogen, carbon, and oxygen impurities have a strong effect on the tensile Ductile–brittle transition temperature (DBTT) of pure Nb (reviewed by Hahn et al.13), with hydrogen levels near 10 ppm increasing the DBTT to 173 K and over 273 K at levels >100 ppm (DBTT of high purity Nb and Nb alloys is near 73 K14,15). The effects of oxygen and carbon were less severe, but influential at levels of 100 ppm and greater. The effect of nitrogen on embrittlement also appears to be as severe as that of oxygen, though some uncertainty exists as to whether solubility limits have been exceeded in the data.1 The effect of interstitial impurities on the irradiated properties of Nb and Nb-base alloys is significant and has been examined, though the overall database for irradiated properties is limited. The interplay between the radiation-created defects and the interstitial impurity elements was investigated by Igata et al.16 for pure Nb (70 wppm oxygen and 30 wppm nitrogen) irradiated to 3.4 1020 n cm2 (E > 1 MeV) at temperatures below 413 K and postirradiation annealed up to 973 K. Increases in yield strength over the as-irradiated values following annealing were measured at 473 and 673 K, attributed to the interplay of the defect clusters trapping oxygen and nitrogen atoms, respectively. Above 773 K, no difference between the annealed and as-irradiated yield stress was observed. Hautoja¨rvi et al.17 and Naidu et al.18 examined the interaction between vacancies and interstitial
Radiation Effects in Refractory Metals and Alloys
impurities in irradiated Nb through positron annihilation studies. In high-purity Nb, vacancy clustering within the collision cascades is observed, starting as low as 160 K, with vacancy migration peaking around 250 K, but in materials with higher hydrogen content, the vacancy migration stage shifts to temperatures close to 400 K.17 The irradiation exposure at intermediate temperatures (0.3–0.6 Tm) can lead to void swelling, irradiation creep, and helium embrittlement through processes involved in (n,a) reactions or impurity gas atoms. Naidu et al.18 examined the effect of He and its interaction with vacancies in pure Nb, leading to the development of bubbles through a-irradiated specimens. At temperatures between 623 and 1023 K, bubble growth occurs through the addition of He atoms and vacancies, followed by migration and coalescence at higher temperatures, eventually leading to the annealing out of the He bubbles and vacancy complexes above 1173 K.18 The irradiation-induced swelling of pure Nb generally appears at temperatures between 673 and 1323 K with peak swelling near 873 K (0.32 Tm), though these limits are not clearly defined and are based on the very limited data available, compiled by Wiffen19 and Pionke and Davis.1 A maximum swelling of 4.8% following irradiation to 2.5 1022 n cm2 at 858 K was reported.1 However, the magnitude of swelling shows considerable scatter in the literature, possibly reflecting the influence of impurity concentrations and differences in irradiation conditions and microstructural interpretation of the materials.19 Fischer20 reported that void concentration increased four to seven times for a fourfold increase in flux for the same total fluence. This produced a reduction in void size with flux and therefore a reduction in the total swelling. Loomis and Gerber21–23 examined the influence of oxygen and substitutional binary alloy additions on the swelling of 3 MeV 58Ni+ ion-irradiated Nb up to 50 dpa. Void formation and characteristics in size and morphology were found to be dependent on temperature, oxygen concentration, and the type of substitutional alloy addition. The average void diameter was found to increase with temperature as well as oxygen up to 0.02 at.%. Higher oxygen concentrations resulted in a decrease in void diameter to 0.1 at.% O, above which void diameters showed no significant changes. The number density of voids was found to decrease with temperature, but increase with oxygen concentration to 0.06 at.%, above which the number density showed no significant change. As the volume fraction of swelling (DV/V) is proportional to both the void number and the cube of the
void diameter, the volume fraction is observed to increase with temperature and oxygen concentration to 0.04 at.%, followed by a decrease and plateau of the volume fraction above 0.1 at.%. The dependence of DV/V on temperature and oxygen concentration is illustrated in Figure 1. Microstructural examination revealed an ordering of the voids into a latticetype structure in the material irradiated at 1050 K to 40 dpa and oxygen concentration 0.039 at.% oxygen. The higher temperature of the maximum swelling as compared to the neutron irradiation data is believed to be associated with the higher displacement damage rate of the ion-bombarded material,19 though the higher impurity levels may also provide an influence. The effect of dilute (2.4 at.%) substitutional alloy addition on the swelling of 0.06 at.% oxygendoped Nb was also examined for 3 MeV 58Ni+ ion irradiation at 1225 K. The DV/V was determined to increase through the addition of Ta, but decreased with increasing effectiveness by the addition of Ti, Zr, V, and Hf. The addition of the reactive alloying elements to Nb suppresses void formation through the gettering of interstitial impurities that act as void nucleation sites. The DV/V was determined to be unaffected by the addition of Ni or Fe. The dependence of DV/V on temperature, oxygen, and substitutional addition is also shown in Figure 1.
14
30 MeV 58Ni+ irradiation Tirr =1225 K, 50 dpa
12 10 DV/V (%)
184
8 1240 K
6
1225 K
4
1240 1200 K
2 0 0.00 0.04
0.08 0.12 0.16 0.20 Oxygen concentration (at.%)
Nb + 2.4 at.% Ta Nb + 2.4 at.% Ni Nb + 2.3 at.% Fe Nb + 2.3 at.% Ti
Tirr (K)
1220 1200
Nb + 2.4 at.% Zr Nb + 2.4 at.% V Nb + 2.4 at.% Hf Nb + 2.4 at.% Mo
Figure 1 The dependence of void volume fraction (DV/V ) in 3 MeV 58Ni+ ion-irradiated Nb on the concentration of oxygen and dilute solute additions. Reproduced from Loomis, B. A.; Gerber, S. B. J. Nucl. Mater. 1983, 17, 224–233.
Radiation Effects in Refractory Metals and Alloys
Swelling in Nb–1Zr has been examined, though only scattered data are available in the examination of temperature and flux dependence. The available swelling data on Nb–1Zr, compiled by Powell et al.24 and Watanabe et al.25 presented in Figure 2, show the lack of data on the temperature range in which peak swelling appears. The swelling data shown in the figure were measured through electron microscopy, with the exception of the data by Powell et al.24 and Wiffen.26 Alloy impurity chemistry, in addition to interpretation and measurement error, may account for the scatter associated with the lower temperatures. The work of Watanabe et al.25 and Garner et al.27 indicates that irradiation-induced swelling is dependent on the thermomechanical history of the material. In that material, cold-working followed by solution anneal and aging exhibited swelling, while material not given the preirradiated cold-working showed some densification. The changes in density of the material are dependent on the phase-related transformations involving precipitation. Swelling in Nb–1Zr appears to be centered over a more narrow temperature range than in Nb, with a peak near 1073 K that is higher than that of the pure metal. While the addition of Zr to Nb appears to delay nucleation of voids to higher temperatures, the voids that form are of larger size than those appearing in pure Nb under comparable conditions. For example,
2.5
DV/V (%)
2.0 1.5 1.0 0.5 Irr 400 ad 600 iat ion 800 te 1000 m pe 1200 ra tu 1400 re (K )
20 0
Powell et al.24 Jang and Moteff145 Sprague et al.146 Wiffen28
40 dpa
60
80
100
Wiffen26 Michel and Smith147 Watanabe et al.25 Garner et al.27
Figure 2 Swelling as a function of irradiation temperature and dose for neutron-irradiated Nb–1Zr from available literature compiled by Powell et al.24 and Watanabe et al.25
185
following irradiation to 2.5 1022 n cm2 (E > 0.1 MeV) at 1063 K, the diameter, concentration, and volume fraction of voids in Nb–1Zr was 57.5 nm, 1.8 1020 m3, and 2.2%, respectively,1 whereas under similar conditions, the same void parameters in pure Nb were 18.6 nm, 2.8 1021 m3, and 1.04%. While void formation and swelling in Nb and Nb–1Zr occurs, the total swelling is generally <5% and within engineering limits, even for high neutron exposures >10 dpa.3 The addition of Ti to Nb was found to increase void resistance and has been found to suppress void formation in V at concentrations as low as 3%.29 The combination of reactive alloy elements and Nb in the C-103 alloy may suggest a greater void formation resistance than in pure Nb and Nb–1Zr. 4.06.2.3 Mechanical Properties of Irradiated Nb and Nb Alloys Little coverage of the changes in mechanical properties following irradiation has been given to Nb and Nb alloys, with the majority of the data for temperatures below 800 K. Some preliminary experimental work on the irradiated mechanical properties of Nb alloys Cb-752 (Nb–10W–2.5Zr)30 and FS-85 (Nb– 10W–28Ta–1Zr)31 is available. However, these alloys are not commercially produced and have shown indications of thermal aging instabilities, leading to grain boundary embrittlement.12,32,33 The irradiated mechanical properties of these alloys show similar radiation hardening as in the pure metal, but with mechanical properties more sensitive to thermal aging conditions. The bulk of the irradiated mechanical properties data is for the Nb–1Zr alloy as well as the pure metal, and is covered in this review. The irradiated mechanical properties of Nb and Nb–1Zr are strongly governed by irradiation temperature, determining whether the mechanical properties are controlled by dislocation loops or a combination of loops and voids in the microstructure. As cavity formation can be delayed or suppressed by higher irradiation temperatures in Nb–1Zr, mechanical property comparisons between the alloy and the base metal will reflect their irradiated microstructure. For Nb and Nb–1Zr irradiated to 3 1022 n cm2 at 728 K, the pure metal contains both dislocation loops and voids, while the alloy exhibits no void formation.19 A comparison of the tensile properties of Nb and Nb–1Zr irradiated under similar conditions is shown in Figure 3. The irradiated strength of both materials shows an increase in tensile strength above the unirradiated
Radiation Effects in Refractory Metals and Alloys
120
Niobium
Nb–1Zr
Stress (1000 psi)
100 80 60 40 20
Elongation (%)
0 60
Control 3.7 ⫻ 1022 n cm–2 At 450 ⬚C
Control 3.0 ⫻ 1022 n cm–2 At 460 ⬚C Ultimate
800
600
Yield and ultimate
Yield
400 Ultimate
Ultimate
Yield
Yield
Stress (MPa)
186
200
0 Total
Total
40 20 0 0
Uniform
Uniform Total
Total Uniform 200
400
600 0 200 Test temperature (⬚C)
400
600
Figure 3 Comparison of tensile properties between Nb and Nb–1Zr tested under similar irradiation conditions. Reproduced from Wiffen, F. W. In Refractory Alloy Technology for Space Nuclear Power Applications, CONF-8308130; Cooper, R. H., Jr, Hoffman, E. E., Eds.; Oak Ridge National Laboratory: Oak Ridge, TN, 1984; pp 252–277.
800
Niobium Ttest = 298 K 3.0 ⫻ 1026 n cm–2 Tirr = 733 K
600 Stress (MPa)
condition, with Nb–1Zr showing a greater sensitivity to irradiation. As the mechanical properties of Nb–1Zr are dominated by the dislocation loop structures, yield instability is observed in the material, leading to the early onset of necking. This results in <0.2% uniform elongation, though total elongation near 10% is still achieved. The yield instability is associated with dislocation channeling, in which deformation dislocations will create defect-free channels along their slip plane, following the annihilation of the loop structures. This occurs only after enough applied stress is achieved to overcome the obstacles, but the net effect is a nonuniform plastic deformation through channels that allow for the movement of deformation dislocations at reduced stress. The irradiated Nb samples whose properties are shown in Figure 3 contain, in addition to dislocation loops, voids that limit dislocation channeling by providing added obstacles to deformation, resulting in some measure of uniform elongation and work hardening upon yielding. The microstructure dependence on the tensile properties can best be illustrated by the comparison shown in Figure 4 of Nb irradiated at 328 and 733 K. The higher irradiation temperature results in the development of microstructural voids and thus the significant differences in the tensile curves. The lower irradiation temperature results in dislocation channeling following yield and the
400
7.5 ⫻ 1024 n cm–2 Tirr = 328 K
200
0 0
4
12 8 Elongation (%)
16
Figure 4 Comparison of tensile curves between Nb irradiated at 328 and 733 K. Yield instability is seen at 328 K due to channeling of deformation dislocations through the irradiated dislocation loop structures. The higher irradiation temperature resulted in the development of small voids providing a barrier to dislocation movement. Reproduced from Wiffen, F. W. In Refractory Alloy Technology for Space Nuclear Power Applications, CONF-8308130; Cooper, R. H., Jr, Hoffman, E. E., Eds.; Oak Ridge National Laboratory: Oak Ridge, TN, 1984; pp 252–277.
associated work softening during necking to failure at around 11% total elongation. While the higher irradiation temperature sample was irradiated to a higher total fluence, the effect of dose is observed only on the
Radiation Effects in Refractory Metals and Alloys
relative strength increase over the unirradiated condition. The higher irradiation temperature produced voids in the microstructure, providing additional obstacles to deformation and higher uniform elongations and modest work hardening. Little is known with regard to the aging properties of Nb–1Zr or the combined thermal and radiation effects. The addition of 1 wt% Zr to Nb creates a dispersion-strengthened alloy, in which the Zr combines with interstitial impurities creating fine precipitates throughout the material. The development of these fine precipitates on aging at 1098 K can increase the tensile strength between 50 and 100 MPa over the annealed condition and provide an effective strengthening greater than that observed through modest irradiation31 (Table 1). Irradiation of Nb–1Zr to 0.9 dpa at 1098 K showed a modest increase in yield and ultimate tensile strength to 135 and 192 MPa, respectively, over the annealed condition. This increase in tensile strength either through aging or irradiation results in a corresponding decrease in uniform elongation from 15% to 3.5% and total elongation from 25% to 15%. Aging at temperatures above 1098 K produces little effective hardening as the precipitates coarsen in the microstructure.33 Irradiation to 0.9 dpa at 1248 and 1398 K of Nb–1Zr showed only a modest increase in the yield strength over the aged and annealed specimens, though ultimate tensile strength and elongation were unchanged or less. Irradiation to 1.88 dpa at 1223 K resulted in weaker tensile properties Table 1 of Nb–1Zr
187
compared to the 0.9 dpa sample, believed to be due to further precipitate coarsening. The time under irradiation conditions for the 1.88 dpa sample was near 1100 h and produced similar tensile properties as that of the aged-only material. As discussed in the preceding paragraphs, the irradiated properties of Nb and Nb–1Zr are governed by their microstructure and are influenced by temperature, displacement damage rate, and neutron spectrum. The tensile properties of neutron-irradiated Nb–1Zr for damage levels between 0.1 and 5 dpa (Horak et al.34 and Wiffen35) summarized by Zinkle and Wiffen3 are shown in Figure 5. At temperatures below 800 K, a large increase in the tensile strength from irradiation is observed with the corresponding low uniform elongations. At higher temperatures, uniform elongation increases because of the presence of voids in the microstructure. However, the data plotted in Figure 5 show uniform elongations remaining low up to 1100 K, while radiation hardening is relatively moderate, suggesting that impurities are the source of the reduced elongation values. No irradiated fracture toughness data exist for Nb or Nb–1Zr, though comparisons can be made from the larger irradiated vanadium alloy database, in which fracture toughness embrittlement becomes a concern when tensile strength exceeds 600–700 MPa and therefore at temperatures below 400 K for Nb–1Zr.36 However, if a conservative value is assigned to the critical stress to induce cleavage fracture of 400 MPa (40% lower than that observed in vanadium alloys),
Tensile property comparison illustrating the effects of aging and irradiation on the mechanical properties
Test/aged/irradiated temperature (K)
Yield strength (MPa)
As-annealed condition 298 185.00 1073 102.33 1223 89.00 1373 83.00 1100 h aged 1073 185.67 1223 111.67 1373 82.50 Irradiated: 2.04 1021 n cm2 (E > 0.1 MeV), 0.93 dpa 1073 134.50 1223 149.50 1373 102.00 Irradiated: 4.13 1021 n cm2 (E > 0.1 MeV), 1.88 dpa 1223 104.50
Ultimate tensile strength (MPa)
Uniform elongation (%)
Total elongation (%)
281.00 191.00 215.00 157.00
18.60 16.03 15.05 10.50
34.60 23.90 23.70 44.80
248.67 156.67 130.50
8.07 8.07 6.80
13.27 13.27 28.65
192.50 182.00 127.50
7.10 3.55 9.55
18.55 17.80 33.65
166.00
8.10
20.90
Source: Busby, J. T.; Leonard, K. J.; Zinkle, S. J. Effects of neutron irradiation on refractory metal alloys, ORNL/LTR/NR-PROM1/05-38; Oak Ridge National Laboratory: Oak Ridge, TN, Dec 2005.
Radiation Effects in Refractory Metals and Alloys
Ultimate tensile strength (MPa)
700
35 Irrad. UTS
600
30 Irrad. eU 25
500
20
400 300
Unirrad. UTS
15
200
10
100
5
0 200
400
600
800 1000 1200 Temperature (K)
1400
Uniform elongation (%)
188
0 1600
Figure 5 Unirradiated and irradiated (0.1–5 dpa) ultimate tensile strengths (UTS) and uniform elongation (eU) of Nb–1Zr. Irradiated data represented by solid symbols and unirradiated by open symbols. Figure reprinted with permission from Zinkle, S. J.; Wiffen, F. W. Radiation effects in refractory alloys. In STAIF 2004, AIP Conference Proceedings; El-Genk, M. S., Ed.; 2004; Vol. 699, pp 733–740. Copyright 2004, American Institute of Physics.
fracture toughness becomes a concern at temperatures below 800 K for Nb–1Zr.3 While irradiated tensile strength above 800 K is close to the unirradiated values, uniform elongation values remain low until irradiation temperatures >1000 K. Therefore, a conservative approach towards engineering design needs to be taken with this alloy. The mechanical properties of irradiated refractory alloys can be influenced by the formation of He developed through the (n,a) reactions, leading to the grain boundary formation of bubbles and the eventual embrittlement of the material. Some scoping investigations on the effect of He on the irradiated mechanical properties of Nb–1Zr have been performed. Wiffen37 investigated the high-temperature mechanical properties of 50 MeV a-irradiated Nb–1Zr. In tensile tests conducted at 1273 and 1473 K, no significant effect of He on the strength or ductility of Nb–1Zr was observed for samples containing 2–20 appm He. Later analysis of the creep ductility reductions was found to be dependent on the observed precipitate phase development through the pick-up of oxygen during implantation.38 He-implanted Nb–1Zr through 100 MeV a-irradiations at 323 and 873 K by Sauges and Auer39 found no significant effect on ductility up to 80 appm He. Wiffen19 observed that uniform elongations stayed around 1% between test temperatures of 723 and 1073 K on 130 appm 10B doped Nb–1Zr irradiated in a fast reactor between 723 and 1223 K up to 6 1022 n cm2. These were slightly higher than those observed in undoped material; this is believed to be due to the formation of He bubbles in the grains of the material
acting similar to voids in generating obstacles to dislocation channeling. In general, no detrimental effects on mechanical properties were reported for accelerator-injected He between 1273 and 1473 K for He concentrations <200 appm.37,40
4.06.3 Tantalum and Ta-Base Alloys 4.06.3.1 Introduction and History of Ta and Ta Alloys Tantalum and its alloys have historically been examined for high-temperature nuclear applications, particularly in the various space reactor programs. For reasons similar to those of Nb and its alloys, various alloying combinations of Ta were examined, particularly in the late 1950s to 1960s. Much of this effort emphasized the development of solid solution (W and Re additions) and dispersion-strengthened (Hf addition) alloys. While Ta-alloys pay a penalty in higher density over, for example, Nb, and decreases the low temperature density-compensated strength to comparable values on Nb-base alloys. The higher melting temperature of Ta (3290 K) results in better strength retention above 1000 K and in density-compensated creep strength.12,41 Early work on substitutional solid solutionstrengthened Ta–10W for aerospace applications42 led to limited examination of this alloy for irradiation environments. The improved strengthening by addition of a maximum of 10 wt% allows the retention of suitable nonirradiated ductility and weldability.43,44
Radiation Effects in Refractory Metals and Alloys
However, the use of Ta–10W in space reactor applications where liquid alkali coolants are considered was unacceptable because of the lack of oxide gettering elements such as Hf that form stable dispersionstrengthened structures. The T-111 (Ta–8%W–2% Hf) alloy, with its demonstrated compatibility with liquid alkali metals and improved strength over pure Ta while retaining ductility and weldability, has been a lead candidate alloy in space reactor systems since the 1960s.45 Though a considerable effort has been made on the Ta–10W and T-111 alloys, the irradiation properties database is very small. Irradiated mechanical property behavior follows typical bcc alloys in which radiation hardening effects including limit ductility appear and are expected at temperatures 0.3 Tm (987 K).3 4.06.3.2 Irradiation-Induced Swelling of Ta and Ta-Base Alloys Swelling data for Ta and its alloys are limited to a few studies.19 Void formation in pure Ta was experimentally observed through TEM examination of material irradiated to 2.5 1022 n cm2 (E > 0.1 MeV) at temperatures between 673 and 1273 K.46 An empirical estimation of the bulk swelling taken from microstructural void size density data of that study is shown in Figure 6. Void concentrations in the material were highest at the peak swelling temperature and decreased with higher irradiation temperature with an associated increase in cavity size. Ordering of the voids at the peak 3.0 Bates and Pitner47 Wiffen46 2.5
ΔV/V (%)
2.0
Neutron fluence 2.5 ⫻ 1022 n cm–2 (E > 0.1 MeV)
1.5
1.0
0.5
0.0 200
400
600
800
1000 1200 1400 1600
Irradiation temperature (K) Figure 6 Swelling data for pure Ta measured through microstructural void density measurements by Wiffen46 and from immersion density measurements by Bates and Pitner.47
189
swelling condition was reported to occur along the {110} planes in the bcc structure. A subject of considerable theoretical debate, the mechanisms of void ordering that have appeared in bcc and fcc metals have been examined,48–50 since the first reported occurrence in irradiated Mo.51 Disordered void structures in the microstructure of the higher temperature irradiated Ta appear as the size of the voids increase, though some rafting, or grouping, was reported.46 The swelling data of Wiffen46 derived from microstructural analysis correlate well with the immersion density data of Bates and Pitner47 (Figure 6), from which an empirical equation for percent swelling as a function of temperature, T (K), and fluence, F (in units of 1022 n cm2, E > 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 <0.3% uniform but with total elongation near 10%. The observed plastic instability, attributed to the lack of uniform elongation following yielding, resulted from dislocation channeling. Some recovery of ductility is observed following irradiation to 913 K, which correlates with temperatures approximating the maximum swelling temperature (Figure 6) and a change in the dominating microstructural features influencing deformation behavior in the metal. The tensile data are presented in Figure 8, along with the irradiated properties of T-111, which are discussed later. The recent work of Byun and Maloy56 investigated tensile behavior as a function of fluence for pure Ta, Ta–1W, and Ta–10W, establishing deformation mode maps for pure Ta and Ta–1W that outline the conditions in which brittle failure and uniform and unstable plastic deformation occur. Following fast-reactor exposures at temperatures <373 K, a progressive hardening and gradual loss in ductility are observed in the tensile properties of pure Ta, leading to a near doubling of the yield stress by 0.14 dpa over the unirradiated value (Figure 9(a)).
4.06.3.3 Mechanical Properties of Irradiated Ta and Ta-Base Alloys The overall mechanical property data for irradiated Ta and Ta-base alloys are very limited, with most studies involving irradiation at temperatures <1073 K. In general, the behavior of Ta and its alloys is similar to that of other bcc materials in that radiation hardening is observed with significant reductions in elongation at temperatures <0.3 Tm (Tm ¼ 3290 K, pure Ta). As is discussed in this section, the addition of solute strengthening elements creates an increased sensitivity to radiation hardening of the material. In addition to the lack of high-temperature irradiation behavior, impact and fracture toughness data for irradiated Ta and Ta alloys are also limited. As with all refractory metals, the mechanical behavior of pure Ta is highly dependent on the impurity levels in the material. This may explain the observed differences between the work of Brown et al.54 and Chen et al.,55 of 800 MeV proton irradiations up to 11 dpa at temperatures <673 K (Figure 7). While chemical analysis quantifying the purity of Ta was not reported in the former, irradiation to 0.26 dpa resulted in a yield strength increase from 350 to 525 MPa over the unirradiated value with a corresponding drop in ductility below 2%. Flow instability following yield was characteristic of samples irradiated to 0.26 and 2.9 dpa.54 Tensile properties of high-purity Ta irradiated to 0.6–11 dpa tested at room temperature and 523 K showed similar increases in tensile strength, while the uniform elongation remained near 8% following irradiation to 0.6 dpa or higher.55 The tensile properties of neutron-irradiated Ta were reported by Claudson and Pessl,30 Wiffen,19
Stress (MPa)
800 600
(a)
Tirr 658-673 K Ttest = 298 K
2.9 dpa
(b)
0.26 dpa
200
0
Control
Ta
5
High-purity Ta (99.99%)
0.6 dpa
Control
400
Tirr < 473 K Ttest = 298 K
11 dpa
10
15
20
0 Strain (%)
10
20
30
Figure 7 Stress–strain curves of pure Ta irradiated by 800 MeV protons (a) for lower purity Ta, and (b) higher purity Ta. (a) Reproduced from Brown, R. D.; Wechsler, M. S.; Tschalar, C. In Influence of Radiation on Material Properties: 13th International Symposium, ASTM STP 956; Garner, F. A., Henager, C. H., Jr, Igata, N., Eds.; ASTM: Philadelphia, PA, 1987; pp 131–140 and (b) Reproduced from Chen, J.; Ullmaier, H.; Floßdorf, T.; et al. J. Nucl. Mater. 2001, 298, 248–254.
Radiation Effects in Refractory Metals and Alloys
240
Test temperature (K) 900
700
Stress (psi)
Tantalum Control 200 1.97 dpa at 663 K 2.5 dpa at 913 K 160
500
700
900
T-111
1600 Control 2.5 dpa at 688 K 2.5 dpa at 913 K
Ultimate Yield
1200 Ultimate Yield
120
800
Yield and ultimate Ultimate
80
Stress (MPa)
500
191
Ultimate Yield
Ultimate
40
Yield
400
Yield
Elongation (%)
0 60
0 Total
40
Total
Uniform
Uniform
20
Total
Total
Total
Uniform
Total
Uniform
0 0
200
400
0 200 600 Test temperature (⬚C)
400
600
Figure 8 Comparison of tensile properties of neutron-irradiated Ta and T-111. Uniform elongations of <0.3% for the 663 K irradiations are not shown in the figure. Reproduced from Wiffen, F. W. In Refractory Alloy Technology for Space Nuclear Power Applications, CONF-8308130; Cooper, R. H., Jr, Hoffman, E. E., Eds.; Oak Ridge National Laboratory: Oak Ridge, TN, 1984; pp 252–277.
An early onset of necking or plastic instability was observed in Ta at doses above 0.0004 dpa. The lower elongation strains in the pure Ta compared with the work by Chen et al.55 is believed to be due to the higher oxygen content in the material.56 The introduction of 1 wt% W resulted in a near30% increase in unirradiated strength over pure Ta (Figure 9(b)). The Ta–1W alloy showed greater sensitivity to radiation hardening than the pure metal. The tensile properties as a function of dose were similar to those of the pure Ta. However, above 0.004 dpa, plastic instability becomes more predominant in the Ta–1W alloy and occurs immediately following yielding. For Ta–1W irradiated from 0.7 to 7.5 dpa in a mixed proton and neutron irradiation from the same study, hardening was saturated with little change in ductility (insert shown in Figure 9(b)). Macroscopic deformation mode maps produced for Ta and Ta–1W by Byun and Maloy56 are a graphical way of predicting the performance of a material in an irradiation environment. The deformation mode map for pure Ta is shown in Figure 10(a), while that
of Ta–1W is shown in Figure 10(b). The yield and plastic instability stress were directly obtained from tensile data, while the fracture stress was calculated through a linear strain hardening model for necking deformation, assuming that during instable deformation, the strain hardening rate remains unchanged and is approximately the plastic instability stress. The fracture and plastic instability stresses are independent of dose, with a ratio between the stresses of 2 for the materials studied. The fracture strength decreases with dose if the material becomes embrittled, for example, through interstitial segregation or secondary phase precipitation at grain boundaries, though this was not observed in their work. The yield strength is highly dose dependent, though the yield stress was significantly lower than the fracture strength in Ta–1W, suggesting that the material may show limited ductility to even higher displacement doses. The effect of increasing test temperature for each material further increases the boundaries for uniform deformation behavior. This increase was found to be greater in pure Ta.
Radiation Effects in Refractory Metals and Alloys
Engineering stress (MPa)
1000
Ta Neutron irradiated Tirr 333–373 K
800
600
Engineering stress (MPa)
192
0.14 dpa
0.04
400
0.004 0.0004
200
Unirr.
0.00004
1000
Ta–1W Neutron irradiated Tirr 333–373 K
800
Proton and neutron irradiated Tirr 323–433 K
800
7.5 dpa 4.6
600
0.14 dpa
400
600
0.04
0
400
200
2.0
0.004
Unirr. 0.7
200
0
10
20
30
0.00004
40
0.0004 Unirr.
0 (a)
0
10
20 30 Elongation (%) 1600
40
10
0
20 30 Elongation (%)
(b)
25.2 dpa
40
50
Ta–10W Proton and neutron irradiated Tirr 323–433 K
1400 Engineering stress (MPa)
0
1200 7.5 1000 800
4.4
Unirr.
0.7 2.0
600 400 200 0
(c)
0
5
15 10 Elongation (%)
20
Figure 9 Room temperature tensile curves for irradiated (a) pure Ta, (b) Ta–1W, and (c) Ta–10W. Reproduced from Byun, T. S.; Maloy, S. A. J. Nucl. Mater. 2008, 377, 72–79.
The room temperature unirradiated tensile strength of Ta–10W is nearly double the value of the Ta–1W and triple that of pure Ta in the material investigated by Byun and Maloy,56 and also shows an increased sensitivity in radiation hardening over the pure metal (Figure 9(c)). This sensitivity is also clearly apparent at higher irradiation temperatures near 673 K, as shown in the comparison of tensile curves that were compiled by Ullmaier and Carsughi58 of earlier work (Figure 11). Near room temperature irradiation of Ta–10W to the mixed proton and spallation neutron exposure by Byun and Maloy56 to doses between 2 and 25.2 dpa showed prompt necking following yielding. Total elongation values of <3% were observed for doses between 2 and 7.5 dpa, with near-zero ductility observed at 25.2 dpa. Fast neutron irradiation studies of Ta–10W by Gorynin et al.59 observed brittle failure after 0.13 dpa in materials irradiated and tested near 600 K. Less than 5% total elongation was measured following 1.97 dpa irradiation at 700 K, despite a near
doubling of the yield stress over the unirradiated material. Limited ductility was also observed following 2.63 dpa exposure in materials irradiated and tested at 1073 K, with a yield strength increase from 240 to 315 MPa over the unirradiated control. While lowtemperature embrittlement following exposure to 0.13 dpa was reported in the neutron-irradiated materials59 and limited ductility following mixed proton and neutron exposure,56 the interstitial concentrations on the behavior of these materials may be more influential than the irradiation spectrum. Similar to Ta and Ta–10W, very limited data exist on the irradiated properties of T-111. The most referenced base-line study is that by Wiffen,19 shown as in Figure 8. Large increases in yield and ultimate tensile strengths are observed following irradiations to 1.9 1022 n cm2 (E > 0.1 MeV), 2.5 dpa, at 688 and 913 K. The increase in radiation hardening is substantially greater than that observed in pure Ta irradiated under similar conditions. Yield and
193
Radiation Effects in Refractory Metals and Alloys
2000
1000
1500 True stress (MPa)
True stress (MPa)
Fracture region
Fracture region
800
600 Plastic instability region 400 Elastic region
200
1000
500
Uniform plasticity (523 K) Uniform plasticity (Trm)
0 0.0
0.0001
0.001 0.01 Dose (dpa)
(a)
Plastic instability region
Elastic region Uniform plasticity (523 K) Uniform plasticity (Trm)
0.1
0 0.0
1 (b)
0.0001 0.001 0.01 0.1 Dose (dpa)
1
10
Figure 10 Deformation mode map of (a) pure Ta and (b) Ta–1W for room temperature irradiations, illustrating fracture, plastic instability, uniform plasticity, and elastic regions as a function of stress and displacement dose. The increases in the uniform plasticity region for temperatures of 523 K are superimposed. Reproduced from Byun, T. S.; Maloy, S. A. J. Nucl. Mater. 2008, 377, 72–79.
400
Tirr~673 K
360 0.39 dpa
320 Stress (MPa)
280
Ta-10W Control
240 200
Ta Control
160
0.13 dpa
120 80 40 0
0
2
4 6 Strain (%)
8
Figure 11 Comparison of the radiation hardening of Ta and Ta–10W irradiated at 673 K to displacement doses of <0.39 dpa. Adapted from Ullmaier, H.; Casughi, F. Nucl. Instr. Methods Phys. Res. B 1995, 101, 406–421.
ultimate tensile strengths of around 1250 MPa are reported for irradiation at 688 K, with uniform and total elongation of <0.3% and 4.5%, respectively. Irradiation at 913 K improves uniform and total elongation values only slightly to 2.5% and 8%. These values represent more than a 50% reduction in ductility over the unirradiated values. No known irradiated property data for T-111 exist for temperatures above
913 K. As tensile strengths of both Ta–10W and T-111 exceed 1000 MPa at temperatures below 1000 K19,56,59 and are well above the stresses that produce brittle behavior in vanadium alloys for which more data are available, it is likely that these Ta alloys are embrittled under these conditions.3 Further expansion of the irradiated materials database including fracture toughness data for Ta and Ta alloys irradiated near and above 1000 K is much needed to ascertain the upper temperature limitations. However, based on this preliminary data, temperatures below1000 K may need to be avoided for Ta and Ta-base alloys. Low-fluence irradiations to 1.2 1015 n cm2 at room temperature and 623 K have been performed to evaluate the performance of T-111 and Ta–10W for use in radioisotope power applications.60 These low-dose irradiations produced little change in the tensile properties of the two alloys. Some variations in the total elongation were observed in T-111, which may be related to the distribution and make-up of the Hf-rich compounds in the material as well as the effects of radiation. Thermal stability of T-111 can be an issue, as a brittle behavior following 1100 h aging at 1398 K has been reported,41 due to precipitation of Hf-rich compounds along grain boundaries. It is not known how the combination of long-term thermal aging under irradiation affects the structure– property relationships or how the detrimental precipitation of the interstitial elements with Hf can be controlled.
194
Radiation Effects in Refractory Metals and Alloys
4.06.4 Molybdenum and Mo-Base Alloys 4.06.4.1 Introduction and History of Mo and Mo Alloys Molybdenum and its alloys are the perennial candidates for refractory metal alloy use in irradiation environments, due in part to their high melting temperature (2896 K), good thermal properties, high-temperature strength, and lower induced radioactivity (as compared to tantalum). The density of molybdenum (10.28 g cm3) is also significantly lower than that of Ta and W, though greater than Nb. But like other refractory metal alloys, Mo can present difficulties in fabrication, low-temperature ductility, and low-temperature embrittlement from radiation damage. The TZM (Mo–0.5%Ti–0.1% Zr) and Mo–Re alloys were examined as part of the SP-100 and JIMO/Prometheus space reactor programs, respectively, and offer additional benefits of improved high-temperature strength over the pure metal.5,19 Molybdenum and its alloys have also been examined for plasma facing and diverter components in fusion reactor designs due to the relatively low sputter yield, high thermal conductivity, and thermal compatibility with other structural materials.5–27,29–63 In addition, because of these benefits, Mo has also been examined for use as a grazing incident metal mirror in fusion diagnostic port designs.64,65 As in all other refractory metals, the mechanical properties are influenced by impurity concentrations, particularly through grain boundary weakening. However, improvements in Mo ductility are achievable through grain refinement, impurity control, and the addition of Re or reactive elements such as Ti and Zr. An upper limit to the acceptable level of C was also found to improve grain boundary strength. Low-carbon arc-cast molybdenum (LCAC-Mo) is one such example, in which oxygen impurities are reduced to tens of ppm, nitrogen to <10 ppm, and carbon to <100 ppm.66 Higher levels of C will result in reduced fracture toughness, unless additional reactive alloy additions are present in the alloy. The TZM alloy also incorporates a small level of carbon to produce Ti- and Zr-carbide strengthening. Improvements in ductility and toughness through the ‘rhenium effect’ have been observed in Mo for some time,67–69 and generally occurs when Group VIa metals are alloyed with elements from Group VIIa and VIIIa metals.70,71 Explanations for this phenomenon range from enhanced mechanical twinning, reduced resistance to dislocation glide, reduction of
oxygen at grain boundaries, and increased interstitial oxygen solubility.67,71–75 Critical evaluation76,77 of the initial work that had suggested a maximum tensile ductility near 11–13 wt% Re78,79 was found to be inconclusive because of inadequate control of O and C impurity levels in the earlier studies. Higher concentration alloys with 40–50 wt% Re have also been examined for use in the radiation environments. Alloys with Re concentrations up to 41–42% are single-phase solid-solution a-Mo, while those at higher levels incorporate the s-Re2Mo phase. Commercially available alloys include Mo–41Re and Mo–47.5Re (sometimes referred to as Mo–50Re). Recently, introduction of oxide dispersion strengthened (ODS)-Mo through the incorporation of lanthanum oxide particles has been examined.80–82 These alloys show great resistance to recrystallization and high-temperature deformation while maintaining low ductile-to-brittle transition temperatures (DBTT) partly because of their refined grain structure.83–85 The radiation effects database for Mo and its alloys is limited to scattered scoping examinations, which show little overlap in the experimental variables such as material purity, alloying level, material thermomechanical history, irradiation conditions, and postirradiation test conditions. Where available, information on the physical and mechanical property changes to LCAC-Mo, TZM, Mo–Re alloys, and ODS-Mo will be reviewed. 4.06.4.2 Irradiation-Induced Swelling and Physical Property Changes in Mo and Mo-Base Alloys Two earlier reviews of the irradiation-induced properties of Mo and TZM have been presented as part of the UWMAK-III fusion reactor study86 and for the SP-100 space nuclear power program.19 Much of the known swelling data on irradiated Mo is contained in these reviews, with the majority of data for irradiations <10 dpa and temperatures below 1073 K. The swelling data available are considerably scattered, with little coherence to examinations on the swelling as a function of temperature or dose. Swelling in Mo is expected to begin around 573– 673 K and continue to temperatures near 1573 K.19 Maximum swelling in pure Mo remains below 4% for fluences up to 1 1023 n cm2 (E > 0.1 MeV), 50 dpa, with peak swelling at irradiation temperatures near 900 K. Attempts at consolidating the reported swelling data as a function of irradiation temperature through normalizing the fluences proved
Radiation Effects in Refractory Metals and Alloys
195
at 1176 K.91 Swelling is expected to reach a maximum of 3–4% on the development of the void lattice structure, based on an attainment of an equilibrium ratio of void diameter to void superlattice parameter.92 At temperatures >1423 K, void lattice formation is no longer observed, leading to the high values of swelling observed in the material ion irradiated to high doses.88 The onset of void growth in neutron-irradiated material appears to be accelerated in cold-worked materials compared to annealed materials, reaching a maximum in swelling at doses near 40 dpa for temperatures below 873 K and 20 dpa at higher temperatures.89 At higher doses, swelling decreases through void shrinkage, with swelling values approaching those of annealed materials. Void shrinkage has also been reported by Bentley et al.93 and Evans53 to occur because of changes in the void sink bias89 presumably due to the segregation of transmuted species at the void surfaces, making them more attractive for interstitials. Irradiation-induced swelling in TZM has been reported53,94–97 and generally shows similar temperature dependence as the pure metal. The fluence and temperature dependence of swelling of TZM was examined by Powell et al.95 and Gelles et al.,94 with
to be inaccurate in determining the upper bound limit for maximum swelling.19 The swelling data collected from numerous sources,53,87–89 including those contained in the review work of Brimhall et al.90 for irradiated Mo as a function of dose and irradiation temperature, are provided in Figure 12. Void swelling was found to be <1% in Mo irradiated to 8 1022 n cm2, E > 0.1 MeV at temperatures between 673 and 1173 K by Evans.53 Void swelling studied by Stubbins et al.88 in 3.1 MeV 51V+ ion-irradiated Mo between 1173 and 1393 K up to 50 dpa remained below 4%, while irradiations between 1523 and 1780 K were near 10%. Void ordering has been observed in both neutronirradiated28,89,91 and ion-irradiated Mo88 at temperatures between 700 and 1373 K. Garner and Stubbins89 examined the irradiation and material conditions that contribute to void ordering. Irradiation temperatures near 700 K delineate the lower boundary temperature for void lattice formation at irradiations above 20 dpa. At lower doses, void lattice formation was not observed. The void superlattice constant, measured as the distance between void centers along the <100> direction in the material, is found to increase with temperature from 2.4 nm at 700 K to 4.5 nm
4 Garner and Stubbins89 (neutron) Stubbins et al.88 (ion) Evans53 (neutron) Lee et al.87 (neutron) Brimhall et al.90 (neutron) Brimhall et al.90 (ion)
DV/ V (%)
3
2
1 100 0
10 400 800 ion tem 1200 per atu re
diat
0.1 (K)
1600
dp a
1
Irra
0.01
Figure 12 Irradiation-induced swelling (DV/V ) as a function of irradiation temperature and displacement damage (dpa) for pure Mo. The irradiation source is marked in the key. Reproduced from Lee, F.; Matolich, J.; Moteff, J. J. Nucl. Mater. 1976, 62, 115–117; Evans, J. H. J. Nucl. Mater. 1980, 88, 31–41; Stubbis, J. F.; Moteff, J.; Taylor, A. J. Nucl. Mater. 1981, 101, 64–77; Garner, F. A.; Stubbins, J. F. J. Nucl. Mater. 1994, 212–215, 1298–1302; Brimhall, J. L.; Simonen, E. P.; Kissinger, H. E. J. Nucl. Mater. 1973, 48, 339–350.
196
Radiation Effects in Refractory Metals and Alloys
results from the latter shown in Figure 13. Peak swelling in TZM following irradiation to 1.78 1023 n cm2 and 873 K remained below 4%, though the data are limited to irradiation temperatures below 923 K. Only limited data are available on direct comparisons between TZM and pure Mo, with Bentley and Wiffen96 reporting 1% swelling in Mo–0.5%Ti and TZM alloys and 0.6% swelling in pure Mo under the same irradiation conditions. Similarly, 4% swelling was observed in TZM and 3% in pure Mo following irradiation to 5.4 1022 n cm2 at 923 K.97 In examining Mo and TZM of different preirradiated material conditions, Evans53 observed equal or greater swelling in TZM compared to Mo following irradiation to 3.5 1022 n cm2 (E > 0.1 MeV) at 823 and 873 K. However, in the materials irradiated at 723 K for the same fluence, the TZM alloy showed lower swelling, except in the carburized condition. The Ti and Zr atoms not tied up as carbides are assumed to have played a role in reducing void size in the material at the lower temperature. There is little information on the swelling behavior of Mo–Re alloys. Measured swelling of 0.44% in Mo–50Re irradiated to 5.3 1022 n cm2 (E > 0.1 MeV) at temperatures which rose during irradiation from 1128 to 1329 K was reported.26 For irradiated Mo–Re alloys, radiation-induced segregation (RIS) and transmutation can lead to precipitation of
equilibrium or nonequilibrium phases, which can be detrimental to mechanical properties. This is examined in the next section. Electrical resistivity changes to 5.4 dpa irradiated Mo at 733 K were examined by Zakharova et al.98 using single crystal samples. Increases in resistivity of 10–14% and 92–110% were measured at postirradiation test temperatures of 298 and 77 K, respectively. The largest resistivity changes were measured in the [100] direction. A residual 10% increase in resistivity was measured following annealing above 0.6 Tm associated with the accumulation of transmuted radionuclides. The changes in electrical resistivity of LCAC-Mo over a 353–1373 K irradiation temperature range up to 3.3 dpa were examined by Li et al.99 and Cockeram et al.,100 with the latter examining the recovery of resistivity following isochronal anneals. The room temperature resistivity for 353 K irradiated LCACMo rapidly increases between 0.01 and 0.1 dpa saturating near 0.2 dpa for an 42% increase over the unirradiated value.99 Increases in room temperature resistivity of 10–12% were reported following 0.5–1.2 dpa irradiation at 543 K, and 3.3–5.3% after 1.4–2.4 dpa at 878 K. At irradiation temperatures 1208 K, little (<3%) to no net increase in resistivity was observed for irradiations up to 3.3 dpa. This is reflected in the higher mobility of vacancies and
4.0 3.5
DV/V (%)
3.0
Gelles et al.94 (11-20 dpa) Gelles et al.94 (36-53 dpa) Gelles et al.94 (47-65 dpa) Evans53
2.5 2.0 1.5 1.0 0.5 80
0.0 600
60 700
40
800 Ti
rr
(K)
900
20 1000
a
dp
1100 Figure 13 The swelling dependence on temperature and fluence for neutron-irradiated TZM. Adapted from Gelles, D. S.; Peterson, D. T.; Bates, J. F. J. Nucl. Mater. 1981, 103–104, 1141–1146; Evans, J. H. J. Nucl. Mater. 1980, 88, 31–41.
Radiation Effects in Refractory Metals and Alloys
interstitials formed during irradiation to diffuse to sinks where annihilation occurs, reducing the electrical scattering effects that these defects have at lower irradiation temperatures. The small increases measured at the higher irradiation temperatures were primarily due to transmutation products. As is shown in the next section, the changes in electrical resistivity with increasing irradiation temperatures also correlate with changes in measured hardness, though at a greater level of sensitivity. This is controlled by microstructural changes, as the small dislocation loops and voids of high distribution density appearing at the lower irradiation temperatures coarsen into larger and fewer defects that have less interaction with deformation dislocations. 4.06.4.3 Mechanical Properties of Irradiated Mo and Mo Alloys The mechanical property performance of pure Mo is strongly controlled by the grain size, oxygen, nitrogen, and carbon concentrations as well as alloy additions. This is true for unirradiated as well as irradiated properties.71,76,83,84 The sensitivity to embrittlement at irradiation temperatures 873 K can be mitigated through a reduction in oxygen and nitrogen while keeping the carbon-to-oxygen ratio high to reduce the segregation of oxygen and nitrogen to the grain boundaries. A reduction in the grain size can further increase the number of sinks and reduce the mean distance that irradiation-induced defects must travel at temperatures at which mobility is limited. Irradiated mechanical properties of wrought LCAC-Mo in both the recrystallized and stressrelieved conditions have been examined over several decades.81,82,84–86,99–106 In general, LCAC-Mo undergoes significant increases in tensile strength through radiation hardening 873 K, which produces reductions in ductility and high DBTT values. A summary of tensile properties as a function of irradiation temperature and dose is shown in Figure 14. Irradiated stress-relieved LCAC-Mo shows less radiation embrittlement than as-crystallized materials at temperatures <1208 K.100 However, at higher irradiation temperatures, recrystallization of the stressrelieved material occurs, leading to large changes in the microstructure and less desirable properties. Increases in hardness of 56% and 112% for stressrelieved and recrystallized LCAC-Mo, respectively, are reported following irradiation to 1.2 dpa at 543 K.100 The increase in hardness decreases slightly following 878 K irradiations to 2.4 dpa, but returns to
197
values near those for the unirradiated material for irradiation at temperatures 1208 K. Materials irradiated in the stress-relieved condition at temperatures 1173 K can result in softening compared to unirradiated materials because of recrystallization.107 A comparison of the changes in irradiated material hardness as a function of dose and temperature is plotted in Figure 15(a) for LCAC-Mo, TZM, and ODS-Mo. The two latter alloys are discussed in detail later. In the three materials investigated by Cockeram et al.,107 the largest increase in hardness was measured for irradiations at 873 K, counter to what is observed in tensile properties for irradiation between 573 and 873 K.81,85 However, tensile failure in the lower temperature irradiated samples generally occurs before the samples yield because of the elevation of the flow stress above the fracture stress for these test conditions. Recovery of the hardness, measured through 1 h anneals at increasing temperatures, is plotted in Figure 15(b) for the 573 K irradiated material. The start of recovery is near 800 Kwith full recovery of the LCAC-Mo hardness by 1253 K. For LCAC-Mo irradiated at a higher temperature of 873 K, recovery of hardness begins near 1163 K and is completed near 1463 K. Substantial increases in the DBTT to values >773 K are observed for irradiated LCAC-Mo over a range of fluences for irradiation temperatures 873 K.26,82,85,103,105,108 A summary of DBTT values for LCAC-Mo is presented in Figure 16, along with data from high-purity grain-refined LCAC-Mo, TZM, and ODS-Mo, which is discussed next. In general, recovery in the DBTT of LCAC-Mo is not observed until irradiation temperatures above 873– 973 K, depending on material conditions. A reduced sensitivity to low-temperature irradiation embrittlement of LCAC-Mo is observed in materials with reduced levels of impurities. A high-purity form of LCAC-Mo (HP-LCAC-Mo) was developed through the use of 1873 K hydrogen atmosphere annealing of LCAC-Mo plates prior to further arc casting, extrusion, and rolling into sheet stock.82 Levels of oxygen and nitrogen were <4 wppm each, with carbon at 20 wppm. Average grain diameters of 1.3 and 452 mm lengths were produced, and represent a considerable change in aspect ratio compared to LCAC-Mo values of 4–5 mm diameter and 78–172 mm lengths.100 The DBTT of 573 K irradiated HP-LCAC-Mo showed no increase over the unirradiated value (123 K) for irradiations up to 0.11 dpa, and increased to 723 K by 1.29 dpa. For 873 K irradiations, the DBTT remained below 223 K up to 1.29 dpa.
198
Radiation Effects in Refractory Metals and Alloys
2000 HP-LCAC (0.11-1.29 dpa, 573 K)
LCAC (unirr.)
LCAC (0.6 dpa, 573 K)
Yield stress (MPa)
1500
LCAC (13.1 dpa, 833-1057 K) LCAC (12.3 dpa, 567 K)
HP-LCAC (unirr.)
EB (0.23 dpa, 873 K)
1000
LCAC (0.9-1.4 dpa, 878 K)
500 LCAC (3.3 dpa, 1373 K)
25
LCAC (3.3 dpa, 1373 K)
Total elongation (%)
20
LCAC (unirr.)
15
EB (0.23 dpa, 873 K)
10
LCAC (0.9 -1.4 dpa, 878 K)
HP-LCAC (unirr.)
5
HP-LCAC (0.11-1.29 dpa, 573 K)
LCAC (13.1 dpa, 833-1057 K) LCAC (12.3 dpa, 567 K)
0 0
200
400
600
800
1000
1200
Test temperature (K) Figure 14 Yield stress and total elongation as a function of test temperature for LCAC-Mo, HP-LCAC-Mo and electron beam (EB) melted Mo. Irradiation dose and temperature (dpa, K) is marked. Unirradiated samples (unirr.) are also presented. Reproduced from Cockeram, B. V.; Smith, R. W.; Leonard, K. J.; Byun, T. S.; Snead, L. L. J. Nucl. Mater. 2008, 382, 1–23; Cockeram, B. V.; Smith, R. W.; Snead, L. L. J. Nucl. Mater. 2005, 346, 165; Abe, K.; Takeuchi, T.; Kikuchi, M.; Morozumi, S. J. Nucl. Mater. 1981, 99, 25–37.
The majority of the irradiated mechanical property database for TZM is limited to displacement damage <5 dpa and irradiation temperatures <1000 K, with much of the testing conducted at temperatures below that used in the irradiation. The database is also sparse, with little connectivity between
different experimental examinations. Furthermore, significant differences are observed in the nonirradiated tensile values of TZM because of variations in material processing leading to differences in grain size, impurity level, distribution of the strengthening phase, and differences in testing procedure. In general,
Radiation Effects in Refractory Metals and Alloys
400
LCAC-Mo (Tirr = 573 K)
LCAC-Mo (Tirr = 873 K)
LCAC-Mo (Tirr = 1173 K)
ODS-Mo (Tirr = 573 K)
ODS-Mo (Tirr = 873 K)
ODS-Mo (Tirr = 1173 K)
TZM (Tirr = 573 K)
TZM (Tirr = 873 K)
TZM (Tirr = 1173 K)
199
Change in hardness (MPa)
300 873 K
200 573 K
100
0 1173 K -100 0.00
2.00
4.00
(a)
1700
Recovery of hardness
6.00 8.00 Displacement dose (dpa)
10.00
12.00
14.00
>1573 K
75-100% 1500
50-75% 25-50%
1 h anneal temperature (K)
0-25% 1300
No change
1100
900
700
500 Dose 0.6 dpa 12.3 dpa 1.4 dpa 13.1 dpa 12.3 dpa 3.9 dpa 13.1 dpa 12.3 dpa 3.9 dpa 13.1 dpa Tirr 543 K 567 K 877 K 833 K 573 K 877 K 833 K 573 K 882 K 833 K LCAC LCAC LCAC LCAC ODS ODS ODS TZM TZM TZM (b)
Figure 15 (a) Change in hardness as a function of neutron dose for LCAC-Mo, TZM, and ODS-Mo irradiated between 573 and 1173 K up to 13.1 dpa. (b) Recovery of hardness as a function of isochronal annealing temperature for material irradiated 573 K. Adapted from Cockeram, B. V.; Smith, R. W.; Byun, T. S.; Snead, L. L. J. Nucl. Mater. 2009, 393, 12–21.
200
Radiation Effects in Refractory Metals and Alloys
800
LCAC-Mo Wiffen26 (13.3-23.4 dpa) Chakin and Kazakov108 (2.1-3 dpa) Hasegawa et al.106 (10.5-50 dpa) Abe et al.104,105 (0.6-1.4 dpa) Webster et al.148 (7.4-10 dpa) Smith and Michel149 (5.3 dpa) Scibetta et al.109 (0.15-0.19 dpa) Cockeram et al.85 (12.3-13 dpa) HP-LCAC-Mo Cockeram et al.82 (1.29 dpa) TZM Cockeram et al.85 (12.3-13 dpa) Wiffen26 (16 dpa) ODS-Mo Cockeram et al.85 (12.3-13.1 dpa)
700 600
DBTT (°C)
500 400 300 200 100 0 -100 0
100 200 300 400 500 600 700 800 900 1000 1100 1200 1300 Irradiation temperature (°C)
Cockeram et al.82 (1.29 dpa)
Figure 16 Summary of DBTT values as a function of neutron irradiation temperature for LCAC-Mo and TZM. Adapted from Cockeram, B. V.; Smith, R. W.; Leonard, K. J.; Byun, T. S.; Snead, L. L. J. Nucl. Mater. 2008, 382, 1–23; Cockeram, B. V.; Smith, R. W.; Snead, L. L. J. Nucl. Mater. 2005, 346, 165.
displacement damage to the TZM alloy produces a large increase in the yield strength of the material with a corresponding drop in total elongation, with the level of change increasing with dose and/or lower irradiation temperatures. Some recovery of properties begins to be observed at test conditions above 973 K for materials irradiated at lower temperatures. A compilation of earlier and more recent tensile data for irradiated TZM is provided in Figure 17. Irradiation 1.2 dpa and temperatures >800 K showed little effective strengthening above the unirradiated values, for tensile tests below the irradiation temperature, 85,109,110 though higher displacement doses resulted in a significant increase.81 Total elongation in material irradiated <1000 K was limited, with uniform elongation values <1%.111 Plastic instability following yielding was also observed in irradiated TZM as well as LCAC-Mo. The plastic instability stress, defined as the maximum true stress in which plastic necking occurs, is strongly dependent on test temperature but nearly independent of fluence and irradiation temperature.81 Increases in the DBTT for irradiated TZM are significant and do not diminish until irradiation temperatures slightly higher than that of LCAC-Mo,81 but comparisons may be difficult because of differences in materials and test methods. The DBTT value for unirradiated TZM is 200 K and increases with increasing irradiation damage. An increase of 230 K in the DBTT over unirradiated values was reported for TZM irradiated to 2–4.8 dpa at 644–661 K.111 DBTTs
of 750 K for Mo–0.5Ti irradiated to 16 dpa at 727 K26 and 1073 K for TZM irradiated to 12.3 dpa at 573 K81,85 were also reported. A summary of DBTT as a function of irradiation temperature for LCAC-Mo and TZM is shown in Figure 16 for irradiation doses <50 dpa.85 Excessive embrittlement is observed for irradiations at <773 K. A reduction in the DBTT for LCAC-Mo appears near 873 K, while the DBTT of TZM remains high. Recovery of DBTT to near unirradiated values occurs for irradiations at >1073 K.81,85 The stage V recovery temperature for vacancy diffusion in Mo is 873 K; however, the kinetics for microstructural changes to occur are still relatively slow at this temperature. Therefore, embrittlement issues can be present until 1073 K. Very little fracture toughness data exist for irradiated TZM. For precracked compact tension specimens irradiated to 0.29–0.35 dpa at 313–695 K, a 4 MPa√m decrease in fracture toughness (15– 20 MPa√m at Trm, unirradiated84,109) was observed up to the irradiation temperature.109 Kitsunai et al.112 examined the impact toughness of irradiated TZM and alloys, incorporating 0.1–1 wt% TiC additions to Mo. The Mo–TiC alloys showed dramatically improved toughness levels over TZM with increasing TiC concentration in samples irradiated to 0.08 dpa between temperatures of 573 and 773 K. Shown in Figure 18 is the shift in DBTT to lower temperatures for the irradiated TiC-containing Mo over the TZM alloy. For the 1 wt% containing sample,
Radiation Effects in Refractory Metals and Alloys
1974 Kasakov et al.: unirradiated
201
2005 Cockeram et al./Byun et al. (unirradiated)
1974 Kasakov et al. (1.6 dpa, T irr = 823 K)
1974 Kasakov et al. (1.2 dpa, T irr = 1223 K)
2005 Cockeram et al. (3.9 dpa, T irr = 833 K)
2005 Cockeram et al. (12.3 dpa, T irr = 567 K)
2008 Byun et al. (3.9 to 13.1 dpa, T irr = 567 - 833 K)
2008 Byun et al. (3.9 to 13.1 dpa, T irr = 1057 - 1209 K)
1400
Yield strength (MPa)
1200
1000
800
600 Unirradiated range 400
200 200
400
600
(a)
800
1000
1200
1400
Test temperature (K) 2005 Cockeram et al./Byun et al. (unirradiated) 1974 Kasakov et al. (1.2 dpa, T irr = 1223 K)
1974 Kasakov et al.: unirradiated 1974 Kasakov et al. (1.6 dpa, T irr = 823 K) 2005 Cockeram et al. (3.9 dpa, T irr = 833 K)
2005 Cockeram et al. (12.3 dpa, T irr = 567 K) 1976 Steichen (4.79 dpa, T irr = 661 K)
1976 Steichen (2.12 dpa, T irr = 644 K)
20 18
Total elongation (%)
16 14 12 Unirradiated range
10 8 6 4 2 0 200
(b)
300
400
500
600
700
800
900
1000
1100
1200
Test temperature (K)
Figure 17 Neutron-irradiated tensile data for TZM (a) yield stress versus test temperature and (b) total elongation versus test temperature. Reproduced from Byun, T. S.; Li, M.; Cockeram, B. V.; Snead, L. L. J. Nucl. Mater. 2008, 376, 240–246; Cockeram, B. V.; Smith, R. W.; Snead, L. L. J. Nucl. Mater. 2005, 346, 165; Kasakov, V. A.; Kolesnikov, A. N.; Krassnoselov, V. A.; et al. Effect of neutron irradiation on properties of potential structural material for thermonuclear reactors, USSR-US Exchange on CTR Materials, Nov 1974; Steichen, J. M. J. Nucl. Mater. 1976, 60, 13–19.
the DBTT remained unchanged with irradiation, despite a 50% increase in Vickers microhardness. More surprisingly, the Mo–1%TiC sample increased in toughness following 0.8 dpa irradiation attributed to grain boundary strengthening by (Ti, Mo)C radiation-enhanced precipitates. Developed to improve the low-temperature ductility and weld characteristics of unalloyed Mo, the Mo–Re alloys have gained considerable attention over the past decade for use in nuclear applications. Single-phase solid solution a-Mo phase field extends
up to 42 wt% Re, above which the s-MoRe2 phase precipitates. At higher Re concentrations, the w-MoRe3 phase is present. However, the exact phase boundaries are not well delineated at temperatures below 1773 K113,114 mainly because of the slow kinetics in phase development.115 The Mo–Re alloys show a hardening response to irradiation stronger than that of the pure metal and TZM following irradiation.116–118 The hardening response of Mo–Re alloys ranging in composition from 2 to 13 as well as 41 wt% Re following
202
Radiation Effects in Refractory Metals and Alloys
0.9 TZM Total absorbed energy (J mm-3)
0.8
Mo-0.1%TiC Mo-0.5%TiC
0.7
Mo-1%TiC 0.6 0.5 0.4 0.3 0.2 0.1 0 200
250
300 350 400 Test temperature (K)
450
500
Figure 18 Total absorbed energy versus test temperature for TZM and 0.1–1.0 wt% TiC additions to Mo irradiated to 0.08 dpa at 573–773 K. Reproduced from Kitsunai, Y.; Kurishita, H.; Narui, M.; Kayano, H.; Hiraoka, Y. J. Nucl. Mater. 1996, 239, 253–260.
irradiation up to 20 dpa at temperatures between 681 and 1072 K was examined by Nemoto et al.116 A linear increase in hardness with Re concentration was observed for the unirradiated controls as well as samples irradiated at temperatures 874 K. For samples irradiated at 1072 K, little variation was observed with increasing Re content, though hardness values remained nearly double that of the unirradiated material. The dependence of hardness on irradiation temperature and fluence for Mo–5Re and Mo–41Re in comparison with LCAC-Mo is presented in Figure 19. The high degree of radiation hardening at temperatures <1000 K exhibited by the Mo–Re alloys is further reflected in low ductility and reported embrittlement. Tensile elongation values of <0.3% were reported for Mo–5Re irradiated to 0.16 dpa and tested at the irradiation temperature of 320 K,118 though higher total elongations of 8% were reported for Mo–5Re fast reactor irradiated to 0.29 dpa and tested near the irradiation temperature of 723 K.109 Unlike LCAC-Mo and TZM, the little deformation occurring in Mo–Re is mostly uniform at temperatures below 1000 K,118,119 while a small degree of work softening has been observed at higher temperatures.120 No evidence of dislocation channeling was found during microstructural examination of tensile tested Mo–5Re irradiated to 0.16 dpa at 373 K.118 For comparable irradiation fluences, dislocation loop concentrations are approximately two times higher than TZM and four to six times higher than Mo, while dislocation loop diameters are smaller in
the Mo–5Re alloy.118 Void development begins to appear in Mo–Re alloys above 623 K118 and shows a slight increase in size with Re concentration (with corresponding decrease in number density) up to 10 wt% with no further increase for the 41% Re alloy.116 Irradiation-induced void swelling in Mo with Re concentrations 13 wt% is 0.5–1.5% for 21 dpa irradiated material at 681–1072 K, while swelling for 41 wt% Re was near 0.1%.116 Radiation-induced precipitation has been reported by Nemoto et al.116 in Mo–(2–41)Re alloys irradiated to 21 dpa between 681 and 1072 K, and by Edwards et al.121 in Mo–41Re irradiated 28–96 dpa at 743– 1003 K. An initial formation of hcp-structured precipitates with a thin plate-like morphology consisting of solid solution Re and Os was observed, appearing with the {110}Mo//{0001}Re, <111>Mo//<2110>Re orientation relationship.121 The precipitation of these plates on dislocation loops resulted in the high density of plates observed, which dominates the microstructure. On further irradiation to higher doses or higher temperatures, these plates develop and coarsen into the w-phase. This nonequilibrium phase development in Re-lean alloys was originally observed by Erck and Rehn122 in Mo–(27–30)Re irradiated by 1.8 MeV He ions at 1023– 1348 K. The s-MoRe2 phase was also reported appearing in all the Mo–Re samples examined by Nemoto and coworkers,116 but was suppressed in the stress-relieved specimens compared to the recrystallized materials. While a limited (<1%) amount of ductility was reported in Mo–5Re alloys irradiated at temperatures <1000 K to 34 dpa,109,118,119 embrittlement of Mo–(1–20)Re irradiated 723–1073 K in a fast reactor up to 5 dpa and Mo–(13 and 47)Re irradiated 373– 673 K in a mixed spectrum reactor to 2 dpa has been reported by Fabritsiev and Pokrovsky.62 The reduction in tensile strength, in many cases below the unirradiated values with no plasticity occurring in the samples, was attributed to the hardening of the material by radiation-created defects along with RIS of oxygen, nitrogen, and transmuted impurities to the grain boundaries. The oxygen and nitrogen content in the embrittled alloys was reported to be near 70 appm. Irradiation hardening above 900 MPa was also observed in Mo–41Re and Mo–47.5Re samples irradiated to 1.46 dpa at temperatures >1073 K.120 Failure of Mo–41Re samples irradiated to 1.46 either prior to yielding or after 5% elongation upon reaching 1600 MPa was observed at 1073 K (Tirr ¼ Ttest). Examples of the tensile curves for the two alloys in the irradiated, 1100 h aged and as-annealed condition tested at 1073 K is shown in Figure 20. Radiation
203
Radiation Effects in Refractory Metals and Alloys
1800 1600 Mo-41Re
Vickers hardness
1400 1200 1000
Mo-5Re
800 Unirrad. 600 400 LCAC-Mo
200 0 200
400
600 800 Irradiation temperature (K)
1000
1200
Mo (Nemoto et al.116 18-21 dpa)
Mo-5Re (Nemoto et al.116 18-21 dpa)
Mo-10Re (Nemoto et al.116 18–21 dpa)
Mo-41Re (Nemoto et al.116 18-21 dpa)
117 6.8-34 dpa)
Mo-5Re (Hasegawa et al.117 6.8-34 dpa)
Mo-5Re (Hasegawa et al.
117 6.8-34 dpa)
Mo-41Re (Hasegawa et al.
Figure 19 Vickers hardness as a function of neutron irradiation temperature and dose for LCAC-Mo, Mo–5Re, and Mo–41Re alloys. Displacement damage levels are provided in the key. Reproduced from Nemoto, Y.; Hasegawa, A.; Satou, M.; Abe, K.; Hiraoka, Y. J. Nucl. Mater. 2004, 324, 62–70; Hasegawa, A.; Ueda, K.; Satou, M.; Abe, K. J. Nucl. Mater. 1998, 258–263, 902–906.
2000 Mo-41Re
0.72 dpa
1.46 dpa
Mo-47.5Re
1500
1500 1.46 dpa Stress (MPa)
0.72 dpa 0.72 dpa
1.46 dpa 1000
1.46 dpa
1000
0.72 dpa
500
500 1100 h aged
1100 h aged
Annealed 0
0
5
10
15 20 25 Strain (%)
30
35
0
5
10
15 20 25 Strain (%)
Annealed 30
35
40
Figure 20 Comparison of stress–strain curves for neutron irradiated, 1100 h and as-annealed Mo–41Re and Mo–47.5 Re samples at 1073 K (Tirr ¼ Ttest). Adapted from Busby, J. T.; Leonard, K. J.; Zinkle, S. J. Effects of neutron irradiation on refractory metal alloys, ORNL/LTR/NR-PROM1/05-38; Oak Ridge National Laboratory: Oak Ridge, TN, Dec 2005; Busby, J. T.; Leonard, K. J.; Zinkle, S. J. J. Nucl. Mater. 2007, 366, 388–406.
204
Radiation Effects in Refractory Metals and Alloys
Total elongation (%)
Ultimate tensile strength (MPa)
2500 Mo-41 (Busby) Mo-Re (Fabrietsiev) Mo-47.5Re (Busby) Mo-5Re (Hasegawa) Open symbols: unirrad. Closed symbols: irrad.
2000
Irrad. UTS 1500 Unirrad. UTS
1000
500
Irrad. UTS (brittle fracture)
15
10
5 Brittle fracture 0 600
700
800
900 1000 1100 Temperature (K)
1200
1300
1400
Figure 21 Tensile data comparisons of Mo–Re alloys detailing the upper limits for irradiation embrittlement. Adapted from Busby, J. T.; Leonard, K. J.; Zinkle, S. J. Effects of neutron irradiation on refractory metal alloys, ORNL/LTR/NR-PROM1/ 05-38; Oak Ridge National Laboratory: Oak Ridge, TN, Dec 2005; Busby, J. T.; Leonard, K. J.; Zinkle, S. J. J. Nucl. Mater. 2007, 366, 388–406.
hardening to levels over twice the as-annealed condition was observed for the alloys irradiated at 1223 and 1373 K; however, total elongation was between 4 and 12%. Analysis of the fractured surfaces of these samples revealed intergranular failure, with the severity increasing with irradiation temperature. A comparison of mechanical property data of Mo–Re samples from the sources discussed is shown in Figure 21. The degree of RIS influencing the properties of Mo–Re alloys varies with temperature, dose rate, and total fluence. At temperatures <0.3 Tm, the recombination of vacancies and interstitials generated by displacement damage dominates because of the limited defect mobility, and therefore RIS is not a factor. At temperatures >0.5 Tm, a reduced driving force for segregation occurs because of the high thermal defect concentrations. At intermediate temperatures (850–1430 K for Mo–Re), the radiation generated point defects diffuse to defect sinks such as grain
boundaries or dislocations. Any preferential coupling of vacancies or interstitial defects fluxes with solute atoms, including transmuted species, will create enrichment at the defect sinks. This is observed in the nucleation of Re-rich phases in the microstructures of neutron-irradiated samples116,121 and the degradation in mechanical properties and transition to intergranular fracture in higher Re concentration alloys.62,120 Further information on RIS can be found in Chapter 1.18, Radiation-Induced Segregation. Through modeling and experimental work, Erck and Rehn123 showed that the degree of segregation per dpa reaches a maximum for Mo–30 at.% Re (45 wt% Re) near 1223 K and that for Mo–7 at.% Re (13 wt% Re) near 1473 K. While the Mo–5Re alloys irradiated up to 20 dpa show some limited ductility,109,118,119 the maximum irradiation temperatures were <1073 K and are therefore at or below the lower temperature limit expected for RIS.
Radiation Effects in Refractory Metals and Alloys
1800 Tirr = 567 K, 12.3 dpa Tirr = 882 K, 3.9 dpa Tirr = 1143 K, 1.2 dpa Tirr = 1143 K, 3.4 dpa Unirradiated: longitudinal, stress-relieved
1600 1400
Yield stress (MPa)
1200 1000 800 600 400 200
20
Total elongation (%)
The Mo–(41 and 47.5)Re alloys irradiated at 1073– 1373 K120 showed indications of RIS even at relatively low damage levels, part of which may have been a contribution of a thermal aging component, which in the unirradiated as-aged Mo–41Re and Mo–47.5Re showed increases in Re at the grain boundaries, leading to precipitation of s- and w-phases at the grain boundaries in the 47.5Re containing alloy.115 Utilizing Mo–Re alloys with a more moderate Re content may improve the irradiation performance of these alloys, especially when considering higher doses and/or longer irradiation times at temperatures at which thermal precipitation effects may further compound RIS influences on mechanical properties. Additional information on the fracture toughness data for Mo–Re alloys is also needed. Preliminary data by Scibetta and coworkers109 on precracked compact tension specimens of Mo–5Re showed reductions in fracture toughness from unirradiated values of 17–23 MPa√m at room temperature and 623 K, to 11 MPa√m for 0.35 dpa irradiated at 313 K, and 15 MPa√m for 0.29 dpa irradiated at 643 K. These low irradiated values, at which no ductile crack growth was observed in the specimens, are a concern. Recent work examining the irradiated properties of wrought, commercially available, ODS-Mo containing lanthanum oxide particles has shown promising results.81,82,124 The fabrication methods produce a microstructure consisting of elongated grains with appreciable texturing and alignment of the oxide particles. The high degree of working associated with fabrication produces a <2 mm grain size, which is stabilized from growth by the ODS particles. Irradiation of the ODS-Mo up to 13.1 dpa at 567 K and 883–882 K produced an increase in yield strength of 57–173%,124 while irradiation at 1143–1273 K produced a 10–34% increase. The irradiated tensile properties of ODS-Mo as a function of irradiation temperature and dose from the work of Cockeram and coworkers124 are shown in Figure 22. The increases in radiation strength are comparable to the higher limits for LCAC-Mo. The most striking result of the ODS work is the improvement in the DBTT for the irradiated samples.82,124 For 567 K irradiation to 12.3 dpa, the DBTT is 1073 K and is comparable to that of LCAC-Mo and TZM (Figure 16). However, the DBTT for ODS-Mo irradiated to 13.1 dpa at 833–882 K is 298 K, while that of LCAC-Mo is 573 and 973 K for TZM. For irradiation to 13.1 dpa at 1143–1209 K, the DBTT is 173 K, unchanged from the nonirradiated material, while those of LCAC-Mo and TZM are between 223–273 K.
205
15
10 Unirradiated 5
0 100
300
500
700 900 1100 Test temperature (K)
1300
1500
Figure 22 Yield stress and total elongation as a function of test temperature for lanthanum oxide ODS-Mo neutron irradiated up to 13 dpa at temperatures between 567 and 1209 K. Adapted from Cockeram, B. V.; Smith, R. W.; Snead, L. L. J. Nucl. Mater. 2005, 346, 165–184.
The reduced susceptibility to irradiation embrittlement of ODS-Mo is in part due to the grain size reduction and presence of the oxide particles. Reducing the distance of possible defect sinks such as grain boundaries and offering additional sites such as the oxide/matrix interface are particularly critical at lower irradiation temperatures at which defect mobility is limited. In addition, Cockeram and coworkers81,124 describe the fine but elongated grain structure as enhancing the plain strain condition acting on each plane of the lamina-shaped grains formed during the fracture process, inducing larger plastic deformation in irradiation-hardened material. This is also true for the HP-LCAC-Mo containing the high aspect ratio grains compared to other forms of LCAC-Mo produced. Currently, no fracture data on the irradiated
206
Radiation Effects in Refractory Metals and Alloys
ODS-Mo are available. Unirradiated fracture toughness values are between 23 and 38 MPa√m, depending on the grain orientation tested.83
4.06.5 Tungsten and W-Base Alloys 4.06.5.1 Introduction and Irradiated Properties Database for W and W Alloys Despite the recurring interest in the use of tungsten as a structural material for very-high-temperature applications and for use as plasma facing components in fusion devices, the database on irradiated properties is very limited and based primarily on fast neutron irradiation experiments. Similar to all bcc materials, tungsten is susceptible to low-temperature embrittlement at T 0.3 Tm (Tm ¼ 3695 K) for fluences >1 1024 n m2, which makes this material even more limited in toughness and ductility. Improvements in the unirradiated mechanical properties of tungsten are observed with the addition of Re, which is found to increase ductility at elevated temperatures125 as well as fracture toughness126 through the reduction in the DBTT. However, as discussed in this section, the gains in performance through added Re content do not necessarily hold for irradiated materials.
with a less pronounced recovery at 0.22 Tm through divacancy and impurity migration, which was followed by vacancy migration above 0.31 Tm. The residual resistivity, not recovered following anneals above 0.4 Tm after irradiation to fluences >3.3 1019 n cm2, was due to the development of Re in the tungsten from transmutation reactions with thermal neutrons. Very little data exist on irradiation-induced swelling in Wand its alloys. Data on pure Ware restricted to two reported series of experiments concerning the temperature dependence of swelling. The irradiationinduced swelling measured by Matolich et al.131 and Wiffen19 using immersion density methods is shown in Figure 23. It should be noted that there is an order of magnitude difference in fluences between the two studies. No other systematic examination of the swelling dependence on fluence and temperature is available. Though swelling data for W–Re alloys are also limited, work by Matolich et al.131 for W–25Re irradiated to 5.5 1022 n cm2 revealed no significant amount of swelling. The data are also shown in Figure 23. Microstructural examination of W–Re alloys with concentrations of 5%, 11%, and 25%Re showed no cavity formation for fluences between 4.3
1.8 1.6
4.06.5.2 Irradiation-Induced Swelling and Physical Property Changes in W and W Alloys
1.2 DV/ V (%)
Early investigations into the behavior of irradiated tungsten59,125,127–129 examined the defect formation and recovery defects. Much of this initial work was through the examination of electrical resistivity following irradiation. Increases in electrical resistivity of pure annealed tungsten of up to 24% following irradiation to 2 1022 n cm2 in a fast reactor and 14% in a mixed spectrum reactor to 1021 n cm2 have been reported.59 Keys et al.127,128 examined the recovery of neutronirradiated tungsten through isochronal resistivity studies following irradiation at 343 K to dose levels of 1.5 1021 n cm2 (E > 1 MeV). The beginning of saturation in resistivity observed in their studies appears just below 1020 n cm2 and correlates with the work by Lacefield et al.130 on the appearance of defect clusters by 2.4 1019 n cm2 identified through TEM examination. The work by Keys et al.127,128 identified distinct stages of recovery such as self-interstitial migration occurring near 0.15 Tm,
1.4
1 0.8 0.6 0.4 0.2 0 600
800
1000 1200 1400 1600 Irradiation temperature (K)
1800
2000
Tungsten: Matolich et al.131 5.5 ´ 1022 n cm–2 (E > 0.1 MeV) Tungsten: Wiffen19 4 to 6 ´ 1022 n cm–2 (E > 0.1 MeV) W-25Re: Matolich et al.131 5.5 ´ 1022 n cm–2 (E > 0.1 MeV)
Figure 23 Irradiation-induced swelling measured through immersion density methods of W and W–25Re by Matolich et al.131 and Wiffen.19
Radiation Effects in Refractory Metals and Alloys
and 6.1 1021 n cm2 (E > 0.1 MeV) at temperatures between 873 and 1773 K,132 while void cavities have been experimentally observed in irradiated pure W over similar fluences.133,134 The effects of increasing Re or Os content in W were experimentally shown to decrease the density and radius of dislocation loops and voids in 0.15 dpa proton and neutronirradiated material by He et al.135 This reduction in size and number density is the result of the restricted mobility of the radiation-induced defects by the lattice dilations from the Re and Os solute. The transmutation of W to Re and Re to Os during irradiation can have an effect on microstructural, physical, and mechanical properties of the material. The transmutation of the material, which results in the shifting of solute concentrations to higher levels, may result in precipitation in alloys that are nominally in a single-phase region. One example of microstructural and physical property changes because of irradiation is the decalibration of type-C (W–3%Re/ W–25%Re or W–5%Re/W–26%Re) thermocouples, such as that used in fuel element centerline temperature measurements. Reviews of early experimental work on W/Re thermocouples and on the dependence of decalibration on the neutron fluence have previously been discussed.136,137 While displacive neutron damage may result in material changes such as vacancy clusters or dislocation loops, the maximum theoretical changes expected in emf output of the thermocouple is 1 mV C1,138 whereas the changes associated with transmutation effects can result in more significant decreases. A 300 C drift in temperature following 6000 h irradiation under 2.7 1022 n cm2 thermal and 8 1021 n cm2 fast fluence was reported for a W–3%Re/W–25%Re couple.139 These changes can also be significant in fast reactor irradiations. Experimental work by Williams et al.132 showed that for a W–5%Re/W–25%Re thermocouple irradiated to 6.1 1021 n cm2 fast fluence at 1173 K, the precipitation-induced changes in the Seebeck coefficient were 6.6 and 0.02 mV C1 for the 5 and 25% Re alloys, respectively. Calculated final compositions following 6.1 1021 n cm2 (14 MeV) irradiation of W–5(wt%)Re produce W–5.130Re– 0.021Os–0.150Ta alloy, while a W–26Re alloy transmutes to W–25.955Re–0.107Os–0.117Ta.140 Postirradiation examination of the microstructure of the irradiated 5, 11, and 25% Re alloys in Williams et al.132 revealed w-phase precipitation at irradiation temperatures above 1373 K, though unidentifiable precipitation was apparent in the alloys at 1173 K.
207
The development of the w-phase over the equilibrium s-phase in irradiated samples, but not in the unirradiated annealed samples, is the result of irradiationinduced solute segregation to defect sinks. The development of the w-phase was also reported in microstructural studies of W–26Re irradiated up to 11 dpa at temperatures between 646 and 1073 K.141 It should be pointed out that the change or temperature shift under irradiation is proportional to the degree of localized transmutation and local temperature gradients and therefore dependent on the profiles of the temperature and irradiation fields to which the thermocouple is exposed. Therefore, experimental work typically involves the irradiation of the entire cable, while in reactor applications, significant variations in temperature and fluence may result. The changes in thermoelectric power (D) as a function of irradiation fluence can be modeled by the following:140 D ¼ 0 for 0 f 0:25 1021 n cm2 D ¼ 100½1 e0:067ð0:25fÞ for 0:25 f 1 1021 n cm2 D ¼ 100½1 e0:104ð0:52fÞ for f > 1 1021 n cm2
½2
Though significant radiation-induced decalibration may occur in fission reactors, this effect may not be readily observed in fusion reactors where the thermal flux is much lower. In addition, typical end-of-life estimates of total neutron fluence of <1021 n cm2 suggest that transmutation effects resulting in decalibration is not an issue.
4.06.5.3 Irradiated Mechanical Properties of W and W Alloys Like all other refractory metal alloys, tungsten is sensitive to embrittlement issues following irradiation, the causes of which are point defect generation, impurity segregation to grain boundaries, and radiation-enhanced precipitation. The strengthening of the metal matrix can raise the deformation stress to levels greater than the cleavage strength of the alloy, resulting in brittle failure. While only a limited number of mechanical property tests have been performed on tungsten and its alloys, the actual composition of the materials reported is not well characterized and therefore may have a significant range of impurity levels that can affect grain boundary cohesion and the mechanical strength of the material. Similarly, preirradiation heat treatments have also shown minor improvements in the postirradiation mechanical
208
Radiation Effects in Refractory Metals and Alloys
properties,59 likely through the development and coarsening of interstitial impurities into precipitate formations that reduce grain boundary sensitivities. For the irradiated tensile properties of tungsten, two works are typically referenced that make up the bulk of the data available. In the work of Steichen,111 the properties of wrought tungsten, stress-relieved at 1273 K, irradiated at 658 K to fluences between 0.4 1022 and 0.9 1022 n cm2 (E > 0.1 MeV), were examined, the results of which are shown in Figure 24. The irradiated yield strength of the material increased to approximately twice that of the unirradiated values,
while significant reduction in ductility was observed. Only in tensile tests well above the irradiation temperature did ductility values approach that of the unirradiated material. In the work by Gorynin et al.,59 pure W consolidated through powder metallurgy was irradiated at temperatures of up to 1073 K in both a mixed and fast reactor up to 2 1022 n cm2 (E > 0.1 MeV). Samples irradiated and tested at temperatures near 573 K showed brittle failures at low stress levels, while some ductility and appreciable hardening were observed for samples tested and irradiated at
1600 1400
0.9 ´ 1022 n cm–2 Tirr = 661 K Steichen111
Yield stress (MPa)
1200 0.5 ´ 1022 n cm–2 Tirr = 644 K Steichen111
1000 800
Unirradiated Steichen111
600 400
Brittle failure
200
1.5 ´ 1022 n cm–2 Tirr = 773 K Gorynin et al.59
1.0 ´ 1022 n cm–2 Tirr = 603–723 K Gorynin et al.59
0
Total elongation (%)
20
15
10
5
0 200
400
600
800 1000 Test temperature (K)
1200
1400
Figure 24 Temperature-dependent tensile properties of irradiated and unirradiated tungsten. Reproduced from Gorynin, I. V.; Ignatov, V. A.; Rybin, V. V.; et al. J. Nucl. Mater. 1992, 191–194, 421–425; Steichen, J. M. J. Nucl. Mater. 1976, 60, 13–19.
Radiation Effects in Refractory Metals and Alloys
1073 K. Limited recovery of strength and ductility was observed in postirradiated material annealed at 1473 K for 1 h. Embrittlement following irradiation due to radiation hardening and loss of grain boundary strength due to impurities resulted in increased DBTT for the aforementioned work. The DBTT is dependent on the test conditions in addition to material conditions prior to irradiation and should be used with caution. The DBTT in samples examined by Steichen111 increases from 333 K in the unirradiated condition to 503 K, following 1–2 dpa irradiation at 653 K, while DBTT values increased from 673 K unirradiated to 873 K after 1 dpa at 373 K for sintered W.59 Increased DBTT values with irradiation were also reported by Krautwasser et al.142 in powder metallurgy to form W, W–10Re, and Densimet 18 (W–3.4Ni–1.6Fe) bend-test bars irradiated between 525 and 575 K up to 5.6 1021 n cm2 (E > 0.1 MeV) (see Figure 25). While the addition of Re to W results in improved nonirradiated mechanical properties,142 the increased DBTT in the irradiated W–10Re is more severe than in pure W. In the case of the former, the possible development of the w-phase may be responsible for the higher DBTT values and the general increased sensitivity to radiation hardening. The w-phase observed in W–26Re irradiated from 2 to 9.5 dpa at temperatures between 373 and 800 C141 is reported as precipitating as plate-like particles on the {110} planes of the W matrix, therefore, restricting slip in the material. 1400 1200
W-pure W-10Re
DBTT (K)
1000 800 600 400
209
It should be noted that due to the limited mechanical property data available for W and W–Re alloys, particularly the lack of irradiated data at elevated temperatures, accurate determination of the DBTT cannot be made. Nonetheless, increases in DBTT between 200 and 500 K for 1 dpa of damage reported for the various grades of pure tungsten create limitations on its use, particularly at low irradiation temperatures.59,109,142 Based on irradiation data for Mo alloys, the minimum irradiation temperature which avoids severe radiation embrittlement is >0.3 Tm or 1300 K for tungsten at neutron fluences >0.03 dpa or 1 1021 n cm2 (E > 0.1 MeV),3 which correlates with the irradiation defect recovery data on tungsten compiled by Keys et al.127 Recent work in the development of ultra-fine grained tungsten incorporating TiC additions has shown promising results in reducing the sensitivity to radiation-induced degradation of properties.143,144 The grain size refinement, in the range of 50–200 nm, depending on TiC additions and process, theoretically reduces the effective size of weak grain boundaries that can act as crack initiators. In addition, significant reductions are observed in the density of void formation in the materials relative to pure W at irradiations conducted at 873 K and 2 1020 n cm2, though interstitial loop densities are unchanged. While unirradiated room temperature tensile properties still show brittle fracture behavior, the fracture stress is up to four times higher in the W–TiC samples than in pure W in addition to showing 100 K lower DBTT in impact testing. In microhardness measurements following irradiation, the W–TiC samples exhibited no radiation hardening compared with pure W. The change in Vickers hardness following irradiation for the W–TiC material of Kurishita et al.143 compared to neutron- and proton-irradiated W and W–Re alloys135 irradiated to similar temperatures and doses is shown in Figure 26. The reduced sensitivity of the W–TiC alloy to radiation hardening offers the potential for further development of these alloys for nuclear applications.
200
4.06.6 Outlook 0
0
5.1
9.2
20.9
44-56
Dose (´1020 n cm–2), Tirr = 523-573 K
Figure 25 Ductile-to-brittle transition temperature (DBTT) as a function of neutron fluence (E > 0.1 MeV) of W and its alloys. Reproduced from Krautwasser, P.; Heinz, D.; Kny, E. High Temp. High Press. 1990, 22, 25–32.
The use of refractory metal alloys in radiation environments can offer high-temperature capabilities not matched in other alloy categories. Refractory metal alloys also offer exceptional compatibility with liquid metal coolants. As described in some detail in this
210
Radiation Effects in Refractory Metals and Alloys
250
Change in Vickers Hardness with irradiation HVirrad – HVunirrad (kg mm–2)
He et al.135 Kurishita et al.143 200
150
100
50
0 Dose 0.15 dpa Type Proton Tirr 773 K Material W-pure
0.15 dpa Neutron 873 K W-pure
0.15 dpa Proton 873 K W-3Re
0.15 dpa Neutron 873 K W-3Re
0.15 dpa Neutron 873 K W-5Re
0.08 dpa 0.08 dpa Neutron Neutron 873 K 873 K W-pure W-0.5TiC
Figure 26 Comparison of the increase in Vickers Hardness for tungsten and tungsten alloys for similar dose and irradiation temperatures. Reproduced from He, J. C.; Hasegawa, A.; Abe, K. J. Nucl. Mater. 2008, 377, 348–351; Kurishita, H.; Kobayashi, S.; Nakai, K.; et al. J. Nucl. Mater. 2008, 377, 34–40.
chapter through mechanical property comparisons, these materials are sensitive to impurity contamination during metallurgical processing as well as in-service exposures that can lead to grain boundary embrittlement issues. The inherent irradiation response of bcc-structured materials also limits refractory metal use at temperatures >0.3 Tm, with significant degradation in material properties with displacive irradiation doses as low as 0.03 dpa.3 Improvements in the irradiated mechanical properties of refractory metal alloys have been observed in recent experimental work, even at low irradiation temperatures. This is in part through improved control over impurity levels and also through thermomechanical processing techniques that result in microstructures with reduced sensitivity to radiation embrittlement. This was discussed with reference to LCAC molybdenum,100 where samples irradiated in the stress-relieved condition showed improvement over material in the recrystallized condition up to the recrystallization temperature. Further development of HP-LCAC molybdenum has resulted in higher aspect ratio grain morphologies that led to plain strain conditions in the grain lamellae during deformation.82 In addition, reduced
grain sizes or higher aspect ratios decrease distances to defect sinks, further reducing irradiation sensitivity. While Mo has traditionally been used to study the behavior of W, the microstructural changes and purity control that have been employed for irradiation studies of Mo have not been incorporated into W. The control over precipitate formation in the preirradiated condition appears to result in changes to some physical material properties, specifically, swelling and densification in Nb–1Zr,25,27 that may lead to variations in mechanical properties. An understanding of the effect of preirradiation thermomechanical processing or in-service microstructural changes that occur during irradiation may lead to improved properties or the ability to avoid dangerous embrittlement issues that can occur through precipitate development. This may be of particular interest in Nb and Ta-base alloys that incorporate Zr or Hf additions that react with impurity elements and produce precipitates. Alloying Mo and W with Re results in improved mechanical properties of unirradiated alloys, increased radiation hardening, and radiation-induced embrittlement.62,120 However, much of this work is
Radiation Effects in Refractory Metals and Alloys
on recrystallized, high Re concentration material, the purity of which may not be ideal. The effect that RIS has on the degradation of properties of Mo–Re alloys is a matter of concern. Further work is needed on higher purity, lower Re (5–20 wt% Re) concentration material with reduced grain size, or that with a tailored aspect ratio similar to that of LCAC-Mo. Initial results show improvements to the irradiated properties of Mo and W through the incorporation of either rare earth oxide124 or TiC additions.112,143 These additions aid in restricting grain growth, provide sinks for radiation-induced defects, and act as obstacles to or deflection points for crack propagation. Though these results are preliminary, they outline the need for further examination of incorporating stable dispersion strengthening particles to refractory metal alloys.
References 1. 2.
3. 4. 5. 6.
7. 8. 9. 10. 11.
12. 13.
14.
Pionke, L. J.; Davis, J. W. Technical assessment of niobium alloys data base for fusion reactor applications. McDonnell Douglas Report C00–4247–2, 1979. Cooper, R. H. In Refractory Alloy Technology for Space Nuclear Power Applications; CONF-8308130; Cooper, R. H., Jr, Hoffman, E. E., Eds.; Oak Ridge National Laboratory: Oak Ridge, TN, 1984; pp 14–17. Zinkle, S. J.; Wiffen, F. W. In STAIF 2004, AIP Conference Proceedings; El-Genk, M. S., Ed.; American Institute of Physics: Melville, NY, 2004; Vol. 699, pp 733–740. El-Genk, M. S.; Tournier, J. M. J. Nucl. Mater. 2005, 340, 93–112. Busby, J. T.; Leonard, K. J. J. Miner. Met. Mater. Soc. 2007, 59(4), 20–26. Wojcik, C. C. In High Temperature Silicides and Refractory Alloys, MRS Symposium Proceedings; Briant, C. L., Ed.; Materials Research Society: Pittsburgh, PA, 1994; Vol. 322, pp 519–530. Goldberg, D. C.; Dicker, G.; Worcester, S. A. Nucl. Eng. Des. 1972, 22, 95–123. Delgrosso, E. J.; Carlson, C. E.; Kaminsky, J. J. J. Less-Common Met. 1967, 12, 173–201. Perkins, R. A. In Proceedings of a Symposium on Advanced Compact Reactor Systems; National Academy Press: Washington, DC, 1983; pp 282–325. English, C. In Proceedings of Symposium on Niobium; Stuart, H., Ed.; The Metallurgical Society of AIME: Warrendale, PA, 1984; pp 239–324. Zinkle, S. J.; Wiffen, F. W.; DiStefano, J. R. Historical basis for selection of Nb–1Zr cladding for space reactor applications, TN Report ORNL/LTR/NR-JIMO/04–05; Oak Ridge National Laboratory: Oak Ridge, TN, 2004. Leonard, K. J.; Busby, J. T.; Hoelzer, D. T.; Zinkle, S. J. Met. Mater. Trans. A 2009, 40(4), 838–855. Hahn, G. T.; Gilbert, A.; Jaffee, R. I. In Refractory Metals and Alloys II, Metallurgical Society Conferences, 17; Perlmutter, I., Semchyshen, M., Eds.; Interscience: New York, NY, 1963; pp 191–221. Begley, R. T.; Platte, W. N. Development of niobium-base alloys, WADC-TR-57-34, Part IV; US Air Force-Wright Air Development Command, 1960.
15.
211
Stephens, J. R. Effects of long-term aging on ductility of the columbium alloys C-103, Cb–1Zr, and CB-752 and molybdenum alloy Mo-TZM, NASA TN-D-8095; National Aeronautics and Space Administration-Lewis Research Center, Oct 1975. 16. Igata, N.; Miyahara, K.; Hakomori, K. J. Nucl. Sci. Technol. 1979, 16(1), 73–75. 17. Hautoja¨rvi, P.; Huomo, H.; Saaiaho, P.; Vehanen, A.; Yli-Kauppila, J. J. Phys. F: Met. Phys. 1983, 13, 1415–1427. 18. Naidu, S. V.; Gupta, A.; Sen, P. J. Nucl. Mater. 1987, 148, 86–91. 19. Wiffen, F. W. In Refractory Alloy Technology for Space Nuclear Power Applications, CONF-8308130; Cooper, R. H., Jr, Hoffman, E. E., Eds.; Oak Ridge National Laboratory: Oak Ridge, TN, 1984; pp 252–277. 20. Fischer, C. Study of vacancy clusters in the form of voids using electron microscopy: Neutron irradiated niobium, quenched nickel, Report FRNC-TH-361; Thesis Universitye Scientifique et Medicale de Grenoble, France; English translation BNWL-TR-108; Feb 1974. 21. Loomis, B. A.; Gerber, S. B. J. Nucl. Mater. 1981, 97(1–2), 113–125. 22. Loomis, B. A.; Gerber, S. B. J. Nucl. Mater. 1981, 103–104, 1193–1198. 23. Loomis, B. A.; Gerber, S. B. J. Nucl. Mater. 1983, 17, 224–233. 24. Powell, R. W.; Peterson, D. T.; Zimmerschied, M. K.; Bates, J. F. J. Nucl. Mater. 1981, 103–104, 969–974. 25. Watanabe, H.; Yasunaga, K.; Muroga, T.; Yoshida, N.; Garner, F. A. J. Nucl. Mater. 1996, 233–237, 577–580. 26. Wiffen, F. W. In Defects and Defect Clusters in bcc Metals and Their Alloys, Nuclear Metallurgy 18; Arsenault, R. J., Ed.; National bureau of Standards: Gaithersburg, MD, 1973. 27. Garner, F. A.; Greenwood, L. R.; Edwards, D. J. J. Nucl. Mater. 1994, 212–215, 426–430. 28. Wiffen, F. W. In Radiation-Induced Voids in Metals, CONF-710601; Corbin, J. W., Ianniello, L. C., Eds.; Albany, NY, 1972; pp 386–396. 29. Gold, R. E.; Harrod, D. L. In International Metals Reviews; Hughes, T. L., Ed.; ASM International: Metals Park, OH, 1980; Vol. 25, pp 234–254. 30. Claudson, T. T.; Pessl, H. J. Irradiation effects on high-temperature reactor structural metals, Technical Report BNWL-23; Pacific Northwest Lab: Battelle-Northwest, Richland, WA, 1965. 31. Busby, J. T.; Leonard, K. J.; Zinkle, S. J. Effects of neutron irradiation on refractory metal alloys, ORNL/LTR/NR-PROM1/05–38; Oak Ridge National Laboratory: Oak Ridge, TN, Dec 2005. 32. Stephens, J. R. Metallography 1977, 10, 1–25. 33. Leonard, K. J.; Busby, J. T.; Hoelzer, D. T.; et al. Microstructural characterization of annealed and aged refractory alloys, ORNL/LTR/NR-PROM/05–35; Oak Ridge National Laboratory: Oak Ridge, TN, Nov 2005. 34. Horak, J. A.; Grossbeck, M. L.; Paxton, M. M. In 11th Symposium on Space Nuclear Power and Propulsion, AIP Conference Proceedings 301; El-Genk, M. S., Hoover, M. D., Eds.; American Institute of Physics: Woodbury, NY, 1994; pp 413–420. 35. Wiffen, F. W. In Alloy Development for Irradiation Performance Quarterly Progress Report; DOE/ET-0058/1, Jan–Mar 1978; Oak Ridge National Laboratory: Oak Ridge, TN, 1978; 142–152. 36. Zinkle, S. J.; Matsui, H.; Smith, D. L.; et al. J. Nucl. Mater. 1998, 258–263, 205–214. 37. Wiffen, F. W. In Radiation Effects and Tritium Technology for Fusion Reactors, CONF-750989; Watson, J. S.,
212
38. 39.
40. 41. 42. 43. 44.
45.
46. 47. 48. 49. 50. 51. 52.
53. 54.
55. 56. 57. 58. 59. 60.
61. 62. 63. 64. 65.
Radiation Effects in Refractory Metals and Alloys Wiffen, F. W., Eds.; Gatlinburg, TN, 1976; Vol. II, pp 344–361. English, C. In Proceedings of the International Symposium on Niobium; Stuart, H., Ed.; The Metallurgical Society of AIME: Warrendale, PA, 1981; p 239. Sauges, A. A.; Auet, J. In Radiation Effects and Tritium Technology for Fusion Reactors, CONF-750989; Watson, J. S., Wiffen, F. W., Eds.; Gatlinburg: TN, 1976; Vol. II, pp 331–343. Remark, J. R.; Johnson, A. B.; Farrar, H.; Atteridge, D. G. Nucl. Technol. 1976, 29, 369–377. Leonard, K. J.; Busby, J. T.; Zinkle, S. J. J. Nucl. Mater. 2007, 366(1–2), 353–368. Torti, M. L. Space/Aeronautics 1961, 36, 87–93. Buckman, R. W. In Alloying; Walter, J. L., Jackson, M. R., Sims, C. T., Eds.; ASM International: Metals Park, OH, 1988; pp 419–445. Lessmann, G. G.; Stoner, D. R. Welding refractory metal for space power system applications. In Paper Presented at the 9th National SAMPLE Symposium on Joining of Materials for Aerospace Systems, Dayton, OH, Nov 1965. Buckman, R. W. In Refractory Alloy Technology for Space Nuclear Power Applications, CONF-8308130; Cooper, R. H., Hoffman, E. E., Eds.; Oak Ridge National Laboratory: Oak Ridge, TN, 1984; pp 86–97. Wiffen, F. W. J. Nucl. Mater. 1977, 67, 119–130. Bates, J. F.; Pitner, A. L. Nuclear Technol. 1972, 16, 406–409. Krishan, K. Phil. Mag. A 1982, 45(3), 401–417. Veshchunov, M. S.; Matveev, L. V. Atomic Energy 1994, 76(1), 29–35. Dubinko, V. Nucl. Instr. Meth. Phys. Res. B 2009, 267(18), 2976–3979. Evans, J. Nature 1971, 229(5248), 403–404. Murgatroyd, R. A.; Bell, I. P.; Bland, J. T. In Properties of Reactor Structure Alloys After Neutron or Particle Irradiation, ASTM STP 570; ASTM: Philadelphia, PA, 1975; pp 421–432. Evans, J. H. J. Nucl. Mater. 1980, 88, 31–41. Brown, R. D.; Wechsler, M. S.; Tschalar, C. In Influence of Radiation on Material Properties: 13th International Symposium, ASTM STP 956; Garner, F. A., Henager, C. H., Jr, Igata, N., Eds.; ASTM: Philadelphia, PA, 1987; pp 131–140. Chen, J.; Ullmaier, H.; Floßdorf, T.; et al. J. Nucl. Mater. 2001, 298, 248–254. Byun, T. S.; Maloy, S. A. J. Nucl. Mater. 2008, 377, 72–79. Zinkle, S. J.; Ghoniem, N. M. Fusion Eng. Design 2000, 51–52, 55–71. Ullmaier, H.; Casughi, F. Nucl. Instr. Meth. Phys. Res. B 1995, 101, 406–421. Gorynin, I. V.; Ignatov, V. A.; Rybin, V. V.; et al. J. Nucl. Mater. 1992, 191–194, 421–425. Barklay, C. D.; Kramer, D. P.; Talnagi, J. In Space Technology and Applications Forum-STAIF 2007; El-Genk, M. S., Ed.; American Institute of Physics, Melville NY, 2007; Vol. 880, pp 224–228. Das, S. K.; Kaminsky, M.; Dusza, P. J. Vac. Sci. Technol. 1978, 15, 710. Fabritsiev, S. A.; Pokrovsky, A. S. J. Nucl. Mater. 1998, 252, 216–227. Kisunai, Y.; Kurishita, J.; Kuwabara, T.; et al. J. Nucl. Mater. 2005, 346, 233. Voitsenya, V. S.; Konovalov, V. G.; Becker, M. F.; Motojima, O.; Narihara, K.; Schunke, B. Rev. Sci. Instrum. 1999, 70(4), 2016–2025. Tang, C. J.; Li, R. H.; Chen, J. L. Plasma Sci. Technol. 2008, 10(4), 412–415.
66. 67. 68. 69. 70. 71. 72. 73. 74.
75. 76. 77. 78.
79.
80. 81. 82. 83. 84. 85. 86. 87. 88. 89. 90. 91. 92. 93.
94. 95. 96.
Standard Specification for Molybdenum and Molybdenum Alloy Plate, Sheet and Foil, ASTM B 386-03; ASTM: West Conshohocken, PA, 2003. Geach, G. A.; Hughes, J. E. In Proceedings of the 2nd Plansee Seminar; Benesovsky, F., Ed.; Pergamon: London, 1956; pp 245–253. Jaffee, R. I.; Sims, C. T.; Harwood, J. J. In Proceedings of the 3rd Plansee Seminar; Benesovsky, F., Ed.; Springer-Verlag: Vienna, Austria, 1959; pp 380–411. Sims, C. T.; Jaffee, R. I. Trans. ASM 1960, 52, 929. Stephens, J. R.; Witzke, W. R. J. Less-Common Met. 1972, 29, 371–388. Klopp, W. D. J. Less-Common Met. 1975, 42, 261–278. Davidson, D. L.; Brotzen, F. R. Acta Metall. 1970, 18, 463–470. Lawley, A.; Maddin, R. Trans. Met. Soc. AIME 1962, 224, 573–583. Korotaev, A. D.; Tyumentsev, A. N.; Manako, V. V.; Pinzhin, Y. P. In Rhenium and Rhenium Alloys; Bryskin, B. D., Ed.; TMS: Warrendale, PA, 1997; pp 671–680. Raffo, P. L. J. Less-Common Met. 1969, 17, 133. Wadsworth, J.; Nieh, T. G.; Stephens, J. J. Scripta Metall. 1986, 20, 637–642. Wadsworth, J.; Nieh, T. G.; Stephens, J. J. Int. Met. Rev. 1988, 33, 131–150. Lundberg, L. B.; Ohriner, E. K.; Tuominen, S. M.; Whelan, E. P.; Shields, J. A. In Physical Metallurgy and Technology of Molybdenum and Its Alloys; AMAX Specialty Metals Corporation: Ann Arbor, MI, 1985; pp 71–80. Klopp, W. D.; Witzke, W. R. Mechanical properties of electron-beam-melted molybdenum and dilute molybdenum–rhenium alloys, NASA TM X-2576; NASA Glenn Research Center: Cleveland, OH, 1972. Baranwal, R.; Burke, M. G. Phil. Mag. 2005, 85(4–7), 519–531. Byun, T. S.; Li, M.; Cockeram, B. V.; Snead, L. L. J. Nucl. Mater. 2008, 376, 240–246. Cockeram, B. V.; Smith, R. W.; Leonard, K. J.; Byun, T. S.; Snead, L. L. J. Nucl. Mater. 2008, 382, 1–23. Cockeram, B. V. Met. Trans. A 2002, 33, 3685–3707. Cockeram, B. V. Met. Trans. A 2005, 36, 1777–1791. Cockeram, B. V.; Smith, R. W.; Snead, L. L. J. Nucl. Mater. 2005, 346, 165. Schmunk, R. E.; Kulcinski, G. L. Survey of Irradiation Data on Molybdenum, University of Wisconsin Report, UWFDM-161, Madison, WI, June 1976. Lee, F.; Matolich, J.; Moteff, J. J. Nucl. Mater. 1976, 62, 115–117. Stubbis, J. F.; Moteff, J.; Taylor, A. J. Nucl. Mater. 1981, 101, 64–77. Garner, F. A.; Stubbins, J. F. J. Nucl. Mater. 1994, 212–215, 1298–1302. Brimhall, J. L.; Simonen, E. P.; Kissinger, H. E. J. Nucl. Mater. 1973, 48, 339–350. Sikka, V. K.; Moteff, J. J. Nucl. Mater. 1974, 54, 325–345. Krishan, K. Radiat. Eff. Defects Solids 1982, 66, 121–155. Bentley, J.; Eyre, B. L.; Loretto, M. In Proceedings of International Conference on the Fundamental Aspects of Radiation Damage in Metals, CONF-751006-P2; Gatlinburg, TN 1975; Vol. 2, pp 925–931. Gelles, D. S.; Peterson, D. T.; Bates, J. F. J. Nucl. Mater. 1981, 103–104, 1141–1146. Powell, R. W.; Peterson, D. T.; Zimmerschied, M. K.; Bates, J. F. J. Nucl. Mater. 1981, 103–104, 969–974. Bentley, J.; Wiffen, F. W. In Proceedings of the Second Topical Meeting on the Technology of Controlled Nuclear Fusion, CONF-760935-P1; Energy Research and
Radiation Effects in Refractory Metals and Alloys
97. 98. 99. 100. 101. 102. 103. 104. 105. 106. 107. 108. 109. 110.
111. 112. 113. 114. 115. 116. 117. 118. 119. 120. 121. 122. 123. 124. 125.
Development Administration: Washington, DC, 1976, pp 209–218. Sprague, J. A.; Smidt, F. A.; Reed, J. R. J. Nucl. Mater. 1979, 85–86, 739–743. Zakharova, M. I.; Artemov, N. A.; Bogdanov, V. V. Inorg. Mater. 2001, 37(8), 786–789. Li, M.; Eldrup, M.; Byun, T. S.; Hashimoto, N.; Snead, L. L.; Zinkle, S. J. J. Nucl. Mater. 2008, 376, 11–28. Cockeram, B. V.; Hollenbeck, J. L.; Snead, L. L. J. Nucl. Mater. 2005, 336, 299–313. Singh, B. N.; Evans, J. H.; Horsewell, A.; Toft, P.; Mu¨ller, G. V. J. Nucl. Mater. 1998, 258–263, 865–872. Li, M.; Byun, T. S.; Snead, L. L.; Zinkle, S. J. J. Nucl. Mater. 2008, 377, 409–411. Cox, B. L.; Wiffen, F. W. J. Nucl. Mater. 1979, 85–86, 901–905. Abe, K.; Takeuchi, T.; Kikuchi, M.; Morozumi, S. J. Nucl. Mater. 1981, 99, 25–37. Abe, K.; Kikuchi, M.; Tate, K.; Morozumi, S. J. Nucl. Mater. 1984, 122–123, 671–675. Hasegawa, H.; Abe, K.; Satou, M.; Ueda, K. J. Nucl. Mater. 1996, 233–237, 565–569. Cockeram, B. V.; Smith, R. W.; Byun, T. S.; Snead, L. L. J. Nucl. Mater. 2009, 393, 12–21. Chakin, V.; Kazakov, V. J. Nucl. Mater. 1996, 233–237, 570–572. Scibetta, M.; Chaouadi, R.; Puzzolante, J. L. J. Nucl. Mater. 2000, 283–287, 455–460. Kasakov, V. A.; Kolesnikov, A. N.; Krassnoselov, V. A.; et al. Effect of neutron irradiation on properties of potential structural material for thermonuclear reactors, USSR-US Exchange on CTR Materials, Nov 1974. Steichen, J. M. J. Nucl. Mater. 1976, 60, 13–19. Kitsunai, Y.; Kurishita, H.; Narui, M.; Kayano, H.; Hiraoka, Y. J. Nucl. Mater. 1996, 239, 253–260. Massalski, T. B. Ed. Binary Alloy Phase Diagrams; ASM International: Metals Park, OH, 1986; Vol. 2. Carlen, J. C.; Bryskin, B. D. J. Mater. Eng. Perform. 1994, 3(2), 282–291. Leonard, K. J.; Busby, J. T.; Zinkle, S. J. J. Nucl. Mater. 2007, 366, 369–387. Nemoto, Y.; Hasegawa, A.; Satou, M.; Abe, K.; Hiraoka, Y. J. Nucl. Mater. 2004, 324, 62–70. Hasegawa, A.; Ueda, K.; Satou, M.; Abe, K. J. Nucl. Mater. 1998, 258–263, 902–906. Sigh, B. N.; Evans, J. H.; Horsewell, A.; Toft, P.; Edwards, D. J. J. Nucl. Mater. 1995, 223, 95–102. Hasegawa, A.; Abe, K.; Satou, M.; Namba, C. J. Nucl. Mater. 1995, 225, 259–266. Busby, J. T.; Leonard, K. J.; Zinkle, S. J. J. Nucl. Mater. 2007, 366, 388–406. Edwards, D. J.; Garner, F. A.; Gelles, D. S. J. Nucl. Mater. 2008, 375, 370–381. Erck, R. A.; Rehn, L. E. Phil. Mag. A 1990, 62(1), 29–51. Erck, R. A.; Rehn, L. E. J. Nucl. Mater. 1989, 168, 208–219. Cockeram, B. V.; Smith, R. W.; Snead, L. L. J. Nucl. Mater. 2005, 346, 165–184. Garfinkle, M.; Witzke, W. R.; Klopp, W. D. Trans. Metall. Soc. AIME 1969, 245(2), 303–309.
126.
213
Mutoh, Y.; Ichikawa, K.; Nagata, K.; Takeuchi, M. J. Mater. Sci. 1995, 30, 770–775. 127. Keys, L. K.; Smith, J. P.; Moteff, J. Phys. Rev. 1968, 176(3), 851–856. 128. Keys, L. K.; Moteff, J. J. Nucl. Mater. 1970, 34, 260–280. 129. Eyre, B. L. J. Phys. F: Metal Phys. 1973, 3(2), 422–470. 130. Lacefield, K.; Moteff, J.; Smith, J. P. Phil. Mag. 1966, 13(125), 1079–1081. 131. Matolich, J.; Nahm, N.; Moteff, J. Scripta Met. 1974, 8, 837–842. 132. Williams, R. K.; Wiffen, F. W.; Bentley, J.; Stiegler, J. O. Met. Trans. A 1983, 14, 655–666. 133. Rau, R. C.; Ladd, R. L.; Moteff, J. J. Nucl. Mater. 1969, 33, 324–327. 134. Sikka, V. K.; Moteff, J. J. Appl. Phys. 1972, 43(12), 4942–4944. 135. He, J. C.; Hasegawa, A.; Abe, K. J. Nucl. Mater. 2008, 377, 348–351. 136. Heckelman, J. D.; Kozar, R. P. Measured drift of irradiated and unirradiated W3%Re/W25%Re thermocouples at a nominal 2000K, NASA Technical Report, NASA TM X-67818; Lewis Research Center: Cleveland, OH, 1971. 137. Vitanza, C.; Stein, T. E. J. Nucl. Mater. 1986, 139, 11–18. 138. Tyler, W. W. Electron theory of thermoelectric effects, Report number KAPL-M-WWT-1; Knolls Atomic Power Laboratory: Schenectady, NY, 1951. 139. Coobs, J. H.; Kasten, P. R. High-temperature gas-cooled reactor base-technology program progress report, for date ending June 30 1975, Report number ORNL-5108; Oak Ridge National Laboratory: Oak Ridge, TN. 140. Van Nieuwenhove, R.; Vermeeren, L. Rev. Sci. Instrum. 2004, 75(1), 75–82. 141. Nemoto, Y.; Hasegawa, A.; Satou, M.; Abe, K. J. Nucl. Mater. 2000, 283–287, 1144–1147. 142. Krautwasser, P.; Heinz, D.; Kny, E. High Temp. High Press. 1990, 22, 25–32. 143. Kurishita, H.; Kobayashi, S.; Nakai, K.; et al. J. Nucl. Mater. 2008, 377, 34–40. 144. Kitsunai, Y.; Kurishita, H.; Kayano, H.; Hiraoka, Y.; Igarashi, T.; Takida, T. J. Nucl. Mater. 1999, 271–272, 423–428. 145. Jang, H.; Moteff, J. The influence of neutron irradiation temperature on the void characteristics of niobium and niobium–1% zirconium alloys. In Radiation Effects and Tritium Technology for Fusion Reactors, CONF-750989, Gatlinburg, TN; Watson J. S., Wiffen, F. W., Eds.; 1976; pp 1106–1121. 146. Sprague, J. A.; Smith, F. A. Jr.; Reed, J. R. J. Nucl. Mater. 1979, 85–86, 739–743. 147. Michel, D. J.; Smith, H. H. In Effects of Radiation on Structural Materials, ASTM-STP-683; Sprague, J. A., Kramer, D., Eds.; ASTM: Philadelphia, PA, 1979; pp 107–124. 148. Webster, T. H.; Eyre, B. L.; Terry, E. A. In Proceedings of BNES Conference on Irradiation Embrittlement and Creep in Fuel Cladding and Core Components; British Nuclear Energy Society: London, 1972; pp 61–80. 149. Smith, H. H.; Michel, D. J. J. Nucl. Mater. 1977, 66, 125–142.
4.07
Radiation Effects in SiC and SiC–SiC
L. L. Snead and Y. Katoh Oak Ridge National Laboratory, Oak Ridge, TN, USA
T. Nozawa Japan Atomic Energy Agency, Rokkasho, Aomori, Japan
ß 2012 Elsevier Ltd. All rights reserved.
4.07.1 4.07.2 4.07.3 4.07.4 4.07.4.1 4.07.4.2 4.07.4.3 4.07.4.4 4.07.5 4.07.6 References
Introduction Irradiation-Induced Swelling and Microstructure of Pure SiC Irradiation-Induced Thermal Conductivity Degradation of Monolithic SiC Effect of Irradiation on the Mechanical Properties of Monolithic SiC Elastic Modulus of Monolithic SiC Hardness of Monolithic SiC Fracture Toughness of Monolithic SiC Strength and Statistical Variation in Strength for Monolithic SiC Irradiation Creep of SiC Silicon Carbide Composites Under Irradiation
Abbreviations ATR BSD BSR CVD CVI dpa DuET ETR FB fcc HFBR HFIR HFIR-METS HNLS HP JMTR Kd Kgb Kirr Knonirr Krd Ku PLS PS
Advanced test reactor Black spot dot Bend stress relaxation Chemical vapor deposition Chemical vapor infiltration Displacement per atom Dual-beam facility for energy science and technology Engineering test reactor Fluidized bed Face-centered cubic High-flux beam reactor High Flux Isotope Reactor High-flux isotope reactor – mapping elevated temperature swelling Hi Nicalon Type S Hot-pressing Japan materials testing reactor Thermal conductivity by defect scattering Thermal conductivity by grain boundary scattering Irradiated thermal conductivity Nonirradiated thermal conductivity Thermal conductivity by radiation Thermal conductivity by Umklapp scattering Proportional limit stress Pressureless sintering
PW PyC SEM SiC/SiC composite SW TEM Tirr TRISO TySA UTS
215 216 221 224 224 226 226 227 233 234 239
Plain weave Pyrolytic carbon Scanning electron microscopy Silicon carbide fiber reinforced silicon carbide matrix composite Satin weave Transmission electron microscopy Irradiation temperature TRIstructural ISOtropic Tyranno SA Ultimate tensile stress
4.07.1 Introduction Silicon carbide (SiC) has been studied and utilized in nuclear systems for decades. Its primary use was, and still is, as the micro pressure vessel for hightemperature gas-cooled reactor fuels. For these so-called TRI-ISOtropic (TRISO) fuels, the SiC is deposited via a gas-phase decomposition process over two layers of pyrolytic graphite surrounding the fuel kernel. In addition to being strong enough to withstand the pressure buildup from the fission product gas liberated, this SiC layer must also withstand chemical attack from metallic fission products such as palladium and the mechanical loads derived from irradiation-induced dimensional changes occurring in the pyrolytic graphite. More recent nuclear applications of SiC include its use as structural composites 215
216
Radiation Effects in SiC and SiC–SiC
support of nuclear fuel coating1–9 and more recently, for various nuclear applications such as structural SiC composites.10 Before proceeding, it is important to distinguish neutron-induced effects on high-purity materials, such as single crystal and most forms of chemical vapor deposited (CVD) SiC, from those on lower purity forms such as sintered with additives, reaction-bonded, or polymer-derived SiC. It is well understood that the presence of significant second phases and/or poorly crystallized phases in these materials leads to unstable behavior under neutron irradiation,11–14 as compared to stoichiometric materials, which exhibit remarkable radiation tolerance. Discussion and data for this section refer only to high purity, stoichiometric, near-theoretical density SiC, unless otherwise specified. Rohm and Haas (currently Dow Chemicals) CVD SiC is an example of such material. The irradiation-induced microstructural evolution of CVD SiC is roughly understood and has been reviewed recently by Katoh et al.15 An updated version of the microstructural evolution map is shown in Figure 1. However, the contribution of the defects themselves to the swelling in SiC is less understood. Below several hundred Kelvin, the observable
(i.e., SiC/SiC) for high-temperature gas-cooled reactors and for fusion power systems. The possibility of using composite and monolithic SiC thermal insulators for both fusion and fission systems is also being investigated. Moreover, both monolithic and composite forms of SiC are being investigated for use in advanced sodium fast, advanced liquid salt-cooled, and advanced light water reactors. In this chapter, the effects of neutron irradiation on relatively pure, radiation resistant forms of SiC are discussed. This chapter has been limited to the effects of irradiation on the microstructure, and the mechanical and thermal properties of SiC, although it is recognized that environment aspects such as oxidation and corrosion will also be factors in eventual nuclear application. These areas are not discussed here.
4.07.2 Irradiation-Induced Swelling and Microstructure of Pure SiC The neutron-induced swelling of SiC has been well studied for low and intermediate temperatures (293– 1273 K). Originally, this material was investigated in 1600
1400
6
7
7
6 6
6 7
Irradiation temperature (⬚C)
1200
6
1000
800
7
3
Black spot defects(BSD) and/or unidentified small loops
6
6
BSD and/or unidentified small loops
7
Frank loops
Frank loops
6 1
Unfaulted loops and/or network Voids
1
1
6
6
4
600
6
6 1
1
6 3
Larger loops dislocation network voids
1
1
7
5
6
5
2
1 6
1
6
2 2
(1973)4
400
200 0.1
1. Price 2. Yano (1998)17 3. Senor (2003)18 4. Iseki (1990)19 5. Katoh, neutron (2006)15 6. Katoh, ion (2006)15 7. Snead (2007)16
1
2
5
10 Fluence (dpa)
100
1000
Figure 1 Updated summary of the microstructural development in cubic SiC during neutron and self-ion irradiation. Reproduced from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377.
Radiation Effects in SiC and SiC–SiC
Tirr = 300 °C, 6 dpa
Tirr = 800 °C, 7.7 dpa Loop number density » 3.3e + 23 m-3 Mean loop diameter = 3.0 nm
Dot number density » 2.2e + 24 m-3 Mean dot diameter = 1 nm
g = 200
217
40 nm g = 200
50 nm
Figure 2 Microstructure for CVD neutron irradiated at 573 and 1073 K.
microstructure of neutron-irradiated SiC is described as containing ‘black spots, which are most likely tiny clusters of self-interstitial atoms in various indeterminate configurations. For irradiation temperatures less than about 423 K, accumulation of strain due to the irradiation-produced defects can exceed a critical level above which the crystal becomes amorphous. This has been shown in the case of both self-ion irradiation and fast neutron irradiation.20–22 As shown by Katoh et al.,23 the swelling at 323 K under self-ion irradiation increases logarithmically with dose until amorphization occurs. The swelling of neutron- and ion-amorphized SiC has been reported to be 10.8% for 343 K irradiation.22 However, there is evidence that the density of amorphous SiC will depend on the conditions of irradiation (dose, temperature, etc.)24 For temperatures above the critical amorphization temperature (423 K), the swelling increases logarithmically with the dose until it approaches saturation, with a steady decrease in the saturation swelling level with increasing irradiation temperature. The dose exponents of swelling during the logarithmical period are in many cases close to twothirds, as predicted by a kinetic model assuming planar geometry for interstitial clusters.25 This temperature regime is generally referred to as the pointdefect swelling regime and can be roughly set between 423 and 1273 K. As an example of how these ‘black spot’ defects mature in the point-defect swelling regime, Figure 2 shows neutron-irradiated microstructures at 573 and 1073 K for doses
consistent with a saturation in density. While these microstructural features are generically classified as ‘black spots,’ the defects formed at 1073 K are clearly coarser compared to those formed under 573 K irradiation. The approach to saturation swelling is shown for High Flux Isotope Reactor (HFIR) neutron irradiated Rohm and Haas CVD SiC in Figure 3. In this figure, the swelling is depicted in both logarithmic (Figure 3(a)) and linear (Figure 3(b)) plots. In addition to the approach to saturation, this figure highlights two other characteristics of neutroninduced swelling of SiC. First, the swelling of SiC is highly temperature dependent. For the data given in Figure 3, the 1 dpa and saturation values of swelling at 473 K are approximately five times that for 1073 K irradiation. This reduced swelling with increasing irradiation temperature is primarily attributed to enhanced recombination of cascadeproduced Frenkel defects due to lower interstitial clustering density at higher temperatures. The second characteristic swelling behavior to note is that the swelling saturates at a relatively low dose. For damage levels of a few dpa (typically months in a fission power core), the swelling in the point-defect recombination range has found its saturation value. At higher temperatures such as 1173–1673 K,4,18,26 Frank faulted loops of the interstitial type become the dominant defects observed by transmission electron microscopy (TEM). It has also been reported that Frank faulted loops appear for lower temperature neutron irradiation at extremely high doses.27
218
Radiation Effects in SiC and SiC–SiC
3 CVD SiC
CVD SiC 2
200 °C
1
400 °C
2.5
200 °C 600 °C
Swelling (%)
Swelling (%)
2
650 °C
300 °C
1.5 400 °C
800 °C
1 600 °C
800 °C
0.1
0.5
0 0 (a)
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3 4 5 6 Neutron dose (dpa)
7
8
0.01 (b)
0.1 1 Neutron dose (dpa)
Tirr = 200 °C
Tirr = 300 °C
Tirr = 400 °C
Tirr = 500 °C
Tirr = 600 °C
Tirr = 800 °C
10
30
Figure 3 Swelling of SiC in the intermediate temperature point defect swelling regime. Reproduced from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377.
Under silicon ion irradiation at 1673 K, the development of Frank loops into dislocation networks through unfaulting reactions at high doses is reported.26 The volume associated with dislocation loops in irradiated SiC has been estimated to be on the order of 0.1%.28 At temperatures where vacancies are sufficiently mobile, vacancy clusters can be formed. Threedimensional (3D) cavities (or voids) are the only vacancy clusters known to commonly develop to large sizes in irradiated SiC. The lowest temperature at which void formation was previously reported under neutron irradiation is 1523 K.4 Senor reported the lack of void production after neutron irradiation to 0.9 dpa at 1373 K, although voids were observed after subsequent annealing at 1773 K for 1 h.18 Under silicon ion irradiation, voids start to form at 1273 K at very low density and become major contributors to swelling at irradiation conditions of 1673 K at >10 dpa.29 Positron annihilation and electron paramagnetic resonance studies have shown that the silicon vacancy in cubic SiC becomes mobile at 1073–1173 K.30,31 Therefore, it would not be surprising for void swelling to take place at as low as 1273 K at high doses, particularly for low damage rate irradiations.
As previously mentioned, data on swelling of SiC in the high-temperature ‘void swelling’ regime has been somewhat limited. Recently, however, work has been carried out in the 1173–1773 K range for Rohm and Haas CVD SiC irradiated in HFIR. Of particular significance to that experiment is the confidence in irradiation temperature owing to the melt-wire passive thermometry.32 Recent TEM imaging by Kondo28 clearly shows the evolution of complex defects. As an example, Figure 4 indicates sparse void formation on stacking faults for material irradiated at 1403 K. Significant growth of voids commences at 1723 K. The well-faceted voids appeared to be tetrahedrally bounded by planes, which likely provide the lowest surface energy in cubic SiC. In many cases, voids appeared to be aligned on stacking faults at all temperatures. However, intragranular voids unattached to stacking faults were also commonly observed at 1723 K. The evolution of dislocation microstructures at 1403–1723 K is shown in Figure 5. In this temperature range, dislocation loops are identified to be Frank faulted loops of interstitial type. Evolution of the dislocation loops into dislocation networks was confirmed for irradiation at 1723 K.
Radiation Effects in SiC and SiC–SiC
20 nm
(a)
1280 ⬚C, 5.0 dpa
1130 ⬚C, 8.5 dpa
1130 ⬚C, 1.8 dpa
(b)
(c) 1450 ⬚C, 8.5 dpa
1450 ⬚C, 5.0 dpa
1450 ⬚C, 1.8 dpa
(d)
219
(e)
(f)
Figure 4 Evolution of voids in high-temperature irradiated CVD SiC.
1130 ⬚C, 8.5 dpa
1130 ⬚C, 1.8 dpa
30 nm
(a)
(b)
(c) 1450 ⬚C, 5.0 dpa
1450 ⬚C, 1.8 dpa
(d)
1280 ⬚C, 5.0 dpa
g
(e)
1450 ⬚C, 8.5 dpa
g
(f)
g
Figure 5 Evolution in dislocation networks for high-temperature irradiated CVD SiC.
Figure 6 plots both historical data, recently published, and unpublished data on the swelling behavior of SiC over a wider range of temperature.16,33 This plot is limited to literature data on high-purity CVD SiC. A divergence from point-defect ‘saturated’ swelling to unsaturated swelling is observed in the 1273–1473 K range, although additional data in this temperature range as a function of fluence would be required to precisely define such behavior. Above 1373 K, there exists a clear unsaturated swelling behavior for CVD SiC. The three divergent curves
drawn in Figure 6 represent data taken at nominally 1.75, 5.0, and 8.5 1025 n m2 (E > 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.
Radiation Effects in SiC and SiC–SiC
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 (<423 K), the saturable point-defect swelling regime (423–1073 K) range, and the unsaturated void swelling regime,
221
which occurs for irradiation temperature >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
Radiation Effects in SiC and SiC–SiC
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 <10 µm
Price (1973)
FB
Grain size <5 µm
Price (1969)3
FB
Grain size <5 µm
Shaffer (1991)51
PS
Li et al. (1998)37
-
Thermal conductivity (W m–1 K–1)
Collins et al. (1990) Collins et al. (1990) 34
300
Calculated value
PS: pressureless sintering, HP: hot-pressing CVD: chemically vapor deposition, FB: fluidized bed N/R: not reported Poly-crystal, large grains
200
Kth = [–0.0003 + 1.05⫻10–5 T]–1 Poly-crystal, small grains Porous poly-crystal, small grains
100
0 0
500
1000
1500
2000
Temperature (K) Figure 7 Thermal conductivity of various forms of SiC as a function of temperature. Reproduced from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377.
and grain boundary scattering, the temperaturedependent phonon scattering cannot be altered and tends to dominate at high temperature (above about 673 K for SiC). The effect of irradiation on SiC in the temperature range of 423–1073 K (the point defect regime) is to produce simple defects and defect clusters that very effectively scatter phonons. For ceramics possessing high thermal conductivity, the irradiationinduced defect scattering quickly dominates, with saturation thermal conductivity typically achieved by a few dpa. Moreover, as the irradiation-induced defect scattering exceeds the phonon–phonon scattering, the temperature dependence of thermal conductivity is much reduced or effectively eliminated. The rapid decrease as well as saturation in thermal conductivity of CVD SiC upon irradiation in the point-defect regime has been reported by several authors.8,34,45,52,53 Figure 8 shows this rapid decrease in thermal conductivity for fully dense CVD SiC, including new data, previous data from the authors,52,53 and that of Rohde.45 It is noted that the data of Thorne is omitted as the material was of exceptionally low density for a CVD SiC material. Moreover, the data of Price34 is published with a range of fluence that is not valuable in the figure.
In recent papers by Snead on the effects of neutron irradiation on the thermal conductivity of ceramics,53 and specifically on SiC,16,52 the degradation in thermal conductivity has been analyzed in terms of the added thermal resistance caused by the neutron irradiation. The thermal defect resistance is defined as the difference between the reciprocals of the irradiated and nonirradiated thermal conductivity (1/Krd ¼ 1/Kirr – 1/Knonirr). This term can be related directly to the defect type and concentration present in irradiated ceramics.53 Moreover, this term can be used as a tool to compare the thermal conductivity degradation under irradiation of various ceramics or, for example, various forms of SiC. It has been shown that, for certain high purity forms of alumina, the accumulation of thermal defect resistance is very similar even though the starting thermal conductivities of the materials are substantially different. Similarly, CVD SiC was shown to have a similar accumulation of thermal defect resistance as a hot-pressed form of SiC with substantially lower (90 W m1 K1) unirradiated thermal conductivity. The utility of this finding is that if the thermal defect resistance is mapped as a function of irradiation temperature and dose for a form of high-purity CVD SiC, it can be applied to determine the thermal
Radiation Effects in SiC and SiC–SiC
Room temperature thermal conductivity (W m−1 K–1)
300
223
Nonirradiated thermal conductivity 381 ± 26 W m−1 K−1 (except Rohde, 108 W m−1 K−1)
Tirr = 200 °C Tirr = 800 °C
100
50 Tirr = 80 °C
30 Tirr = 750-850 °C Tirr = 600 °C Tirr = 480-550 °C Tirr = 400-480 °C
10
Tirr = 375 °C Tirr = 200 °C
CVD SiC 5 0.001
0.01
0.1
1
10
Fast neutron dose (´ 1025 n m-2 E > 0.1 MeV) Figure 8 Degradation in room-temperature thermal conductivity for CVD SiC. (Rohde data designated as .) Reproduced from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377.
conductivity of any high-purity CVD SiC, independent of the starting thermal conductivity. The accumulation in thermal defect resistance generated from the data of Figure 8 is shown in Figure 9. Another result of the previously reported analysis on irradiated CVD SiC16,52 is that the thermal defect resistance appears to be directly proportional to the irradiation-induced swelling, although the data-set for making the previous assertion was somewhat limited. A compilation plot including the previous dataset as well as the new data of Figure 9 is shown in Figure 10. It is clear from this plot that a linear relationship exists between swelling and thermal defect resistance. Moreover, there does not appear to be any effect of irradiation temperature on this result. The fact that the thermal defect resistance is proportional to the irradiation-induced swelling allows a rough estimate of thermal conductivity. As measurement of thermal conductivity for the SiC TRISO shell is not practical, while measurement of density is routine, this finding allows an indirect determination of thermal conductivity by measurement of the density change in the TRISO SiC shell by means of a density gradient column or some other technique.
The thermal conductivity degradation discussed up to this point has been for irradiation temperature associated with the point defect regime. For irradiation above this temperature (the nonsaturable void swelling regime), the thermal properties are not expected to saturate (at least at low dpa). The primary reason for this is that the formation of voids and other complex defects in the high-temperature regime (which contributes to the unsaturated swelling as seen in Figure 6) contributes to phonon scattering, and these defects will not saturate. Moreover, it has been shown that the linear relationship that existed between swelling and thermal defect resistance (as seen in Figure 10) does not exist in this elevated temperature irradiation regime.16,52 This underlines the fact that the phonon scattering and swelling are not controlled by the same defects in the lower temperature ‘saturable,’ and elevated temperature ‘nonsaturable’ irradiation regimes. A compilation plot of room-temperature thermal conductivity as a function of irradiation temperature for the saturable and nonsaturable temperature regimes is given in Figure 11. By comparison to the unirradiated roomtemperature conductivity value of 280 W m1 K1,
224
Radiation Effects in SiC and SiC–SiC
Thermal defect resistance (m K W–1)
Tirr = 80 °C
Tirr = 200 °C Tirr = 375 °C
0.1
Tirr = 400-480 °C Tirr = 480-550 °C Tirr = 600 °C
Rhode, Tirr = 80 °C Tirr = 750-850 °C
0.01
Tirr = 80 °C Tirr = 200 °C Tirr = 375 °C Tirr = 400-480 °C Tirr = 480-550 °C Tirr = 600 °C Tirr = 750-800 °C
0.001 0.001
0.01
0.1
1
10
Fast neutron dose (´ 1025 n m-2 E > 0.1 MeV) Figure 9 Thermal defect resistance for stoichiometric CVD SiC as a function of neutron dose. (Rohde data designated as .) Reproduced from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377.
it is clear that the thermal conductivity degradation in the highest temperature regimes is less dramatic, even though the swelling is rapidly increasing (see Figure 6). This is opposite to the behavior in the lower temperature, saturable regime, where high swelling corresponds to extreme reduction in thermal conductivity. Unfortunately, the data on thermal conductivity reduction in the nonsaturable regime is limited, and given the lack of knowledge of the specific defects governing the phonon scattering, it is not possible to accurately predict behavior outside of the data-set of Figure 6. Data presented thus far has been limited to measurement of thermal conductivity at room temperature. As described in Figure 7, there is a dramatic dependence of thermal conductivity on measurement temperature. The temperature dependence of irradiated materials can be found by applying the temperature dependence of unirradiated SiC (the Umklapp thermal resistance term) to the as-neutron-degraded room-temperature values. This approximation (dashed lines) is compared to actual
data (solid lines) in Figure 12 and shows fair correspondence. However, it is clear that such a treatment systematically underestimated the thermal conductivity degradation. This implies temperature dependence on the defect scattering that is not presently understood.
4.07.4 Effect of Irradiation on the Mechanical Properties of Monolithic SiC 4.07.4.1
Elastic Modulus of Monolithic SiC
Figure 13 summarizes the irradiation temperature dependence of the elastic modulus change. Irradiation generally reduces modulus to a greater extent for lower temperature irradiation. The modulus reduction becomes negligible when irradiation temperature reaches or exceeds 1273 K. There seems to be an indistinct stage between 1073 and 1273 K. As expected, the elastic modulus trends with ‘point defect swelling’ of SiC. Point defect swelling is an isotropic volume
Radiation Effects in SiC and SiC–SiC
Tirr = 200 °C
Tirr = 400-480 °C
Tirr = 300 °C
Tirr = 480-550 °C X Tirr = 80 °C
Tirr = 375 °C
Tirr = 600 °C
225
Tirr = 750-850 °C
0.12
X
100
0.1 Open symbols Right 0.08
0.06
0.04
Thermal defect resistance (m K W-1)
Room temperature thermal conductivity (W m-1 K–1)
0.14
10 0.02
Closed symbols Left X
0 0
0.5
1
1.5
2
2.5
Swelling (%) Figure 10 The room-temperature thermal conductivity and thermal defect resistance as a function of irradiation-induced density change. Reproduced from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377.
expansion that is believed to occur by lattice relaxation due to accumulated isolated point defects and small point defect clusters during irradiation at temperatures where vacancies are not readily mobile. In SiC irradiated in the point defect swelling regime, a fairly good agreement between dimensional expansion and lattice spacing has been confirmed by X-ray diffractometry studies. In contrast, the data in the nonsaturable swelling regime is somewhat limited, although the data suggest that there is little reduction in elastic modulus in spite of the swelling being relatively large. However, in this regime, the defects responsible for swelling are voids and other relatively large defects, which would have less of an effect on elastic modulus as compared to point defects. An estimation of the influence of lattice relaxation on elastic modulus was attempted using the Tersoff potential.54 The result predicted a linear lattice
swelling of 1% causing approximately 10% reduction in elastic modulus (Figure 14). The predicted elastic modulus change could be varied by more than 10% depending on a selection of interatomic potential, with the Tersoff potential giving a relatively high sensitivity of modulus to the interatomic distance among various empirical interatomic potential functions. Therefore, the measured elastic modulus changes observed in this experiment are generally greater than the theoretical prediction. It is noted that the methods applied for generating the data of Figure 14 are various and of differing quality. Typically, elastic modulus as measured by nanoindentation, which sometimes is the only alternative for miniature specimens, tends to give widely scattered and less reliable data than the mechanical or sonic modulus methods. Nonetheless, it is clear that the lattice expansion is a major cause of the irradiation-induced elastic modulus reduction in SiC.
226
Radiation Effects in SiC and SiC–SiC
Amorphization regime
Saturable regime
Unsaturable regime
120 Nonirradiated conductivity ~280 W m-1 K-1 Rohm and Haas CVD SiC
Room temperature thermal conductivity (W m-1 K-1)
100 ~1.75 dpa 80
~5 dpa
60
~8.5 dpa 40
20
0 0
200
400
600 800 1000 Irradiation temperature (°C)
1200
1400
1600
Figure 11 Room-temperature thermal conductivity in the saturable and nonsaturable regime. Reproduced from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377.
4.07.4.2
Hardness of Monolithic SiC
The irradiation effect on nanoindentation hardness of Rohm and Haas CVD SiC in a fluence range of 0.1–18.7 dpa is summarized in Figure 15. It is interesting to note that the nanoindentation hardness exhibits relatively small scatter for the individual experiments, and the trend in data as a function of temperature is uniform. This observation is in contrast to both the flexural strength and the indentation fracture toughness data, which indicate a broad peak at an intermediate temperature and a relatively large scatter. It is worth noting that nanoindentation hardness of brittle ceramics is, in general, determined primarily by the dynamic crack extension resistance in the near surface bulk material, and therefore should be more relevant to fracture toughness than to plastic deformation resistance. However, surface effects of the original sample affect the nanoindentation hardness less, as the samples are generally polished prior to testing. 4.07.4.3 Fracture Toughness of Monolithic SiC The effect of irradiation on the fracture toughness of Rohm and Haas CVD SiC is summarized in Figure 16. This compilation plots data using the Chevron notched beam technique, although the bulk of the
data sets report Vicker’s or nanoindentation generated data.55–57 The general trend is that the irradiation-induced toughening seems to be significant at 573–1273 K for the indentation fracture toughness data, in spite of the decrease in elastic modulus, which confirms the increase in fracture energy caused by irradiation. The scatter of the indentation fracture toughness data among different experiments is likely caused by both the condition of the sample surface and the lack of standardized experimental procedures. Typically, indentation should be applied on the polished surfaces, but conditions of polishing are not always provided in literature. Moreover, the crack length measurements are done using optical microscopy, conventional scanning electron microscopy (SEM), or field emission SEM, all of which may give very different crack visibility. In addition, a few different models have been used for derivation of the fracture toughness. In conclusion, indentation fracture toughness techniques can be used only for qualitative comparison within a consistent set of experiments. It is noted that the experiment employing the Chevron notched beam technique also indicates the irradiation-induced toughening, although scatters of toughness values were even greater. These results lead to the conclusion that, in the intermediate irradiation temperature range, the increase of the fracture toughness of SiC exists.
Radiation Effects in SiC and SiC–SiC
227
120 500
Rohm and Haas CVD SiC irradiated in HFIR
Rohm and Haas CVD SiC irradiated in HFIR at 800 ˚C
400 300
100 Tirr = 1500 ⬚C
100
Thermal conductivity (W m-1 K–1)
Thermal conductivity (W m-1 K–1)
200
Nonirradiated
80 60 50
0.05 dpa ++
40
+ + +
+
30
+
0.5 dpa +
1.94 dpa 4.3 dpa
20 Calculated based on Umklapp scattering
1.75 dpa 80
60
Tirr = 1500 ⬚C 7.5 dpa
40
20
Calculation based on Umklapp scattering
Fit to data
Fit to data
10 0
(a)
200 400 600 800 1000 Irradiation and measurement temperature ( ˚C)
0 0
400 500 600 700 100 200 300 Irradiation and measurement temperature ( ⬚C)
(b)
800
40 Rohm and Haas CVD SiC irradiated in HFIR 35
Thermal conductivity (W m-1 K–1)
T = 1020 ⬚C irr 1.75 dpa
30
+ Tirr = 1050 ⬚C
25
5 dpa
+
+ + +
+
1.75 dpa
+
+
5 dpa
+ + +
Tirr = 1060 ⬚C
20
+
7.5 dpa
8.5 dpa
15
10
Calculation based on Umklapp scattering Fit to data
5
0 (c)
0
100 200 300 400 500 600 700 800 Irradiation and measurement temperature ( ˚C)
Figure 12 Effect of temperature on the conductivity of irradiated SiC. (a) Tirr ¼ 1073 K, (b) Tirr ¼ 1773 K, and (c) Tirr ¼ 1293–1333 K. Reproduced from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377.
4.07.4.4 Strength and Statistical Variation in Strength for Monolithic SiC There have been several studies on the effect of neutron irradiation on the strength of various types of SiC forms including reaction-bonded, sintered, pressureless sintered, and CVD SiC materials.1,11,13,14,58,60–65
The strength of SiC depends significantly on stoichiometry under neutron irradiation. Both the sintered SiC and the reaction-bonded SiC forms exhibit significant deterioration in strength by neutron irradiation (Figure 17).13 The presence of impurities such as sintering additives for sintered SiC and excess
228
Radiation Effects in SiC and SiC–SiC
1.20
Relative Young’s modulus
1.10
1.00
0.90
0.80
0.70
0.60 200
Osborne (1999)55,HFIR, 2 dpa Nogami (2002)56,HFIR + HFBR, 0.15–7.7 dpa Park (2003)57,DuET 5.1 MeV Si, 3 dpa Katoh (2005)58,HFIR – 14J, 6.0–7.7 dpa Snead (2007)16, HFIR – METS, 0.7–8.6 dpa Price(1977)59,ETR, 2.8–12.2 dpa Snead (2007)16, HFIR, 0.7–4.2 dpa
400
600
Nanoindentation
4pt. bend Sonic resonance
800 1000 1200 1400 Irradiation temperature (K)
1600
1800
2000
Figure 13 Irradiation temperature dependence of irradiated elastic modulus of CVD SiC, at ambient temperature, normalized to unirradiated values. The error bars are showing standard deviations for all the neutron data points and ranges of data scatter for the ion data points. Reproduced from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377.
1.20 Nanoindentation
Relative Young’s modulus
1.10
4pt. bend Sonic resonance
Osborne (1999)55,HFIR, 2 dpa Nogami (2002)56,HFIR + HFBR, 0.15–7.7 dpa Park (2003)57,DuET 5.1MeV Si, 3 dpa Katoh (2005)58,HFIR - 14J, 6.0–7.7 dpa Snead (2007)16, HFIR - METS, 1.7–8.6 dpa Price (1977)59, ETR, 2.8–12.2 dpa Snead (2007)16, HFIR, 0.7–4.2 dpa Model (Tersoff potential)
1.00
0.90
0.80
0.70 0.0
0.5
1.0
1.5 2.0 2.5 Volumetric swelling (%)
3.0
3.5
Figure 14 Irradiation-induced change of elastic modulus versus swelling of CVD SiC. An estimation of the influence of lattice relaxation on elastic modulus is calculated using Tersoff potential. Reproduced from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377.
Radiation Effects in SiC and SiC–SiC
229
1.6
Irradiated nano-indentation hardness normalized to unirradiated values
1.4 1.2 1 0.8 0.6 0.4 0.2 0 200
Katoh (2005)58, HFIR-14J, 6.0–7.7 dpa Nogami (2002)56, HFIR+HFBR, 0.2–7.7 dpa Osborne (1999)55, HFIR, 0.1–5.0 dpa Park (2003)57, DuET 5.1 MeV Si, 3 dpa
400
600
800
1000
1200
1400
1600
Irradiation temperature (K) Figure 15 Nanoindentation hardness of CVD SiC at ambient temperature as a function of irradiation temperature. Reproduced from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377.
2 1.8
Irradiated fracture toughness normalized to unirradiated values
1.6 1.4 1.2 1 0.8 0.6 0.4 0.2 0 200
Nogami (2002)56, HFIR+HFBR, 0.2–7.7 dpa Osborne (1999)55, HFIR, 0.1–5.0 dpa Park (2003)57, DuET 5.1 MeVSi, 3 dpa Snead (2007)16, HFIR, 0.7–4.2 dpa
400
600
800
1000
1200
1400
1600
1800
Irradiation temperature (K) Figure 16 Indentation fracture toughness of CVD SiC at ambient temperature as a function of irradiation temperature. Reproduced from Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377.
Si for reaction-bonded SiC, which typically segregate to grain boundaries during sintering, tends to have a significant influence on strength under neutron irradiation. For the case of sintered SiC with boron
compounds as sintering additives, the reaction of B(n, a)7Li causes the accumulation of helium bubbles at and near the grain boundary phases under neutron irradiation.60–63 In contrast, unmatched
10
230
Radiation Effects in SiC and SiC–SiC
1.5
Normalized strength
Hot pressed and sintered SiC forms
1.0
Matheney (1979)64 Matthews (1974)11 Iseki (1990)19 RB(reaction bonded) Iseki (1990)19 PLS(pressureless sintered)
0.50
Iseki (1990)19 HP(hot pressed) Price (1982)63 NC-430 Price (1982)63 Carborundum Dienst (1991)62 0.75% B Dienst (1991)62 1.9% B Corelli (1983)60 NC430
0.0
0
0.1
1 Dose (dpa)
10
100
Figure 17 Fluence dependence of irradiated flexural strength of hot-pressed and sintered SiC normalized to unirradiated strength. Irradiation is variable but in the saturable swelling regime. Reproduced from Newsome, G. A.; Snead, L. L.; Hinoki, T.; Katoh, Y.; Peters, D. J. Nucl. Mater. 2007, 371, 76–89.
2.0 CVD silicon carbide
Normalized strength
1.5
1.0
Snead Price Dienst Newsome 300 ⬚C Newsome 500 ⬚C Newsome 800 ⬚C
0.50
0.0 0.1
1
10
100
Dose (dpa) Figure 18 Flexural strength of CVD SiC at ambient temperature as a function of irradiation dose. Reproduced from Newsome, G. A.; Snead, L. L.; Hinoki, T.; Katoh, Y.; Peters, D. J. Nucl. Mater. 2007, 371, 76–89.
swelling between Si and SiC for reaction-bonded SiC causes disruption at the grain boundary, severely reducing the strength.11,19,60–64 Meanwhile, the highpurity materials such as CVD SiC exhibit superior irradiation resistance. The irradiation effect on flexural strength of Rohm and Haas CVD SiC in a fluence range of 0.15–30 dpa is summarized in Figure 18. In comparing Figure 18
with Figure 17, it is clear that CVD SiC retains stability in strength to a much higher dose than the sintered and reaction-bonded forms of SiC. It is to be noted that in Figure 18, the data of Dienst does indicate a significant as-irradiated degradation in strength around 15 dpa. However, such degradation is not seen for the 30 dpa irradiation of Snead. It is speculated that the degradation in the Dienst data may
Radiation Effects in SiC and SiC–SiC
have been due to statistical limitations of the study and/or due to issues with sample handling postirradiation. This issue is discussed in the Dienst reference.65 A compilation of strength data as a function of irradiation temperature is given in Figure 18, indicating no apparent correlation for the dose and temperature ranges studied. However, as with the fracture toughness data, irradiation-induced strengthening seems to be significant at 573–1273 K. The large scatter in flexural strength of brittle ceramics is inevitable, as the fracture strength is determined by the effective fracture toughness, morphology, and characteristics of the flaw that caused the fracture. Irradiation possibly modifies both the flaw characteristics and the fracture toughness through potential surface modification, relaxation of the machining-induced local stress, modifications of elastic properties, and fracture energy. A typical means of describing the failure of ceramics is through the use of Weibull statistics, which is a departure from the analysis of data that is assumed to follow a normal Gaussian distribution. In the twoparameter Weibull formalism, sometimes referred to as a weakest-link treatment, the failure probability F is described as F ðxÞ ¼ 1 e ðx
m
=x0 Þ
where m is the Weibull modulus and xo is the distribution size parameter. A change in the Weibull statistics, indicating a higher scatter in as-irradiated flexural strength has been observed by previous authors, although the point could not be made convincingly because of limitations in the number of tests observed. In the earliest work known to the authors, Sheldon66 noted a 14% decrease in crushing strength of highly irradiated CVD SiC shells with an increase of the coefficient of variation from 8% to 14%. Price63 went on to a 4-point bend test using relatively thin (0.6 mm) strips of CVD SiC deposited onto a graphite substrate. In his work, the flexural strength following a 9.4 1025 n m2 (E > 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
Radiation Effects in SiC and SiC–SiC
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 <0.7 dpa in the 673– 1303 K temperature range, whereas a major transition at higher doses likely exists between 1223 and 1353 K. Below 1223 K, the creep strain appeared highly nonlinear with neutron fluence because of the
234 Table 1 Tirr ( C)
CVD SiC 400 600 600 640 700 750 1030 1080 3C–SiC 640 700 1030 1080
Radiation Effects in SiC and SiC–SiC Irradiation creep data for CVD SiC from bend stress relaxation experiments Fluence (dpa)
Reactor
Initial/final bend stress (MPa)
Initial/final bend strain (104)
Creep strain (104)
BSR ratio m
Average creep compliance 106 (MPa dpa)1
0.6 0.2 0.6 3.7 0.7 0.6 0.7 4.2
JMTR JMTR JMTR HFIR HFIR JMTR HFIR HFIR
82/60 81/57 81/46 87/36 102/72 80/55 85/61 101/8
1.80/1.39 1.80/1.31 1.80/1.05 1.95/0.83 2.27/1.64 1.80/1.27 1.94/1.42 2.29/0.19
0.41 0.49 0.75 1.12 0.63 0.53 0.52 2.10
0.77 0.73 0.58 0.42 0.72 0.71 0.73 0.08
0.97 3.5 2.0 0.50 1.1 1.3 0.97 0.91
3.7 0.7 0.7 4.2
HFIR HFIR HFIR HFIR
87/30 102/90 86/57 101/1
1.94/0.68 2.27/2.06 1.94/1.31 2.29/0.02
1.26 0.21 0.63 2.27
0.35 0.87 0.67 0.01
0.59 0.34 1.2 1.1
early domination of the transient irradiation creep. The transient creep is speculatively caused by the rapid development of defect clusters and the structural relaxation of as-grown defects during early stages of irradiation damage accumulation. At 1353 K, irradiation creep mechanisms, which are common to metals, are likely operating. In metals, steady-state irradiation creep rates are generally proportional to the applied stress and neutron (or other projectiles) flux, f,75,76 and therefore, irradiation creep compliance, B, has been conveniently introduced75: eic ¼ sðBf þ DSÞ where S is void swelling and D is a coefficient of swelling–creep coupling. Ignoring the swelling– creep coupling term (valid in the saturable swelling regime), preliminary estimations of the steady-state irradiation creep compliance of CVD SiC were given to be 2.7 2.6 10–7 and 1.5 0.8 10–6 (MPa dpa)1 at 873–1223 K and 1353 K, respectively. If linear-averaged, creep compliances of 1–2 10–6 (MPa dpa)1 were obtained for doses of 0.6–0.7 dpa at all temperatures. Monocrystalline 3C–SiC samples exhibited a significantly smaller transient creep strain by 0.7 dpa and a greater subsequent deformation when loaded along <011> direction. To better define the irradiation creep behavior of SiC and the underlying physical mechanisms, it will be essential to further examine the stress dependence, dose dependence, effect of crystallographic orientation, microstructures of the crept samples, and the coupling between irradiation creep and swelling.
4.07.6 Silicon Carbide Composites Under Irradiation SiC composites are a family of materials of varied constituents and architectures. Up to the point of writing this chapter, nuclear-grade SiC composites (those specifically developed for application in fast neutron environments and exhibiting neutron irradiation damage resistance) are more precisely defined as continuous fiber-reinforced ceramic composites. The history of development for these materials has been reviewed in a number of publications.29,77–79 The primary constituents of these nuclear-grade composites are the continuous SiC fiber, a fiber/matrix interphase material that can be SiC or pyrolytic graphite or a combination of the two, and a matrix of SiC infiltrated into the woven fiber preform. The most common matrix material is derived from chemical vapor infiltration (CVI), and is essentially identical in structure, properties, and irradiation response to the CVD SiC discussed in previous sections. While there has been little direct study on the effects of irradiation on the material properties of the SiC interphase, it can be assumed that it would also behave in a similar manner to the SiC matrix. However, the effect of neutron irradiation on pyrolytic graphite interphase (if used) will be substantially different from that on both matrix and fiber. While the effect of irradiation on the underlying properties of graphite interphase has not been well studied, it can be assumed that the interphase will behave in a similar manner to nuclear graphite (discussed in Chapter 4.05, Radiation Damage of Reactor Pressure Vessel Steels).
Radiation Effects in SiC and SiC–SiC
Y
(a)
235
(b) X
Z Y (c)
300 mm
SiC matrix
SiC/PyC multilayer
PyC 1 mm
Hi-Nicalon type-S fiber
Figure 21 Example of braided nuclear-grade SiC/SiC composite. Fiber: Hi-Nicalon™ Type-S; Interphase: Multilayer SiC with pyrolytic carbon; Matrix: CVI SiC deposited through an isothermal process. Reproduced from Nozawa, T.; Lara-Curzio, E.; Katoh, Y.; Shinavski, R. J. Tensile properties of advanced SiC/SiC composites for nuclear control rod applications. Wiley: 2007; pp 223–234.
An example of an SiC/SiC composite that has been developed for high-temperature gas-cooled reactor control rod applications is shown in Figure 21. The basic textile weaving of the composite is evident on inspection of Figure 21(a). In this case, a 55 weave is depicted. For the polished section of Figure 21(b), large voids, which are an unavoidable characteristic of chemical vapor infiltrated materials and also the primary reason why it is difficult to produce gasimpermeable SiC/SiC composite, are clearly observed. In Figure 21(c), the complicated structure of the interphase is seen. In this case, alternating layers of SiC and pyrolytic graphite have been applied. The pyrolytic graphite layer between the SiC layers is quite thin (tens of nanometer), with a relatively thick graphite layer in contact with the fiber itself. From the earliest study of SiC/SiC composites under irradiation, it was clear that the fiber was the key to performance. As with the impure forms of SiC monolithic ceramics (cf. Figure 17), the impure and oxygen-rich early grades of SiC fiber (trade name Nicalon™) were quite unstable under neutron irradiation.12,80,81 Researchers were able to directly link an irradiation-induced shrinkage of the SiC-based
fibers with debonding of the fiber–matrix interface that severely compromised the ability to load the high-strength fibers.80 Composite mechanical properties such as strength suffered appreciably. With continued evolution of the fiber systems to increasingly pure, stoichiometric materials, the irradiation stability improved significantly. Presently, there are two commercial fiber systems used in nuclear-grade composites, both of which have relatively low impurity contents and are approaching a 1:1 stoichiometry. Specifically, the 11 micron Hi-Nicalon™ Type-S fiber has the nominal chemistry of SiC1.05, 0.2%-O, while the 7.5 and 10 mm Tyranno™ SA-3 fibers have the nominal chemistry of SiC1.07, 0.5% Al. Study has revealed that these ‘near stoichiometric’ fibers exhibit irradiationinduced swelling similar to that of CVD,82 thus avoiding the debonding phenomenon mentioned in the previous paragraph. For this reason, composites fabricated from these materials are superior under irradiation to their predecessors. Consistent with the discussion of properties of irradiated monolithic SiC, the following discussion will be limited to the more pure, near stoichiometric fiber materials.
236
Radiation Effects in SiC and SiC–SiC
Weibull mean strength (GPa)
4
280
3
300
470
800
500
470 770
2 280
950 950
1
Offset for nonirradiated Hi-Nicalon™ type-S, Katoh (2010)67 Tyranno™-SA3, Katoh (2010)67 Hi-Nicalon™ type-S, Nozawa (2004)83
0 0
2
4 6 Neutron dose (dpa)
8
10
Figure 22 Effect of neutron irradiation on fiber strength. Data labels indicate the nominal irradiation temperature (in C) for Hi-Nicalon™ Type-S (upright) and Tyranno™ SA-3 (oblique) fibers. Reproduced from Katoh, Y.; Snead, L. L.; Nozawa, T.; Kondo, S.; Busby, J. T. J. Nucl. Mater. 2010, 403, 48–61.
The effect of neutron irradiation on the Weibull mean strength of individual ‘near stoichiometric’ fibers is given in Figure 22.83,84 Within inherent statistical scatter, no change in strength is observed for either the Hi-Nicalon™ Type-S or the Tyranno™ SA-3 bare fibers. The numbers inset to the figure indicate the irradiation temperature of the SiC fibers, with no apparent function of irradiation temperature on strength observed. From the same study, the effect of irradiation on composite properties is also observed. Figure 2367 gives the proportional limit stress for which the load departs from elastic behavior and the ultimate tensile strength. As with the fiber data, and the data for monolithic CVD SiC (Figure 18), the composite strength does not exhibit any statistically meaningful change. Supporting studies14,82,83,85–87 on the strength in tension or bending of neutron-irradiated stoichiometric fiber composites support the fact that at least up to 40 dpa, composite strength is not significantly affected by irradiation. A recent study88 on the fracture toughness of irradiated and unirradiated Hi-Nicalon™ Type-S composites also reports no appreciable change. However, a minor difference in the fracture surface (length of fiber pull out) and a trend in the fiber–matrix interphase properties are reported,89 suggesting that mechanical property evolution may occur at higher doses.
In the unirradiated state, the thermal conductivity of SiC composites is dependent on variables including the fibers and matrix constituents, processing, and the level of porosity. For the nuclear composite considered here, there is considerable thermal conductivity anisotropy and temperature dependence typical of all ceramics. This is demonstrated in Figure 24, which gives the measured and calculated thermal conductivity for the two nuclear-grade SiC composites.90 Presented are Hi-Nicalon™ Type-S fiber and Tyranno™ SA fiber composites, each matrix infiltrated through CVI.58 Architectures included balanced (1:1:1 for x:y:z) and unbalanced (1:1:4) 3D forms and 2D laminates (SW: satin weave, PW: plain weave.) In each case, a pyrolytic graphite interphase was applied. The conductivity for all materials is presented in the through thickness direction (perpendicular to the plate and the fabric for the 2D composite.) This typically represents the lowconductivity direction. As evident from Figure 24 and the supporting analysis by Katoh,90 the fiber makes a significant contribution to the thermal conductivity of these highly stoichiometric fiber composites, and this conductivity is a fairly strong function of temperature. However, the absolute conductivity is only a fraction of that for the highest thermal conductivity CVD SiC (cf. Figure 7.)
Radiation Effects in SiC and SiC–SiC
237
400 Hi-Nicalon-S CVI UTS Tyranno-SA3 CVI UTS Hi-Nicalon-S CVI PLS
Nonirradiated
420
300 400
Tensile stress (MPa)
Tyranno-SA3 CVI PLS
460 480 690
760
780
610
UTS
530 380
200
220
1000 480
570 610
PLS
100
350
0 0
1
2
4 3 Neutron dose (dpa)
5
6
Figure 23 Effect of neutron dose on tensile proportional limit and ultimate tensile stresses for composites. Data labels indicate the nominal irradiation temperature in C for Hi-Nicalon™ Type-S (upright) and Tyranno™ SA-3 (oblique) composites. Reproduced from Katoh, Y.; Snead, L. L.; Nozawa, T.; Kondo, S.; Busby, J. T. J. Nucl. Mater. 2010, 403, 48–61.
80 3D 1:1:4 TySA/PyC, through-thickness, model 3D 1:1:1 TySA/PyC, through-thickness, model 2D-PW TySA/PyC, through-thickness, model 5HSW HNLS/PyC, through-thickness, model 3D 1:1:4 TySA/PyC through-thickness, experiment 3D 1:1:1 TySA/PyC through-thickness, experiment 2D-PW TySA/PyC through-thickness, experiment 5HSW HNLS/PyC through-thickness, experiment
Thermal conductivity (W m-1 K–1)
70 60 50 40 30 20 10 0 0
200
400
800 600 Temperature (⬚C)
1000
1200
Figure 24 Thermal conductivity of representative nuclear-grade SiC/SiC composite in unirradiated condition. Reproduced from Katoh, Y.; Nozawa, T.; Snead, L. L.; Hinoki, T.; Kohyama, A. Fus. Eng. Des. 2006, 81, 937–944.
As with the CVD SiC discussed in section 4.07.3, silicon carbide composite also undergoes significant degradation in thermal conductivity because of neutron irradiation. The data is somewhat limited;
however, Figure 25 gives the ambient throughthickness thermal conductivity for a plain weave Hi-Nicalon™ Type-S, multilayer SiC interphase, and CVI SiC matrix composite. It is noted that, in
238
Radiation Effects in SiC and SiC–SiC
Knonirr = 10.1 ± 2.2
Thermal conductivity at ambient (W m-1 K–1)
10
8
6 Tirr ~ 200 ⬚C Tirr ~ 800 ⬚C 4 ~ 600 ⬚C ~ 400 ⬚C
2 Nicalon Type-S fiber composite 0
0.001
0.01
0.1 Dose (dpa)
1
10
Figure 25 Effect of neutron irradiation on the through-thickness thermal conductivity of Hi-Nicalon™ Type-S, CVI matrix composite.
1
Thermal defect resistance (W m-1 K-1)-1
~ 800 ⬚C composite
0.1 Tirr ~ 200 ⬚C composite
0.01
0.001
Tirr ~ 800 ⬚C CVD SiC
Tirr ~ 200 ⬚C CVD SiC
0.001
0.01
0.1 Dose (dpa)
1
10
Figure 26 Comparison of the thermal defect resistance for neutron irradiated CVD SiC and Hi-Nicalon™ Type-S, CVI matrix composite.
Radiation Effects in SiC and SiC–SiC
comparison to the conductivity shown in Figure 24 (second from lowest curve), the ambient throughthickness thermal conductivity for the material of Figure 25 is relatively low (10.2 2.2 W m1 K1). This is mostly ascribed to the large porosity for that composite. Nevertheless, the figure clearly shows a significant, irradiation temperature-dependent reduction in thermal conductivity as a function of irradiation dose. The fact that this is temperature dependent suggests that the degradation is due to the production of stable point defects and clusters, as discussed in Section 4.07.3, although this may not be the sole factor determining the degradation. Figure 26 provides the accumulated thermal defect resistance at the lowest and highest irradiation temperature for the composite materials of Figure 25, compared with high-conductivity CVD SiC. It is interesting to note that the thermal defect resistance for the composite, while accumulating in the same manner as that of the CVD SiC, is about an order of magnitude greater than that of CVD SiC at a given dose (at least prior to saturation.) This greater accumulation of thermal defect resistance has been recently observed by Katoh.67 The reason for this is unclear, although it is plausible that, in addition to defect production, propagation of internal interfaces (e.g., cracks) in the composite is occurring under irradiation. It is also possible that the defects population responsible for phonon scattering for the composite material is stabilized at a higher level than that of the highly pure CVD SiC.90
13. 14. 15. 16. 17. 18. 19. 20. 21. 22. 23. 24. 25. 26.
27. 28. 29. 30. 31. 32.
References 1. 2. 3. 4. 5. 6. 7.
8. 9. 10. 11. 12.
CEGA NP-MHTGR Material Models of Pyrocarbon and Pyrolytic Silicon Carbide; CEGA-002820, Rev 1; July 1993. Blackstone, R.; Voice, E. H. J. Nucl. Mater. 1971, 39, 319–322. Price, R. J. J. Nucl. Mater. 1969, 33, 17–22. Price, R. J. J. Nucl. Mater. 1973, 48, 47–57. Primak, W.; Fuchs, L. H.; Day, P. P. Phys. Rev. 1956, 103(5), 1184–1192. Balarin, M. Phys. Stat. Sol. 1965, 11, K67–K71. Pravdyuk, N. F.; Nikolaenko, V. A.; Kapuchin, V. I.; Kusnetsov, V. N. In Properties of Reactor Materials and the Effects of Radiation Damage Proceedings, Littler, D. J., Ed.; Butterworths: 1962; p 57. Thorne, R. P.; Howard, V. C.; Hope, B. Proc. Brit. Ceramic Soc. 1967, 7. Stevens, R. Phil. Mag. 1972, 25, 523–528. Senor, D. J.; Youngblood, G. E.; Moore, C. E.; Trimble, D. J.; Newsome, G. A.; Woods, J. J. Fus. Tech. 1996, 30, 943–955. Matthews, R. J. Nucl. Mater. 1974, 51, 203–208. Hollenberg, G. W.; Henager, C. H., Jr.; Youngblood, G. E.; Trimble, D. J.; Simonson, S. A.; Newsome, G. A.; Lewis, E. J. Nucl. Mater. 1995, 219, 70–86.
33. 34. 35. 36. 37. 38. 39. 40.
41. 42. 43.
239
Snead, L. L.; Hinoki, T.; Katoh, Y. Strength of neutron irradiated SiC carbide and silicon carbide composite; DOE/ER-0313/33; 2002; pp 49–57. Newsome, G. A.; Snead, L. L.; Hinoki, T.; Katoh, Y.; Peters, D. J. Nucl. Mater. 2007, 371, 76–89. Katoh, Y.; Hashimoto, Y.; Kondo, S.; Snead, L. L.; Kohyama, A. J. Nucl. Mater. 2006, 351, 228–240. Snead, L. L.; Nozawa, T.; Katoh, Y.; Byun, T-S.; Kondo, S.; Petti, D. A. J. Nucl. Mater. 2007, 371, 329–377. Yano, T.; Miyazaki, H.; Akiyoshi, M.; Iseki, T. J. Nucl. Mater. 1998, 253, 78–86. Senor, D. J.; Youngblood, G. E.; Greenwood, L. R.; et al. J. Nucl. Mater 2003, 317, 145–159. Iseki, T.; Maruyama, T.; Yano, T.; Suzuki, T.; Mori, T. J. Nucl. Mater. 1990, 170, 95–100. Katoh, Y.; Kishimoto, H.; Kohyama, A. J. Nucl. Mater. 2002, 307–311, 1221–1226. Snead, L. L.; Zinkle, S. J.; Hay, J. C.; Osborne, M. C. Nucl. Instrum. Meth. Phys. Res. B 1998, 141, 123–132. Snead, L. L.; Zinkle, S. J. Nucl. Instrum. Meth. Phys. Res. 2002, 191B, 497–503. Katoh, Y.; Kishimoto, H.; Kohyama, A. Mater. Trans. 2002, 43(4), 612–616. Golubov, S. I. Phys. Met. Metall. 1985, 60(3), 7–13. Kuwabara, T.; Kurishita, H.; Ukai, S.; Narui, M.; Mizuta, S.; Yamazaki, M.; Kayano, H. J. Nucl. Mater. 1998, 258–263, 1236–1241. Snead, L. L.; Zinkle, S. J. In Microstructure Evolution During Irradiation, Robertson, I. M., Was, G. S., Hobbs, L. W., de la Rubia, T. D., Eds.; Materials Research Society: Pittsburgh, 1997; Vol. 439, pp 595–606. Kondo, S.; Park, K. H.; Katoh, Y.; Kohyama, A. Fus. Sci. Tech. 2003, 44, 181–185. Kondo, S.; Katoh, Y.; Snead, L. L. J. Nucl. Mater. 2009, 386–388, 222–226. Katoh, Y.; Snead, L. L.; Henager, C. H.; et al. J. Nucl. Mater. 2007, 367, 659–671. Itoh, H.; Hayakawa, N.; Nashiyama, I.; Sakuma, E. J. Appl. Phys. 1989, 66, 4529–4531. Kawasuso, A.; Itoh, H.; Morishita, N.; et al. Appl. Phys. 1998, 67A, 209–212. Snead, L. L.; Burchell, T. D.; Katoh, Y. J. Nucl. Mater. 2008, 381, 55–61. Snead, L. L.; Katoh, Y.; Connery, S. J. Nucl. Mater. 2007, 367–370, 677–684. Price, R. J. J. Nucl. Mater. 1973, 46, 268–272. Snead, L. L.; Scholz, R.; Hasegawa, A.; Rebelo, A. F. J. Nucl. Mater. 2002, 307–311, 1141–1145. Olesinski, R. W.; Abbaschian, G. J. Bull. Phase Alloy Diagrams 1984, 5. Li, J.; Porter, L. J.; Yip, S. J. Nucl. Mater. 1998, 255, 139–152. Huang, H. C.; Ghoniem, N. M.; Wong, J. K.; Baskes, M. I. Modelling . Simul.Mater. Sci. Eng. 1995, 3, 615–627. Bockstedte, M.; Mattausch, A.; Pankratov, O. Phys. Rev. B 2003, 68, 205201-1–205201-17. Lam, C. H.; Ling, C. C.; Beling, C. D.; Fung, S.; Weng, H. M.; Hang, D. S. Vacancies in electron irradiated 6H silicon carbide studied by positron annihilation spectroscopy. In Materials Research Society Symposium Proceedings, 2004; pp R3.19.1–R3.19.6. Bockstedte, M.; Heid, M.; Pankratov, O. Phys. Rev. B 2003, 67, 193102–1–4. de Sousa Balona, L. A.; Loubser, J. H. N. J. Phys. C Solid State Phys. 1970, 3, 2344–2351. Lee, C. W.; Pineau, F. J.; Corelli, J. C. J. Nucl. Mater. 1982, 108–109, 678–684.
240
Radiation Effects in SiC and SiC–SiC
44. Snead, L. L.; Burchell, T. D. Stored Energy in Irradiated Silicon Carbide; Fusion Reactor Materials Semiannual Progress Report, DOE ER-0313/21; Dec 31, 1996, 21. 45. Rohde, M. J. Nucl. Mater. 1991, 182, 87–92. 46. Taylor, R.E.; Groot, H.; Ferrier, J. Thermophysical Properties of CVD SiC. Thermophysical Properties Research Laboratory Report, TRPL 1336, School of Mechanical Engineering, Purdue University, 1993. 47. Product sheet from Rohm and Haas Co. http://www.dow. com/assets/attachments/business/gt/advanced_ceramics/ cvd_silicon_carbide/tds/cvd_silicon_carbide.pdf. 48. Goaebner, J. E.; Altman, H.; Balzaretti, N. M.; et al. Diam. Relat. Mater. 1998, 7, 1589–1604. 49. Pcikering, M. A.; Tayler, R. L.; Keeley, J. T.; Graves, G. A. Nucl. Instrum. and Meth. A 1990, 291, 95–100. 50. Collins, A. K.; Pickering, M. A.; Taylor, R. L. J. Appl. Phys. 1990, 68, 6510–6512. 51. Shaffer, P.T.B. Engineering Properties of Carbides. In Engineering Materials Handbook, 1991; pp 804–811. 52. Snead, L. L. J. Nucl. Mater. 2004, 329–333, 524–529. 53. Snead, L. L.; Zinkle, S. J.; White, D. P. J. Nucl. Mater. 2005, 340, 187–202. 54. Tersoff, J. Phys. Rev. B 1989, 39(8), 5566–5568. 55. Osborne, M. C.; Hay, J. C.; Snead, L. L.; Steiner, D. J. Am. Ceram. Soc. 1999, 82(9), 2490–2496. 56. Nogami, S.; Hasegawa, A.; Snead, L. L. J. Nucl. Mater. 2002, 307–311, 1163–1167. 57. Park, K. H.; Kondo, S.; Katoh, Y.; Kohyama, A. Fus. Sci. Technol. 2003, 44, 455–459. 58. Katoh, Y.; Snead, L. L. J. ASTM Int. 2005, 2(8), 12377-1–12377-13. 59. Price, R. Nucl. Technol. 1977, 35, 320–336. 60. Corelli, J. C.; Hoole, J.; Lazzaro, J.; Lee, C. W. J. Am. Ceram. Soc. 1983, 66(7), 529–537. 61. Harrison, S. D.; Corelli, J. C. J. Nucl. Mater. 1984, 122(1–3), 833–839. 62. Dienst, W. Fus. Eng. Design 1991, 16, 311–316. 63. Price, R. J.; Hopkins, G. R. J. Nucl. Mater. 1982, 108–109, 732–738. 64. Matheny, R. A.; Corelli, J. C. J. Nucl. Mater. 1979, 83, 313–321. 65. Dienst, W.; Fett, T.; Heidinger, R.; Roehrig, H. D.; Schulz, B. J. Nucl. Mater. 1990, 174, 102–109. 66. Sheldon, B. E. UKAEA Report, AERE-R8025; 1975. 67. Katoh, Y.; Snead, L. L.; Nozawa, T.; Kondo, S.; Busby, J. T. J. Nucl. Mater. 2010, 403, 48–61. 68. Byun, T. S.; Lara-Curzio, E.; Lowden, R. A.; Snead, L. L.; Katoh, Y. Miniturized Fracture Stress Tests for
69. 70. 71. 72. 73. 74. 75. 76. 77. 78. 79. 80. 81. 82. 83. 84.
85. 86. 87. 88. 89. 90.
Thin-Walled Tubular SiC Specimens. In 12th International Conference on Fusion Reactor Metals, Santa Barbara, CA, 2005. Hong, S. G.; Byun, T. S.; Lowden, R. A.; Snead, L. L.; Katoh, Y. J. Am. Ceram. Soc. 2007, 90(1), 184–191. Scholz, R. J. Nucl. Mater. 1998, 258–263, 1533–1539. Scholz, R. J. Nucl. Mater. 1998, 254, 74–77. Scholz, R.; Mueller, R.; Lesueur, D. J. Nucl. Mater. 2002, 307–311, 1183–1186. Scholz, R.; Youngblood, G. E. J. Nucl. Mater. 2000, 283–287, 372–375. Katoh, Y.; Snead, L. L.; Hinoki, T.; Kondo, S.; Kohyama, A. J. Nucl. Mater. 2007, 367–370, 758–763. Ehrlich, K. J. Nucl. Mater. 1981, 100, 149–166. Matthews, J. R.; Finnis, M. W. J. Nucl. Mater. 1988, 159, 257–285. Snead, L. L.; Osborne, M. C.; Lowden, R. A.; et al. J. Nucl. Mater. 1998, 253, 20–30. Katoh, Y.; Kohyama, A.; Hinoki, T.; Snead, L. L. Fus. Sci. Technol. 2003, 44(1), 155–162. Jones, R. H.; Steiner, D.; Heinisch, H. L.; Newsome, G. A.; Kerch, H. M. J. Nucl. Mater. 1997, 245, 87–107. Snead, L. L.; Steiner, D.; Zinkle, S. J. J. Nucl. Mater. 1992, 191–194, 566–570. Snead, L. L.; Osborne, M.; More, K. L. J. Mater. Res. 1995, 10(3), 736–746. Hinoki, T.; Snead, L. L.; Katoh, Y.; Kohyama, A. J. Nucl. Mater. 2002, 307–311, 1157–1162. Nozawa, T.; Hinoki, T.; Snead, L. L.; Katoh, Y.; Kohyama, A. J. Nucl. Mater. 2004, 329–33, 544–548. Nozawa, T.; Lara-Curzio, E.; Katoh, Y.; Shinavski, R. J. Tensile properties of advanced SiC/SiC composites for nuclear control rod applications. In Proceedings of the 31st Annual International Conference on Advanced Ceramics & Composites, 2007; pp 223–234. Nozawa, T.; Snead, L. L.; Katoh, Y.; Miller, J. H. J. Nucl. Mater. 2007, 371, 304–313. Ozawa, K.; Nozawa, T.; Katoh, Y.; Hinoki, T.; Kohyama, A. J. Nucl. Mater. 2007, 367–370, 713–718. Katoh, Y.; Nozawa, T.; Snead, L. L.; Hinoki, T. J. Nucl. Mater. 2007, 367–370, 774–779. Ozawa, K.; Katoh, Y.; Nozawa, T.; Snead, L. L. J. Nucl. Mater. 2011 (in press). Nozawa, T.; Katoh, Y.; Snead, L. L. J. Nucl. Mater. 2009, 384(3), 195–211. Katoh, Y.; Nozawa, T.; Snead, L. L.; Hinoki, T.; Kohyama, A. Fus. Eng. Des. 2006, 81, 937–944.
4.08
Oxide Dispersion Strengthened Steels
S. Ukai Hokkaido University, Sapporo, Japan
ß 2012 Elsevier Ltd. All rights reserved.
4.08.1
Introduction
242
4.08.2 4.08.2.1 4.08.2.2 4.08.3 4.08.3.1 4.08.3.2 4.08.3.2.1 4.08.3.2.2 4.08.3.2.3 4.08.3.3 4.08.3.3.1 4.08.3.3.2 4.08.3.4 4.08.4 4.08.4.1 4.08.4.2 4.08.4.3 4.08.4.4 4.08.5 4.08.5.1 4.08.5.2 4.08.5.3 4.08.5.4 4.08.6 4.08.7 4.08.7.1 4.08.7.2 4.08.7.3 4.08.7.4 4.08.8 4.08.8.1 4.08.8.2 4.08.8.3 4.08.8.3.1 4.08.8.3.2 4.08.8.3.3 4.08.9 References
Nanoscale Oxide Particle Control Dissociation and Precipitation Structure and Coherency Martensitic 9Cr-ODS Steels Chemical Composition and Microstructure Residual Ferrite Formation and Strength Characterization Mechanically alloyed powder characterization Pinning of a–g interface by oxide particles Strength characterization Cladding Manufacturing Continuous cooling transformation diagram Manufacturing process Creep and tensile properties Ferritic 12Cr-ODS Steels Strength Anisotropy Recrystallization Tests Cold-Rolling Cladding Manufacturing Internal Creep Rupture Property Al-Added 16Cr-ODS Steels Application and Technical Issues Thermal Aging Embrittlement Due to High Cr Content Mechanical Properties Cladding Manufacturing Existing ODS Steel Cladding Corrosion and Oxidation Sodium Compatibility LBE Compatibility SCPW Compatibility Oxidation Irradiation Simulated Irradiation Neutron Irradiation of Materials Fuel Pin Irradiation 9Cr- and 12Cr-ODS steel cladding in BOR-60 12Cr-ODS steel cladding in EBR-II DT2203Y05 in Phe´nix Summary
242 242 243 244 244 245 245 245 247 248 248 249 250 250 250 252 252 253 255 255 256 256 257 258 259 259 259 262 262 263 263 264 268 268 269 269 269 270
Abbreviations CTT Continuous cooling transformation CEN-SCK Centre d’Etude de l’e´nergie Nucleaire – Studiecentrum voor Kernenergie
CVN EFTEM EPMA
Charpy V-notch Energy-filtered transmission electron microscopy Electron probe microanalysis
241
242
Oxide Dispersion Strengthened Steels
FFT HIP HRTEM INCO JAEA LBE LFR LMP MA MOX ODS PMW PRW SCPW SEM SFR TIG UTS
Fast Fourier transformation Hot isostatic pressing High-resolution transmission electron microscopy International Nickel Company Japan Atomic Energy Agency Lead–bismuth eutectic Lead fast reactor Larson–Miller parameter Mechanical alloying Mixed oxide Oxide dispersion strengthened Pulse magnetic Welding Pressurized resistance welding Super critical pressurized water Secondary electron microscopy Sodium fast reactor Tungsten inert gas welding Ultimate tensile strength
4.08.1 Introduction Recent progress in oxide dispersion strengthened (ODS) steels produced by mechanical alloying (MA) techniques allows them to be used as fuel cladding in sodium-cooled fast reactors (SFR). The thermally stable oxide particles dispersed in the ferritic matrix improve the radiation resistance and creep resistance at high temperature. As a result, ODS steels have a strong potential for high burnup (long-life) and hightemperature applications typical for SFR fuels. The attractiveness of ODS steels is due not only to the nanosize oxide particles composed of Y–Ti–O atoms but also to their controlled micron-size grain morphology. We review existing knowledge on the crystalline structure and lattice coherency of these nanosize particles with their surrounding matrix, since these factors dominate the dispersion and strength-determining mechanism through dislocation interaction. The development of manufacturing processes is a principal issue for hardened ODS steels to realize long, thin-walled ODS steel cladding on production scales. There was the long-standing problem in low hoop strength due to the extremely elongated fine grains parallel to the rolling direction. To soften hardened cold-rolled products and modify their grain morphology, martensitic 9Cr-ODS steels and ferritic 12Cr-ODS steels have been developed. Current progress in the development of these ODS steel claddings, including their relevant mechanical properties, for
example, tensile and creep rupture strengths in the hoop directions, and irradiation performance, is reviewed. The development of Al-added high CrODS steel cladding is also addressed, with a focus on superior resistance to oxidation and corrosion in a lead– bismuth eutectic (LBE), and supercritical pressurized water (SCPW) in the international Generation IV advanced nuclear power system. Nanocluster ODS steels,1 for example, 14YWT, etc., for fusion blanket structure materials, are not addressed in this chapter.
4.08.2 Nanoscale Oxide Particle Control 4.08.2.1
Dissociation and Precipitation
The fine distribution of Y2O3 particles, which is essential to improving the high temperature strength of ODS steels, is attained by the dissociation of oxide particles during MA processing.2 The thermodynamically stale Y2O3 particles are forcedly decomposed into the ferritic steel matrix during the MA process. Subsequent annealing induces oxide particles to precipitate finely at elevated temperature of around 1000 C. The co-addition of Ti during MA processing promotes the decomposition of Y2O3 and then the precipitation of Y–Ti complex oxide particles through an annealing heat treatment.3,4 A field emission ion micro-probe (FIM) analysis confirmed that this type of complex oxide is constituted of several nanometer-sized Y–Ti–O compounds.5–7 The precipitation process of the decomposed Y2O3 was investigated by means of a small angle neutron scattering (SANS) experiment.8 The neutron-scattering cross-section (dS/dO) versus scattering vector (q2) plots for the milled U14YWT(Fe–14Cr–0.4Ti–3W– 0.25Y2O3) are shown in Figure 1(a). They indicate that the hot isostatic pressing (HIP) of U14YWT at 850 C leads to the precipitation of a high number density of nanoclusters, as designated by Odette. Figure 1(b) shows the effects of HIP (filled symbols) and powder annealing (open symbols) at temperatures of 700, 850, 1000, and 1150 C. The increase in magnitude and decrease in slope of the dS/dO versus q2 curves indicate that the radius of nanoclusters decreases and their number density increases with decreasing temperature at HIP and powder annealing. Annealing at 700 C produces the highest scattering and lowest sloping, which indicates that the smallest-sized nanoclusters precipitate with the highest number density at lower temperatures. In terms of an X-ray diffraction experiment using Super Photon
Oxide Dispersion Strengthened Steels
10
10 HIPed materials
850 ⬚C HIP (U14YWT)
1
As-MA (U14YWT) As-MA (No Y2O3) 0.1
(a)
0
2
4
6
8
q2 (nm−2)
1 Decreasing Increasing Nd
1000 ⬚C 850 ⬚C
0.1
0.01
10
Powder anneals
1150 ⬚C
dΣ/dΩ (cm srad)−1
dS/dW (cm srad)−1
850 ⬚C HIP (No Y2O3)
0.01
243
700 ⬚C
1150 1000 850 ⬚C Controls (No Y2O3)
0
2
4
6
8
10
q2 (nm−2)
(b)
Figure 1 Results of a SANS experiments for as-mechanically alloyed powders and after HIP and annealing in U14YWT (Fe–14Cr–0.4Ti–3W–0.25Y2O3). (a) As-MA, 850 C HIP and annealing and (b) HIP (filled) and annealing (annealing) at 1150, 1000, 850, and 700 C . Reproduced from Alinger, M. J.; Odette, G. R.; Hoelzer, D. T. J. Nucl. Mater. 2004, 382, 329–333.
ring-8 eV (Spring-8) constructed in Japan, Kim et al. recently reported that nanoclusters could be in a noncrystalline state and can be transformed to nanocrystalline oxide particles at around 1000 C.9 4.08.2.2
− (0 1 1) 3.06 Å 70.5 ⬚
Structure and Coherency
With regard to ODS steels without Ti, high resolution (HR) TEM investigations were performed by Klimiankou to investigate the structure of Y2O3.10 The crystallographic lattice of the metal matrix corresponds to a-Fe with a bcc structure and a lattice constant of a0 ¼ 0.287 nm.11 The Y2O3 has a crystalline bcc structure with a 1.06 nm lattice constant.11 Figure 2 shows an HRTEM image taken from an Y2O3 particle that is surrounded by the matrix (M) lattice. This image was taken from the grain, oriented with [1 1 1]M zone axis to the electron beam. A fast Fourier transformation (FFT) of the image shows the matrix lattice as a hexagonal pattern with diffraction spots of the {1 1 0} type and dM(110) ¼ 0.203 nm distance. In the FFT image, the Y2O3 (YO) lattice is rectangular, with diffraction spots of the {2 2 2} type and a corresponding atomic planes distance of dYO(2 2 2) ¼ 0.306 nm. The angle of 70.5 between diffraction spots of the {2 2 2}YO type marked in Figure 2(a) confirms that the Y2O3 particle is oriented with the [1 1 0] zone axis, and consequently [1 1 0]YO// [1 1 1]M. The orientation correlation of both lattices is (1 1 1)YO//(1 1 0)M. Therefore, the following Kurdjumov–Sachs orientation relationship12 is satisfied: ð11 1 ÞYO ==ð11 0ÞM ; ½110YO ==½111M :
½1
3.06 Å
− (2 2 2) − (1 0 1) − (2 2 2) − (1 1 0)
− (1 1 0) −− (2 2 2 ) − − − (1 0 1 ) (2 2 2 )
− (0 1 1 ) (a)
(b)
Figure 2 HRTEM micrograph of the Y2O3 particle with surrounding matrix (a) and FFT image of micrograph (b). The diffraction spots from Y2O3 particle of {2 2 2} type form the rectangle, whereas diffraction spots from the matrix of {1 1 0} type form the hexagon at the [1 1 0] zone axis and [1 1 1] of matrix. Reproduced from Kliniankou, M.; Lindau, R.; Moslang, A. J. Nucl. Mater. 2004, 329–333, 347–351.
The interfacial coherency between the Y2O3 particle and the ferritic matrix can be estimated by Klimiankou as follows: 3dMð1 1 0Þ 2dYOð2 2 2Þ 0:5%: 2dYOð2 2 2Þ
½2
This result suggests that a coherency could be satisfied for the Y2O3 particles surrounded by a ferritic matrix. Concerning the Y–Ti–O complex oxide particles formed in Ti-added ODS steels, Figure 3(a) shows an energy-filtered (EF) TEM micrograph from a Y–Ti–O particle in which two atomic planes are simultaneously visible.10 Figure 3(b) shows the (0 0 4) and (2 2 2) atomic planes of Y2Ti2O7 cubic (a0 ¼ 1.01 nm) phases with the [1 1 0] zone axis. In fact, the measured
244
Oxide Dispersion Strengthened Steels
(a)
(b) −− 222
− 222
Residual ferrite
004
5 nm Figure 3 EFTEM images of Y2Ti2O7 particles. Reproduced from Kliniankou, M.; Lindau, R.; Moslang, A. J. Nucl. Mater. 2004, 329–333, 347–351.
data are equal to the following data calculated from the Y2Ti2O7 structure: d2 2 2 ¼ 0.29 nm and d0 0 4 ¼ 0.25, and an angle between the (0 0 4) and (2 2 2) atomic planes of 54.7 . The analysis of EFTEM results definitively shows that Y–Ti–O particles have a Y2Ti2O7 composition. These findings suggest that nano-oxide particles precipitate from the ferritic matrix, maintaining crystalline coherency or partial-coherency with a ferritic matrix. In general, the nucleation and growth of precipitates proceeds, as both interfacial and strain energies become minimal. In the case of ODS steels, interfacial coherency could be maintained between thermodynamically stable nanoparticle precipitates and the ferritic matrix in order to decrease the free energy in the system from the extremely high energy state induced by MA. Elucidation of the details of the nanoscale precipitation is important not only as basic materials science research but also as the development of high-strength engineering materials.
4.08.3 Martensitic 9Cr-ODS Steels 4.08.3.1 Chemical Composition and Microstructure 9Cr-ODS steels are being developed by the JAEA ( Japan Atomic Energy Agency) for application to SFR fuel cladding. Their standard chemical composition is 9Cr–0.13C–0.2Ti–2W–0.35Y2O3 (wt%). The chromium concentration was determined to be 9 wt% in terms of ductility, fracture toughness, and corrosion resistance based on a series of irradiation data of ferrite steels. The addition of titanium produces the nanoscale dispersion of oxide particles, which leads to a markedly improved high-temperature strength. If titanium is added to excess, however, it creates too much strength, which negatively impacts
Tempered martensite
20 μm
Figure 4 Microstructure of 9Cr-ODS steel showing residual ferrite and tempered martensite.
cold rolling and cold workability. To achieve a balance between strength and workability, a value of 0.2 wt% was selected. Tungsten of 2 wt% is also added in order to improve high-temperature strength by means of solid solution hardening. The microstructure of 9Cr-ODS steels13–18 can be easily controlled by a reversible a–g transformation with a remarkably high driving force of a few hundred MJ m3, as compared with a driving force of irreversible recrystallization with a few MJ m3 for 12Cr-ODS steels. By inducing reversible a–g transformations, 9Cr-ODS steel cladding for fast-reactor fuel elements is currently being manufactured at the JAEA. The microstructure of 9Cr-ODS steel cladding is basically tempered martensite. However, it has been recognized that 9Cr-ODS steel cladding manufactured in an engineering process possesses a dual-phase structure that comprises both tempered martensite and ferrite phases. An example of their microstructure is shown in Figure 4. The ferrite phase appears white, and the elongated phase is indicated by arrows. Their size is about 30–60 mm in length 3–10 mm in width. The formation of a ferrite phase in 9Cr-ODS steel is somewhat unusual, because only the full martensite phase can be expected in 9Cr-ferritic steel without yttria under normalizing and air-cooling conditions. Moreover, the high-temperature strength of manufactured 9Cr-ODS steel is significantly improved by the presence of the ferrite phase.19–21 This is obvious from the creep rupture data shown in Figure 5.20
Oxide Dispersion Strengthened Steels
200
1500
700 ⬚C
L
δ
1400
γ+δ
1300 1200
100 90 80
Solid: with residual ferrite Open: without residual ferrite
70 60 50 10
Temperature (⬚C)
Uni-axial stress (MPa)
245
1100
γ+δ +TiC
1000 900
α +TiC
800
100 1000 Time to rupture (h)
A1
α +TiC +Laves
700 600 500
Mechanically milled without Y2O3
γ + TiC A3
γ + TiC + M23C6 α + M23C6 + TiC
α + M23C6 + laves + TiC
400 10 000
Figure 5 Uni-axial creep rupture strength of 9Cr-ODS steels at 700 C after the normalizing-and-tempering (1050 C 1 h, Ar-gas cooling (AC) = > 780 C 1 h, AC) with and without residual ferrite. Reproduced from Ohtsuka, S.; Ukai, S.; Fujiwara, M.; Kaito, T.; Narita, T. Mater. Trans. 2005, 46, 487.
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 Dispersion Strengthened Steels
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
Oxide Dispersion Strengthened Steels
ODS steels, the a–g-phase transformation could induce the coarsening of oxide particles by disturbing the interface coherency between these particles and the g-phase matrix. 4.08.3.3
Cladding Manufacturing
4.08.3.3.1 Continuous cooling transformation diagram
The preparation of a CCT (continuous cooling 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
Oxide Dispersion Strengthened Steels
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
Oxide Dispersion Strengthened Steels
consists of normalizing at 1050 C for 1 h, followed by tempering at 800 C for 1 h. 4.08.3.4
Creep and tensile properties
The lifetime of a fast reactor fuel pin is most strongly determined by the internal creep rupture strength of the cladding induced by the internal pressure of the fission gas at a temperature of around 700 C. For 9Cr-ODS steel cladding, internal creep rupture data at 650, 700, and 750 C are shown in Figure 15.30 Additionally, the best fit lines for hoop stress versus rupture time at each temperature are shown by solid lines. These results confirmed that creep rupture strengths in the hoop and longitudinal directions of cladding are almost the same, due to their equi-axed grains. The corresponding creep rupture curves for HT931 and austenitic PNC31632 are also presented for comparison. PNC316 is a typical austenitic 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.
Oxide Dispersion Strengthened Steels
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
15 Uniform elongation (%)
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
Oxide Dispersion Strengthened Steels
there is strong strength anisotropy, and the bi-axial creep rupture strength is considerably lower than that of the uni-axial direction. Microstructure 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 <0.25 mass % to attain the recrystallized structure in the ODS ferritic steels. The two types of 12Cr-ODS steels in the chemical composition of 0.23Y2O3–2W–0.4Ti (A3) and 0.34Y2O3–3W– 0.4Ti (A15) were extruded at 1150 C, followed by 60% cold rolling and annealed at 1200 C for 1 h. From these heat-treated bars, the internal creep rupture specimens were machined and tested at 650 C. Figure 1835 exhibits a comparison of the creep rupture strength of recrystallized (A3) and unrecrystallized (A15) 12Cr-ODS steels at 650 C, where
von Mises’ equivalent stress was estimated for the internal hoop stress. The unrecrystallized specimen shows significant strength anisotropy in uni-axial and internal creep rupture strength, whereas the recrystallized specimens reveal decrease of anisotropy, where uni-axial creep rupture strength decreases and internal strength approaches the uni-axial strength. These results demonstrate that the recrystallization process adequately improves the creep rupture strength in the internal hoop direction. Furthermore, softening by recrystallization makes it possible to manufacture cladding by cold-rolling processing. 4.08.4.3 Cold-Rolling Cladding Manufacturing Based on the aforementioned finding, cladding manufacturing tests were conducted using coldrolling pilger mill in 12Cr-ODS ferritic steels with the limited Y2O3 content <0.25 mass % to induce the recrystallized structure, and their internal creep rupture properties were evaluated at 700 C, not at 650 C. The chemical composition of the manufactured cladding is listed in Table 1, where the specimens are denoted as F1 to F4. The four levels of Y2O3 contents were selected in the range of <0.25 mass %, and the titanium content ranged from 0.13 to 0.31 mass %. Cold rolling by a pilger mill was
Solid: uni-axial Open: internal
600
Equivalent stress (MPa)
500 400
300
200 A3-2(1123 K extrusion) A3-1(1423 K extrusion) A15
100 10
100
1000
10 000
Time to rupture (h) Figure 18 Creep rupture strength of recrystallized and unrecrystallized 12Cr-ODS steel claddings by using pressurized specimens at 650 C. Reproduced from Ukai, S.; Nishida, T.; Okada, H.; Okuda, T.; Fujiwara, M.; Asabe, K. J. Nucl. Sci. Technol. 1997, 34(3), 256–263.
253
Oxide Dispersion Strengthened Steels
Table 1 steels
Chemical composition of F1–F4 specimens with different titanium and yttria contents (mass %) in 12Cr-ODS
Specimen no.
F1 F2 F3 F4
Chemical composition (mass %) C
Cr
W
Ni
Ti
Y2O3
O
N
Ar
0.065 0.054 0.065 0.056
11.8 11.8 11.8 11.7
1.92 1.94 1.93 1.92
0.03 0.03 0.03 0.03
0.13 0.13 0.22 0.31
0.08 0.13 0.22 0.24
0.08 0.05 0.09 0.04
0.010 0.010 0.012 0.010
0.005 0.004 0.005 0.004
Source: Reproduced from Ukai, S.; Okuda, T.; Fujiwara, M.; Kobayashi, T.; Mizuta, S.; Nakashima, H. J. Nucl. Sci. Technol. 2002, 39(8), 872–879.
F2
F3
F4
Longitudinal
F1
Transverse
20 mm
20 mm
Figure 19 Optical microstructure of the F1, F2, F3, and F4 specimens in the final claddings. Reproduced from Ukai, S.; Okuda, T.; Fujiwara, M.; Kobayashi, T.; Mizuta, S.; Nakashima, H. J. Nucl. Sci. Technol. 2002, 39(8), 872–879.
repeated twice with a reduction ratio of about 50% per rolling. The intermediate heat-treatment to soften the cold-rolled cladding was performed at 1100 C for 30 min, and the final heat-treatment was performed at 1150 C for 0.5 h. Figure 19 shows the optical microstructures of the manufactured claddings in the longitudinal and transverse directions.36 All of the specimens seem to be recrystallized. However, the extent of recrystallization depends on the yttria and titanium contents. In the transverse cross-section, the grain size becomes slightly finer with increasing yttria and titanium contents from F1 to F4. In the F4 specimen, the elongated grains can be still seen and the aspect ratio in the longitudinal (L) and transverse (T) directions is large, whereas the aspect ratio of specimen F1 appears to be nearly unity. These findings show that F4 specimen with higher yttria and titanium contents did not attain the perfectly recrystallized grain structure by the annealing of 1150 C for 0.5 h.
4.08.4.4
Internal Creep Rupture Property
The internal creep rupture properties of the manufactured 12Cr-ODS steels at 700 C are shown in Figure 20.36 Increasing the yttria and titanium contents improves the internal creep rupture strength (F4 > F3 > F2 > F1). The uni-axial creep rupture strength for F4 is also plotted; there is the strength anisotropy between the uni-axial and internal hoop directions. This strength anisotropy can be associated with the slightly elongated grain structure shown in Figure 19. The stress–strain rate relationship was 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
Oxide Dispersion Strengthened Steels
1000
Hoop stress (MPa)
F1 F2 F3 F4 F4 (uni-axial, this work)
Uni-axial
100
Internal, bi-axial direction
10
100 1000 Time to rupture (h)
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
Oxide Dispersion Strengthened Steels
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.
Oxide Dispersion Strengthened Steels
at an operating temperature of 600 C. It is known that 9Cr-ODS steels have superior compatibility with sodium, but their corrosion resistance is not adequate for SCPW and LBE at a temperature >600 C. Thus, the most critical issue for the application of 9Cr-ODS steels to 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 <16 wt%. 4.08.5.3
Mechanical Properties
The addition of Al to ODS ferritic steels sometimes softens their creep and tensile properties. Figure 2544 shows the effects of Al addition and Cr content on the
0.8 Aged at 500 ⬚C
0.7 Absorbed energy (J)
256
0.6
As-received
0.5 0.4 0.3 0.2 0.1 0.0 12
100 h
1000 h 4300 h 10 000 h 14
16
18
20
22
Cr content (mass%) Figure 24 Aging embrittlement of high Cr-ODS steels with respect to Cr content. Absorbed fracture energy was measured at room temperature with the use of miniaturized Charpy V-notch (CVN) specimen which measures 1.5 mm square with 20 mm length. Reproduced from Kimura, A.; Kasada, R.; Iwata, N.; et al. In Proceedings of ICAPP ’09, Tokyo, Japan, May 10–14, 2009; Paper 9220.
tensile strength of high Cr-ODS steels. A decrease in UTS versus Al content is obvious in 16Cr-ODS steels; this dependency becomes weaker at higher temperature. The effect of Cr concentration on the UTS is not so obvious between 13.7 and 17.3 wt% at 450 and 700 C. In the case of 9Cr–12Cr-ODS steels, hightemperature strength is considerably enhanced by the uniform dispersion of Y–Ti complex oxide (Y2Ti2O7) particles. In Al-added high Cr-ODS steels, however, Y–Al complex oxides and/or Al oxides are formed rather than Y–Ti complex oxides, which leads to larger oxide particles, causing a degradation of hightemperature mechanical strength. Therefore, Hf or Zr, which form thermodynamically stable oxides, were added to form Y–Hf or Y–Zr complex oxide particles rather than Y–Al complex oxide. The process of manufacturing these materials is exactly the same as that used for 9Cr-ODS ferritic steels: MA by an attrition type ball mill and hot extrusion at a nominal temperature of 1150 C. The extruded bars were provided for mechanical tests. The creep rupture properties of Al-added high Cr-ODS steels are summarized in Figure 26.44 The creep strength of the standard steel is generally lower than that of Al-free steel. On the other hand, the addition of Zr and Hf induces an improved creep strength which approaches that of Al-free steel. Furthermore, it was found that their
Oxide Dispersion Strengthened Steels
1400
1000 Fe–16Cr–0W–0.1–0.35Y2O3
1200
700 ⬚C
700 ⬚C
450 ⬚C
800 UTSave. (MPa)
1000
UTSave. (MPa)
257
800 600 400
Fe–4Al–0W–0.1–0.35Y2O3
450 ⬚C
600 400 200
200
0
0 0
3 1 2 Al concentration (mass %)
4
12
13 14 15 16 17 18 Cr concentration (mass %)
19
Figure 25 The UTS with respect of Al and Cr content in Al added high Cr-ODS steels at 450 C and 700 C. Reproduced from Furukawa, T.; Ohtsuka, S.; Inoue, M.; et al. In Proceedings of ICAPP ’09, Tokyo, Japan, May 10–14, 2009; Paper 9221.
103
Stress (MPa)
700 ⬚C
102 0Al (16Cr–0.1Ti–0.35Y2O3) Standard (15Cr–4Al–2W–0.1Ti–0.35Y2O3) 0.63Zr (15Cr–4Al–2W–0.63Zr–0.1Ti–0.35Y2O3) 0.62Hf (15Cr–4Al–2W–0.62Hf–0.1Ti–0.35Y2O3)
10 101
103 102 Time to rupture (h)
104
105
Figure 26 Creep rupture strength of various Al added high Cr-ODS steels in hoop direction by using pressurized specimens at 700 C. Reproduced from Furukawa, T.; Ohtsuka, S.; Inoue, M.; et al. In Proceedings of ICAPP ’09, Tokyo, Japan, May 10–14, 2009; Paper 9221.
fracture elongation and reduction of area are slightly higher than those of Al-free steel. Under microstructural observation by TEM, oxide particles consisting of Y3A5O12, YAlO3, and Al2O3, were observed in typical Al-added ODS steel, whereas this Y–Al complex oxide can be changed to Y2Hf2O7 in Hf-added ODS steel. The improved creep rupture strength in Hf-added ODS steel could be attributed to the nanosize dispersion of the Y2Hf2O7 complex oxide. 4.08.5.4
Cladding Manufacturing
16Cr–4Al-ODS steels exhibit a full ferrite structure without a–g-phase transformation, and the manufacture of their cladding is effected by means of cold-rolling
and recrystallization annealing, which is the same as 12Cr-ODS steel described in Section 4.08.4. From mother tubes of 16Cr–2W–0.1Ti–4Al– 0.35Y2O3 with an 18 mm outer diameter and 3 mm thickness, cold-rolling was repeated four times using a pilger mill; the reduction rate for each rolling reached about 45%. Annealing is required to soften the deformed structure in order to make possible the next round of cold rolling. A recrystallized structure cannot be reproduced by the final heat treatment once recrystallization has taken place in the intermediate annealing.45 A heat treatment test was therefore conducted to determine an appropriate annealing temperature to induce softening by dislocation recovery but without recrystallization. The temperature for the final heat treatment was selected in order to realize recrystallized grains. Figure 27 illustrates an orientation image map (OIM) measured after heat treatment at 900 C for 1 h, 1000 C for 1 h, and 1150 C for 1 h, from which an evolution of recrystallization can be clearly identified in the cladding.46 The primary recrystallization begins with an {111}<112> orientation, which is designated by the color blue at 900 C annealing. When the temperature increases to 1000 C, a {110}<100> Goss orientation, designated by green, partially appears. The primary recrystallization orientation of {111}<112> is almost replaced by a {110}<100> Goss orientation when the temperature is elevated to 1150 C, which corresponds to a secondary recrystallization. Thus, the final heat treatment was conducted at 1150 C for 1 h in order to create a perfectly recrystallized structure. The manufactured cladding has secondary recrystallized grains with a {110}<100> Goss orientation, which was formed by annealing from the cold rolled {112}<110> orientation through primary recrystallization of the {111}<112> orientation.47
258
Oxide Dispersion Strengthened Steels
4.08.6 Existing ODS Steel Cladding The basic chemical composition of the representative ODS steels is summarized in Table 2. ODS steels are divided into two groups which have either been
200 μm
200 μm
200 μm
(a)
(b)
(c)
Figure 27 OIM for 16Cr–4Al–ODS steels heat-treated at 900 C for 1 h (a), 1000 C for 1 h (b), and 1150 C for 1 h (c). Reproduced from Ukai, S.; Ohnuki, S.; Hayashi, S.; et al. In Proceedings of ICAPP’09, Tokyo, Japan, May 10–14, 2009; Paper 9232.
Table 2
commercialized or are under development. The first group includes Incoloy MA956 and PM2000. The former is produced by what was formerly the International Nickel Company (INCO) and is now the Special Metals (SM) Company. The latter is a product of the Plansee Company of Austria. MA956 and PM2000 are 20% Cr-ODS steels containing 5% Al, which exhibit superior resistance to oxidation and corrosion in hot gases at temperatures >1000 C. Tubes, sheets, and bars made from these steels are commercially used in various stationary and 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’e´nergie 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’e´nergie Nucleaire – Studiecentrum voor Kernenergie.
Development
JAEA/Japan KU/Japan
Oxide Dispersion Strengthened Steels
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
100 1000 Time to rupture (h)
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
Oxide Dispersion Strengthened Steels
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
Oxide Dispersion Strengthened Steels
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
Oxide Dispersion Strengthened Steels
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)
Oxide Dispersion Strengthened Steels
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
Oxide Dispersion Strengthened Steels
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 (<10%) than that of the unirradiated ones at all test temperatures, due to irradiation hardening. Figure 40 plots uniform elongation before and after irradiation as a function of test temperature.67 Uniform elongation for unirradiated F94 and F95 cladding is almost the same at all test temperatures, and that of M93 is lower in relation to strength. Uniform elongation in the hoop direction for all three claddings is more than 3% at these test temperatures, though that of 1DS was particularly low (<1%) due to its microstructural anisotropy, as shown in Figure 17. Figure 40 indicates that there is no significant degradation in uniform elongations for F94, F95, and M93, due to irradiation. This indicates that the microstructural improvement by recrystallization or a–g-phase transformation is quite effective in maintaining well-balanced mechanical properties for ODS steel cladding, especially those of strength and ductility, not only for as-received conditions but also following irradiation. In-pile creep rupture tests were conducted in JOYO using the Material Testing Rig with Temperature Control (MARICO-2) as a new irradiation test device.68 The test specimens were prepared from the claddings of 9Cr-ODS steel (Mm14) and 12Cr-ODS steel (F14). Both end-plugs of the specimens were joined by means of pressurized resistance welding (PRW). The hoop stress was set by adjusting the pressure of the enclosed helium gas. To identify the rupture of time and specimens, a unique blend of stable
20
265
40 Fraction (%)
Unirradiated condition
15 10 5 0 20
Unirradiated condition
30 20 10 0
T = 500 ⬚C, dose = 5 dpa
15
40 Fraction (%)
10 5 0 20 T = 500 ⬚C, dose = 50 dpa
15
20 10 0
5
40
0 20 T = 500 ⬚C, dose = 150 dpa
15 10 5
T = 600 ⬚C, dose = 5 dpa
30
10 Fraction (%)
Fraction (%)
Fraction (%)
Fraction (%)
Fraction (%)
Oxide Dispersion Strengthened Steels
T = 600 ⬚C, dose = 50 dpa
30 20 10 0
0 0
2
4
6
8 10 12 14 16 18 20 22 Particle size (nm) Fraction (%)
40
0
2
4
6
8
10 12 14 16 18 20 22
Particle size (nm)
Unirradiated condition
30 20 10
Fraction (%)
0 40 T = 700 ⬚C, dose = 5 dpa
30 20 10
Fraction (%)
0 40 T = 700 ⬚C, dose = 50 dpa
30 20 10
Fraction (%)
0 40 T = 700 ⬚C, dose = 150 dpa
30 20 10 0 0
2
4
6
8 10 12 14 16 18 20 22 Particle size (nm)
Figure 38 Particle size (diameter) distribution for 9Cr-ODS steel irradiated at 500, 600, and 700 C to doses of 0, 5, 50, and 150 dpa. Reproduced from Allen, T. R.; Gan, J.; Cole, J. I.; et al. J. Nucl. Mater. 2008, 375, 26–37.
266
Historical survey of yttrium–titanium-oxides reported to be stable under radiation
Author
Material
Irradiation particle (dpa)
Pareige et al.74 Asano et al.75
12YWT MA957
Hide et al.76
MA957
150 keV Fe 1 MeV He (þ) 4 MeV Ni 42 keV He at 25 C (þ) 200 keV C
Hide et al.76
MA957
Little77
DT2203YO5
Saito et al.61
13Cr–0.5TiO2 –0.2Y2O3
220 keV He at 25 C (+)3 MeV Niþ 52 MeV Cr6 þ (þ) 4 MeV He 1 MeV electron
Kinoshita et al.78
13Cr-ODS (+) Nb, V, Zr
1 MeV electron
Akasaka et al.79
9Cr and 12Cr-ODS
JOYO
Mathon et al.80
MA957
Monnet et al.63
DY EM10þY2O3EM10 þ MgO
Thermal neutrons (OSIRIS) 1 MeV Helium
Kimura et al.81
(13–19)Cr–4Al-ODS
a
Dose (dpa)
Dose rate (dpa s1)
Result
300 450 650 475 525 575 625 475 525 475
0.7 150
1.9 104 2 103
Stable dispersions Stable oxides
200
1.0–1.4 102
Stable oxides
150
3.0 103
Stable oxides
50
3.0 104
Stable oxides
400 500 350 450 330 400 450 500 325
12
2.2 103
Stable oxides
15
2 103
Stable oxides
7.0 2.5 14.0 15.0 0.8, 2.0 3.5, 5.5 0.05
Not reporteda
Stable oxides
1 1014 n cm2 (E > 1 MeV) Not reported
Stable dispersions No change in oxide particles
20
1 104
No reported change in oxide size
Temperature ( C)
400 300 –500
Typical fast reactor displacement rates in the driver fuel portion of the core are 1 106 dpa s1. Source: Reproduced from Allen, T. R.; Gan, J.; Cole, J. I.; et al. J. Nucl. Mater. 2008, 375, 26–37.
Oxide Dispersion Strengthened Steels
Table 3
Oxide Dispersion Strengthened Steels
267
Historical survey of yttrium–titanium-oxides reported change size under radiation
Table 4 Author
Material
Irradiation
Temperature ( C)
Dose (dpa)
Dose rate (dpa s1)
Result
Yamashita et al.64
IDS (11Cr)
JOYO
450–561
21
Not reporteda
IDK (13Cr) DT2203YO5
Small particles disappear. Average particles increase slightly with increasing temperature or dose.
Phenix
400–580
81
Not reporteda
Monnet et al.63
DT2203YO5
Phenix
400–580
81
Not reporteda
Monnet et al.63
DY EM10 + Y2O3 EM10 + MgO
1 MeV and 1.2 MeV Electron
300–550
100
3–6 103
Oxide particles are totally dissolved (small oxides) or reduced in size and were surrounded by a halo of smaller oxides (large oxides). Disappearance of small oxides and significant halo of smaller oxides at higher temperatures and doses. Oxides stable at 400 C under 1.0 MeV electrons but dissolve under 1.2 MeV.
Dubuisson et al.62
a Typical fast reactor displacement rates in the driver fuel portion of the core are about 1 106 dpa s1. Source: Reproduced from Allen, T. R.; Gan, J.; Cole, J. I.; et al. J. Nucl. Mater. 2008, 375, 26–37.
(Fluence; ⫻ 1026 n m–2 (E > 0.1 MeV)
Yield strength (MPa)
(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.
10 Uniform elongation (%)
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
Oxide Dispersion Strengthened Steels
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
Oxide Dispersion Strengthened Steels
(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 Phe´nix
Fuel pins with DT2203Y05 cladding were irradiated in an experimental capsule placed in a special subassembly in Phe´nix. 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 <10 nm in size and with a density higher than the original oxide
density, was observed in the lower part of the fuel pin at temperatures <500 C. These precipitates were found to be a0 -phase, as shown in Figure 43(a).63 At irradiation temperatures above 500 C, precipitation of w-phase was uniformly distributed throughout all grains, as shown in Figure 43(b). Their chemical composition was slightly different from that of the intergranular w-phase that was present before irradiation. At a high temperature and low dose, w-phase is replaced by the thermal precipitation of Laves phase, as shown in Figure 43(c). From the results of tensile tests at levels corresponding to the fissile column, rupture occurred without striction, and uniform and total elongations were equal. The elongation values reached 0.2% close to the maximum dose. These results indicate that DT2203Y05 cladding was highly embrittled by irradiation. At the bottom of the fuel pin, where the temperature is below 500 C, a0 precipitation, oxide redistribution and dislocation loops are the main features of the microstructure. Dubuisson63 pointed out that dislocation glides on the dislocation denuded bands in the hardened materials, that the deformation is all localized in these bands, and that this heterogeneous shear could nucleate cracks that then propagate along these channels. At higher temperatures, w precipitation induces a loss of ductility.
4.08.9 Summary The formation process of nanosized oxide particles through decomposition by MA and subsequent precipitation by annealing was reviewed. Based on
270
Oxide Dispersion Strengthened Steels
information concerning the irradiation embrittlement of DT2203Y05 cladding produced by CEN-SCK Mol due to the formation of a0 -phase below 500 C and w-phase above 500 C, 9Cr-ODS and 12Cr-ODS steels containing low Cr and low Ti were developed. The manufacture of cladding and improvement of the creep rupture strength in the hoop direction were successfully achieved by introducing a–g reverse transformation or recrystallization. 9Cr-ODS steel has a unique structure consisting of tempered martensite and residual ferrite that induces superior strength through finely dispersed oxide particles, which are promising candidates for advanced SFR fuel cladding. 16Cr–4Al–ODS steels present an advantage due to their superior resistance to corrosion and oxidation in LBE and SCPW environments. There is still uncertainty concerning the irradiation performance of ODS steels, such as the oxide particle dissolution due to their diffusion-based mechanism. In order to substantiate the use of cladding materials as advanced fast reactor fuels, abundant ODS cladding fuel pins should be irradiated, and their results should provide feedback contributing to the further improvement of ODS steels.
References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16.
Odette, G. R.; Alinger, M. J.; Wirth, B. D. Annu. Rev. Mater. Res. 2008, 38, 471–503. Okuda, T.; Fujiwara, M. J. Mater. Sci. Lett 1995, 14, 1600. Ukai, S.; Harada, M.; Okada, H.; Inoue, M.; Nishid, T.; Fujiwara, M. J. Nucl. Mater. 1993, 204, 65–73. Ukai, S.; Fujiwara, M. J. Nucl. Mater. 2002, 307–311, 749. Okuda, T.; Ukai, S.; Miyahara, K. Materia (Japan) 1999, 39, 954. Larson, D. J.; Maziasz, P. J.; Kim, I.-S.; Miyahara, K. Scr. Mater. 2001, 44, 359. Miller, M. K.; Kenik, E. A.; Russell, K. F.; Heatherly, L.; Hoelzer, D. T.; Maziasz, P. J. J. Nucl. Mater. 2003, A353, 140. Alinger, M. J.; Odette, G. R.; Hoelzer, D. T. J. Nucl. Mater. 2004, 382, 329–333. Kim, S.-W.; Shobu, T.; Ohtsuka, S.; Kaito, T.; Inoue, M.; Ohnuma, M. Mater. Trans. 2009, 50(4), 917–921. Kliniankou, M.; Lindau, R.; Moslang, A. J. Nucl. Mater. 2004, 329–333, 347–351. Smith, J. V. American Society for Testing Materials, Powder diffraction file, 13. Inorganic; Philadelphia, PA, 1962–1965. Luo, C. P.; Dahmen, U. Acta. Mater. 1998, 46, 2063. Ukai, S.; Nishida, T.; Okuda, T.; Yoshitake, T. J. Nucl. Sci. Technol. 1998, 35, 294. Ukai, S.; Nishida, T.; Okuda, T.; Yoshitake, T. J. Nucl. Mater. 1998, 1745, 258–263. Ukai, S.; Muzuta, S.; Fujiwara, M.; Okuda, T.; Kobayashi, T. J. Nucl. Sci. Technol. 2002, 39, 778. Ukai, S.; Mizuta, S.; Fujiwara, S.; Okuda, T.; Kobayashi, T. J. Nucl. Mater. 2002, 307–311, 758.
17. 18. 19. 20. 21. 22. 23. 24. 25. 26. 27. 28. 29.
30. 31. 32. 33. 34. 35. 36. 37. 38. 39. 40. 41. 42. 43. 44. 45. 46. 47. 48. 49.
Ukai, S.; Kaito, T.; Otsuka, S.; Narita, T.; Fujiwara, M.; Kobayashi, T. ISIJ Int. 2003, 43, 2038. Ukai, S.; Narita, T.; Alamo, A.; Pamentier, P. J. Nucl. Mater. 2004, 329–333, 356. Ohtsuka, S.; Ukai, S.; Fujiwara, M.; Kaito, T.; Narita, T. J. Nucl. Mater. 2004, 329–333, 372–376. Ohtsuka, S.; Ukai, S.; Fujiwara, M.; Kaito, T.; Narita, T. Mater. Trans. 2005, 46, 487. Ohtsuka, S.; Ukai, S.; Fujiwara, M.; Kaito, T.; Narita, T. J. Nucl. Mater. 2006, 351, 241. Yamamoto, M.; Ukai, S.; Hayashi, S.; Kaito, T.; Ohtsuka, S. Mater. Sci. Eng. A 2010, 527, 4418–4423. Ukai, S.; Ohtsuka, S. Energy Mater. 2007, 2(1), 26–35. Nishizawa, T.; Ohnuma, I.; Ishida, K. Mater. Trans. JIM 1997, 38(11), 950–956. Martion, J. W.; Doherty, R. D. Stability of Microstructure in Metallic Systems; Cambridge University Press: Cambridge, 1976; p 173. Ukai, S.; Ohtsuka, S.; Kaito, T.; Sakasegawa, H.; Chikata, N.; Hayashi, S.; Ohnuki, S. Mater. Sci. Eng. A 2009, 510–511, 115–120. Tabor, D. The Hardness of Metals; Oxford university press: Oxford, 1951. Yamamoto, M.; Ukai, S.; Hayashi, S.; Kaito, T.; Ohtsuka, S. J. Nucl. Mater. 2011, 237–240, 417. Lambard, V. Development of ODS ferritic-martensitic steels for application to high temperature and irradiation environment; Rapport CEA-R-5918; France 2000. Allen, T.; Burlet, H.; Nanstad, R. K.; Samaras, M.; Ukai, S. Mater. Res. Soc. Bull. 2009, 34(1), 20–27. Patriarca, P. In Proceedings of the Topical Conference on Ferritic Alloys for Use in Nuclear Energy Technologies, Snowbird, Utah, June 19–23 1983; p 107. Shibahara, I.; Ukai, S.; Onose, S.; Shikakura, S. J. Nucl. Mater. 1993, 204, 131. Ukai, S.; Harada, M.; Okada, H.; et al. J. Nucl. Mater. 1993, 204, 74–80. Fischer, J. J. US Patent 4075010, 1978. Ukai, S.; Nishida, T.; Okada, H.; Okuda, T.; Fujiwara, M.; Asabe, K. J. Nucl. Sci. Technol. 1997, 34(3), 256–263. Ukai, S.; Okuda, T.; Fujiwara, M.; Kobayashi, T.; Mizuta, S.; Nakashima, H. J. Nucl. Sci. Technol. 2002, 39(8), 872–879. Arzt, E. Res. Mech. 1991, 31, 399–453. Srolovitz, D. J.; Perkovic-Luton, R. A.; Luton, M. J. Phil. Mag. A 1983, 48, 795. Scattergood, R. O.; Bacon, D. J. Phil. Mag. A 1975, 31, 170. Foreman, A. J. E.; Makin, M. J. Phil. Mag. A 1966, 14, 911. Stoller, R. E.; Zinkle, S. J. J. Nucl. Mater. 2000, 283–287, 349. Guerin, Y.; Was, G. S.; Zinkle, S. J. Mater. Res. Soc. Bull. 2009, 34(1), 10–14. Kimura, A.; Kasada, R.; Iwata, N.; et al. In Proceedings of ICAPP ’09, Tokyo, Japan, May 10–14, 2009; Paper 9220. Furukawa, T.; Ohtsuka, S.; Inoue, M.; et al. In Proceedings of ICAPP ’09, Tokyo, Japan, May 10–14, 2009; Paper 9221. Narita, T.; Ukai, S.; Kaito, T.; Otsuka, S.; Kobayashi, T. J. Nucl. Sci. Technol. 2004, 41(10), 1008. Ukai, S.; Ohnuki, S.; Hayashi, S.; et al. In Proceedings of ICAPP’09, Tokyo, Japan, May 10–14, 2009; Paper 9232. Dorner, D.; Zaefferer, S.; Raabe, D. Acta. Mater. 2007, 55, 2519. Seran, J. L.; Levy, V.; Dubuisson, P.; et al. In Effects of Radiation on Materials: 15th International Symposium; Philadelphia, PA 1992; pp 1209–1233, ASTM STP 1125. Lippens, M.; Ehrlich, K.; Levy, V.; Brown, C.; Calzabini, A. In Proceedings of the International Conference on Ferritic Alloys for Use in Nuclear Energy Technologies, Utah, 1983; pp 329–334.
Oxide Dispersion Strengthened Steels 50. DeWilde, L.; Gedopt, J.; DeBurbure, S.; Delbrassibe, A.; Driesen, C.; Kazimierzak, B. In Proceeding in Materials for Nuclear Reactor Core Applications; BNES: London, 1987; 271–276. 51. Huet, J. J.; Coheur, L.; De Bremaecker, A.; et al. Nucl. Technol. 1985, 70, 215–219. 52. Hamilton, M. L.; Gelles, D. S. PNNL-13168, Feb 2000. 53. Yoshida, E.; Kato, S. J. Nucl. Mater. 2004, 329–333, 1393–1397. 54. Machut, M.; Sridharan, K.; Li, N.; Ukai, S.; Allen, T. J. Nucl. Mater. 2007, 371, 134–144. 55. Furukawa, T.; Muller, G.; Schumacher, G.; et al. J. Nucl. Sci. Technol. 2004, 41(3), 265–270. 56. Takaya, S.; furukawa, T.; Aoto, K.; et al. Fall meeting of Japan Atomic Energy Society, Sept 2008. 57. Lee, J. H.; Kimura, A.; Kasada, R.; et al. In Proceedings of ICAPP ’09, Tokyo, Japan, May 10–14, 2009; Paper 9223. 58. Kaito, T.; Narita, T.; Ukai, S.; Matsuda, Y. J. Nucl. Mater. 2004, 329–333, 1388–1392. 59. Chen, Y.; Sridharan, K.; Ukai, S.; Allen, T. R. J. Nucl. Mater. 2007, 371, 118–128. 60. Allen, T. R.; Gan, J.; Cole, J. I.; et al. J. Nucl. Mater. 2008, 375, 26–37. 61. Saito, J.; Suda, T.; Yamashita, S.; et al. J. Nucl. Mater. 1998, 258–263, 1264. 62. 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, 1999; p 882, ASTM STP 1325. 63. Monnet, I.; Dubuisson, P.; Serruys, Y.; Ruault, M. O.; Kaitasov, O.; Jouffrey, B. J. Nucl. Mater. 2004, 335, 311. 64. Yamashita, S.; Oka, K.; Ohnuki, S.; Akasaka, N.; Ukai, S. J. Nucl. Mater. 2002, 307–311, 283. 65. Rechtin, M. D.; Wiedersich, H. Radiat. Eff. 1977, 31, 181. 66. Zinkle, S. J.; Kinoshita, C. J. Nucl. Mater. 1997, 251, 200.
67. 68. 69. 70. 71. 72. 73. 74. 75. 76.
77.
78. 79. 80. 81.
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Yoshitake, T.; Abe, Y.; Akasaka, N.; Ohtsuka, S.; Ukai, S.; Kimura, A. J. Nucl. Mater. 2004, 329–333, 342–346. Kaito, T.; Ohtsuka, S.; Inoue, M.; et al. J. Nucl. Mater. 2009, 386–388, 294–298. Gelles, D. S. Fusion Materials, Semiannual Progress Report for Period Ending March 31, DOE/ER-0313/16 1994; pp 146–160. Seki, M.; Hirako, K.; Kono, S.; Kihara, Y.; Kaito, T.; Ukai, S. J. Nucl. Mater. 2004, 329–333, 534–1538. Ukai, S.; Kaito, T.; Seki, M.; Mayorshin, A. A.; Shishalov, O. V. J. Nucl. Sci. Technol. 2005, 42(1), 109–122. Kaito, T.; Ukai, S.; Povstyanko, A. V.; Efimov, V. N. J. Nucl. Sci. Technol. 2009, 46(6), 529–533. Bottcher, J.; Ukai, S.; Inoue, M. Nucl. Technol. 2002, 138, 238–245. Pareige, P.; Miller, M. K.; Stoller, R. E.; Hoelzer, D. T.; Cadel, E.; Radiguet, B. J. Nucl. Mater. 2007, 360, 136. Asano, K.; Kohno, Y.; Kohyama, A.; Suzuki, T.; Kusanagi, H. J. Nucl. Mater. 1988, 155–157, 928. Hide, K.; Sekimura, N.; Fukuya, K.; et al. In Effects of Radiation on Materials: 14th International Symposium, vol. I.; Packan, N. H., Stoller, R. E., Kumar, A. S. Eds.; American Society for Testing and Materials: Philadelphia, PA, 1988; p 61, ASTM STP 1046. Little, E.A. In Effects of Radiation on Materials: 17th International Symposium; Gelles, D.S., Nanstad, R.K., Kumar, A.S., Little, E.A., Eds.; American Society for Testing and Materials: Philadelphia, PA, 1996; p 739, ASTM STP 1270. Kinoshita, H.; Akasaka, N.; Takahashi, H.; Shibahara, I.; Onose, S. J. Nucl. Mater. 1992, 191–194, 874. Akasaka, N.; Yamashita, S.; Yoshitake, T.; Ukai, S.; Kimura, A. J. Nucl. Mater. 2004, 329–333, 1053. Mathon, M.-H.; de Carlan, Y.; Averty, X.; Alamo, Ch.-H.; de Novion. J. ASTM Int. 2005, 2(8), Paper ID JAI 12381. Kimura, A.; Cho, H.; Toda, N.; et al. J. Nucl. Sci. Technol. 2007, 44(3), 323.
4.09
Welds for Nuclear Systems
G. A. Young, M. J. Hackett, J. D. Tucker, and T. E. Capobianco Knolls Atomic Power Laboratory, Schenectady, NY, USA
ß 2012 Elsevier Ltd. All rights reserved.
4.09.1
Welding Defects
273
4.09.1.1 4.09.1.1.1 4.09.1.1.2 4.09.1.1.3 4.09.1.2 4.09.1.2.1 4.09.1.2.2 4.09.1.3 4.09.2 4.09.2.1 4.09.2.1.1 4.09.2.1.2 4.09.3 4.09.3.1 4.09.3.2 4.09.3.3 4.09.4 4.09.4.1 4.09.4.2 4.09.4.3 4.09.4.4 References
Supersolidus Cracking Solidification cracking Liquation cracking Hot tearing Subsolidus Cracking Precipitation-induced cracking Segregation-induced cracking Other Welding Defects Stresses and Strains in Welds Quantification of Residual Stresses and Strains Elastic stress Plastic strain In-Service Performance Environmentally Assisted Cracking Microchemical Changes Microstructural Changes Weldability of Specific Alloy Systems Low-Alloy Steels Austenitic Stainless Steels Nickel-Based Alloys Zirconium Alloys
274 274 278 278 279 279 282 285 285 287 287 287 289 289 291 293 294 294 294 295 295 296
Abbreviations AMIS CTE EAC GTAW HAZ PIC PMZ SCC SIC SMAW SS Zr-4
Average intragrain misorientation Coefficient of thermal expansion Environmentally assisted cracking Gas tungsten arc welding Heat-affected zone Precipitation-induced cracking Partially melted zone Stress corrosion cracking Segregation-induced cracking Shielded metal arc welding Stainless steel Zircaloy-4
4.09.1 Welding Defects Freedom from cracks or other sharp discontinuities is of primary concern in weld quality and performance.
Crack-like defects can degrade component lifetime by eliminating the initiation stage of phenomena such as fatigue or stress corrosion. Similarly, other flaws (e.g., lack of fusion defects, gas porosity, and inclusions) can act to magnify global stresses, produce locally aggressive environments via their occluded geometry or composition, and initiate cracking. In order to mitigate these flaws, it is critical to differentiate between defect types. The first step in the prevention of cracking is understanding the temperature range over which the cracking occurs. The primary measure is to determine whether defects are ‘hot cracks’ or ‘cold cracks,’ that is, whether they form above or below the solidus temperature of the alloy. Secondly, the location of the crack in the weld (composite region, unmixed zone, partially melted zone, heat-affected zone) and in the microstructure (solidification boundary, crystallographic boundary, etc.) must be determined.1 Once these distinctions are made, 273
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strategies to eliminate cracking can be developed via changes to the welding process, weld parameters, filler metal, joint design, fixturing, and/or postweld heat treatment. The supersolidus/subsolidus distinction, combined with the unifying concept of homologous temperature, illustrates the commonality of welding defects and degradation mechanisms across alloy systems as shown in Figure 1. Supersolidus ‘hot cracking’ defects include solidification cracking, liquation cracking, and hot tearing. Subsolidus ‘cold cracking’ includes precipitation, transformation, and segregationinduced cracking (SIC) mechanisms. 4.09.1.1
Supersolidus Cracking
Cracks that occur between the liquidus and the solidus are commonly termed ‘hot cracks.’ Hot cracks can be further differentiated depending on whether they occur in the composite region of the weld bead on cooling (solidification-type) or in the partially melted or heat-affected zone where a composition gradient or low melting phase acts to locally depress the solidus of the alloy (liquation-type). Additionally, a third type of ‘hot cracking,’ that is, hot tearing, can be
distinguished from solidification and liquation cracks. Hot tears are primarily mechanical in nature, driven by the geometry and stresses in the weldment. A schematic of the locations and representative micrographs of the different types of supersolidus cracking are shown in Figure 2. 4.09.1.1.1 Solidification cracking
Solidification cracks occur in the mushy zone of a weld bead on cooling, as the strains that develop exceed the ductility of the (solid þ liquid) mixture. The modern theory of solidification cracking was developed by Borland,2 who highlighted the importance of the quantity and distribution of the liquid near the terminal phase of solidification, as well as the stresses that act on that liquid. The primary factors that affect hot cracking are summarized in texts by Kou, Messler, and others.3–6 These factors are listed below: 1. The solidification temperature range : The larger the solidification temperature range, the more extensive the solid þ liquid mushy zone, which is susceptible to cracking. While large solidification temperature ranges may promote crack healing via
Homologous temperature (T/Tmelt)
0.1
0.2
0.3
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0.7
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0.9
1.0
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Supersolidus
Subsolidus Segregation-induced cracking: hydrogen
Precipitation-induced cracking: also known as ductility dip cracking, strain-age cracking, reheat cracking, subsolidus cracking
Liquation -type
‘Hot tearing’ (mechanical)
Solidification -type
Environmental degradation Hydrogen embrittlement (low temp. crack propagation, hydriding)
Impurity segregation via diffusion Ordering reactions/brittle second phases precipitate
Radiation-induced segregation Radiation hardening
Creep-rupture
Liquid and solid metal embrittlement Figure 1 Comparison of the typical temperature ranges for the different types of weld cracking (top) and forms of environmental degradation common to nuclear power systems (bottom). All temperature ranges are approximate, based on the homologous temperature of the alloy under consideration.
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Figure 2 Illustration of the different forms of hot cracking. Liquation cracks occur in the partially melted and/or heat-affected zone (HAZ) of the weld bead being deposited (top). Solidification-type cracks occur in the composite region of the weld during solidification (middle), and hot tears are dominated by mechanical forces and occur at macroscopic notches (bottom). Note that the cracks are all from nickel–chromium alloy welds but are not all from the same weldment.
backfilling, solute-rich ‘backfill’ may have degraded properties relative to the bulk weld metal. The approximate solidification temperature ranges of several alloys used in nuclear construction are shown in Figure 3. 2. The solidification path : Solidification crack susceptibility is markedly influenced by the type and distribution of solid phases, for example, initial d-ferrite formation from the liquid in austenitic stainless steel welds imparts hot crack resistance
by breaking up the solidification structure and by scavenging tramp elements (e.g., sulfur and phosphorous). Conversely, eutectic-type reactions during terminal solidification (e.g., liquid ! g þ Laves in nickel-based alloys) are notably detrimental to solidification cracking resistance.7 Figure 4 illustrates the calculated solidification path and hot cracking resistance of two nickel– chromium filler metals. The more solute-rich filler metal that forms Ni2(Nb, Mo)-type Laves phase is
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1900 Note: Scheil calculations are approximate and overpredict the actual solidification temperature range
Calculated solidification T range (⬚C )
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Figure 3 Comparison of the calculated solidification temperature ranges of some common materials used in nuclear power systems (JMatPro, Version 4.1). For a given alloy class, hot cracking is promoted by larger solidification temperature range and low solidus temperature. Note that the compositions used are ‘typical’ values and significant variability exists within each alloy’s specification range.
much more prone to hot cracking than the Lavesfree alloy. 3. The surface tension of the terminal liquid : Low surface tension liquids wet solidification boundaries and promote cracking by increasing the amount of interface incapable of supporting appreciable tensile strains. 4. The metallurgical structure of the weld: Large columnar dendritic grains are more susceptible to solidification cracking than finer equiaxed structures. Coarse solidification structures result in longer crack paths and less grain boundary area to distribute elements that lower the solidus and/or embrittle the boundary. Columnar grains may exacerbate hot cracking by promoting wetting of the grain faces and can result in linear solidification boundaries near the centerline of the weld bead where tensile stresses are often the highest. 5. The mechanical forces that act on the weld: High tensile strains during the terminal stages of solidification promote cracking. Given the complexity of the factors that contribute to solidification cracking, it is difficult to predict its
occurrence in production welds. However, weldability tests such as the transvarestraint test enable a quantitative ranking of alloys with respect to solidification cracking susceptibility and offer a standardized methodology to optimize welding parameters.8–10 Results from transvarestraint tests on several corrosion-resistant alloys are shown in Figure 5, which compares the maximum crack distance (i.e., the extent of the mushy zone when the crack distance becomes insensitive to the applied strain) and hence the intrinsic susceptibility of the alloy to solidification cracking. A representative transvarestraint sample is shown in Figure 6, which illustrates the locations of solidification- and liquation-type cracks. Note that solid-state cracks can also be produced in this test.10 In general, alloying additions that are rejected into the liquid (i.e., whose equilibrium segregation coefficient, k, is <1) lower the solidus and increase the solidification temperature range. This increases the extent of the solid þ liquid ‘mushy’ zone and increases the susceptibility to solidification cracking. For example, niobium and molybdenum have segregation
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Figure 4 Illustration of the effect of solidification path on cracking resistance. (a) Two Ni–30Cr filler metals show markedly different hot crack susceptibilities. (b) Scheil modeling predicts that the more solute-rich alloy has a larger solidification temperature range and can form s, Laves and d in the terminal solid and (c) SEM investigation confirms Nb, Mo-rich Laves near solidification cracks.
coefficients <1 in austenitic stainless steels and nickelbased alloys, which explains the longer crack lengths in 347 SS than in 308L SS (see Figures 5 and 7). Similarly, in nickel alloys, the susceptibility to solidification
cracking is Alloy 625 > EN52MSS > EN52i > EN82H > 68HP, which is a direct result of the decreasing alloying content of the strong melting point depressants molybdenum and niobium.7,10–15
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Transvarestraint test data GTAW weld buildups 180 amps, 12.7 volts, 5 ipm ~45.7 kJ in.–1
3.00
Alloy 625
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Maximum crack distance (mm)
EN52MSS
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EN52i 1.50 EN82H SANICRO 68HP 1.00 347 SS
0.50 308L SS Different symbols of the same color denote different heats
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4
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6 7 8 9 Applied strain (%)
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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.
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Autogenous weld bead
Liquation cracks
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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
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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 <50 wt ppm sulfur, <100 wt ppm P in the bulk alloy) nor grain boundary sliding plays a significant role in cracking. For example, if sulfur is migrating to grain boundaries at 870 C (1600 F), ductility would not be expected to recover after short hold times (10 s as shown in Figure 9). Similarly, the temperature dependence of the ductility minimum must be explained, as the effect of embrittling agents such as sulfur should persist to low homologous temperatures.18 Similarly, if grain boundary sliding were contributing to the intergranular fracture, the fastest quenched sample in Figure 9 should show the most embrittlement, that is, where sliding would be favored by a microstructure without carbides to pin the grain boundaries.29,30 A relative grain-by-grain map of the plastic strains from samples strained to 5 and 10% in the ductility dip temperature range (Figure 11) shows direct evidence against the
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Figure 8 Illustration of intergranular liquation-type cracks in a low-alloy steel. The cracks occurred in both the partially melted zone and heat-affected zone along prior austenite grain boundaries (top) and subsequent analysis of crack surfaces via Auger electron microscopy identified bands of MnS-type sulfide inclusions (bottom).
notion that solid-state cracking is caused by sliding, that is, the uniformity of slip with some strain accumulation (yellow and red areas) near the ductility dip cracks. If grain boundary sliding played a role in ductility dip cracking, there would be less strain contrast near the cracks, not more, as this technique is sensitive only to the diffraction pattern rotation produced by dislocations. The scenario of weld cracking occurring in the time/temperature regime of precipitation of a partially or fully coherent second phase is also well
recognized as controlling strain-age or ‘reheat’-type cracking in g0 /g00 -strengthened alloys.10,18,20,21,33 While susceptibility to reheat cracking is often plotted as a function of the aluminum and titanium content of the alloy, a more fundamental correlation based on the transformation kinetics and precipitate/ matrix mismatch of the alloy is possible.18 As shown in Figure 12, alloys with high susceptibility to strainage cracking display fast transformation kinetics and a large negative precipitate/matrix mismatch (i.e., tension develops between precipitates).
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1100
Temperature (⬚C)
1000
Finish of grain boundary
Start of grain boundary M23 C6 precipitation
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900 1
2
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700 M23 C6 TTT kinetics Estimated from JMatPro 4.1 Ni–29Cr–9Fe–0.01C alloy
600
500 0
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35 0 (b)
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Hold time at 870 ⬚C prior to straining (s)
Figure 9 Comparison of on-cooling hot ductility tests on EN52 (b) with the estimated M23C6 time–temperaturetransformation behavior (a). The local ductility minimum correlates with the onset of grain boundary M23C6 precipitation and ductility recovery with the completion of precipitation.
As transformation stresses are displaced to longer times and precipitation-induced stresses become compressive, weldability is improved. However, while weldability is increased by displacing precipitationinduced stresses to longer times, hot workability is often degraded for the same reasons. It is notable that some titanium alloys undergo a tensile ductility loss when tested near the b ! a transformation.23,34,35 While the authors do not know of equivalent research on zirconium alloys, these alloys could also be susceptible to this form of precipitationinduced cracking (PIC). Mechanistically, this could be caused by the nucleation of a from the b-phase
if the (110)bk(0001)a is significant, or from some other phase with partial coherency (e.g., HCP Laves on HCP-a). 4.09.1.2.2 Segregation-induced cracking
The most technologically important manifestation of SIC is hydrogen-induced cracking (also known as ‘cold cracking’) caused by the numerous sources of hydrogen in both fabrication and service, the relative ease of hydrogen entry into metals, its high diffusivity, and its ability to weaken metallic bonds or form brittle second phases.36–39 Hydrogen cracking in steels and hydride-type cracking of zirconium alloys
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30 Cr alloy a ≈3.58
Partially coherent grain boundary Cr23C6
Tension between carbides
Cr23C6 a≈10.66
Grain boundary
(Cr,Fe)23C6
(a)
0.1μm
(Cr,Fe)23C6
Cracking between carbides (b) Small cracks coalesce
1μm
2 mm
Acc.V Spot Magn Det WD 20.0 kV 3.5 10000x SE 12.7
Grain boundary
Grain 1
‘Ductility Dip’ Crack μm μm 100
(c)
Grain 2
0.1%
Estimated strain from M23C6 precipitation
0.0%
Alloy 600 ~15 wt% Cr
–0.1% –0.2% –0.3% –0.4%
EN82H ~20 wt% Cr
–0.5% Alloy 690 ~29 wt% Cr
–0.6% –0.7% 3.550 (d)
3.555
3.565 3.560 3.570 Lattice parameter of alloy (Å)
3.575
3.580
Figure 10 Illustration of the mechanism of ductility dip cracking in Ni–Cr alloys. (a) Partially coherent (Cr,Fe)23C6 carbides form in reheated weld metal, which have misfit with the grain boundary. (b) Precipitation generates local grain boundary tensile stresses between carbides which (c) results in ductility dip cracking when sufficiently large global stresses are present during welding or are applied during tensile testing. (d) The increased misfit between the carbide and the matrix with increasing chromium concentration helps explain the susceptibility of alloy 690/EN52 and the resistance of A600/EN82 to DDC.
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1
2
(a)
=10 mm; LocMis2; Step = 0.1 mm; Grid274 x 207
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DDC
0
2
1
3
(c)
(d) =20 mm; LocMis2; Step = 0.3333 mm; Grid350 x 372
=20 mm; BC; Step = 0.3333 mm; Grid350 x 372
Figure 11 Comparison of the local misorientation (left) and band contrast (right) images for EN52 strained to 5% ((a) and (b), respectively) and to 10% strain ((c) and (d)) during cooling. Note the generally uniform plasticity with some strain accumulation at the grain boundary.
0.8
X-750
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0.6
718 0.4 0.2
Compression between precipitates
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Astroloy IN 100 – 0.6 – 0.8 0.1
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10 100 Calculated nose of the TTT curve (s)
1000
Figure 12 Correlation of the subsolidus cracking susceptibility of selected superalloys with the misfit and kinetics of second-phase precipitation (g0 or g00 ). Adapted from Young, G. A., et al. Welding J. Res. Suppl. 2008, 31S–43S; Prager, M.; Shira, C. S. Welding Res. Council Bull. 1968, 128; calculations done with JMatPro, Version 4.1.
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have been treated recently in the literature and more extensive reviews are found elsewhere.40,41 However, it should be highlighted that lowstrength austenitic alloys are resistant but not immune to hydrogen-induced cracking. Figure 13 shows a hydrogen-induced crack in a Ni–20Cr–3Mn– 2.5Nb–1Fe weld metal (EN82) that was produced by the combination of poor welding practice and the use of hydrogen-bearing shield gas. The 95%Ar–5%H2 shield gas helps minimize surface oxides and interpass grinding but results in 12 wt ppm hydrogen dissolved in the filler metal. ‘Refuse welding’ or remelting beads in an attempt to improve the tie-in and contour increases the plastic strain in the joint and can trigger cold cracking.42,43
lost if they oxidize prior to solidification (e.g., the grain nucleating effect of Ti(C,N)-type particles). Recent research shows that control of oxygen is critical to the weld puddle flow and wetting in nickelbased filler metals.10 In practice, this often translates into careful wire drawing practice so as to minimize the extent of embedded oxides or wire drawing lubricants into the filler metal. Figure 15 shows variability in the bead contour and tie-in of two filler metals welded under identical conditions, which was later traced to wire cleanliness. Additionally, separate testing shows that 100 wt ppm levels of oxides can have macroscopic detriment on the regular flow and contour of Ni–30Cr-type filler metals.10
4.09.1.3
4.09.2 Stresses and Strains in Welds
Other Welding Defects
In addition to cracking, defects such as lack of fusion between weld beads or the weld bead and the sidewall, variable penetration, or second-phase inclusions can degrade weld quality. Lack of fusion defects is a notable concern when welding high-alloy nickelbased materials, which have notably ‘sluggish’ weld pools and are difficult to wet and tie into adjacent material (Figure 14(a)). Inclusion-type defects are another concern and can be grouped into at least two types: (1) those that result from alloying additions or slag and (2) those that form via reaction with the environment. An example of the first type is given in Figure 14(b), which shows an unmelted iron–niobium Laves phase that is an intentional alloying addition to the flux coating of a shielded metal arc electrode. While alloying in this manner is a cost-effective way to tailor the composition of the electrode, it can lead to brittle second phases that also affect the local composition. In this case, the Nb-rich Laves phase is a strong melting point suppressant which can lead to either solidification- or liquation-type cracking. The second type of inclusion is generally oxide or nitride-type particles that form via reaction with air. The corrosion-resistant alloys used in nuclear power systems (i.e., Fe-based stainless and Ni-based alloys) are especially prone to oxide-type defects as the nature of their corrosion resistance depends on the formation of stable, tenacious oxide films. An example of an aluminum–titanium-rich eutectictype oxide that formed in a poorly shielded Alloy 690 fusion weld is shown in Figure 14(c). Another consideration of this oxide formation is that whatever metallurgical effect these alloying elements have is
Fusion welding leads to residual stresses and strains (distortion) via thermally induced stresses, solidification shrinkage, and phase transformations. Thermal stresses arise from the large temperature gradients inherent to fusion welding and from differences in the coefficients of thermal expansion (CTE) between materials that make up the weldment (Table 1). Thermal contraction generates stress and distortion during on-cooling, with the maximum residual stress often being the flow stress at which the lowest temperature distortion occurs.5,45 Dissimilar metal welds are regions of special concern for nuclear power systems as the residual stresses are often higher than for similar metal welds.18 For example, in ‘safe end’-type welds, the CTE difference between a pressure vessel low-alloy steel and a corrosion-resistant austenitic alloy leads to higher stresses in these welds and in fact, these locations are known to be at increased risk of stress corrosion cracking for this reason.46–48 Solidification shrinkage is a lesser effect, as the liquid cannot support appreciable stress but can affect the local bead contour, the deformation at the weld face, and lead to stress raisers (e.g., concave beads or cracks).5 Phase transformations can also have appreciable effects on the sign, magnitude, and distribution of residual stresses in welds. Becker et al. have shown that for accurate prediction of the residual stresses in pressure vessel-type steels, it is critical to account for the on-cooling phase transformations that occur from welding.49 Furthermore, phase transformations during postweld heat treatments and from service exposure must also be considered. For example, nickel alloys can be susceptible to the development of
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M-3, 8:1 Phos.
=1500 mm; BC + GB; Step = 5 mm; Grid800 ´ 100
=1500 mm; E123; Step = 5 mm; Grid800 ´ 100
=1500 mm; BC + GB; Step = 5 mm; Grid800 ´ 100
=1500 mm; E123; Step = 5 mm; Grid800 ´ 100
=1500 mm; BC + GB; Step = 5 mm; Grid800 ´ 100
0.5 cm
=1500 mm; E123; Step = 5 mm; Grid800 ´ 100
524´D
25.0 kV
100 mm
AMRAY
#0000*
2,540´D
25.0 kV
10 mm
AMRAY
#0000*
Figure 13 Example of a hydrogen crack produced in EN82H from the use of 95%Ar–5%H2 shielding gas and abusive welding practice (refuse welding). The top figures show the crack in cross-section and the corresponding electron backscatter diffraction strain maps. The bottom fractographs show the intergranular/interdendritic nature of the hydrogen crack.
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(a)
Acc.V Spot Magn Det WD Exp 20.0 kV 5.0 500x BSE 4.0 6001
Acc.V Spot Magn Det WD Exp
(b)
50 μm
500 μm
20 mm
50 μm Acc.V Spot Magn Det WD Exp 20.0 kV 4.0 1200x BSE 4.8 6001
JP1-100 10 mil
287
Acc.V Spot Magn Det WD Exp 20.0 kV 5.0 150x SE 5.2 6001
200 μm
20 mm
(c)
50 μm
50 μm
Figure 14 Illustration of some welding defects in commercially available Ni–30 wt% Cr alloys. (a) Lack of fusion defects in an SMAW (left) and a GTAW (right), (b) unmelted NbFe2 alloying addition in an SMAW (left) and a slag inclusion from an SMAW (right) and (c) surface (left) and internal (right) oxides from a laser weld.
short- and long-range order. The ordered structure typically has a smaller lattice parameter than the bulk alloy and can lead to increased residual stresses with service exposure.10,50–54 Another point to note is that the welds typically have considerable texture that can be a significant factor in both the macroscopic and microscopic (intergrain) stresses and strains in welds. 4.09.2.1 Quantification of Residual Stresses and Strains 4.09.2.1.1 Elastic stress
Stresses in welds can be determined via several computational and experimental techniques. Computational methods are generally based on finite element methods, while experimental techniques include X-ray and neutron diffraction, hole drilling, and surface deformation mapping (e.g., slitting). Details of the application of these techniques can be found in several research proceedings55–57 and recent books.58,59
4.09.2.1.2 Plastic strain
The evolution of automated electron backscatter diffraction analysis has made the mapping and quantification of plastic strains in welds accessible via the scanning electron microscope.43,60–64 Strains can be visualized qualitatively via the intragrain misorientation (Figures 11 and 19(c)) of the diffraction pattern or quantitatively (Figure 13) via the average intragrain misorientation (i.e., the ‘AMIS’ parameter) of many grains and an appropriate calibration curve. Calibration curves from uniaxial tensile samples for several nickel-based alloys are given in Figure 16. For reference, the measured plastic strain in several different welds and a heat-affected zone are compared in Table 2. Appreciable plastic strains (2–4%) occur even in unconstrained bead-on-plate welds and a wide range of strains (2% to almost 30%) may be found, depending on the precise weld geometry, constraint, and welding practice. An example of the effect of welding practice is
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2 cm Ca
SEM
11/28/101 F
10.0 keV
900.0X
50.0 mm
11/28/101 F
10.0 keV
900.0X
C
11/28/101 F
10.0 keV
900.0X
50.0 mm
50.0 mm
O
11/28/101 F
10.0 keV
900.0X
50.0 mm
Figure 15 An example of the effect of wire cleanliness on weld quality. The top picture shows two GTAW welds made under identical conditions. The right hand weld displayed poor flow and tie-in and was later traced back to impurities (likely drawing lubricant) embedded in the weld wire. The bottom figure shows an SEM image of the wire surface and Auger maps of embedded calcium, carbon, and oxygen contamination.
Table 1
Comparison of some physical properties for elements of interest to nuclear power systems
Metal or alloy
Melting point (K)
Thermal conductivity (W mK1)
Coefficient of thermal expansion, 293–373 K (106 K1)
Fe Ni Zr
1809 1726 2125
80.4 74.9 21.1
11.8 13.3 5.0
given in Figure 17 for a 2 in. thick, Alloy 690 narrow groove weld made with EN82H filler metal via automatic gas tungsten arc welding (A-GTAW).42,43 If welded with no ‘repairs’ (i.e.,
Reference
[44]
autogenous remelting of beads to improve beadto-bead tie-in), it shows 5.5% plastic strain near the weld root. This plastic strain increases if the beads above the weld are remelted as shown in the
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40 A600 (as-received plate) A690 (as-received plate)
35
y = 2.8332x2 - 3.9127x + 6.3447 R2 = 0.9844 (A690)
HP nickel (annealed)
y = 1.9331x2 - 1.498x + 1.3342
EN82H (annealed)
30
R2 = 0.9482 (HP nickel - annealed)
Tensile strain (%)
304L and 316L (GE data) y = 0.5754x2 + 3.3861x - 0.9728 R2 = 0.9949 (A600)
Linear (304L and 316L (GE data))
25 20
y = 6.7568x - 2.1824 R2 = 1 (304L and 316L)
y = 4.9776x – 2.8443 R2 = 1 (EN82H - annealed)
15 10 5 0 0
1
2
3 4 AMIS parameter (degrees)
5
6
Figure 16 Tensile data used to calibrate the ‘AMIS’ parameter for several austenitic alloys. The ‘GE Data’ are from Angeliu, T. In Tenth International Conference on Environmental Degradation of Materials in Nuclear Power Systems; NACE: Lake Tahoe, NV, 2000.
Table 2 Comparison of the experimentally measured ‘AMIS’ parameter and the calculated plastic strain for several nickel-alloy welds and a heat-affected zone Weld
Range of AMIS measured (degrees)
EN82H, unconstrained bead-on-plate, A-GTAW A600/EN82H pipe weld, near root, M-GTAW A690/EN82H narrow groove weld near root, A-GTAW Best practice A690/EN82H narrow groove weld, A-GTAW Abusive weld practice A600 HAZ (unconstrained, E-182 SMAW)
0.9–1.3 1.1–1.9 1.0–3.2
2–4 3–8 2–13
3.0–6.2
12–28
1.0–1.5
3–6
graph with strains of 11.5%, 15.0%, and 16.5% with 1, 2, and 3 simulated ‘repairs’ above the weld. As expected, high levels of plastic strain lead to increased yield strength, decreased ductility, and increased susceptibility to stress corrosion cracking.
4.09.3 In-Service Performance 4.09.3.1 Environmentally Assisted Cracking Welds and their heat-affected zones have long been known to be the areas of concern for environmentally assisted cracking (EAC) because of their propensity
Approximate plastic strain in weld or HAZ (%)
for as-fabricated flaws, high residual stresses, elevated plastic strains, chemical heterogeneity, and microstructural differences relative to base metals (Chapter 5.04, Corrosion and Stress Corrosion Cracking of Ni-Base Alloys; Chapter 5.05, Corrosion and Stress Corrosion Cracking of Austenitic Stainless Steels; Chapter 5.06, Corrosion and Environmentally-Assisted Cracking of Carbon and Low-Alloy Steels; Chapter 5.02, Water Chemistry Control in LWRs; and Chapter 5.08, Irradiation Assisted Stress Corrosion Cracking). Common EAC concerns in the nuclear industry include corrosion fatigue of low-alloy steels, hydride-induced
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18
Plastic strain near weld root (%)
16
2⬙ thick, A690 Narrow groove weld, A-GTAW Weld filler metal is EN82H. ‘Repairs’ are autogenous remelting of welds beads.
14 12 10 8 6 4 2 0
Best practice
1 repair
2 repairs
3 repairs
Figure 17 Illustration of the effect of simulated repairs (i.e., autogenous remelting of beads) on the plastic strain developed in a 2 in. thick Alloy 690/EN82H narrow groove weld.
cracking of zirconium alloys, and stress corrosion cracking of corrosion-resistant structural alloys. Specifically, stress corrosion of austenitic stainless steel65–68 and nickel-alloy welds and their heataffected zones has been a topic of considerable research.9,66,69–72 In austenitic stainless steels, sensitization- and welding-induced plastic strains are the key factors in stress corrosion resistance. In nickel-alloy welds, the bulk chromium concentration, the solidification segregation, and the as-fabricated plastic strain are critical factors for understanding their stress corrosion performance. The stress corrosion cracking growth rate of nickelalloy filler metals in high-temperature, high-purity water is shown as a function of their bulk chromium concentration in Figure 18(a). Note the strong decrease in stress corrosion crack growth rate at bulk chromium levels near 22 wt%. This decrease is likely associated with a change in crack tip oxide from NiO-type to a more stable spinel (NiCr2O4) or corundum (Cr2O3) structure.73–75 However, the bulk chromium concentration does not explain the extreme resistance of EN52 weld metal compared to other 30 wt% bulk chromium alloys (Figure 18(a)). Consideration of how solidification segregation affects the grain boundary chromium concentration is critical to understanding the stress corrosion resistance of the high-alloy weld metals.76 Specifically, niobium- and molybdenum-bearing alloys tend to deplete the solidification grain boundaries in
chromium, while the Nb- and Mo-free EN52 grain boundaries are enriched in chromium as shown in the graph in Figure 18(b). In Alloy 600 heat-affected zones, the increased susceptibility to SCC in high-temperature deaerated water is due, in large part, to the lack of intergranular chromium carbides.77,78 Figure 19(a) shows a crosssection of a stress corrosion crack grown in an Alloy 600 heat-affected zone and the flat grain boundary topography (GBT) in the HAZ, which is an indication of a low degree of intergranular chromium carbide precipitation.78 Additionally, Figure 19 shows the different chromium concentration profiles in the HAZ and base metal (Figure 19(b)), the increased strain in the weld and HAZ relative to the base metal (Figure 19(c)), and the transmission electron micrographs of the grain boundaries in the HAZ (showing sparse (M7C3- and M23C6-type carbides) versus the large continuous Cr7C3 carbides in the unaffected base metal (Figure 19(d)). The diffraction patterns in Figure 19(d) identify the M23C6 (left) and M7C3 (right) carbides. Stress corrosion crack growth rate predictions for Alloy 600 heat-affected zones are shown in Figure 20, which illustrates the strong temperature dependence as well as the effects of the applied stress intensity factor and the electrochemical potential. Figure 20 is based on eqn [1], which describes the crack growth rate of A600-type alloys exposed to high-temperature, high-purity water,77 in which A0, n, m, b, x0, and c are the
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102
100
10–8
360 ⬚C, Initial KI = 50 MPa m½ 20 SCC H2 per kg H2O Adjusted to 100% engagement
E-182*
10–9
EN82H*
100X
10–10 10–1 10–11 10–2
EN52i : SCC benefit with weldability of EN82H
10–3
SCC rate (m s–1)
SCC rate (mils day–1)
101
10–12
EN52
10–13 10–4 10
15
20 25 Bulk chromium (wt%)
30
35
102
100X
10–9
EN82H*
10–10 10–1 10–11 10–2
EN52i : SCC benefit with weldability of EN82H
SCC rate (m s–1)
SCC rate (mils day–1)
100
10–8
360 ⬚C, Initial KI = 50 MPa m½ 20 SCC H2 per kg H2O Adjusted to 100% engagement
E-182*
101
–12 EN52 10
10–3 10–13 10–4
10
15 20 25 30 Grain boundary chromium (wt%)
35
Figure 18 (a) The effect of chromium concentration on the stress corrosion crack growth rates of nickel–chromium weld metals. While a dramatic improvement in SCC resistance is seen near 22 wt% bulk chromium, high chromium (30 wt%) alloys appear to show variable resistance. Understanding how solidification segregation affects the grain boundary chromium level (b) is key to understanding the different stress corrosion resistance of these alloys. E-182 and EN82H SCC rates are from model predictions, E-182 grain boundary chromium concentration estimated.
experimentally determined constants, KI is the stress intensity factor, sYS is the yield strength, DEcP is the electrochemical potential relative to the Ni/NiO phase transition, Q Effective is the apparent activation energy, R is the gas constant, and T, the temperature. The appropriate parameters for Alloy 600 HAZs are given in Table 3. a_ ¼ A0 KIn smYS ( " #) ðEcPNi=NiO x0 Þ 2 1 þ b exp 0:5 c Q Effective ½1 exp RT
4.09.3.2
Microchemical Changes
In addition to solidification segregation, the local composition of welds can change in-service via thermal exposure and via radiation-induced transmutation and radiation-induced segregation (RIS). Thermally induced embrittlement is most notable in low-alloy steels that are high in tramp elements, in locations where welding produces local compositional enrichment, and in higher nickel grades (e.g., A508 Gr4N), which are intrinsically more susceptible because of the cosegregation of nickel and phosphorous.79–82 RIS is a phenomenon in which irradiation-created defects cause spatial redistribution of alloying elements as they diffuse to, and get trapped in, sinks
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End of precrack
A600GBT heat Flat affected zone SCC
EN82H Weld
Chromium concentration (wt%)
(a)
(b)
17 16 15 14 13 12 11 10 9 8 7 6 5
0.5 mm
Far from weld
A600 HAZ
EN82H GTA weld
A600 Base Metal A600 Heat Affected Zone
–2
Hardness indents
0 1 –1 Distance from grain boundary (μm)
2
High strain
Low strain
(c)
HAZ
Misorientation (°)
Base
1 mm 01 1378 6 018000
01 1352 6 010500
1 mm
(d)
Figure 19 Comparison of Alloy 600 heat-affected zone and base metal structure, chemistry, and strain: (a) cross-section of stress corrosion sample showing the location of the cracking, (b) grain boundary chromium profiles for base metal (blue) and the HAZ (red), (c) qualitative strain map for the base metal, HAZ/weld interface, and (d) typical grain boundary microstructure for the HAZ (sparse M23C6 and M7C3) and base metal (extensive M7C3).
(e.g., voids and grain boundaries) (Chapter 1.18, Radiation-Induced Segregation). RIS can occur when a point defect flux interacts preferentially with a certain elemental species in the alloy, causing
that element to be enriched or depleted near the defect sinks. This preference can be driven kinetically (migration barriers) and/or thermodynamically (binding/ordering). RIS is segregation that occurs
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1.0E – 8
Predicted crack growth rate (m s–1)
A600 HAZ
360 ⬚C 1.0E – 9 305 ⬚C 1.0E – 10 250 ⬚C
1.0E – 11
8 7 0 60 0
1.0E – 12 75 50
50 4 ½) 0 –25 30 0 20 am –50 ΔEcP MP 10 –75 ( (mV) KI –100
25
Figure 20 Illustration of the predicted crack growth rates for Alloy 600 HAZ material as a function of potential, stress intensity factor, and temperature. Table 3 Figure 20
Fitted parameters for the Alloy 600 heat-affected zone stress corrosion crack growth rate data presented in
Best estimate 95% confidence
ln (A0 )
n
m
B
x0 (mV)
c (mV)
Q (kJ mol 1)
22.607 3.729
0.869 0.349
0a –
3.604 –
15.61 20.71
42.79 19.25
136.0 18.2
a
Insufficient data were available to determine this dependence.
in addition to thermal segregation. Like thermal segregation, this local change in composition can result in detrimental changes to mechanical and corrosion properties.83 RIS has been a concern in the nuclear industry for over 30 years and is considered one of the many factors that lead to irradiation-assisted stress corrosion cracking.6–9 This phenomenon was first predicted by Anthony84 and has been observed in a number of different alloys and steels used in nuclear reactors.85–87 In both the iron- and nickel-based face-centered cubic Fe–Ni–Cr alloys, experimental RIS trends at grain boundaries are generally chromium depletion, nickel enrichment, and possible compensation through iron enrichment or depletion.85,86 In body-centered cubic ferritic–martensitic steels, both chromium enrichment and depletion have been reported.88,89 RIS in welds has not been extensively researched but the possibility of grain boundary depletion of chromium in corrosion-resistant
alloys that would act in addition to the chromium depletion that occurs during solidification segregation (e.g., Ni–Cr–Nb and Ni–Cr–Mo alloys) is of significant concern. 4.09.3.3
Microstructural Changes
In addition to the relatively short-time microstructural changes that can occur on-cooling or with postweld heat treatment (typically <10 h) and induce PIC as discussed earlier, long-time microstructural changes can also occur. While most nuclear alloys have had sufficient vetting to preclude these concerns, recently developed high-alloy nickel-based filler metals raise the concern that topologically close-packed (TCP) phases (e.g., sigma) or long-range ordering could occur in Ni–Cr–Mo welds.10,50 The degradation of toughness by the formation of TCP phases is well established in the superalloy literature and the
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formation of long-range order can lead to increased residual stresses and decreased resistance to EAC, most notably to hydrogen embrittlement.50,52,53,90
4.09.4 Weldability of Specific Alloy Systems 4.09.4.1
Low-Alloy Steels
Low-alloy steels generally have good weldability with the main concern being liquation cracking near impurities (Figure 8), the propensity of some grades (esp. Cr–Mo steels and some pressure vessel grades) to reheat-type cracking (discussed earlier) and to hydrogen-induced cracking.1,91,92 The susceptibility to hydrogen-induced cracking is controlled by four major considerations: 1. 2. 3. 4.
The composition of the steel The mobile hydrogen concentration The stresses in the weldment The thermal management of the weld
In general, the more hardenable the steel (i.e., the more easily martensite is formed), the more susceptible it is to cracking. Since hardenability generally increases with carbon content and alloying additions, several parameters have been developed to gauge susceptibility to hydrogen-induced cracking, including the carbon equivalent (Ceq) in eqn [2] and the ItoBessyo ‘cold cracking’ parameter (Pcm) (eqn [3]).93 The concentrations in eqns [2] and [3] are for weight percentage. Ceq ¼ C þ
Cr þ Mo þ V Mn Ni þ Cu þ þ 5 6 15
Pcm ¼ 5 B þ C þ
½2
V Mo Mn þ Cr þ Cu þ þ 10 15 20
Si Ni þ ½3 30 60 The hydrogen concentration in the weldment can be minimized by proper cleanliness, preheat, and postweld heat treatment.92,94 Additionally, microstructural hydrogen traps can provide significant benefit by preventing hydrogen redistribution to regions of high stress.37 As the mobile hydrogen concentration is most detrimental, significant work has gone into standards to accurately assess the amount of diffusible hydrogen in steels.37,92,94 Minimizing residual stresses and avoiding geometric stress concentrators (e.g., notches) in the weldment also impart resistance to hydrogen-induced cracking. þ
Thermal management of the weldment (e.g., preheating, interpass temperature, bead tempering, and postweld heat treatment) is also critical to the mitigation of hydrogen-induced cracking. Preheating of components and interpass temperature control act to outgas hydrogen or hydrogen-bearing compounds (e.g., water) and lower the cooling rate. Additionally, careful control of heat input can produce hydrogenresistant microstructures (i.e., bead tempering). Postweld heat treatment can act to lower residual stresses, produce beneficial hydrogen traps, and remove dissolved hydrogen from the weld.3,4,95,96 Several sources provide guidelines to mitigate hydrogen cracking in specific grades of steel.95–97 4.09.4.2
Austenitic Stainless Steels
The main concerns in austenitic stainless steel welding are solidification, liquation, and PIC. Additionally, avoiding grain boundary chromium depletion via carbide precipitation (i.e., sensitization or ‘knifeline’ attack) is critical to maintaining in-service corrosion resistance (Chapter 2.09, Properties of Austenitic Steels for Nuclear Reactor Applications). A key factor in solidification cracking resistance is proper control of the weld chemistry to form delta ferrite in the initial solid.98–100 Delta ferrite formation breaks up long, linear solidification boundaries and acts to scavenge tramp elements from the liquid and prevent their concentration at interdendritic boundaries during terminal solidification.101,102 Delta ferrite levels are typically controlled to 5–10% by volume to impart solidification cracking resistance but retain the mechanical properties of a face-centered cubic alloy. Specifically, the bodycentered cubic ferrite is susceptible to cleavage at low temperatures and to spinodal decomposition of the ferrite into iron-rich (a) and chromium-rich (a0 ) phases at intermediate temperatures. Control of delta ferrite levels for different alloys is given in handbooks and can be predicted via the experimentally based Schaeffler or Delong diagrams, or computationally, by multicomponent phase diagrams, as shown in Figure 21.103–107 As a practical example, it is often difficult to weld over fully austenitic metals (e.g., nickel-based alloys) with austenitic filler metals (e.g., 308), as the increased nickel (from dilution) promotes the primary austenite solidification mode. As with most fusion welds, liquation cracking can be controlled by minimizing solidification segregation (e.g., faster cooling rates and finer, more equiaxed structures) and by using lower heat input.
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1500 Primary austenite forms
1475
Primary ferrite forms Liquid
Temperature (⬚C)
Liquid+bcc_A2
1450 Liquid+fcc_A1
1425
1400
fcc_A1
fcc_A1+bcc_A2
bcc_A2
1375 70Fe-XCr-(30-X)Ni
1350 14
15
16
17 18 19 20 21 22 23 Chromium concentration (wt%)
24
25
26
Figure 21 Illustration of the effect of chromium concentration on the solidification behavior of model ‘austenitic’ stainless steels. Formation of primary (d) ferrite provides hot cracking resistance. This pseudo-binary phase diagram was generated with Pandat Version 8.1, after Lippold, J. C.; Savage, W. F. Welding J. 1979, 58, 362s–374s.
As discussed earlier, stainless steels can be susceptible to PIC via partially coherent M23C6 carbide precipitation. One distinction relative to nickel-based alloys is that cracking may more likely occur on heating or with longer in-service exposures, as carbon may be in solution as-welded, and precipitation kinetics are usually slower than for high-chromium nickel-based alloys.108
Like austenitic stainless steels, single-phase nickelbased alloys are susceptible to solidification and liquation cracking (Chapter 2.08, Nickel Alloys: Properties and Characteristics). Common melting point suppressants that must be controlled to avoid cracking include tramp elements such as sulfur and intentional alloying additions such as boron, zirconium, niobium, and molybdenum.7,10–15,18,109,110 Additionally, both low-strength and precipitation hardenable grades can be susceptible to PIC mechanisms (ductility dip and strain-age cracking) as discussed previously.
properties and the corrosion performance of the weld (Chapter 2.07, Zirconium Alloys: Properties and Characteristics).92,111,112 Vacuum or inert gases (argon, helium, or Ar–He mixtures) can be used to shield zirconium, but again, care needs to be taken to ensure sufficient vacuum level or gas purity to prevent contamination.92,111 Zirconium alloys can be susceptible to both supersolidus and subsolidus (i.e., hydride-type) cracking. Supersolidus cracking is typically solidification-type and many common alloying elements and/or potential contaminants promote susceptibility. For example, iron, nickel, chromium, and copper all stabilize low-temperature eutectic reactions, and small concentrations can greatly increase the solidification temperature range. An example of this is shown in Figure 22, which compares the maximum crack length in transvarestraint tests for a Zr–Cr alloy, Zircaloy-4 (Zr-4), and Zr–2.5Nb welded under identical conditions. As shown, the Zr–Cr alloy exhibits the most susceptibility to solidification cracking.
4.09.4.4
Acknowledgments
4.09.4.3
Nickel-Based Alloys
Zirconium Alloys
Zirconium alloys are readily weldable but as with all reactive metals, need special precautions to prevent pickup of interstitial elements such as oxygen, carbon, and nitrogen that can degrade both the mechanical
The Authors are indebted to the welders, technicians, specialists, and engineers of the Welding & Materials Process Development Unit at Knolls Atomic Power
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2.5 Transvarestraint tests Wrought plate GTAW, 8 kJ in.–1
Max crack length (mm)
2.0
Zr-1.0Cr-0.5Sn-0.5Fe
1.5
1.0 Zr-4 0.5 Zr-2.5Nb 0.0
0
1
2
3 4 Applied strain (%)
5
6
7
Figure 22 Comparison of transvarestraint test data for three zirconium alloys. The strong effects of Cr and Fe on the solidification temperature range likely explain the increased susceptibility of the Zr–Cr alloy relative to Zr-4 and Zr–2.5Nb.
Laboratory, whose dedication and expertise made this work possible. They are also grateful to Dr. David S. Knorr of General Electric for his important contributions to the manuscript.
15. 16. 17. 18.
References 1.
2. 3. 4. 5. 6. 7. 8.
9. 10. 11. 12. 13. 14.
Savage, W. F. In Weldments: Physical Metallurgy and Failure Phenomena, Proceedings of the Fifth Bolton Landing Conference; General Electric Company: Bolton Landing, NY, 1978. Borland, J. C. Br. Welding J. 1960, 7, 508–512. Kou, S. Welding Metallurgy, 2nd ed.; Wiley: Hoboken, NJ, 2003. Messler, R. W. Principles of Welding; Wiley: New York, 1999. Linnert, G. E. Welding Metallurgy: Carbon and Alloy Steels; American Welding Society: New York, 1967; Vol. II, p 674. Easterling, K. Introduction to the Physical Metallurgy of Welding, 2nd ed.; Butterworth-Heinemann: Boston, 1992; p 270. Cieslak, M. J. The Solidification Behavior of an Alloy 625/718 Variant in Superalloys 718, 625 and Various Derivatives; TMS: Pittsburgh, PA, 1991. Hanninen, H.; et al. Hot Cracking and Environmentally Assisted Cracking Susceptibility of Dissimilar Metal Welds, 2399; Helsinki University of Technology: Helsinki, 2007; pp 1–186. Hanninen, H.; et al. In 13th International Conference on Environmental Degradation of Materials in Nuclear Power Systems; TMS/ANS/NACE: Whistler, BC, CA, 2007. Young, G. A. In Welding & Repair Technology; EPRI: Fort Myers, FL, 2008; pp 1–50. Cieslak, M. J. In The Metal Science of Joining; TMS: Cincinnati, OH, 1991. Cieslak, M. J. Welding J. 1991, 78, 49S–56S. Cieslak, M. J.; Headley, T. J.; Frank, R. B. Welding J. 1989, 68(12), 473s–482s. DuPont, J. N.; Robino, C. V.; Marder, A. R. Met. Trans. A 1998, 29A, 2797–2806.
19. 20. 21. 22. 23. 24.
25. 26. 27. 28. 29. 30. 31. 32. 33. 34. 35.
DuPont, J. N.; et al. Met. Trans. A 1998, 29A, 2785–2796. Pepe, J. J.; Savage, W. F. Weld. J. Res. Suppl. 1970, 49, 545s–553s. Pepe, J. J.; Savage, W. F. Weld. J. Res. Suppl. 1967, 46, 411s–422s. Young, G. A.; et al. Welding J. Res. Suppl. 2008, 87(2), 31S–43S. Prager, M.; Shira, C. S. Welding Res. Council Bull. 1968, 128. Berry, T. F.; Hughs, W. P. Welding J. 1969, 48(11), 505s–513s. Hughes, W. P.; Berry, T. F. Welding J. 1967, 46(8), 361s–370s. Swift, R. A.; Rodgers, H. C. Welding J. 1973, 52(4), 145s–153s, 172s. Rath, B. B.; et al. High temperature ductility loss in titanium alloys – A review, DE-AC04–94AL85000; Sandia National Laboratories, 1994. Capobianco, T. E.; Hanson, M. In Proceedings of the 7th International Conference Trends in Welding Research; ASM International: Callaway Gardens, GA, 2005. Collins, M. G.; Lippold, J. C. Welding J. 2003, 82(10), 288s–295s. Collins, M. G.; Ramirez, A. J.; Lippold, J. C. Welding J. 2003, 82(12), 348s–354s. Collins, M. G.; Ramirez, A. J.; Lippold, J. C. Welding J. 2004, 83(2), 39s–49s. Nishimoto, K.; Saida, K.; Okauchi, H. Sci. Technol. Welding Joining 2006, 11(4), 455–479. Noecker, F. F.; DuPont, J. N. Welding J. 2009, 88, 62s–77s. Noecker, F. F.; DuPont, J. N. Welding J. 2009, 88, 7s–20s. Ramirez, A. J.; Lippold, J. C. Mat. Sci. Eng. A 2004, 380, 259–271. Ramirez, A. J.; Lippold, J. C. Mat. Sci. Eng. A 2004, 380, 245–258. Franklin, J. E.; Savage, W. F. Welding J. 1974, 53(7), 380s–387s. Hayduk, D.; et al. Welding J. 1986, 65(7), 251s–260s. Damkroger, B. K.; Edwards, G. R.; Rath, B. B. Welding J. 1989, 68(7), 290s–302s.
Welds for Nuclear Systems 36. 37. 38. 39. 40. 41. 42. 43. 44. 45. 46.
47. 48.
49. 50. 51. 52. 53. 54. 55. 56. 57.
58. 59. 60.
61. 62.
Oriani, R. A.; Josephic, P. H. Acta Met. 1974, 22, 1065–1074. Hirth, J. P. Met. Trans. A 1980, 11A, 861–890. Birnbaum, H. K. In Hydrogen Effects on Material Behavior; TMS: Jackson Hole, WY, 1990. Wriedt, H. A.; Oriani, R. A. Acta Met. 1970, 18, 753–760. Strasser, A.; Adamson, R.; Garzarolli, F. The Effect of Hydrogen on Zirconium Alloy Properties 1, ZIRAT-13; ANT International 2009. Strasser, A.; et al. The Effect of Hydrogen on Zirconium Alloy Properties 2, ZIRAT-13; ANT International, 2009. Young, G. A.; et al. In Trends in Welding Research; ASM International: Pine Mt., GA, 2002. Young, G. A. In Trends in Welding Research; ASM International: Pine Mt., GA, 2002. Gale, W. F.; Totemeier, T. C., Eds. Smithells Metals Reference Book, 8th ed.; Elsevier: Boston, MA, 2004. Diehl, M. J.; Messler, J. R. W. Welding J. 1995, 74(4), 109S. McMinn, A.; Nelson, J. L. In Third International Symposium on Environmental Degradation of Materials in Nuclear Power Systems-Water Reactors; ANS/TMS/ NACE: Miami, FL, 1987. Csontos, A.; et al. In 13th International Conference on Environmental Degradation of Materials in Nuclear Power Systems. ANS/TMS/NACE: Whistler, BC, 2007. Herrera, M.; et al. Materials reliability program: Evaluation of the effect of weld repairs on dissimilar metal butt welds, MRP-114; EPRI: Palo Alto, CA, 2004; pp 1–100. Becker, M.; et al. In Trends in Welding Research; Pine Mt., GA, 2004. Wong, F. Aging and phase stability of waste package outer barrier, ANL-EBS-MD-000002 REV 02, Sept 2004; Bechtel: Las Vegas, NV. Martinovitch, M.; et al. Nickel base welds in nuclear components, 98NB00014, July 1997, Electricite de France, pp 7–14. Marucco, A.; Nath, B. J. Mat. Sci. 1988, 23, 2107–2114. Tucker, J. D. Investigation of alloying effects on the kinetics of Ni2Cr formation in Ni–Cr alloys. In TMS Annual Meeting; TMS: San Francisco, CA, 2009. Young, G. A.; et al. In 13th Conference on Environmental Degradation of Materials in Nuclear Power Systems; CNS: Whistler, BC, 2007. David, S. A.; et al., Eds. Trends in Welding Research; ASM International: Pine Mt., GA, 2002; pp 1–1022. David, S. A.; et al., Eds. Trends in Welding Research; ASM International: Pine Mt., GA, 2005; pp 1–1025. Hall, J. F.; et al. In Conference on the Contribution of Materials Investigation to the Resolution of Problems Encountered in Pressurized Water Reactors; Societe Francaise d’Energie Nucleare: Paris, 1994. Feng, Z., Eds. Processes and Mechanisms of Welding Residual Stress and Distortion; Woodhead Publishing: Abington, 2005; p 364. Radaj, D. Welding Residual Stress and Distortion; DVSVerlag, 2003; p 397. Young, G. A.; Lewis, N. In Conference on Vessel Heat Penetration Inspection, Cracking, and Repairs; US NRC and Argonne National Laboratory: Gaithersburg, MD, 2003. Sutliff, J. A. Micros. Microanal. Proc. 1999, 5(Suppl. 2), 236–237. Angeliu, T. In Tenth International Conference on Environmental Degradation of Materials in Nuclear Power Systems; NACE: Lake Tahoe, NV, 2000.
63.
64.
65. 66. 67.
68.
69.
70.
71. 72.
73. 74. 75.
76.
77. 78. 79. 80.
81.
82.
297
Lehockey, E. M.; Brennenstuhl, A. M. Characterization of plastic strains and crystallographic properties surrounding defects in steam generator tubes by orientation imaging microscopy. In 4th CNS International Steam Generator Conference, Toronto, ON, 2002. Lehockey, E. M.; Lin, Y. P.; Lepik, O. E. In Electron Backscatter Diffraction in Materials Science; Schwartz, A. J., Kumar, M., Adams, B. L., Eds.; Kluwer/Plenum: New York, 2000, pp 247–264. Ford, F. P. In Environment-Induced Cracking of Metals; NACE: Kohler, WI, 1988. Andresen, P. L.; Morra, M. M. Corrosion 2008, 64(1), 15–29. Kilian, R.; et al. In 10th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors; NACE: Lake Tahoe, NV, 2001. Lu, Z.; et al. In 13th International Conference on Environmental Degradation of Materials in Nuclear Power Systems; TMS/ANS/NACE: Whistler, BC, CA, 2007. Vaillant, F.; et al. In 13th International Conference on Environmental Degradation of Materials in Nuclear Power Systems; TMS/ANS/NACE: Whistler, BC, CA, 2007. Bruemmer, S.; Vetrano, J. S.; Toloczko, M. B. In 13th International Conference on Environmental Degradation of Materials in Nuclear Power Systems; TMS/ANS/NACE: Whistler, BC, CA, 2007. Kaneda, J.; et al. In 13th International Conference on Environmental Degradation of Materials in Nuclear Power Systems; TMS/ANS/NACE: Whistler, BC, CA, 2007. Jacko, R. J.; Gold, R. E.; Kroes, A. In 11th International Conference on Environmental Degradation in Nuclear Power Systems – Water Reactors; TMS: Stevenson, WA, 2003. Lewis, N.; et al. In Chemistry and Electrochemistry of Corrosion and Stress Corrosion Cracking; TMS: New Orleans, LA, 2001. Legras, L.; et al. In 13th International Conference on Environmental Degradation of Materials in Nuclear Power Systems; TMS/ANS/NACE: Whistler, BC, CA, 2007. Morton, D. S.; et al. In 13th International Conference on Environmental Degradation of Materials in Nuclear Power Systems; TMS/ANS/NACE: Whistler, BC, CA, 2007. Young, G. A.; et al. In 14th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors; ANS: Virginia Beach, VA, 2009. Young, G. A.; et al. In 12th International Conference on Environmental Degradation of Materials in Nuclear Power Systems; TMS: Salt Lake City, UT, 2005. Young, G. A.; et al. In Tenth International Conference on Environmental Degradation of Materials in Nuclear Power Systems; NACE: Lake Tahoe, NV, 2001. Krauss, G. Steels: Heat Treatment and Processing Principles; ASM International: Materials Park, OH, 1990; p 497. Naudin, C.; Pineau, A.; Frund, J. M. In 10th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors; ANS/NACE/ TMS: Lake Tahoe, NV, 2001. Li, Y. Y.; Sofanak, R. J.; Beggs, W. J. In 11th International Conference on Environmental Degradation of Materials in Nuclear Power Systems; ANS/NACE/TMS: Stevenson, WA, 2003. Knorr, D. B.; McGee, J. J. Welding J. Res. Suppl. 2009, 88(10).
298 83. 84. 85. 86. 87. 88. 89. 90. 91. 92. 93. 94. 95. 96. 97.
Welds for Nuclear Systems Was, G. S. Fundamentals of Radiation Materials Science; Springer-Verlag: Berlin, 2007. Anthony, T. R. In International Conference on RadiationInduced Voids in Metals, Jun 9–11, 1971; p 630. Damcott, D. L.; Allen, T. R.; Was, G. S. J. Nucl. Mat. 1995, 225(97), 97–107. Allen, T. R. Acta Mat. 1998, 46(10), 3679–3692. Okamoto, P. R.; Rehn, L. E. J. Nucl. Mat. 1979, 83(2), 2–23. Allen, T. R. In Proceedings of the 22nd Symposium on Effects of Radiation in Materials; ASTM: Amsterdam, The Netherlands, 2006. Lu, Z.; et al. Scripta Mat. 2008, 58, 878–881. Sims, C. T.; Stoloff, N. S.; Hagel, W. C., Eds. Superalloys II; Wiley: New York, NY, 1987; p 615. Suzuki, K. Reactor pressure vessel materials, in neutron irradiation effects in pressure vessel steels and weldments. IWG-LMNNP – 98/3; IAEA, 1998. Kearns, W. H., Ed. Metals and Their Weldability, 7th ed.; Welding Handbook; American Welding Society: Miami, 1982; Vol. 4; pp 1–582. Ito, Y.; Bessyo, K. Weldability of high strength steels related to heat affected zone cracking, IX-567–68 1968; International Institute of Welding. Kotecki, D. J. Welding J. 1992, 71(8), 35–43. Okuda, N.; et al. Welding J. 1987, 66(5), 140S–146S. Stout, R. D.; Doty, W. Weldability of Steels, 3rd ed.; Welding Research Council: New York, 1978. Structural Welding Code – Steel, AWS D1.1/D1.1M:2008; American Welding Society: Miami, FL, 2008.
98.
Lippold, J. C.; Savage, W. F. Welding J. 1979, 58, 362s–374s. 99. Lippold, J. C.; Savage, W. F. Welding J. 1980, 59, 48s–58s. 100. Cieslak, M. J.; Savage, W. F. Welding J. 1980, 59, 136s–146s. 101. Takalo, T.; Suutala, N.; Moisio, T. Met. Trans. A 1979, 10A, 1173–1181. 102. Kujanpaa, V. P.; David, S. A.; White, C. L. Welding J. 1986, 65, 203s–211s. 103. Schaeffler, A. L. Metal Progress. 1949, 56(11), 680–680B. 104. Chang, A. H. Pandat; Computherm: Madison, WI, 2007. 105. DeLong, W. T. Welding J. 1974, 53, 273s–286s. 106. Kotecki, D. J. Welding J. 1999, 78, 180s–192s. 107. Kotecki, D. J.; Siewert, T. A. Welding J. 1992, 71, 171s–178s. 108. Parker, T. D. In Source Book on Materials for ElevatedTemperature Applications; Bradley, E. F., Ed.; American Society for Metals: Metals Park, OH, 1979; pp 178–197. 109. Cieslak, M. J.; Stephens, J. J.; Carr, M. J. Met. Trans. A 1988, 19A, 657–667. 110. MacPherson, B. M., Ed. Effects of Minor Elements on the Weldability of High Nickel Alloys; Welding Research Council: Houston, TX, 1967. 111. Rudling, P.; Strasser, A.; Garzarolli, F. Welding of Zirconium Alloys, van Swann, L., Reviewer, Special Topic Report 2007; ANT International: Skultuna, Sweden, Oct 2007. 112. Zircadyne Zirconium Welding; ATI Wah Chang: Albany, OR, 2009.
4.10
Radiation Effects in Graphite
T. D. Burchell Oak Ridge National Laboratory, Oak Ridge, TN, USA
Published by Elsevier Ltd.
4.10.1
Introduction
300
4.10.2 4.10.3 4.10.4
Nuclear Graphite Manufacture Graphite-Moderated Reactors Displacement Damage and Induced Structural and Dimensional Changes in Graphite Neutron-Induced Property Changes Wigner Energy Mechanical and Physical Properties Irradiation Creep The Relevance of Creep to Reactor Design and Operation The Irradiation-Induced Creep Mechanism (In-Crystal) Review of Prior Creep Models Linear viscoelastic creep model The UK creep model The Kennedy model The Kelly and Burchell model The M2 model Deficiencies in Current Creep Models at High Neutron Doses Outlook
302 303
4.10.5 4.10.5.1 4.10.5.2 4.10.6 4.10.6.1 4.10.6.2 4.10.6.3 4.10.6.3.1 4.10.6.3.2 4.10.6.3.3 4.10.6.3.4 4.10.6.3.5 4.10.6.4 4.10.7 References
Abbreviations AGR ASTM
Advanced gas-cooled reactor American Society for Testing and Materials CEN Centre European Nuclear CP-1 Chicago Pile No. 1 CTE Coefficient of thermal expansion DWNTs Double-walled carbon nanotubes Esu Elastic strain unit HFR High flux reactor HOPG Highly oriented pyrolytic graphite HRTEM High-resolution transmission electron microscope IV Interstitial–vacancy MHTGR Modular high-temperature gas-cooled reactor NGNP Next Generation Nuclear Plant ORNL Oak Ridge National Laboratory PGA Pile grade A PKA Primary knock-on atom SKA Secondary knock-on atom STM Scanning tunneling microscope
TEM USSR
305 310 310 311 315 315 316 317 317 317 317 318 319 321 323 323
Transmission electron microscope Union of Soviet Socialist Republics
Symbols a a b b B c c C Cp d«c/dg E E0
Constant in viscoelastic creep model Crystallographic a-direction (within the basal planes) Burger’s vector Constant in viscoelastic creep model Empirical fitting parameter, analogous to the steady-state creep coefficient Crystallographic c-direction (perpendicular to the basal planes) Flaw size Specific heat Specific heat at constant pressure Initial secondary (steady-state) creep rate Elastic modulus Initial (preirradiated) Young’s modulus
299
300
Radiation Effects in Graphite
Ed
Ep Eg Fx Fx0 G gx
gx0
h k k k0 (g) k0 k1 k2 La Lc R S(g) T W (1/Xa) (dXa/dg) (1/Xc) (dXc/dg) XT Z a aa ac ax a´ x a(f)
Displacement energy for a carbon atom from its equilibrium lattice position Young’s modulus after initial increase due to dislocation pinning Young’s modulus at dose g Pore generation term Pore generation term for a crept specimen Fracture toughness or strain energy release rate Rate of change of dimensions in the x-direction with respect to neutron dose Rate of change of dimensions in the x-direction for a crept specimen with respect to neutron dose (dimensional change component) Planck’s constant Boltzmann’s constant Steady-state creep coefficient Modified steady-state creep coefficient Initial secondary creep coefficient Primary creep dose constant Recoverable creep dose constant Mean graphite crystal dimensions in the a-direction Mean graphite crystal dimensions in the c-direction Gas constant (8.314 J mol1K) Structure factor (E/Ep) Temperature Oxidation change factor Rate of change of crystallite dimensions perpendicular to the hexagonal axis Rate of change of crystallite dimensions parallel to the hexagonal axis Crystal shape parameter Atomic number Coefficient of thermal expansion Crystal coefficient of thermal expansion in the a-direction Crystal coefficient of thermal expansion in the c-direction Coefficient of thermal expansion in the x-direction Coefficient of thermal expansion of a crept specimen in the x-direction Crystal coefficient of thermal expansion at angle f to the c-direction
«˙ «c «´ c «d «e «p «s «t «Total f g l m n uD r r0 st s v V j
Strain rate Creep strain Apparent creep strain Dimensional change strain Elastic strain Primary creep strain Secondary creep strain Thermal strain Total strain Fast neutron flux Fast neutron fluence Empirical fitting parameter Empirical constant (0.75) Dislocation velocity Debye temperature Density True density Tensile strength Tensile stress Frequency of vibrational oscillations Mobile dislocation density Empirical fitting parameter
4.10.1 Introduction There are many graphite-moderated, powerproducing, fission reactors operating worldwide today.1 The majority are in the United Kingdom (gas-cooled) and the countries of the former Soviet Union (watercooled). In a nuclear fission reactor, the energy is derived when the fuel (a heavy element such as 92U235) fissions or ‘splits’ apart according to the following reaction: 235 92 U
þ0 n1 !92 U236 ! F1 þ F2 þ n þg energy
½I
An impinging neutron usually initiates the fission reaction, and the reaction yields an average of 2.5 neutrons per fission. The fission fragments (F1 and F2 in eqn [I]) and the neutron possess kinetic energy, which can be degraded to heat and harnessed to drive a turbinegenerator to produce electricity. The role of graphite in the fission reactor (in addition to providing mechanical support to the fuel) is to facilitate the nuclear chain reaction by moderation of the high-energy fission neutron. The fission fragments (eqn [I]) lose their kinetic energy as thermal energy to the uranium fuel mass in which fission occurred by successive collisions with the fuel atoms. The fission neutrons (n in eqn [I]) give up
Radiation Effects in Graphite
their energy within the moderator via the process of elastic collision. The g-energy given up in the fission reaction (eqn [I]) is absorbed in the bulk of the reactor outside the fuel, that is, moderator, pressure vessel, and shielding. The longer a fission neutron dwells in the vicinity of a fuel atom during the fission process, the greater is its probability of being captured and thereby causing that fuel nucleus to undergo fission. Hence, it is desirable to slow the energetic fission neutron (E 2 MeV), referred to as a fast neutron, to lower thermal energies (0.025 eV at room temperature), which corresponds to a velocity of 2.2 1015 cm s1. The process of thermalization or slowing down of the fission fast neutron is called ‘moderation,’ and the material in a thermal reactor (i.e., a reactor in which fission is caused by neutrons with thermal energies) that is responsible for slowing down the fast fission neutrons is referred to as the moderator. Good nuclear moderators should possess the following attributes:
301
target elements. Low atomic number (Z) is thus a prime requirement of a good moderator. The density (number of atoms per unit volume) of the moderator and the likelihood of a scattering collision taking place must also be accounted for. Frequently used ‘Figures of merit’ for assessing moderators are the ‘slowing down power’ and the ‘moderating ratio.’ Figure 1 shows these Figures of merit for several candidate moderator materials. The slowing down power accounts for the mean energy loss per collision, the number of atoms per unit volume, and the scattering cross-section of the moderator. The tendency for a material to capture neutrons (the neutron capture cross-section) must also be considered. Thus, the second figure of merit, the moderation ratio, is the ratio of the slowing down power to the neutron absorption (capture) cross-section. Ideally the slowing down power is large, the neutron capture cross-section is small, and hence the moderating ratio is also large. Practically, the choices of moderating materials are limited to the few elements with atomic number <16. Gasses are of little use as moderators because of their low density, but can be combined in chemical compounds such as water (H2O) and heavy water (D2O). The available materials/compounds reduce to the four shown in Figure 1 (beryllium, carbon (graphite), water, and heavy water). Water is relatively unaffected by neutron irradiation, is easily contained, and inexpensive. However, the moderating ratio is reduced by the neutron absorption of hydrogen, requiring the use of enriched (in 235U) fuels to maintain the neutron economy. Heavy water is a good moderator because 1H2 and 8O16 do not absorb neutrons, the slowing down power is large, and the moderating ratio is therefore very large. Unfortunately,
do not react with neutrons (because if they are captured in the moderator the fission reaction cannot be sustained); should efficiently thermalize (slowdown) neutrons with few (elastic) collisions in the moderator; should be inexpensive; compatible with other materials in the reactor core; meet the core structural requirements; and ideally do not undergo any damaging chemical or physical changes when bombarded with neutrons. In the fast neutron thermalization process, the maximum energy lost per collision occurs when the target nucleus has unit mass, and tends to zero for heavy
1.E + 05
1.6 1.4 1.2
Moderating ratio
Slowing-down power (cm–1)
1.8
1 0.8 0.6 0.4
1.E + 04
1.E + 03
1.E + 02
0.2 0
1.E + 01 5
10 15 20 Mass number (M)
25
5
10 15 20 Mass number (M)
25
Figure 1 Moderator Figures of merit for several candidate moderators: beryllium (M ¼ 9); graphite, C (M ¼ 12); light water, H2O (M ¼ 18); heavy water, D2O (M ¼ 20).
302
Radiation Effects in Graphite
the cost of separating the heavy hydrogen isotope is large. Beryllium and beryllium oxide are good moderators but are expensive, difficult to machine, and suffer toxicity problems. Finally, graphite (carbon) is an acceptable moderator. It offers a compromise between nuclear properties, utility as a core structural material, and cost. It also has the advantage of being able to operate at very high temperatures (in the absence of oxygen). Unfortunately, the properties of graphite are markedly altered by neutron irradiation and this has to be considered in the design of graphite reactor cores.
Raw petroleum or pitch coke Calcined at 1300 ⬚C Calcined coke Crushed, ground, and blended Blended particles
Coal tar binder pitch
Mixed Cooled Extruded, molded, or isostatically pressed Green artifact
4.10.2 Nuclear Graphite Manufacture The invention of an electric furnace2 capable of reaching temperatures approaching 3000 C by Acheson in 1895 facilitated the development of the process for the manufacture of artificial polygranular graphite. Detailed accounts of the manufacture of polygranular graphite may be found elsewhere.2–4 Figure 2 summarizes the major processing steps in the manufacture of nuclear graphite. Nuclear graphite consists of two phases: a filler material and a binder phase. The predominant filler material, particularly in the United States, is a petroleum coke made by the delayed coking process. European nuclear graphites are typically made from a coal-tar pitch-derived coke. In the United Kingdom, Gilsonite coke, derived from naturally occurring bitumen found in Utah, USA, has been used. Both coke types are used for nuclear graphite production in Japan. The coke is usually calcined (thermally processed) at 1300 C prior to being crushed and blended. Typically, the binder phase is a coal-tar pitch. The binder plasticizes the filler coke particles so that they can be formed. Forming processes include extrusion, molding, vibrational molding, and isostatic pressing. The binder phase is carbonized during the subsequent baking operation (800–1000 C). Frequently, engineering graphites are pitch impregnated to densify the carbon artifact, followed by rebaking. Useful increases in density and strength are obtained with up to six impregnations, but two or three are more typical. The final stage of the manufacturing process is graphitization (2500–3000 C) during which, in simplistic terms, carbon atoms in the baked material migrate to form the thermodynamically more stable graphite lattice. Nuclear graphites require high chemical purity to minimize neutron absorption.
Baked at 800–1000 ⬚C Baked artifact Impregnated to densify (petroleum pitch) Rebaked and reimpregnated artifact Graphitized 2500–2800 ⬚C Graphite Purified Nuclear graphite
Figure 2 The major processing steps in the manufacture of nuclear graphite.
Moreover, certain elements catalyze the oxidation of graphite and must be reduced to an acceptable level. This is achieved by selecting very pure cokes, utilizing a high graphitization temperature (>2800 C), or by including a halogen purification stage in the manufacture of the cokes or graphite. Recently, comprehensive consensus specifications5,6 were developed for nuclear graphites. The electronic hybridization of carbon atoms (1s2, 2 2s , 2p2) allows several types of covalent bonded structure. In graphite, we observe sp2 hybridization in a planar network in which the carbon atom is bound to three equidistant nearest neighbors 120 apart in a given plane to form the hexagonal graphene structure. Covalent double bonds of both s-type and p-type are present, causing a shorter bond length than in the case of the tetrahedral bonding (s-type sp3 orbital hybridization only) observed in diamond. Thus, in its perfect form, the crystal structure of graphite (Figure 3) consists of tightly bonded (covalent) sheets of carbon atoms in a hexagonal lattice network.7 The sheets are
Radiation Effects in Graphite
A
B
c 0.670 nm
303
(grain sizes <100 mm) formed via isostatic pressing often exhibit complete isotropy in their properties. In response to the recent renewed interest in hightemperature gas-cooled reactors, many graphite vendors have introduced new nuclear graphites grades. Table 1 summarizes some of the grades available currently, although this list is not exhaustive. The graphite manufacturer is listed along with the coke type and comments related to the given graphite grade.
4.10.3 Graphite-Moderated Reactors
A
a 0.246 nm Figure 3 The crystal structure of graphite showing the ABAB stacking sequence of graphene planes in which the carbon atoms have threefold coordination. Reproduced from Burchell, T. D. In Carbon Materials for Advanced Technologies; Burchell, T. D., Ed.; Elsevier Science: Oxford, 1999, with permission from Elsevier.
weakly bound with van der Waals type bonds in an ABAB stacking sequence with a separation of 0.335 nm. The crystals in manufactured polygranular graphite are less than perfect, with approximately one layer plane in every six constituting a stacking fault. The graphite crystals have two distinct dimensions, the crystallite size La measured parallel to the basal plane and the dimension Lc measured perpendicular to the basal planes. In a coke-based nuclear graphite, values of La 80 nm and Lc 60 nm are typical.8 A combination of crystal structure bond anisotropy and texture resulting from forming imparts anisotropic properties to the filler coke and the manufactured nuclear graphite. The coke particles become preferentially aligned during forming, either with their long axis parallel to the forming axis in the case of extrusion, or with their long axis perpendicular to the forming axis in the case of molding or vibrational molding. Consequently, the graphite artifacts are often attributed with-grain and against-grain properties as in the American Society for Testing and Materials (ASTM) specifications.5,6 The degree of isotropy in manufactured graphite can be controlled through the processing route. Factors such as the nature of the filler coke, its size and size distribution, and the forming method contribute to the degree of isotropy. Nuclear graphites are typically medium or fine grain graphites (filler coke size <1.68 mm)5,6 and are considered near-isotropic. Fine grain graphites
Graphite has been used as a nuclear moderator in nuclear fission reactors since the very beginning of the nuclear age.1 Indeed, the Chicago Pile No. 1 (CP-1), constructed by Enrico Fermi under the stands at Stagg Field, University of Chicago, used The National Carbon Company’s AGOT grade graphite. On 2 December 1942, Enrico Fermi and his research team achieved the world’s first nuclear chain reaction in CP-1. Subsequently, the early weapons materials production reactors constructed in the United States, United Kingdom, France, and the former Union of Soviet Socialist Republics (USSR) were all graphitemoderated reactors, as were the first commercial power generating fission reactors. The core of a graphite-moderated reactor is comprised of stacks of graphite blocks that are usually keyed to one another to facilitate transmission of mechanical loads throughout the core. Vertical channels penetrate the core into which fuel stringers are placed via the reactor charge face using a refueling machine. The nuclear fuel, which may be natural uranium (a mixture of 238U, 235U, and 234U) or enriched uranium, is usually sheathed (clad) in a metallic cladding. Typically, the cladding is a light alloy (aluminum or magnesium), but it can also be stainless steel (requiring enriched fuel) if a higher fuel temperature is desired (>600 C). The fuel stringer and cladding material may be one and the same, as in the United Kingdom designed Magnox reactor,1 or the fuel may be in the form of stainless steel clad ‘pins’ arranged in graphite fuel sleeves, which are joined to one another and form the fuel stringer as in the UK Advanced Gas-Cooled Reactor1 (AGR). The metallic fuel clad retains the gaseous fission products that migrate from the fuel during the fission reaction and prevents contact between the gaseous coolant and the fuel. An alternative core layout uses integral fuel/moderator elements in which the uranium fuel is placed directly into cavities in the graphite moderator block,
304 Table 1
Radiation Effects in Graphite Currently available nuclear grade graphites
Grade
Manufacturer
Coke type
Comments
IG-430
Toyo Tanso
Pitch coke
IG-110
Toyo Tanso
Petroleum coke
NBG-10
SGL
Pitch coke
NBG-17
SGL
Pitch coke
NBG-18
SGL
Pitch coke
PCEA
GrafTech International
Petroleum coke
PGX 2020
GrafTech International Carbone of America
Petroleum coke Petroleum coke
2191
Carbone of America
Petroleum (sponge) coke
Isostatically molded, candidate for high-dose regions of NGNP concepts Isostatically molded, candidate for high-dose regions of NGNP concepts Extruded, candidate for high-dose regions of NGNP pebble bed concepts; PBMR core graphite Vibrationally molded, candidate for high-dose regions of NGNP prismatic core concepts Vibrationally molded, candidate for high-dose regions of NGNP pebble bed concepts; PBMR core graphite Extruded, candidate for high-dose regions of NGNP prismatic core concepts Large blocks for permanent structure in a prismatic core Isostatically molded, candidate for permanent structures in a prismatic core Isostatically molded, candidate for permanent structures in a prismatic core
and the entire block is discharged from the reactor when the fuel is spent. Fuel elements of this design typically utilize ceramic (UO2 or UC2) rather that metallic fuel so as to be capable of reaching higher fuel temperatures. The ceramic fuel kernel is over coated with layers of SiC and pyrolytic carbon to provide a fission product barrier and to negate the use of a metallic fuel clad (see Chapter 3.07, TRISOCoated Particle Fuel Performance), allowing the reactor core to operate at very high temperatures (>1000 C).1 The coated particle fuel is usually formed into fuel pucks or compacts but may be consolidated into fuel balls, or pebbles.1 The US designed modular high-temperature gas-cooled reactor (MHTGR) and Next Generation Nuclear Plant (NGNP), and the Japanese high-temperature test reactor (HTTR) are examples of gas-cooled reactors with high-temperature ceramic fuel. Additional vertical channels in the graphite reactor core house the control rods, which regulate the fission reaction by introducing neutron-adsorbing materials to the core, and thus reduce the number of neutrons available to sustain the fission process. When the control rods are withdrawn from the core, the self-sustaining fission reaction commences. Heat is generated by the moderation of the fission fragments in the fuel and moderation of fast neutrons in the graphite. The heat is removed from the core by a coolant, typically a gas, that flows freely through the core and over the graphite moderator. The coolant is forced through the core by a gas circulator and passes into a heat exchanger/boiler (frequently referred to as a steam generator).
The primary coolant loop (the reactor coolant) is maintained at elevated pressure to improve the coolant’s heat transfer characteristics and thus, the core is surrounded by a pressure vessel. A secondary coolant (water) loop runs through the heat exchanger and cools the primary coolant so that it may be returned to the reactor core at reduced temperature. The secondary coolant temperature is raised to produce steam which is passed through a turbine where it gives up its energy to drive an electric generator. Some reactor designs, such as the MHTGR, are direct cycle systems in which the helium coolant passes directly to a turbine. The reactor core and primary coolant loop are enclosed in a concrete biological-shield, which protects the reactor staff and public from g radiation and fission neutrons and also prevents the escape of radioactive contamination and fission product gasses that originate in the fuel pins/blocks. The charge face, refueling machine, control rod drives, discharge area, and cooling ponds are housed in a containment structure which similarly prevents the spread of any contamination. Additional necessary features of a fission reactor are (1) the refueling bay, where new fuel stringers or fuel elements are assembled prior to being loaded into the reactor core; (2), a discharge area and cooling ponds where spent fuel is placed while the short-lived isotopes are allowed to decay before the fuel can be reprocessed. The NGNP, a graphite-moderated, helium cooled reactor, is designed specifically to generate electricity and produce process heat, which could be used for the production of hydrogen, or steam generation for the recovery of tar sands or oil shale.
Radiation Effects in Graphite
Two NGNP concepts are currently being considered, a prismatic core design and a pebble bed core design. In the prismatic core concept, the TRISO fuel is compacted into sticks and supported within a graphite fuel block which has helium coolant holes running through its length.1 The graphite fuel blocks are discharged from the reactor at the end of the fuel’s lifetime. In the pebble bed core concept, the TRISO fuel is mixed with other graphite materials and a resin binder and formed into 6 cm diameter spheres or pebbles.1 The pebbles are loaded into the core to form a ‘pebble bed’ through which helium coolant flows. The pebble bed is constrained by a graphite moderator and reflector blocks which define the reactor core shape. The fuel pebbles migrate slowly down through the reactor core and are discharged at the bottom of the core where they are either sent to spent fuel storage or returned to the top of the pebble bed. Not all graphite-moderated reactors are gascooled. Several designs have utilized water cooling, with the water carried through the core in zirconium alloy tubes at elevated pressure, before being fed to a steam generator. Moreover, graphite-moderated reactors can also utilize a molten salt coolant, for example, the Molten Salt Reactor Experiment (MSRE)1 at Oak Ridge National Laboratory (ORNL). The fluid fuel in the MSRE consisted of UF4 dissolved in fluorides of beryllium and lithium, which was circulated through a reactor core moderated by graphite. The average temperature of the fuel salt was 650 C (1200 F) at the normal operating condition of 8 MW, which was the maximum heat removal capacity of the air-cooled secondary heat exchanger. The graphite core was small, being only 137.2 cm (54 in.) in diameter and 162.6 cm (64 in.) in height. The fuel salt entered the reactor vessel at 632 C (1170 F) and flowed down around the outside of the graphite core in the annular space between the core and the vessel. The graphite core was made up of graphite bars 5.08 cm (2 in.) square, exposed directly to the fuel which flowed upward in passages machined into the faces of the bars. The fuel flowed out of the top of the vessel at a temperature of 654 C (1210 F), through the circulating pump to the primary heat exchanger, where it gave up heat to a coolant salt stream. The core graphite, grade CGB, was specially produced for the MSRE, and had to have a small pore size to prevent penetration of the fuel salt, a long irradiation lifetime, and good dimensional stability. Moreover, for molten salt reactor moderators, a low permeability (preferably <108 cm2 s1) is desirable in order to prevent the build up of
305
unacceptable inventories of the nuclear poison Xe in the graphite. At ORNL, this was achieved by sealing the graphite surface using a gas phase carbon deposition process.1
135
4.10.4 Displacement Damage and Induced Structural and Dimensional Changes in Graphite The discovery of Fullerenes9 and carbon nanotubes,10 and other nanocarbon structures, has renewed interest in high-resolution microstructural studies of carbon nanostructures and the defects within them.11 This in turn has given new insight to the nature of displacement damage and the deformation mechanisms in irradiated graphite crystals. The binding energy of a carbon atom in the graphite lattice12 is about 7 eV. Impinging energetic particles such as fast neutrons, electrons, or ions can displace carbon atoms from their equilibrium positions. There have been many studies of the energy required to displace a carbon atom (Ed), as reviewed by Kelly,13 Burchell,14 Banhart,11 and Telling and Heggie.15 The value of Ed lies between 24 and 60 eV. The latter value has gained wide acceptance and use in displacement damage calculations, but a value of 30 eV would be more appropriate. Moreover, as discussed by Banhart,11 Hehr et al.,16 and Telling and Heggie,15 an angular dependence of the threshold energy for displacement would be expected. The value of Ed in the crystallographic c-axis is in the range 12–20 eV,11,17 while the in-plane value is much greater. The primary atomic displacements, primary knock-on carbon atoms (PKAs), produced by energetic particle collisions produce further carbon atom displacements in a cascade effect. The cascade carbon atoms are referred to as secondary knock-on atoms (SKAs). The displaced SKAs tend to be clustered in small groups of 5–10 atoms and for most purposes it is satisfactory to treat the displacements as if they occur randomly. The total number of displaced carbon atoms will depend upon the energy of the PKA, which is itself a function of the neutron energy spectrum, and the neutron flux. Once displaced, the carbon atoms recoil through the graphite lattice, displacing other carbon atoms and leaving vacant lattice sites. However, not all of the carbon atoms remain displaced and the temperature of irradiation has a significant influence on the fate of the displaced atoms and lattice vacancies. The displaced carbon atoms easily diffuse between the graphite layer planes
306
Radiation Effects in Graphite
in two dimensions and a high proportion will recombine with lattice vacancies. Others will coalesce to form C2, C3, or C4 linear molecules. These in turn may form the nucleus of a dislocation loop – essentially a new graphite plane. Interstitial clusters may, on further irradiation, be destroyed by a fast neutron or carbon knock-on atom (irradiation annealing). Adjacent lattice vacancies in the same graphitic layer are believed to collapse parallel to the layers, thereby forming sinks for other vacancies which are increasingly mobile above 600 C, and hence can no longer recombine and annihilate interstitials. The migration of interstitials along the crystallographic c-axis is discussed later. Banhart11 observed typical basal plane defects in a graphite nanoparticles using high-resolution transmission electron microscopy (HRTEM). These defects can be understood as dislocation loops which form when displaced interstitial atoms cluster and form less mobile agglomerates. Other interstitials condense onto this agglomerate which grows into a disk, pushing the adjacent apart. Further agglomeration leads to the formation of a new lattice planes (Figure 4). Other deformation mechanisms have been proposed for irradiated graphite. Wallace18 proposed a mechanism whereby interstitial atoms could facilitate sp3 bonds between the atomic basal planes, this mechanism allowing the stored energy (discussed in Section 4.10.5.1) to be explained. Jenkins19 argued that the magnitude of the increase in shear
1 nm Figure 4 A high-resolution electron micrograph showing the basal planes of a graphitic nanoparticle with an interstitial loop between two basal planes, the ends of the inserted plane are indicated with arrows. Reproduced from Banhart, F. Rep. Prog. Phys. 1999, 62, 1181–1221, with permission from IOP Publishing Ltd.
modulus (C44) with low dose irradiation could not be explained by interstitial clusters pinning dislocations, but that a few sp3 type covalent bonds between the planes could easily account for the observed changes. More recently, Telling and Heggie,15 in their ab-initio calculations of the energy of formation of the ‘spirointerstitial,’ advocate this mechanism to explain the stored energy characteristics of displacement damaged graphite, particularly the large energy release peak seen at 473 K (discussed in Section 4.10.5.1). The first experimental evidence of the interlayer interstitial–vacancy (IV) pair defect with partial sp3 character in between bilayers of graphite was recently reported by Urita et al.20 in their study of doublewalled carbon nanotubes (DWNTs). Jenkins19 invoked the formation of sp3 bonding to explain the c-axis growth observed as a result of displacement damage. If adjacent planes are pinned, one plane must buckle as the adjacent planes shrink due to vacancy shrinkage; buckled planes yield the c-axis expansion that cannot be explained by swelling from interstitial cluster alone. Telling and Heggie15 are very much in support of this position on the basis of their review of the literature and ab-initio simulations of the damage mechanisms in graphite. Their simulations showed how the spiro-interstitial (cross-link) essentially locked the planes together. Additionally, divacancies could lead to the formation of pentagons and heptagons in the basal planes causing the observed bending of graphene layers and c-axis swelling.11,21,22 The predicted c-axis crystal expansion via this mechanism is in closer agreement with the experimentally observed single crystal and highly oriented pyrolytic graphite (HOPG) dimensional change data. The buckling of basal planes as a consequence of irradiation damage has been observed in HRTEM studies of irradiated HOPG by Tanabe21 and Koike and Pedraza.22 In their study, Koike and Pedraza22 observed 300% expansion of thin HOPG samples subject to electron irradiation in an in-situ transmission electron microscope (TEM) study. Their experimental temperatures ranged from 238 to 939 K. They noted that the damaged microstructure showed retention of crystalline order up to 1 dpa (displacements per atom). At higher doses, they observed the lattice fringes break up in to segments 0.5–5 nm in length, with up to 15 rotation of the segments with respect to the original {0001} planes. The evidence in favor of the formation of bonds between basal planes involving interstitials is considerable. However, such bonds are not stable at high temperature. As reported by numerous authors and
Radiation Effects in Graphite
reviewers11,15,19,20 the sp3 like bond would be expected to break and recombine with lattice vacancies with increasing temperature, such that at T >500 K they no longer exist. Indeed, the irradiated graphite stored energy annealing peak at 473 K, and the HRTEM observations of Urita et al.20 demonstrate this clearly. Figure 5 shows a sequential series of HRTEM images illustrating the formation rates of interlayer defects at different temperatures with the same time scale (0–220 s) in DWNTs. The arrows indicate possible interlayer defects. At T ¼ 93 K (Figure 5(a)) the electron irradiation-induced defects are numerous, and the nanotubes inside are quickly damaged because of complex defects. At 300 K (Figure 5(b)), the nanotubes are more resistive to the damage from electron irradiation, yet defects are still viable. At 573 K (Figure 5(c)), defect formation is rarely observed and the DWNTs are highly resistant to the electron beam irradiation presumably because of the ease of defect self-annihilation (annealing). In an attempt to estimate the critical temperature for the annihilation of the IV defect pairs, a systematic HRTEM study was undertaken at elevated temperatures by Urita et al.20 The formation rate of the IV defects that showed sufficient contrast in the HRTEM is plotted in Figure 6. The reported numbers were considered to be an underestimate as single IV pairs may not have sufficient contrast to be
93 K
0s
300 K
307
convincingly isolated from the noise level and thus may have been missed. However, the data was considered satisfactory for indicating the formation rate as a function of temperature. The number of clusters of IV pairs found in a DWNT was averaged for several batches at every 50 K and normalized by the unit area. As observed in Figure 6, the defect formation rate displays a constant rate decline, with a threshold appearing at 450–500 K. This threshold corresponds to the stored energy release peak (discussed in Section 4.10.5.1) as shown by the dotted line in Figure 6. Evidentially, the irradiation damage resulting from higher temperature irradiations (above 473–573 K) is different in nature from that occurring at lower irradiation temperatures. Koike and Pedraza22 studied the dimensional change in HOPG caused by electron-irradiation-induced displacement damage. They observed in situ the growth c-axis of the HOPG crystals as a function of irradiation temperature at damage doses up to 1.3 dpa. Increasing c-axis expansion with increasing dose was seen at all temperatures. The expansion rate was however significantly greater at temperatures ≲473 K (their data was at 298 and 419 K) compared to that at irradiation temperatures ≳473 K (their data was at 553, 693, and 948 K). This observation supports the concept that separate irradiation damage mechanisms may exist at low irradiation temperatures (T <473 K), that is,
573 K
110 s
140 s
220 s
(a)
(b)
(c)
Figure 5 Sequential high-resolution transmission electron microscope images illustrating the formation rates of interlayer defects at different temperatures with the same time scale (0–220 s). (a) 93 K, (b) 300 K, (c) 573 K, in double-walled carbon nanotubes. The arrows indicate possible interlayer defects. Scale bar ¼ 2 nm. Reproduced from Urita, K.; Suenaga, K.; Sugai, T.; Shinohara, H.; Iijima, S. Phys. Rev. Lett. 2005, 94, 155502, with permission from American Physical Society.
308
Radiation Effects in Graphite
Defect formation rate (barns)
200
150
100
50
0
0
200
400 600 Temperature (K)
800
1000
Figure 6 Normalized formation rates of the clusters of interstitial–vacancy pair defects per unit area of bilayer estimated in high-resolution transmission electron microscope images recorded at different temperatures. The dotted line shows the known temperature for Wigner-energy release (473 K). Reproduced from Urita, K.; Suenaga, K.; Sugai, T.; Shinohara, H.; Iijima, S. Phys. Rev. Lett. 2005, 94, 155502, with permission from American Physical Society.
buckling due to sp3 bonded cross linking of the basal planes via interstitials, and at more elevated irradiation temperatures (T ≳ 473 K), where the buckling of planes is attributed to clustering of interstitials which induce the basal planes to bend, fragment, and then tilt. Koike and Pedraza22 also observed crystallographic a-axis shrinkage upon electron irradiation in-situ at several temperatures (419, 553, and 693 K). The shrinkage increased with dose at all irradiation temperatures, and the shrinkage rate reduced with increasing irradiation temperature. This behavior is attributed to buckling and breakage of the basal planes, with the amount of tilting and buckling decreasing with increasing temperature due to (1) a switch in mechanism as discussed above and (2) increased mobility of lattice vacancies above 673 K. Jenkins19,23 also discussed the deformation of graphite crystals in terms of a unit c-axis dislocation (prismatic dislocation), that is, one in which the Burgers vector, b, is in the crystallographic c-direction. The c-axis migration of interstitials can take place by unit c-axis dislocations. The formation and growth of these, and other basal plane dislocation loops
undoubtedly play a major role in graphite crystal deformation during irradiation. Ouseph24 observed prismatic dislocation loops (both interstitial and vacancy) in unirradiated HOPG using scanning tunneling microscopy (STM). Their study allowed atomic resolution of the defect structures. Such defects had previously been observed as regions of intensity variations in TEM studies in the 1960s.25 Telling and Heggie’s15 first principle simulations have indicated a reduced energy of migration for a lattice vacancy compared to the previously established value. Therefore, they argue, the observed limited growth of vacancy clusters at high temperatures (T >900 K) indicates the presence of a barrier to further coalescence of vacancy clusters (i.e., vacancy traps). Telling and Heggie implicate a cross-planer metastable vacancy cluster in adjacent planes as the possible trap. The disk like growth of vacancy clusters within a basal plane ultimately leads to a prismatic dislocation loop. TEM observations show that these loops appear to form at the edges of interstitial loops in neighboring planes in the regions of tensile stress. The role of vacancies needs to be reexamined on the basis of the foregoing discussion. If the energy of migration is considerably lower than that previously considered, and there is a likelihood of vacancy traps, the vacancy and prismatic dislocation may well play a larger role in displacement damage induced incrystal deformation. The diffusion of vacancy lines to the crystal edge essentially heals the damage, such that crystals can withstand massive vacancy damage and recover completely. Regardless of the exact mechanism, the result of carbon atom displacements is crystallite dimensional change. Interstitial defects will cause crystallite growth perpendicular to the layer planes (c-axis direction), and relaxation in the plane due to coalescence of vacancies will cause a shrinkage parallel to the layer plane (a-axis direction). The damage mechanism and associated dimensional changes are illustrated (in simplified form) in Figure 7. As discussed above, this conventional view of c-axis expansion as being caused solely by the graphite lattice accommodating small interstitial aggregates is under some doubt, and despite the enormous amount of experimental and theoretical work on irradiation-induced defects in graphite, we are far from a widely accepted understanding. It is to be hoped that the availability of high-resolution microscopes will facilitate new damage and annealing studies of graphite leading to an improved understanding of the defect structures and of crystal deformation under irradiation.
309
Radiation Effects in Graphite
Collapsing vacancy line
(c)
Contraction
(a)
Interstitial
Vacancy New plane Expansion
Figure 7 Neutron irradiation damage mechanism illustrating the induced crystal dimensional strains. Reproduced from Burchell, T. D. In Carbon Materials for Advanced Technologies; Burchell, T. D., Ed.; Elsevier Science: Oxford, 1999, with permission from Elsevier.
100
80
ced shrinkage
Irradiation-indu
(%)
60
40 20
0 20
Ne
15
ro
ut
2000
10
n do
2500 ⬚C) re ( atu r e mp
5
se
)
pa
(d
Dimensional changes can be very large, as demonstrated in studies on well-ordered graphite materials, such as HOPG that has frequently been used to study the neutron-irradiation-induced dimensional changes of the graphite crystallite.13,26 Price27 conducted a study of the neutron-irradiation-induced dimensional changes in pyrolytic graphite. Figure 8 shows the crystallite shrinkage in the a-direction for neutron doses up to 12 dpa for samples that were graphitized at a temperature of 2200–3300 C prior to being irradiated at 1300–1500 C. The a-axis shrinkage increases linearly with dose for all of the samples, but the magnitude of the shrinkage at any given dose decreases with increasing graphitization temperature. Similar trends were noted for the c-axis expansion. The significant effect of graphitization temperature on irradiation-induced dimensional change accumulation can be attributed to thermally induced improvements in crystal perfection, thereby reducing the number of defect trapping sites in the lattice. Nuclear graphites possess a polycrystalline structure, usually with significant texture resulting from the method of forming during manufacture. Consequently, structural and dimensional changes in polycrystalline graphites are a function of the crystallite dimensional changes and the graphite’s texture. In polycrystalline graphite, thermal shrinkage cracks that occur during manufacture and that are preferentially aligned in the crystallographic a-direction will initially accommodate the c-direction expansion, so mainly a-direction contraction will be observed. The graphite thus undergoes net volume shrinkage. With increasing neutron dose (displacements), the incompatibility of crystallite dimensional changes leads to the generation of new porosity, and the volume shrinkage rate falls, eventually reaching zero. The graphite now begins to swell at an
3000 n te atio
0 3500
itiz aph
Gr
Figure 8 Neutron irradiation-induced a-axis shrinkage behavior of pyrolytic graphite showing the effects of graphitization temperature on the magnitude of the dimensional changes. Reproduced from Burchell, T. D. In Carbon Materials for Advanced Technologies; Burchell, T. D., Ed.; Elsevier Science: Oxford, 1999, with permission from Elsevier.
increasing rate with increasing neutron dose. The graphite thus undergoes a volume change ‘turnaround’ into net growth that continues until the generation of cracks and pores in the graphite, due to differential crystal strain, eventually causes total disintegration of the graphite. Irradiation-induced volume and dimensional change data for H-451 are shown28 in Figures 9–11. The effect of irradiation temperature on volume change is shown in Figure 9. The ‘turn-around’ from volume shrinkage to growth occurs at a lower fluence and
310
Radiation Effects in Graphite
10 H-451 @ 600 ⬚C H-451 @ 900 ⬚C
Volume change (%)
8 6 4 2 0 -2 -4 -6 -8 -10 0
1
2
3
4
Fast fluence 1026 n m-2 [E > 0.1 MeV]
Figure 9 Irradiation-induced volume changes for H-451 graphite at two irradiation temperatures. From Burchell, T. D.; Snead, L. L. J. Nucl. Mater. 2007, 371, 18–27.
H-451 graphite irradiated at 600 ⬚C
Dimensional change (%)
1 0 -1 -2 -3 -4
Perpendicular to extrusion (AG) Parallel to extrusion direction (WG)
-5 -6
0
1
2
3
4
5
Fast fluence 1026 n m-2 [E > 0.1 MeV]
Figure 10 Dimensional change behavior of H-451 graphite at an irradiation temperature of 600 C. From Burchell, T. D.; Snead, L. L. J. Nucl. Mater. 2007, 371, 18–27.
Dimensional change (%)
4.10.5 Neutron-Induced Property Changes
H-451 graphite irradiated at 900 ⬚C
5
4.10.5.1
Perpendicular to extrusion (AG) Parallel to extrusion direction (WG)
4 3 2 1 0 -1 -2 -3
0
0.5 1 1.5 2 2.5 Fast fluence 1026 n m-2 [E > 0.1 MeV]
available at the higher temperatures and the c-axis growth dominates the a-axis shrinkage at lower doses. The irradiation-induced dimensional changes of H-451at 600 and 900 C are shown in Figures 10 and 11, respectively. H-451 graphite is an extruded material and therefore, the filler coke particles are preferentially aligned in the extrusion axis (parallel direction). Consequently, the crystallographic a-direction is preferentially aligned in the parallel direction and the a-direction shrinkage is more apparent in the parallel (to extrusion) direction, as indicated by the parallel direction dimensional change data in Figures 10 and 11. The dimensional and volume changes are greater at an irradiation temperature of 600 C than at 900 C; that is, both the maximum shrinkage and the turnaround dose are greater at an irradiation temperature of 600 C. This temperature effect can be attributed to the thermal closure of internal porosity that is aligned parallel to the a-direction that accommodates the c-direction swelling. At higher irradiation temperatures, a greater fraction of this accommodating porosity is closed and thus the shrinkage is less at the point of turnaround. A general theory of dimensional change in polygranular graphite due to Simmons29 has been extended by Brocklehurst and Kelly.30 For a detailed account of the treatment of dimensional changes in graphite the reader is directed to Kelly and Burchell.31
3
Figure 11 Dimensional change behavior of H-451 graphite at an irradiation temperature of 900 C. From Burchell, T. D.; Snead, L. L. J. Nucl. Mater. 2007, 371, 18–27.
the magnitude of the volume shrinkage is smaller at the higher irradiation temperature. This effect is attributed to the thermal closure of aligned microcracks in the graphite which accommodate the c-axis growth. Hence, there is less accommodating volume
Wigner Energy
The release of Wigner energy (named after the physicist who first postulated its existence) was historically the first problem of radiation damage in graphite to manifest itself. The lattice displacement processes previously described can cause an excess of energy in the graphite crystallites. The damage may comprise Frankel pairs or at lower temperatures the sp3 type bond previously discussed and observed by Urita et al.20 When an interstitial carbon atom and a lattice vacancy recombine, or interplanar bonds are broken, their excess energy is given up as ‘stored energy.’ If sufficient damage has accumulated in the graphite, the release of this stored energy can result in a rapid rise in temperature. Stored energy accumulation was found to be particularly problematic in the early graphite-moderated reactors, which operated at relatively low temperatures. Figure 12 shows the rate of release of stored energy with
Radiation Effects in Graphite
0.7
dS/dT (cal g–1 ⬚C–1)
0.6
Exposures in MWd/at and dpa (approximately)
0.5 0.4
Specific heat
0.3
5700/0.60 1075/0.10
0.2 0.1 0
100/0.01 100
200 300 400 Annealing temperature (⬚C)
500
Figure 12 Stored energy release curves for CSF graphite irradiated at 30 C in the Hanford K reactor cooled test hole. Source: Nightingale, R. E. Nuclear Graphite; Academic Press: New York, 1962. From Burchell, T. D. In Carbon Materials for Advanced Technologies; Burchell, T. D., Ed.; Elsevier Science: Oxford, 1999, with permission from Elsevier.
temperature, as a function of temperature, for graphite samples irradiated at 30 C to low doses in the Hanford K reactor.32 The release curves are characterized by a peak occurring at 200 C. This temperature has subsequently been associated with annealing of interplanar bonding involving interstitial atoms.20 In Figure 12, the release rate exceeds the specific heat and therefore, under adiabatic conditions, the graphite would rise sharply in temperature. For ambient temperature irradiations it was found9 that the stored energy could attain values up to 2720 J g1, which if released adiabatically would cause a temperature rise of some 1300 C. A simple experiment,8 in which samples irradiated at 30 C were placed in a furnace at 200 C and their temperature monitored, showed that when the samples attained a temperature of 70 C their temperature suddenly increased to a maximum of about 400 C and then returned to 200 C. In order to limit the total amount of stored energy in the early graphite reactors, it became necessary to periodically anneal the graphite. The graphite’s temperature was raised sufficiently, by nuclear heating or the use of inserted electrical heaters, to ‘trigger’ the release of stored energy. The release then self-propagated slowly through the core, raising the graphite moderator temperature and thereby partially annealing the graphite core. Indeed, Arnold33 reports that it was during such a reactor anneal that the Windscale (UK) reactor accident occurred in 1957. Rappeneau et al.34 report a second release peak at very high temperatures (1400 C). They studied the energy release up to temperatures of 1800 C
311
of graphites irradiated in the reactors BR2 (Mol, Belgium) and HFR (Petten, Netherlands) at doses between 1000 and 4000 MWd T1 and at temperatures between 70 and 250 C. At these low irradiation temperatures, there is little or no vacancy mobility, so the resultant defect structures can only involve interstitials. On postirradiation annealing to high temperatures, the immobile single vacancies become increasingly mobile and perhaps their elimination and the thermal destruction of complex interstitial clusters or distorted and twisted basal planes contribute to the high-temperature stored energy peak. The accumulation of stored energy in graphite is both dose and irradiation temperature dependent. With increasingly higher irradiation temperatures, the total amount of stored energy and its peak rate of release diminish, such that above an irradiation temperature of 300 C stored energy ceases to be a problem. Accounts of stored energy in graphite can be found elsewhere.1,8,29,32 4.10.5.2 Mechanical and Physical Properties The mechanical and physical properties of several medium-grained and fine-grained nuclear grade graphites currently in production are given in Table 2 (see also Chapter 2.10, Graphite: Properties and Characteristics). The coke type, forming method, and potential uses of these grades are in Table 1. The most obvious difference between the four grades listed in Table 2 is the filler particle sizes. Grade IG-110 is an isostatically pressed, isotropic grade, whereas the others grades shown are near-isotropic and have properties reported either with-grain or against-grain. As discussed earlier (see Section 4.10.2), the orientation of the filler coke particles is a function of the forming method. The mechanical properties of nuclear graphites are substantially altered by radiation damage. In the unirradiated condition, nuclear graphites behave in a brittle fashion and fail at relatively low strains. The stress–strain curve is nonlinear, and the fracture process occurs via the formation of subcritical cracks, which coalesce to produce a critical flaw.35,36 When graphite is irradiated, the stress–strain curve becomes more linear, the strain to failure is reduced, and the strength and elastic modulus are increased. On irradiation, there is a rapid rise in strength, typically 50%, that is attributed to dislocation pinning at irradiation-induced lattice defect sites. This effect is largely saturated at doses >1 dpa. Above 1 dpa, a more gradual increase in strength occurs because of
312
Radiation Effects in Graphite
Table 2
Typical physical and mechanical properties of unirradiated nuclear graphites
Property
Graphite grade IG-110
PCEA
NBG-10
NBG-18
Maximum filler particle size (mm) Bulk density (g cm3) Tensile strength (MPa)
10 1.77 24.5
Flexural strength (MPa)
39.2
Compressive strength (MPa)
78.5
800 1.83 21.9 (WG) 16.9 (AG) 32.4 (WG) 23.3 (AG) 60.8 (WG) 67.6 (AG) 11.3 (WG) 9.9 (AG) 162 (WG) 159 (AG) 3.5 (WG) 3.7 (AG) (30–100 C) 7.3 (WG) 7.8 (AG)
1600 1.79 20.0 (WG) 18.0 (AG) 24.0 (WG) 27.0 (AG) 47.0 (WG) 61.0 (AG) 9.7 (WG) 9.7 (AG) 148 (WG) 145 (AG) 4.1 (WG) 4.6 (AG) (20–200 C) 9.1 (WG) 9.3 (AG)
1600 1.88 21.5 (WG) 20.5 (AG) 28 (WG) 26 (AG) 72.0 (WG) 72.5 (AG) 11.2 (WG) 11.0 (AG) 156 (WG) 150 (AG) 4.5 (WG) 4.7 (AG) (20–200 C) 8.9 (WG) 9.0 (AG)
Young’s modulus (GPa)
9.8
Thermal conductivity (W m1 K1) (measured at ambient temperature) Coefficient of thermal expansion (106 K1) (over given temperature range)
116
Electrical resistivity (mO m)
11
4.5 (350–450 C)
WG, with-grain; AG, against-grain.
ðE=E0 Þirradiated ¼ ðE=E0 Þpinning ðE=E0 Þstructure
½1
where E/E0 is the ratio of the irradiated to unirradiated elastic modulus. The dislocation pinning contribution to the modulus change is due to relatively mobile small defects and is thermally annealable at 2000 C. The irradiation-induced elastic modulus changes for GraphNOL N3M graphite37 are shown in Figure 13. The low dose change due to dislocation pinning (dashed line) saturates at a dose <1 dpa. The elastic modulus and strength are related by a Griffith theory type relationship. Strength; st ¼ ½GE=pc1=2
½2
50
40 Young’s modulus (GPa)
structural changes within the graphite. For nuclear graphites, the dose at which the maximum strength is attained loosely corresponds with the volume change turnaround dose, indicating the importance of pore closure and generation in controlling the high-dose strength behavior, and may be as much as twice the unirradiated value. The strain behavior of nuclear graphites subjected to an externally applied load is largely controlled by shear of the component crystallites. As with strength, irradiation-induced changes in Young’s modulus are the combined result of in-crystallite effects, due to low fluence dislocation pinning, and superimposed structural change external to the crystallite. The effects of these two mechanisms are generally considered separable, and related by
30
20 875 ⬚C 600 ⬚C
10
0
0
5
10
15 20 Fluence (dpa)
25
30
Figure 13 Neutron irradiation-induced Young’s modulus changes for GraphNOL N3M at irradiation temperatures 600 and 875 C. From Burchell, T. D.; Eatherly, W. P. J. Nucl. Mater. 1991, 179–181, 205–208.
where G is the fracture toughness or strain energy release rate ( J m2), E is the elastic modulus (Pa), and c is the flaw size (m). Thus, irradiation-induced changes in st and E (in the absence of changes in [G/c]) should follow st/E1/2. High-dose data reported by Ishiyama et al.38 show significant deviation from this relationship for grade IG-110 graphite, indicating that changes in G and or c must occur.
Thermal conductivity (Wm–1 K–1)
Radiation Effects in Graphite
160 12 dpa 25 dpa 26 dpa Unirradiated
140 120 100 80 60 40 20 0 0
200
400 600 800 Temperature (⬚C)
1000
1200
Figure 14 Temperature dependence of thermal conductivity in the irradiated and unirradiated condition for typical nuclear grade graphite. Irradiation temperature ¼ 600 C.
Graphite is a phonon conductor of heat. Therefore, any reduction in the intrinsic defect population causes a reduction in the degree of phonon-defect scattering, an increase in the phonon mean free path, and an increase in the thermal conductivity. In graphite, such thermally induced improvements are attributable to increases in crystal perfection and a concomitant increase in the size of the regions of coherent ordering upon graphitization. With increasing temperature, the dominant phonon interaction becomes phonon– phonon scattering (Umklapp processes). Therefore, there is a reduction of thermal conductivity with increasing temperature.39 This decrease in the thermal conductivity with increasing temperature can clearly be seen in Figure 14. The mechanism of thermal conductivity and the degradation of thermal conductivity have been extensively reviewed.13,14,26,40 The increase of thermal resistance due to irradiation damage has been ascribed by Taylor et al.41 to the formation of (1) submicroscopic interstitial clusters, containing 4 2 carbon atoms; (2) vacant lattice sites, existing as singles, pairs, or small groups; and (3) vacancy loops, which exist in the graphite crystal basal plane and are too small to have collapsed parallel to the hexagonal axis. The contribution of collapsed lines of vacant lattice sites and interstitial loops, to the increased thermal resistance, is negligible. The reduction in thermal conductivity due to irradiation damage is temperature and dose sensitive. At any irradiation temperature, the decreasing thermal conductivity will reach a ‘saturation limit.’ This limit is not exceeded until the graphite undergoes gross structural changes at very high doses. The ‘saturated’ value of conductivity will be attained more rapidly, and will be lower, at lower irradiation temperatures.42 In graphite, the neutron irradiation-induced degradation of
313
thermal conductivity can be very large, as illustrated in Figure 14. This reduction is particularly large at low temperatures. Bell et al.43 have reported that the room temperature thermal conductivity of pile grade A (PGA) graphite is reduced by more than a factor of 70 when irradiated at 155 C to a dose of 0.6 dpa. At an irradiation temperature of 355 C, the room temperature thermal conductivity of PGA was reduced by less than a factor of 10 at doses twice that obtained at 155 C. Above 600 C, the reduction of thermal conductivity is less significant. For example, Kelly8 reported the degradation of PGA at higher temperatures: at an irradiation temperature of 600 C and a dose of 13 dpa, the thermal conductivity was degraded only by a factor of 6; at irradiation temperatures of 920 and 1150 C, the degradation was minimal (less than a factor of 4 at 7 dpa). For the fine-grained, isomolded graphite shown in Figure 14, the degradation of thermal conductivity at the irradiation temperature (600 C) was only by a factor of 3, but was by a factor 6 at a measurement temperature of 100 C. There are two principal thermal expansion coefficients in the hexagonal graphite lattice; ac, the thermal expansion coefficient parallel to the hexagonal c-axis and aa, the thermal expansion coefficient parallel to the basal plane (a-axis). The thermal expansion coefficient in any direction at an angle f to the c-axis of the crystal is aðfÞ ¼ ac cos2 f þ aa sin2 f
½3
The value of ac varies linearly with temperature from 25 106 K1 at 300 K to 35 106 K1 at 2500 K. In contrast, aa is much smaller and increases rapidly from 1.5 106 K1 at 300 K to 1 106 K1 at 1000 K, and remains relatively constant at temperatures up to 2500 K.39 The large anisotropy in the crystal coefficient of thermal expansion (CTE) values is a direct consequence of the bond anisotropy and the resultant anisotropy in the crystal lattice compliances. The thermal expansion of polycrystalline graphites is controlled by the thermal closure of aligned internal porosity which forms as a result of thermal shrinkage strains on cooling after graphitization. Thus, the c-axis expansion of the graphite crystals is initially, partially accommodated by this internal porosity and a much lower bulk CTE value is observed. On further heating, the graphite crystals fill more of the available internal porosity and more of the c-axis expansion is observed. The bulk CTE thus increases with temperature (Figure 15).
Radiation Effects in Graphite
6
5.0 4.0 3.0 2.0 1.0 0.0 0
200 400 600 800 Measurement temperature (⬚C)
1000
Figure 15 Temperature dependence of the coefficient of thermal expansion for typical nuclear grade graphite.
As the CTE of polycrystalline graphite is dependent on the pore structure, irradiation-induced changes in the pore structure (see discussion of structural changes in Section 4.10.4) can be expected to modify the thermal expansion behavior of carbon materials. Burchell and Eatherly37 report the behavior of GraphNOL N3M (which is typical of many fine-textured graphites), which undergoes an initial increase in the CTE followed by a steady reduction to a value less than half the unirradiated value of 5 106 C1 (Figure 16). Similar behavior is reported by Kelly8 for grade IM1-24 graphite. Heat energy is stored in the crystal lattice in the form of lattice vibrations. The Debye equation therefore gives the specific heat, C, as follows: C ¼ 9R
T yD
y
3 ðT 0
z4 ez dz ðez 1Þ2
½4
Coefficient of thermal expansion a(10-6 ⬚C-1)
6.0
875 ⬚C 600 ⬚C 5
4
3
2
0
5
10
15 20 Fluence (dpa)
25
30
Figure 16 The irradiation-induced changes in coefficient of thermal expansion (25–500 C) for GraphNOL N3M graphite at two irradiation temperatures. From Burchell, T. D.; Eatherly, W. P. J. Nucl. Mater. 1991, 179–181, 205–208.
2400 Specific heat (J Kg–1 K–1)
Average coefficient of thermal expansion (10-6 ⬚C-1)
314
2000 1600 1200 800 Calculated value Experimental data
400 0 0
500
1000 1500 2000 Temperature (K)
2500
3000
where R is the gas constant (8.314 J mol1 K1); T, the temperature; yD , the Debye temperature; and z ¼ ho/2kTp, where o is the frequency of vibrational oscillations; k, the Boltzmann’s constant; T, the temperature; and h is the Plank’s constant. At low temperatures, where (T/yD) <0.1, z in eqn [4] is large, we can approximate eqn [4] by allowing the upper limit in the integral to go to infinity such that the integral becomes (p4/15), and on differentiating we get
Figure 17 The temperature dependence of the specific heat of graphite, a comparison of calculated values and literature data for POCO AXM-5Q graphite. Sources: ASTM C 781. Standard Practice for Testing Graphite and Boronated Graphite Materials for High-Temperature Gas-Cooled Nuclear Reactor Components, Annual Book of Standards. ASTM International: West Conshohocken, PA; Vol. 05.05; Hust, J. G. NBS Special Publication 260–89; US Department of Commerce, National Bureau of Standards, 1984; p 59.
C ¼ 1941ðT =yD Þ3 J mol1 K1
10% of the Debye temperature (0.1yD), the specific heat should rise exponentially with temperature to a constant value at T yD. Figure 17 shows the specific heat of graphite over the temperature range 300–3000 K. The data has been shown to be well represented by the eqn [6],44,45 and is applicable to all nuclear graphites. The release of energy from the thermal annealing of damage accumulated at low irradiation temperature
½5
Thus, at low temperatures, the specific heat is proportional to T3 (eqn [5]). At high temperatures, z is small and the integral in eqn [4] reduces to z2dz; hence, on integrating we get the Dulong–Petit value of 3R, that is, the theoretical maximum specific heat of 24.94 J mol1 K1. As we are typically concerned only with the specific heat at temperatures above
Radiation Effects in Graphite
(Wigner energy) will reduce the effective specific heat (see Section 4.10.4). Cp ¼
11:07T 1:644
1 J Kg1 K1 ½6 þ 0:0003688T 0:02191
The electrical resistivity of graphite is also affected by radiation damage. The mean free path of the conduction electron in unirradiated graphite is relatively large, being limited only by crystallite boundary scattering. Neutron irradiation introduces (1) scattering centers, which reduce charge carrier mobility; (2) electron traps, which decrease the charge carrier density; and (3) additional spin resonance. The net effect of these changes is to increase the electrical resistivity on irradiation, initially very rapidly, with little or no subsequent change to relatively high fluence.14,37 A subsequent decrease at very high neutron doses is attributed to structural degradation.
4.10.6 Irradiation Creep 4.10.6.1 The Relevance of Creep to Reactor Design and Operation Graphite will undergo creep (inelastic strain) during neutron irradiation and under stress at temperatures where thermal creep is generally negligible. The phenomenon of irradiation creep has been widely studied because of its significance to the operation of graphite-moderated fission reactors. Indeed, if irradiation-induced stresses in graphite moderators could not relax via radiation creep, rapid core disintegration would result. The total strain, eTotal, in a graphite component under irradiation in a reactor core is given by the materials constitutive equation: eTotal ¼ ee þ et þ ed þ ec
½7
where ee is the elastic strain; et, the thermal strain; ed, the dimensional change strain; and ec is the creep strain, which is given as ec ¼ ep þ es
½8
where ep and es are the primary and secondary creep strains, respectively. Tsang and Marsden46 concluded that irradiation creep strain is particularly important in reactor design because without creep strain self-induced shrinkage stresses would build up to levels exceeding the graphite component failure strength. The significance of irradiation creep to reactor core design and operation has been the subject
315
of recent work, where it has been shown how uncertainties in the assumed magnitudes of the irradiation-induced creep strains in a graphite reactor core component can substantially impact the predicted stress levels, and hence the predicted failure probabilities of core components.47,48 Li et al.47 assumed the current UK creep law and showed that a 50% decrease in the assumed creep strain resulted in a 50% increase in the magnitude of the predicted hoop stress in a hollow cylindrical core brick. Similarly, a 50% increase in the assumed creep strain yielded a 30% reduction in the predicted brick hoop stress. Wang and Yu48 report the effect of varying creep strain ratio (analogous to Poisson’s ratio, but where the two perpendicular strains are creep strains) on the magnitude of the modified equivalent stress in a graphite component and the associated probability of failure, as a function of neutron dose. In addition, they examined the influence of primary creep in reducing the magnitudes of stresses and associated failure probabilities in graphite core components. Wang and Yu’s results clearly indicate that variations of the creep strain ratio resulted in considerable change in the stress distributions and the corresponding failure probabilities of graphite components. In addition, they showed that the primary creep appears to play the same important role as secondary creep in certain cases. Because of the significance of irradiation-induced creep to the stress levels in graphite core components, accurate models of creep have long been sought. Recently, the breakdown of the currently accepted model(s) of creep at high temperatures and doses has been reported, and possible improvement or alternative models have been postulated.49,50 Analysis of the creep behavior of H-451 at high doses indicated that further modification to the current Kelly and Burchell51 model is required to allow for the generation of new porosity at higher doses and temperatures.50 The extent to which high-dose creep strain behavior differs between the compressive load and tensile load situations is shown in Figure 18, which compares the creep behavior of ATR-2E graphite for the þ5 MPa and the 5 MPa loading cases;52 the dashed lines are polynomial fits to the data. A more rapidly increasing creep rate in the tensile loading case compared to the compressive case is clearly observed. Because of the importance of irradiation-induced creep to the design and operation of graphite reactor cores, the subject is treated here in considerable detail,
Radiation Effects in Graphite
Creep strain (%)
316
2 1.8 1.6 1.4 1.2 1 0.8 0.6 0.4 0.2 0
Tensile creep strain (% MPa) Compressive creep strain (%) Poly tensile creep strain (% MPa) Poly compressive creep strain (%)
0
1 2 3 Neutron dose 1022 n cm-2 [E > 50 keV]
4
Figure 18 A comparison of the tensile (þ5 MPa) and compressive creep (5 MPa) rates of ATR-2E graphite at irradiation temperature of 500–550 C. Source: Haag, G. Properties of ATR-2E Graphite and Property Changes due to Fast Neutron Irradiation; Report No. Jul-4183; Published FZ-J, Germany, 2005; Available at http://juwel.fz-juelich.de.
including a review of the in-crystal creep mechanism and irradiation-induced graphite creep models. 4.10.6.2 The Irradiation-Induced Creep Mechanism (In-Crystal) A mechanism for the irradiation-induced creep of graphite was proposed by Kelly and Foreman53 which involves irradiation-induced basal plane dislocation pinning/unpinning in the graphite crystals. Pinning sites are created and destroyed by neutron irradiation (radiation annealing). Under neutron irradiation, dislocation lines in the basal planes may be completely or partially pinned depending upon the dose and temperature of irradiation. The pinning points were speculated to be interstitial atom clusters 4 2 atoms in size,54,55 that is, the same defects clusters assumed to contribute to the reduction in thermal conductivity. The interstitial clusters are temporary barriers as they are annealed (destroyed) by further irradiation. Thus, irradiation can release dislocation lines from their original pinning site and the crystal can flow as a result of basal plane slip at a rate determined by the rate of pinning and unpinning of dislocations. Kelly and Foreman’s theory assumes that polycrystalline graphite consists of a single phase of true density r0 and apparent density r. The material may be divided into elementary regions in which the stress may be considered uniform and which may be identified as monocrystalline graphite. Significantly, the model excludes porosity. It is further assumed that the only deformation mode is basal plane slip for which the strain rate is determined by e_ xz ¼ kðsxz Þf
½9
and e_ yz ¼ kðsyz Þf
½10
where f is the fast neutron flux; k, the steady-state creep coefficient, and s is the stress in the given direction. The microscopic deformation assumes the usual relationship between the basal plane shear strain rate (e_) and the mobile dislocation density (O), and is given by e_ ¼ Obn ¼ ksf ½11 where b is the Burger’s vector and n is the dislocation velocity as a function of the pinning point concentration in the basal plane as the pins are created and destroyed by neutron flux. The dislocation line flow model used the flexible line approach where the dislocation line is pinned/unpinned and the dislocation line bowing is a function of the line tension and pin spacing. The concentration of pinning sites increases under irradiation from the initial value (from intrinsic defects) to a steady-state concentration. The initial creep rate is high and decreases to a steady-state value as the pinning concentration saturates at a level controlled by the neutron flux and temperature. This saturation would be expected to occur over the same dose scale as the reduction of thermal conductivity to its saturation limit (see Section 4.10.5.2). Thus, a two stage model can be envisioned where the primary creep rate is initially high and falls to a secondary or ‘steady-state’ creep rate. The steady-state creep term should be the dominant term when the dose has reached values at which physical property changes due to dislocation pinning have saturated (see Section 4.10.5.2). Kelly and Foreman state that at higher temperatures the steady-state (secondary) creep rate (k)
Radiation Effects in Graphite
would be expected to increase because of (1) incompatibility of crystal strains increasing the internal stress and thus enhancing the creep rate, and (2) additional effects due to the destruction of interstitial pins by thermal diffusion of vacancies (thermal annealing as well as irradiation annealing). Kelly and Foreman53 further speculate that the nonlinearity of creep strain with stress, which is expected at higher stress levels, may also be related to the high-dose dimensional change behavior of polycrystalline graphite.56 The possibility of other dislocation and crystal deformation mechanisms being involved in irradiation creep must also be considered. For example, prismatic dislocations may play an enhanced role at high temperatures (>250 C) when the graphite lattice is under stress, as suggested by others.57 Are there mechanisms of dislocation climb and glide that need to be explored? Can dislocation lines climb/glide past the assumed interstitial cluster barriers via a mechanism that is active only when structural rearrangements occur during irradiation? This behavior is analogous to carbons and graphites undergoing thermal creep when they undergo structural reorganization, that is, during carbonization and graphitization (thermal relaxation or slumping). 4.10.6.3
Review of Prior Creep Models
4.10.6.3.1 Linear viscoelastic creep model
Irradiation-induced (apparent) creep strain is conventionally defined as the difference between the dimensional change of a stressed specimen and an unstressed specimen irradiated under identical conditions. Early creep data was found to be well described by a viscoelastic creep model1,58–63 where total irradiation creep (ec Þ ¼ primary (transient) creep þ secondary (steady-state) creep. as ½12 ec ¼ ½1 expðbgÞ þ ksg E0 where ec is the total creep strain; s, the applied stress; E0, the initial (preirradiated) Young’s modulus; g, the fast neutron fluence; a and b are constants (a is usually ¼ 1); and k is the steady-state creep coefficient in units of reciprocal neutron dose and reciprocal stress. Equation [13] thus conforms to the Kelly–Foreman theory of creep with an initially large primary creep coefficient, while the dislocation pinning sites develop to the equilibrium concentration, at which time the creep coefficient has fallen to the steady-state or secondary value. Early creep experiments in several
317
countries showed the primary creep saturated at approximately one elastic strain (s/E0) so that the true creep may be represented as s ½13 ec ¼ þ ksg E0 This is often normalized to the initial elastic strain and written as ec ¼ 1 þ kE0 g
½14
in elastic strain units (esu) (esu is defined as the externally applied stress divided by the initial static Young’s modulus), or creep strain per unit initial elastic strain; kE is the creep coefficient in units of reciprocal dose [United Kingdom 0.23 1020 cm2 n1 EDN up to Tirr 500 C]. (EDN – equivalent DIDO nickel dose, a unit of neutron fluence used in the United Kingdom and Europe.) 4.10.6.3.2 The UK creep model
The UK model1,64,65 recognizes that the initial creep coefficient is modified by irradiation-induced structure changes (i.e., changes to the pore structure). Hence, the total creep strain is given by ðg s dec s S 1 ðgÞdg ec ¼ þ dg 0 E0
½15
0
where s is the applied stress; (dec/dg)0, the initial secondary creep rate; g, the fast neutron fluence; S(g), the structure factor, given by S(g) ¼ Eg/Ep the ratio of the Young’s modulus at dose g to the Young’s modulus after the initial increase due to dislocation pinning. The structure factor, S(g), thus attempts to separate those effects due to dislocation pinning occurring within the crystallites and structural effects occurring ex-crystal through changes in the Young’s modulus. However, the effect of creep strain (tensile or compressive) on modulus is not considered when evaluating the structure term. The unstressed Young’s modulus changes are used to establish the magnitude of S(g). 4.10.6.3.3 The Kennedy model
Kennedy et al.66 replaced the structure term in the UK model with a parameter based on the volume change behavior of the graphite: s ½16 ec ¼ þ k0 ðgÞsg E0 where 0
ðg
k ðgÞ ¼ k0 0
1m
DV =V0 ðDV =V0 Þmax
dg
½17
318
Radiation Effects in Graphite
Here, m is an empirical constant equal to 0.75 and k0 is the steady-state creep coefficient established from low dose creep experiments. Although the Kennedy et al.66 model was shown to perform well in the prediction of high-dose tensile creep data, it did not predict the compressive data nearly as well. Moreover, as with the UK model, the sign of the applied stress is not considered when evaluating the influence of structure change (as reflected in volume changes). The quotient in eqn [17] is evaluated solely from unstressed (stressfree) samples irradiation behavior. As discussed by Kelly and Burchell,51 the term (DV/Vmax) does not exist at low irradiation temperatures where graphites expand in volume. 4.10.6.3.4 The Kelly and Burchell model
The Kelly and Burchell50,51 model recognizes that creep produces significant modifications to the dimensional change component of the stressed specimen compared to that of the control and that this must be accounted for in the correct evaluation of creep strain data. The rate of change of dimensions with respect to neutron dose g(n cm2) in appropriate units is given by the Simmons’ theory29 for direction x in the unstressed polycrystalline graphite: ax aa dXT 1 dXa þ þ Fx ½18 gx ¼ ac aa dg Xa dg where ax is the thermal expansion coefficient in the x-direction, and ac and aa are the thermal expansion coefficients of the graphite crystal parallel and perpendicular to the hexagonal axis, respectively, over the same temperature range. The term Fx is a pore generation term that becomes significant at intermediate doses when incompatibilities of irradiationinduced crystal strains cause cracking of the bulk graphite.67 For the purposes of their analysis, Kelly and Burchell ignored the term Fx . The parameters (1/Xc)(dXc/dg) and (1/Xa)(dXa/dg) are the rates of change of graphite crystallite dimensions parallel and perpendicular to the hexagonal axis, and dXT 1 dXc 1 dXa ¼ dg Xc dg Xa dg
½19
The imposition of a creep strain is known to change the thermal expansion coefficient of a stressed specimen, increasing it for a compressive strain and decreasing it for a tensile strain compared to an unstressed control. Thus, the dimensional change
component of a stressed specimen at dose g(n cm2) is given by 0 ax aa dXT 1 dXa 0 þ þ Fx0 ½20 gx ¼ ac aa dg Xa dg where a0x is the thermal expansion coefficient of the crept sample, and Fx0 is the pore generation term for the crept specimen. The difference between these two equations is thus the dimensional change correction that should be applied to the apparent creep strain (the pore generation terms Fx and Fx0 were neglected): 0 a aa dXT gx0 gx ¼ x ac a a dg ax aa dXT ac aa dg 0 a ax dXT ¼ x ½21 ac aa dg The true creep strain rate can now be expressed as 0 de de0 a ax dXT x ½22 ¼ ac aa dg dg dg where e is the true creep strain and e0 is the apparent creep strain determined experimentally in the conventional manner. Thus, the true creep strain (ec) parallel to the applied creep stress is given by ðg 0 ax ax dXT 0 dg ½23 ec ¼ ec ac aa dg 0
where e0c is the induced apparent creep strain, ða0x ax Þ is the change in CTE as a function of dose, ðac aa Þ is the difference of the crystal thermal expansion coefficients of the graphite parallel and perpendicular to the hexagonal axis, XT is the crystal shape change parameter given above, and g is the neutron dose. The apparent (experimental) creep strain is thus given by ðg 0 ax ax dXT 0 dg ½24 ec ¼ ec þ ac aa dg 0
Substituting for ec from eqn [13] gives the apparent (experimental) creep strain e0c as ðg 0 s ax ax dXT 0 dg ½25 þ ksg þ ec ¼ ac aa dg E0 0
with the terms as defined above. The Kelly–Burchell model is unique in that it does take account of the sign of the applied stress in
Radiation Effects in Graphite
predicting creep strain through changes in the CTE of the stressed graphite. While the model gave good agreement between the predicted H-451 graphite apparent creep strain and the experimental data at low doses and high temperatures51 (Figures 19–22), the creep model was shown to be inadequate at doses >0.5 1022 n cm2 [E >50 keV] (3.4 dpa) at an irradiation temperature of 900 C (Figure 23).50 4.10.6.3.5 The M2 model
Based upon the evidence from UK and US creep experiments, Davies and Bradford49,68 suggest the following: The strain induced change in CTE is not a function of secondary creep strain, but saturates after a dose of 30 1020 n cm2 EDN (3.9 dpa). There is evidence, from both thermal and irradiation annealing, for a recoverable strain several
times that of primary creep, and a lower associated secondary creep coefficient that has been previously assumed. The dose at which the recoverable strain saturates bears a striking similarity to that of the saturation of the CTE change. Davies and Bradford49,68 proposed a new creep model (the M2 model) without the term reflecting changes in CTE due to creep, but containing one additional term, recoverable creep: gð1 lk1 s k1 g1 expk1 g dg ec ¼ exp E0 SW
0
gð1 gð1 xk2 s b s k2 g1 k2 g exp dg þ dg exp þ E0 SW E0 SW ½26 0 0
Creep strain (%)
0.5 Experimental creep strain
0
True creep strain
-0.5
CTE correction strain
-1
Predicted apparent creep strain
-1.5 -2 0
0.1
0.2
0.3
0.4
0.5
0.6
Neutron dose 1022 n cm-2 [E > 50 keV] Figure 19 Comparison of predicted apparent creep strain (from eqn [25]) and the experimental creep strain data for irradiation creep at 600 C under a compressive stress of 13.8 MPa. The true creep strain is calculated from eqn [13]. From Burchell, T. D. J. Nucl. Mater. 2008, 381, 46–54.
1
Creep strain (%)
0.5
Experimental creep strain
0
True creep strain
-0.5 -1
CTE correction strain
-1.5
Predicted apparent creep strain
-2 -2.5 -3 0
319
0.1 0.2 0.3 0.4 0.5 Neutron dose 1022 n cm-2 [E > 50 keV]
0.6
Figure 20 Comparison of predicted apparent creep strain (from eqn [25]) and the experimental creep strain data for irradiation creep at 600 C under a compressive stress of 20.7 MPa. The true creep strain is calculated from eqn [13]. From Burchell, T. D. J. Nucl. Mater. 2008, 381, 46–54.
Radiation Effects in Graphite
Creep strain (%)
320
1.5 1 0.5 0 -0.5 -1 -1.5 -2 -2.5 -3 -3.5
Experimental (apparent) creep True creep strain CTE change correction Predicted apparent creep strain
0
0.1 0.2 0.3 0.4 0.5 Neutron dose 1022 n cm-2 [E > 50 keV]
0.6
Figure 21 Comparison of predicted apparent creep strain (from eqn [25]) and the experimental creep strain data for irradiation creep at 900 C under a compressive stress of 13.8 MPa. The true creep strain is calculated from eqn [13]. From Burchell, T. D. J. Nucl. Mater. 2008, 381, 46–54.
2
Creep strain (%)
1 Experimental creep strain
0
True creep strain
-1
Dimensional change correction
-2 -3
Predicted apparent creep strain
-4 -5 0
0.1
0.2
0.3 22
Neutron dose 10
n cm-2
0.4
0.5
0.6
[E > 50 KeV]
Figure 22 Comparison of predicted apparent creep strain (from eqn [25]) and the experimental creep strain data for irradiation creep at 900 C under a compressive stress of 20.7 MPa. The true creep strain is calculated from eqn [13]. From Burchell, T. D. J. Nucl. Mater. 2008, 381, 46–54.
5.0 Apparent (experimental) creep
Creep strain (%)
4.0 3.0
True creep
2.0
Dimensional change correction
1.0 0.0
Predicted apparent creep
-1.0 -2.0
0
0.5 1 1.5 Neutron dose, 1022 n cm-2 [E > 50 KeV]
2
Figure 23 Comparison of predicted apparent creep strain (from eqn [26]) and the experimental creep strain data for irradiation creep at 900 C under a tensile stress of 6 MPa. The true creep strain is calculated from eqn [13]. From Burchell, T. D. J. Nucl. Mater. 2008, 381, 46–54.
Radiation Effects in Graphite
where ec is the total creep strain; s, the applied stress; l, x, and b, are the empirical fitting parameters; k1 and k2, the primary and recoverable dose constants respectively; and W is the oxidation change factor (with respect to Young’s modulus) and is analogous to the structure factor. The terms in the eqn [26] are proportional to esu and the effects of structural changes and radiolytic oxidation (gasification of graphite by an activated species that occurs in CO2 cooled reactors) are also included. The rates of saturation of the primary and recoverable creep components are controlled by the dose constants k1 and k2. The first and last terms in eqn [26] are primary and secondary creep as in the prior UK creep model, with the middle term being recoverable creep. Primary creep is still fast acting, but in the AGR temperature range of 400–650 C, appears to act on a longer fluence scale equivalent to that associated in the United Kingdom with the Young’s modulus pinning,69 k1 ¼ 0.1, and saturates at 1 esu (a ¼ 1). The irrecoverable creep is synonymous with secondary creep, but with a coefficient, b, derived from the irrecoverable strain postthermal anneal, as 0.15 per 1020 n cm2 EDN (1.3 dpa) in the AGR temperature range. The lateral creep strain ratios for primary and recoverable creep are assumed to be the Poisson’s ratio and secondary creep is assumed to occur at constant volume. Figure 24 shows the performance of the M2 models applied to some high dose ATR-2E tensile creep data52 when irradiated at 500 C in high flux reactor (HFR), Petten. The prediction matches the observed data well up to significant fluence of 160 1020 n cm2 EDN (21 dpa). Only beyond
321
this fluence does the new model prediction deviate from the data with a delay in the increase in creep strain at high doses that is often referred to as the ‘tertiary’ creep phase. Figure 25 shows the corresponding compressive creep data,52 irradiated at 550 C. The model over predicts the data slightly but follows the trend remarkably well up to a significant fluence of 160 1020 n cm2 EDN (21 dpa). Beyond this fluence, the compressive prediction also indicates a ‘tertiary’ creep, but the data does not extend into this region. The data52 also indicates a possible difference between tensile and compressive creep (seen more clearly in Figure 18). Saturation of CTE with creep strain as reported by Davies and Bradford49,68 is not however in agreement with other published data. Gray70 reported CTE behavior with creep strain (up to 3%) for three different graphites at irradiation temperatures of 550 and 800 C. Saturation of the CTE in the manner described by Davies and Bradford49,68 for UK AGR graphite was not observed. 4.10.6.4 Deficiencies in Current Creep Models at High Neutron Doses The poor performance of the Kelly and Burchell model (eqn [25]) at predicting the high temperature (900 C) and high dose 6 MPa tensile creep data suggests that the model requires further revision.50,71 H-451 graphite irradiated at 900 C goes through dimensional change turn-around in the dose range 1.3–1.5 1022 n cm2 [E >50] (8.8–10.2 dpa). This behavior is understood to be associated with the 0.02
0.025 M2 model 500 T
0.015 Creep strain
Creep strain
0.02
Model 550 ⬚C 550 ⬚C
0.015 0.01
0.01
0.005
0.005 0
0 0
50
100 150 EDND (1020 n cm-2)
200
250
Figure 24 Comparison of the M2 models prediction and experimental creep strain data for ATR-2E tensile creep data, when irradiation was at 500 C in Petten. Reproduced from Davies, M.; Bradford, M. J. Nucl. Mater. 2008, 381, 39–45.
0
50
100 150 Dose (1020 n cm-2 EDN)
200
250
Figure 25 Comparison of the M2 models prediction and experimental creep strain data for ATR-2E compressive creep data, when irradiation was at 550 C in Petten. Reproduced from Davies, M.; Bradford, M. J. Nucl. Mater. 2008, 381, 39–45.
322
Radiation Effects in Graphite
generation of new porosity due to the increasing mismatch of crystal strains. The Kelly–Burchell model accounts for this new porosity only to the extent to which it affects the CTE of the graphite, through changes in the aligned porosity. Gray70 observed that at 550 C the creep rate was approximately linear. However, at 800 C he reported a marked nonlinearity in the creep rate and the changes in CTE were significant. Indeed, for the two high density graphites (H-327 and AXF-8Q) Gray reports that the 900 C creep strain rate reverses. Gray postulated a creep strain limit to explain this behavior, such that a back stress would develop and cause the creep rate to reduce. Other workers have shown that a back stress does not develop.62 However, Gray further argued that a creep strain limit is improbable as this cannot explain the observed reversal of creep strain rate. Note that a reversal of the creep rate is clearly seen in the 900 C tensile creep strain data reported here for H-451 (Figure 23). Also, a creep strain limit would require that tensile stress would modify the onset of pore generation behavior in the same way as compressive stress, because the direction of the external stress should be immaterial.70 More recent data52 and the behavior reported by Burchell71 show that this is not the case. Gray70 suggests that a more plausible explanation of his creep data is the onset of rapid expansion accelerated by creep strain; that is, net pore generation begins earlier under the influence of a tensile applied stress. Indeed, it has been observed52 that compressive creep appears to delay the turnaround behavior and tensile creep accelerates the turnaround behavior (Figure 18). In discussing possible explanations for his creep strain and CTE observations, Gray70 noted that changes in the graphite pore structure that manifested themselves in changes in CTE did not appear to influence the creep strain at higher doses. The classical explanation of the changes in CTE invokes the closure of aligned porosity in the graphite crystallites. Further crystallite strain can be accommodated only by fracture. A result of this fracture is net generation of porosity resulting in a bulk expansion of the graphite. A requirement of this model is that the CTE should increase monotonically from the start of irradiation. A more marked increase in CTE would be seen when the graphite enters the expansion phase (i.e., all accommodating porosity filled). The observed CTE behavior, reported previously50 and in Gray’s70 work, does not display this second increase in CTE; thus, the depletion of (aligned)
accommodation porosity is not responsible for the early beginning of expansion behavior. The observation by Gray70 and Kennedy63 that creep occurs at near constant volume (up to moderate fluence) indicates that creep is not accompanied by a net reduction of porosity compared to unstressed graphite, but this does not preclude that stress may decrease pore dimensions in the direction of the applied stress and increase them in the other, that is, a reorientation of the pore structure. Pore reorientation could effectively occur as the result of a mechanism of pore generation where an increasing fraction of the new pores are not well-aligned with the crystallites basal planes (and thus they would not manifest themselves in the CTE behavior) or accompanied with the closure of pores aligned with the basal planes. Kelly and Foreman53 report that their proposed creep mechanism would be expected to break down at high doses and temperatures, and thus deviations from the linear creep law (eqn [12]) are expected. They suggest that this is due to (1) incompatibility of crystal strains causing additional internal stress and an increasing crystal creep rate, (2) destruction of interstitial pins by diffusion of vacancies (thermal annealing of vacancies in addition to irradiation annealing), and (3) pore generation due to incompatibility of crystal strains. It is likely that pore generation can manifest itself in two ways: (1) changes in CTE with creep strain – thus, pores aligned parallel to the crystallite basal planes are affected by creep strain – and (2) at high doses, pore generation or perhaps pore reorientation, under the influence of applied and internal stress that must be accounted for in the prediction of high neutron dose creep behavior. Brocklehurst and Brown62 report on the annealing behavior of specimens that had been subjected to irradiation under constant stress and sustained up to 1% creep strain. They observed that the increase in creep strain with dose was identical in compression and tension up to 1% creep strain, and that the CTE was significantly affected in opposite directions by compressive and tensile creep strains. Irradiation annealing of the crept specimens caused only a small recovery in the creep strain, and therefore provided no evidence for a back stress in the creep process, which has implications for the in-crystal creep mechanism. Thermal annealing also produced a small recovery of the creep strain at temperatures below 1600 C, presumably because of the thermal removal of the irradiation-induced defects responsible for dislocation pinning. Higher temperature annealing
Radiation Effects in Graphite
produced a further substantial recovery of creep strain. Most significantly, Brocklehurst and Brown62 reported the complete annealing of the creep induced changes in CTE, in contrast to the total creep strain, where a large fraction of the total creep strain is irrecoverable and has no effect on the thermal expansion coefficient. Brocklehurst and Brown62 discuss two interpretations of their results, but report that neither is satisfactory. One interpretation requires a distinction between changes in porosity that affect the CTE and changes in porosity affecting the elastic deformation under external loads, that is, two distinct modes of pore structure changes due to creep in broad agreement with the mechanism discussed earlier. The modified Simmons model29,30,67 for dimensional changes (eqn [18]) and that for dimensional changes of a crept specimen (eqn [20]) both have pore generation terms which are currently neglected. It now appears necessary to modify the current Kelly–Burchell creep model (eqn [25]) to account for this effect of creep strain on this phenomena; that is, we need to evaluate and take account of the terms Fx and Fx0 as well as include the term (Fx0 –Fx) in eqn [25]. Such a term should account for pore generation and/or reorientation caused by fracture when incompatibilities in crystallite strains become excessive.71 Clearly, further work is needed in the area of irradiation-induced creep of graphite.
4.10.7 Outlook For more than 60 years, nuclear graphite behavior has been the subject of research and development in support of graphite-moderated reactor design and operations. The materials physics and chemistry, as well as the behavior of nuclear graphite under neutron irradiation are well characterized and understood, although new high-resolution characterization tools, such as HRTEM and STM, and other nanoscale characterization techniques, coupled with powerful computer based simulations of crystal deformation and displacement damage, are yielding new insights to the deformation mechanisms that occur in graphite throughout its life in the reactor core. Perhaps the biggest remaining challenge is to gain a fuller understanding of irradiation-induced dimensional change and irradiation creep in graphite. Currently, new creep irradiation experiments are underway at ORNL in the High Flux Isotope Reactor, and at Idaho National Laboratory in the Advanced Test Reactor. Studies of pore structure change from
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unirradiated reference samples, irradiated unstressed samples (controls), and irradiated stressed samples (crept samples), may advance our understanding of pore generation. Work in other countries is directed at reviewing existing creep data and assessing the observable graphite dimensional changes and creep strain in currently operating reactors. A recent Coordinated Research Project initiated by the International Atomic Energy Agency (IAEA) has the goal of bringing these various strands of research together to form a single unified theory of irradiation-induced creep deformation in graphites. The knowledge gained through these many years of work, and 50 years of graphite-moderated reactor operating experience is currently being used to underwrite the safety cases of graphite reactors through out the world. In the first part of the twenty-first century, more knowledge will be gained from the new graphitemoderated reactors in Japan and China that operate at higher temperatures. Several nations (within the Generation IV International Forum) are pursuing high-temperature, graphite-moderated, gas-cooled reactor projects with the goal of developing versatile and inherently safe reactor systems that can efficiently deliver both process heat and electricity. With the realization and acceptance that greenhouse gas emissions from fossil fueled power plants are causing global climate changes, as evidenced by the Kyoto and recent Copenhagen Agreements, the nuclear option may once again become attractive for clean electric power generation. At that time, it is to be hoped that inherently safe, graphite-moderated, gas-cooled reactors may find renewed popularity.
Acknowledgments This work is sponsored by the US Department of Energy, Office of Nuclear Energy Science and Technology under contact DE-AC05-00OR22725 with Oak Ridge National Laboratories managed by UT-Battelle, LLC.
References 1. 2. 3.
Burchell, T. D. In Carbon Materials for Advanced Technologies; Burchell, T. D., Ed.; Elsevier Science: Oxford, 1999; pp 429–484. Eatherly, W. P.; Piper, E. L. In Nuclear Graphite; Nightingale, R. E., Ed.; Academic Press: New York, 1962; pp 21–51. Ragan, S.; Marsh, H. J. Mater. Sci. 1983, 18, 3161–3176.
324 4. 5.
6.
7. 8.
9. 10. 11. 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.
Radiation Effects in Graphite Inagaki, M. In Graphite and Precursors; Delhae`s, P., Ed.; Gordon & Breach Science: The Netherlands, 2001; pp 179–198. ASTM D 7219. Standard Specification for Isotropic and Near-Isotropic Nuclear Graphites, Annual Book of Standards; ASTM International: West Conshohocken, PA, 2010; Vol. 05.05. ASTM D 7301. Standard Specification for Nuclear Graphite Suitable for Components Subjected to Low Neutron Irradiation Dose, Annual Book of Standards; ASTM International: West Conshohocken, PA, 2010; Vol. 05.05. Ruland, W. Chem. Phys. Carbon 1968, 4, 1–84. Kelly, B. T. In Materials Science and Technology: A Comprehensive Treatment: Vol. 10A – Nuclear Materials; Cahn, R. W., Haasen, P., Kramer, E. J., Eds.; Wiley-VCH: Weinheim, 1994; pp 365–417. Kroto, H. W.; Heath, J. R.; O’Brien, S. C.; Curl, R. F.; Smalley, R. E. Nature 1985, 318, 162–163. Iijima, S. Nature 1991, 354, 56–58. Banhart, F. Rep. Prog. Phys. 1999, 62, 1181–1221. Thrower, P. A.; Meyer, R. M. Phys. Status Solidi A 1978, 47, 11–37. Kelly, B. T. Physics of Graphite; Applied Science: London, 1981. Burchell, T. D. In Physical Processes of the Interaction of Fusion Plasmas with Solids; Hofer, W. O., Roth, J., Eds.; Academic Press: San Diego, CA, 1996; pp 341–384. Telling, R. H.; Heggie, M. I. Philos. Mag. 2007, 87, 4797–4846. Hehr, B. D.; Hawari, A. I.; Gillette, V. H. Nucl. Technol. 2007, 160, 251–256. Nakai, K.; Kinoshita, C.; Matsunaga, A. Ultramicroscopy 1991, 39, 361–368. Wallace, P. R. Solid State Commun. 1966, 4, 521–524. Jenkins, G. M. Chem. Phys. Carbon 1973, 11, 189–242. Urita, K.; Suenaga, K.; Sugai, T.; Shinohara, H.; Iijima, S. Phys. Rev. Lett. 2005, 94, 155502. Tanabe, T. Phys. Scripta 1996, T64, 7–16. Koike, J.; Pedraza, D. F. J. Mater. Res. 1994, 9, 1899–1907. Jenkins, G. M. Carbon 1969, 7, 9–14. Ouseph, P. J. Phys. Status Solidi A 1998, 169, 25–32. Amelinckx, S.; Dellavignette, P.; Heerschap, M. Chem. Phys. Carbon 1965, 1, 1–71. Engle, G. B.; Eatherly, W. P. High Temp. High Press. 1972, 4, 119–158. Price, R. J. Carbon 1974, 12, 159–169. Burchell, T. D.; Snead, L. L. J. Nucl. Mater. 2007, 371, 18–27. Simmons, J. W. H. Radiation Damage in Graphite; Pergamon: Oxford, 1965. Brocklehurst, J. E.; Kelly, B. T. Carbon 1993, 31, 155–178. Kelly, B. T.; Burchell, T. D. Carbon 1994, 32, 499–505. Nightingale, R. E. Nuclear Graphite; Academic Press: New York, 1962. Arnold, L. Windscale 1957, Anatomy of a Nuclear Accident; St. Martin Press: London, 1992. Rappeneau, J.; Taupin, J. L.; Grehier, J. Carbon 1966, 4, 115–124. Tucker, M. O.; Rose, A. P. G.; Burchell, T. D. Carbon 1986, 24, 581–602. Burchell, T. D. Carbon 1996, 34, 297–316. Burchell, T. D.; Eatherly, W. P. J. Nucl. Mater. 1991, 170–181, 205–208. Ishiyama, S.; Burchell, T. D.; Strizak, J. P.; Eto, M. J. Nucl. Mater. 1996, 230, 1–7. Burchell, T. D. In Graphite and Precursors; Delhae`s, P., Ed.; Gordon & Breach Science: The Netherlands, 2001; pp 87–109.
40. 41. 42. 43. 44.
45. 46. 47. 48. 49. 50. 51. 52.
53. 54. 55. 56.
57.
58. 59. 60. 61. 62. 63. 64. 65. 66. 67. 68. 69. 70. 71.
Burchell, T. D. MRS Bull. 1997, XXII, 29–35. Taylor, R.; Kelly, B. T.; Gilchrist, K. E. J. Phys. Chem. Solids 1969, 30, 2251–2267. Snead, L. L.; Burchell, T. D. J. Nucl. Mater. 1995, 224, 222–229. Bell, J. C.; Bridge, H.; Cottrell, A. H.; Greenough, G. B.; Reynold, W. N.; Simmons, J. W. H. Phil. Trans. R. Soc. London Ser. A 1962, 245, 361–395. ASTM C 781. Standard Practice for Testing Graphite and Boronated Graphite Materials for High-Temperature Gas-Cooled Nuclear Reactor Components, Annual Book of Standards; ASTM International: West Conshohocken, PA, 2010; Vol. 05.05. Hust, J. G. NBS Special Publication 260–89; U.S. Department of Commerce, National Bureau of Standards: Gaithersburg, MD, USA, 1984; p 59. Tsang, D. K. L.; Marsden, B. J. In Management of Ageing Processes in Graphite Reactor Cores; Neighbour, G., Ed.; RSC: Cambridge, 2007; pp 158–166. Li, H.; Fok, A. S. L.; Marsden, B. J. J. Nucl. Mater. 2006, 372, 164–170. Wang, H.; Yu, S. Nucl. Eng. Des. 2008, 238, 2256–2260. Davies, M.; Bradford, M. J. Nucl. Mater. 2008, 381, 39–45. Burchell, T. D. J. Nucl. Mater. 2008, 38, 146–154. Kelly, B. T.; Burchell, T. D. Carbon 1994, 32, 119–125. Haag, G. Properties of ATR-2E Graphite and Property Changes due to Fast Neutron Irradiation; Report No. Jul-4183; Published FZ-J, Germany, 2005; Available at http://juwel.fz-juelich.de. Kelly, B. T.; Foreman, A. J. E. Carbon 1974, 12, 151–158. Martin, D. G.; Henson, R. W. Philos. Mag. 1967, 9, 659–672. Martin, D. G.; Henson, R. W. Carbon 1967, 5, 313–314. Kelly, B. T.; Simmons, J. H. W.; Gittus, J. H.; Nettley, P. T. In Proceedings of the Third United Nations Conference on Peaceful Uses of Atomic Energy; IAEA: Vienna, 1972; Vol. 10, p 339. Heggie, M. I.; Davidson, C. R.; Haffenden, G. L.; Suarez-Martinez, I.; Campanera, J.-M.; Savini, G. In Proceedings of CARBON 2007; American Carbon Society: Oak Ridge, TN, 2007. Kelly, B. T.; Brocklehurst, J. E. J. Nucl. Mater. 1977, 65, 79–85. Jouquet, G.; Kleist, G.; Veringa, H. J. Nucl. Mater. 1977, 65, 86–95. Oku, T.; Eto, M.; Ishiyama, S. J. Nucl. Mater. 1990, 172, 77–84. Hansen, H.; Loelgen, R.; Cundy, M. J. Nucl. Mater. 1977, 65, 148–156. Brocklehurst, J. E.; Brown, R. G. Carbon 1969, 7, 487–497. Kennedy, C. R. In Conference 901178-1; ORNL: Oak Ridge, TN, 1990. Brocklehurst, J. E.; Kelly, B. T. Carbon 1993, 31, 155–178. Kelly, B. T. Carbon 1992, 30, 379–383. Kennedy, C. R.; Cundy, M.; Kliest, G. In Proceedings of CARBON’88, Newcastle upon Tyne, UK, July 1988; Institute of Physics: London, 1988; pp 443–445. Kelly, B. T.; Martin, W. H.; Nettley, P. T. Philos. Trans. R. Soc. 1966, A260, 51–71. Davies, M. A.; Bradford, M. R. In Management of Ageing Processes in Graphite Reactor Cores; Neighbour, G. B., Ed.; RSC: Cambridge, 2007; pp 100–107. Bradford, M. R.; Steer, A. G. J. Nucl. Mater. 2008, 381, 137–144. Gray, W. J. Carbon 1973, 11, 383–392. Burchell, T. D. Irradiation Induced Creep in Graphite at High Temperature and Dose – A Revised Model; ORNL/TM-2008/098; Oak Ridge National Laboratory: Oak Ridge, TN, Feb 2009.
4.11
Graphite in Gas-Cooled Reactors
B. J. Marsden and G. N. Hall The University of Manchester, Manchester, UK
ß 2012 Elsevier Ltd. All rights reserved.
4.11.1 4.11.2 4.11.2.1 4.11.2.2 4.11.2.3 4.11.2.4 4.11.2.5 4.11.3 4.11.3.1 4.11.4 4.11.5 4.11.5.1 4.11.5.2 4.11.5.2.1 4.11.5.3 4.11.5.4 4.11.5.4.1 4.11.5.5 4.11.5.6 4.11.5.7 4.11.5.8 4.11.5.9 4.11.6 4.11.6.1 4.11.7 4.11.7.1 4.11.7.2 4.11.7.2.1 4.11.7.3 4.11.7.4 4.11.7.5 4.11.7.6 4.11.7.7 4.11.8 4.11.9 4.11.9.1 4.11.9.2 4.11.10 4.11.11 4.11.11.1 4.11.11.2 4.11.11.3 4.11.11.4 4.11.11.5 4.11.11.6 4.11.12
Introduction Graphite Crystal Structures Graphite Crystal Atomic Structure and Properties Coefficient of Thermal Expansion Modulus Thermal Conductivity Microcracking (Mrozowski Cracks) Artificial Nuclear Graphite Microstructure/Property Relationships Graphite Core Fast Neutron Fluence, Energy Deposition, and Temperatures Dosimetry (Graphite Damage Dose or Fluence) Early Activation Measurements on Foils Reactor Design and Assessment Methodology: Fuel Burnup Calder effective dose Equivalent Nickel Flux Integrated Flux and Displacements per Atom DIDO equivalent flux Energy Above 0.18 MeV Equivalent Fission Flux (IAEA) Fluence Conversion Factors Irradiation Annealing and EDT Summary of Fast Neutron Dose (Fluence) Graphite ‘Energy Deposition’ (Nuclear Heating) The Use of Titanium for Installed Sample Holders Radiolytic Oxidation Introduction Ionizing Radiation Energy deposition Radiolytic Oxidation Mechanism Inhibition Internal Porosity Prediction of Weight Loss in Graphite Components Weight Loss Prediction in Inhibited Coolant Graphite Temperatures Variation of Fluence, Temperature, and Weight Loss in a Reactor Core Fuel End Effects Temperature and Weight Loss Distribution of Fluence Within an Individual Moderator Brick Fast Neutron Damage in Graphite Crystal Structures Stored Energy Crystal Dimensional Change Coefficient of Thermal Expansion Modulus Thermal Conductivity Raman Property Changes in Irradiated Polycrystalline Graphite
327 327 327 328 328 329 329 329 330 332 333 334 335 335 336 336 337 338 339 339 339 339 340 341 341 341 341 341 341 342 342 343 343 346 346 347 347 347 348 348 352 353 354 354 354 355 325
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4.11.13 4.11.14 4.11.14.1 4.11.14.2 4.11.14.3 4.11.14.4 4.11.15 4.11.15.1 4.11.15.2 4.11.15.3 4.11.15.4 4.11.16 4.11.16.1 4.11.16.2 4.11.16.3 4.11.16.4 4.11.17 4.11.17.1 4.11.17.2 4.11.17.3 4.11.17.4 4.11.17.5 4.11.17.6 4.11.18 4.11.19 4.11.20 4.11.20.1 4.11.20.2 4.11.20.3 4.11.20.4 4.11.20.4.1 4.11.20.4.2 4.11.20.5 4.11.20.6 4.11.20.6.1 4.11.20.6.2 4.11.20.6.3 4.11.20.6.4 4.11.20.7 4.11.21 References
Averaging Relationships Dimensional Change Pile Grade A Gilsocarbon Effect of Radiolytic Oxidation on Dimensional Change Dimensional Change Rate Coefficient of Thermal Expansion Pile Grade A Gilsocarbon Methodology for Converting Between Temperature Ranges Effect of Radiolytic Oxidation on CTE Thermal Conductivity Pile Grade A Gilsocarbon Thermal Conductivity Temperature Dependence of Irradiated Graphite Predicting the Thermal Conductivity of Irradiated Graphite for Reactor Core Assessments Young’s Modulus Relationship Between Static and Dynamic Young’s Modulus Pile Grade A Gilsocarbon Separation of Structure and Pinning Terms Effect of Radiolytic Weight Loss on Dimensional Change and Young’s Modulus Small Specimen Strength Effect of Radiolytic Oxidation on Thermal Conductivity, Young’s Modulus, and Strength The Use of the Product Rule Irradiation Creep in Nuclear Graphite Dimensional Change and Irradiation Creep Under Load Types of Irradiation Creep Experiments The UKAEA Creep Law Observed Changes to Other Properties Coefficient of thermal expansion Young’s modulus Lateral Changes Creep Models and Theories UKAEA creep law German and US creep model Further modifications to the UKAEA creep law: interaction strain Recent nuclear industry model Final Thoughts on Irradiation Creep Mechanisms Concluding Remark
Abbreviations AG AGR BAF BEPO CPV
Against grain Advanced gas-cooled reactor Bacon anisotropy factor British Experimental Pile Zero Closed pore volume
CTE DFR DSC DYM EDND EDNF
Coefficient of thermal expansion Dounreay Fast Reactor Differential scanning calorimeter Dynamic Young’s modulus Equivalent DIDO nickel dose Equivalent DIDO nickel flux
357 359 360 363 364 366 366 366 367 367 368 368 369 369 369 370 371 372 372 372 374 374 375 375 375 376 378 378 378 380 380 381 381 381 383 385 385 387 387 387 388
Graphite in Gas-Cooled Reactors
EDT FWHM HFR HOPG HRTEM
Equivalent DIDO temperature Full-width, half-maximum High Flux Reactor Highly oriented pyrolytic graphite High-resolution transmission electron microscopy HTR High temperature reactor IAEA International Atomic Energy Agency MTR Materials test reactor NDT Nondestructive testing OPV Open pore volume PGA Pile Grade A RBMK Reaktor Bolshoy Moshchnosti Kanalniy (there are other quoted translations) RPV Reactive pore volume SEM Scanning electron microscopy SYM Static Young’s modulus TEM Transmission electron microscopy TPV Total pore volume UKAEA United Kingdom Atomic Energy Authority WG With grain
327
This chapter aims to address that need by explaining the influence of microstructure on the properties of nuclear graphite and how irradiation-induced changes to that microstructure influence the behavior of graphite components in reactor. Nuclear graphite is manufactured from coke, usually a by-product of the oil or coal industry. (Some cokes are a by-product of refining naturally occurring pitch such as Gilsonite.9) Thus, nuclear graphite is a porous, polycrystalline, artificially produced material, the properties of which are dependent on the selection of raw materials and manufacturing route. In this chapter, the properties of the graphite crystal structures that make up the bulk polycrystalline graphite product are first described and then the various methods of manufacture and resultant properties of the many grades of artificial nuclear graphite are discussed. This is followed by a description of the irradiation damage to the crystal structure, and hence the polycrystalline structure, and the implication of graphite behavior. The influence of radiolytic oxidation on component behavior is also discussed as this is of interest to operators or designers of graphitemoderated, carbon dioxide-cooled reactors, many of which are still operating.
4.11.1 Introduction Nuclear graphite has, and still continues, to act as a major component in many reactor systems. In practice, nuclear graphite not only acts as a moderator but also provides major structural support which, in many cases, is expected to last the life of the reactor. The main texts on the topic were written in the 1960s and 1970s by Delle et al.,1 Nightingale,2 Reynolds,3 Simmons,4 in German, and Pacault5 Tome I and II, in French with more recent reviews on works by Kelly6,7 and Burchell.8 This text is mainly on the basis of the UK graphite reactor research and operating experience, but it draws on international research where necessary. During reactor operation, fast neutron irradiation, and in the case of carbon dioxide-cooled systems radiolytic oxidation, significantly changes the graphite component’s dimensions and properties. These changes lead to the generation of significant graphite component shrinkage and thermal stresses. Fortunately, graphite also exhibits ‘irradiation creep’ which acts to relieve these stresses ensuring, with the aid of good design practice, the structural integrity of the reactor graphite core for many years. In order to achieve the optimum core design, it is important that the engineer has a fundamental understanding of the influence of irradiation on graphite dimensional stability and material property changes.
4.11.2 Graphite Crystal Structures The properties and irradiation-induced changes in graphite crystals have been studied using both ‘naturally occurring’ graphite crystals and an artificial product referred to as highly orientated pyrolytic graphite (HOPG), formed by depositing a carbon substrate using hydrocarbon gas6 followed by compression annealing at around 3000 C. HOPG is considered to be the most appropriate ‘model’ material that can be used to study the behavior of artificially produced polycrystalline nuclear graphite. It has a density value near to that of a perfect graphite crystal structure, but perhaps more appropriately, it has imperfections similar to those found in the structures that make up artificial polycrystalline graphite. A detailed description of the properties of graphite can be found in Chapter 2.10, Graphite: Properties and Characteristics. 4.11.2.1 Graphite Crystal Atomic Structure and Properties In this section, the atomic structure of graphite crystal structures is discussed briefly, along with some of the properties relevant to the understanding of the
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Graphite in Gas-Cooled Reactors
4.11.2.3
irradiation behavior of graphite. Graphite can be arranged in an ABAB stacking arrangement termed hexagonal graphite (see Figure 1). This is the most thermodynamically stable form of graphite and has a density of 2.266 g cm3. The a-spacing is 1.415 A˚ and the c-spacing is 3.35 A˚. However, in both natural and artificial graphite stacking faults and dislocations abound.10
The crystal elastic moduli6 are C11 (parallel to the basal planes) ¼ 1060.0 109 N m2, C12 ¼ 180.0 109 N m2, C13 ¼ 15.0 109 N m2, C33 (perpendicular to the basal planes) ¼ 34.6 109 N m2, and C44 (shear of the basal planes) ¼ 4.5 109 N m2 as defined by the orthogonal co-ordinates given below: 0
4.11.2.2
Modulus
Coefficient of Thermal Expansion
sxx
0
1
C11 C12 C13 0
0
0
C11 C13 0
0
0
C13 C33 0
0
0
0 C44 0
0
B C B B syy C B C12 B C B B C B B szz C B C13 B C B Bt C ¼ B 0 B zx C B B C B Bt C B 0 @ zy A @
The coefficient of thermal expansion (CTE) as measured for natural graphite and HOPG is temperature dependent (Figure 2) and the data from a number of authors has been collated by Kelly.6 The room temperature values of CTE are about 27.5 106 K1 and 1.5 106 K1 in the ‘c’ and ‘a’ directions, respectively.
txy
0
0 0
0
0 C44
0
0
0
0
10 C C C C C C C C C C A
0 1ðC C Þ 12 2 11
exx
1
B C B eyy C B C B C B ezz C B C Be C B zx C B C Be C @ zy A exy ½1
c
c
a
a
Upper layer (A)
a
Lower layer (B)
2
40
1
30
CTE ac (10−6 K−1)
CTE aa (10−6 K−1)
Figure 1 The crystalline structure of graphite.
0 Bailey and Yates Steward et al. Harrison Yates et al.
−1
−2
20 Bailey and Yates Steward et al. Harrison Yates et al. Nelson and Riley
10
0 0
500
1000 1500 2000 Temperature (K) ‘c’ direction
2500
3000
0
500
1000 1500 2000 Temperature (K)
2500
‘a’ direction
Figure 2 Crystal coefficient of thermal expansion. Modified from Kelly, B. T. Physics of Graphite; Applied Science: London, 1981.
3000
Graphite in Gas-Cooled Reactors
The strength of the crystallite is also directly related to the modulus, that is, the strength along the basal planes is higher than the strength perpendicular to the planes, and the shear strength between the basal panes is relatively weak. 4.11.2.4
Thermal Conductivity
The thermal conductivity of graphite along the basal plane ‘a’ direction is much greater than the thermal conductivity in the direction perpendicular to the basal plane ‘c.’ At the temperature of interest to the nuclear reactor engineer, graphite thermal conduction is due to phonon transport. Increasing the temperature leads to phonon–phonon or Umklapp scattering (German for turn over/down). Imperfections in the lattice will lead to scattering at the boundaries.
329
width and many micrometers in length (as seen in Figure 3(b)), appears to be counterintuitive and has led to speculation that these microcracks may contain some low-density carbonaceous structure. The presence of these microcracks is very important in understanding the properties of nuclear graphite as they provide accommodation for thermal or irradiationinduced crystal expansion in the ‘c’ direction. Therefore, two crystal structures are of interest; the ideal, ‘perfect’ structure and the nonperfect structures as may be defined with reference to HOPG. It is of the latter that many of the crystal behaviors and properties have been studied. Definition: In this chapter on nuclear graphite, ‘crystal’ refers to the perfect crystal structure and ‘crystallite’ refers to the nonperfect crystal structures containing Mrozowski-type microcracks (and nanocracks).
4.11.2.5 Microcracking (Mrozowski Cracks)
4.11.3 Artificial Nuclear Graphite
During the manufacture of artificial graphite, very high temperatures (2800–3000 C) are required in the graphitization process. On cooling from these high temperatures, thermoplastic deformation is possible until a temperature of 1800 C is reached. Below this temperature, the large difference in thermal expansion coefficients between the ‘c’ and ‘a’ directions leads to the formation of long, thin microcracks parallel to the basal planes, often referred to as ‘Mrozowski’ cracks.11 These types of cracks are even observed in HOPG (Figure 3). The high density of HOPG when compared to the large number of microcracks, a few nanometers in
The reactor designer requires a high-density, very pure graphite, with a high scattering cross-section, a low absorption cross-section, and good thermal and mechanical properties, both in the unirradiated and irradiated condition. The purity is important to ensure not only a low absorption cross-section but also that during operation the radioactivity of the graphite remains as low as possible for waste disposal purposes. Artificial graphite is manufactured from coke obtained either from the petroleum or coal industry, or in some special cases (such as Gilsocarbon, a UK grade of graphite) from a ‘graphitizable’ coke derived
1 µm 1 µm
(a) (b)
Figure 3 Transmission electron microscopic images of highly orientated pyrolytic graphite. (a) View into the ‘basal’ plane, ‘c’ direction, of HOPG (reproduced from Kelly, B. T. MSc thesis, University of Cardiff, Cardiff, Wales, 1966) and (b) Mrozowski cracks in HOPG as seen along the ‘basal’ planes, ‘a’ direction. Courtesy of A. Jones, University of Manchester.
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Graphite in Gas-Cooled Reactors
from naturally occurring pitch deposits.9 The raw coke is first calcined to remove volatiles and then ground or crushed for uniformity, before being blended and mixed with a pitch binder. (Crushed ‘scrap’ artificial graphite may be added to help with heat removal during the subsequent baking. For nuclear graphite, this should be of the same grade as the final product.) This mixture is then formed into blocks using one of various techniques such as extrusion, pressing, hydrostatic molding, or vibration molding, to produce the so-called ‘green article.’ The ‘green’ blocks are then put into large ‘pit’ or ‘intermittent’ gas or oil-fired furnaces. The blocks are usually arranged in staggers, covered by a metallurgic coke, and baked at around 800 C in a cycle lasting about 1 month to produce carbon blocks. These carbon blocks may be used for various industrial purposes such as blast furnace liners; it has even been used for neutron shielding in some nuclear reactors. (Care must be taken as the carbon blocks are not as pure as graphite and may lead to waste disposal issues at the end of the reactor life.) To improve the properties of the graphite produced from the carbon block, the carbon block is often impregnated with a low-density pitch under vacuum in an autoclave. To facilitate the entry of the pitch into the body of the block, the block surface may be broken by grinding. After impregnation the blocks are then rebaked. This process of impregnation and rebaking may be repeated 2, 3, or 4 times. However, the improvement in the properties by this process is subject to diminishing rewards. The next process is graphitization at about 2800– 3000 C by passing a large electrical current at low voltage through the blocks either in an ‘Acheson furnace’ or using an ‘in-line furnace.’ In both cases, the blocks are covered by a metallurgical coke to prevent oxidation. This graphitization cycle may take about 1 month. If necessary, there may be a final purification step. This involves heating the graphite blocks to around 2400 C in a halogen gas atmosphere to remove impurities. The final product can then be machined into the many intricate components required in a nuclear reactor. For quality assurance purposes, during manufacture the blocks are numbered at an early stage and this number follows the block through the manufacturing process. This is clearly an expensive manufacturing process and therefore, at each stage, quality control is very important. Many samples will be taken from the blocks to ensure that the final batch (or heat) is of appropriate quality compared to previous heats. It is important that the reactor operators
retain this data in electronic form as it may be required to investigate any anomalous behavior as the reactor ages. Samples of ‘virgin’ unirradiated graphite blocks should also be retained for future reference. Records should include information on the batch or heat, property measurements, nondestructive testing (NDT) results, and measurements of impurities. It is not enough just to have the ‘ash’ content after incineration and the ‘boron equivalent’ as some impurities, such as nitrogen, chlorine, and cobalt, will cause significant issues related to reactor operation and final waste disposal. It is important that the reactor operator takes responsibility for these measurements as in the past it has been found that reactor designers and graphite manufacturers close down or merge, and records are lost. Final inspection will uncover issues related to damage, imperfection, quality, etc. Therefore, a ‘concessions’ policy is required to determine what is acceptable and where such components can be used in reactor. Again, the reactor operator will require an electronic record of these concessions. 4.11.3.1 Microstructure/Property Relationships The microstructure of a typical nuclear graphite is described with reference to Gilsocarbon. This product was manufactured from coke obtained from a naturally occurring pitch found at Bonanza in Utah in the United States. To understand the microstructural properties, one has to start with the raw coke. The structure of Gilsonite coke is made of spherical particles about 1 mm in diameter as shown in Figure 4. This structure is retained throughout manufacture and into the final product. In Figure 4(b), the spherical shaped cracks following the contours of the spherical particles are clearly visible. This coke will be carefully crushed in order to keep the spherical structures that form the filler particles and help to give Gilsocarbon its (semi-) isotropic properties. At a larger magnification in a scanning electron microscopy (SEM), the complexity of these cracks is clearly visible, Figure 4(c), and at an even larger magnification, a ‘swirling structure’ made up of graphite platelets stacked together is discernable between the cracks. In essence, the whole structure contains a significant amount of porosity. After graphitization, the Gilsonite coke filler particles are still recognizable (Figure 5(a) and 5(b)). From the polarizing colors, one can see that the main ‘a’ axis orientation of the crystallites follows the
Graphite in Gas-Cooled Reactors
(a)
331
(b)
(c)
(d)
Figure 4 Gilsonite raw-coke microstructure. (a) Photograph of Gilsonite coke, (b) Scanning electron microscopy (SEM) image of polished Gilsonite coke, (c) detail in an SEM image showing the region around cracks that follow the spherical shape of the coke particles, and (d) a higher magnification SEM image showing the intricate, random arrangement of platelets. Courtesy of W. Bodel, University of Manchester.
(a)
(c)
500 µm
(b)
200 µm
(d)
Figure 5 Polarized optical and scanning electron microscopic images of Gilsocarbon graphite. (a) Optical image, (b) optical image, (c) SEM image, (d) SEM image. Courtesy of A. Jones, University of Manchester.
332
Graphite in Gas-Cooled Reactors
spherical particles circumferentially, as does the orientation of the large calcination cracks. The crystallite structures in the binder phase are much more randomly oriented, and this phase contains significant amounts of gas-generated porosity. There are also what appear to be broken pieces of Gilsonite filler particles contained within the binder phase. The bulk properties of polycrystalline nuclear graphite strongly depend on the structure, distribution, and orientation of the filler particles.12 The spherical Gilsonite particles and molding technique give Gilsocarbon graphite semi-isotropic properties, whereas in the case of graphite grades such as the UK pile grade A (PGA), the extrusion process used during manufacture tends to align the ‘needle’ type coke particles. Thus, the crystallite basal planes that make up the filler particles tend to align preferentially, with the ‘c’ axis parallel to the extrusion direction and the ‘a’ axis perpendicular to the extrusion direction. The long microcracks are also aligned in the extrusion direction. The terms ‘with grain (WG)’ and ‘against grain (AG)’ are used to describe this phenomenon, that is, WG is equivalent to the parallel direction and AG is equivalent to the perpendicular direction. Thus, the highly anisotropic properties of the crystallite are reflected in the bulk properties of polycrystalline graphite (Table 1). A graphite anisotropy ratio is usually defined by the AG/WG ratio of CTE values. For needle coke graphite, this ratio can be two or more, while for a more randomly orientated structure, values in the region of 1.05 can be achieved by careful selection of material and extrusion settings. A more scientific way of defining anisotropy ratio is by use of the Bacon anisotropy factor (BAF).13 Other forming methods are usually used to produce isotropic graphite grades such as the Gilsocarbon grade described above. In this case, it was found that Gilsocarbon graphite produced by extrusion was not isotropic enough to meet the advanced gascooled reactor (AGR) specifications. Therefore, a Table 1
Relative properties–grain direction relationships
Property Coefficient of thermal expansion (CTE) Young’s modulus Strength Thermal conductivity Electrical resistivity
With grain (WG)
Against grain (AG)
Lower
Higher
Higher Higher Higher Lower
Lower Lower Lower Higher
‘molding’ method where the blocks were formed by pressing in two directions was used. This had the effect of slightly aligning the grains such that the AG direction was parallel to the pressing direction and the WG was perpendicular to the pressing direction. However, Gilsocarbon has proved to be one of the most isotropic graphite grades ever produced, even in its irradiated condition. Another approach is to choose an ‘isotropic coke’ crushed into fine particles and then produce blocks using ‘isostatic molding’ process. The isostatic molding method involves loading the fine-grained coke binder mixture into a rubber bag which is then put under pressure in a water bath. In this way, high quality graphite can be produced mainly for use for specialist industries such as the production of electronic components. This type of graphite (such as IG110 and IG-11) has been used for high-temperature reactor (HTR) moderator blocks, fuel matrix, and reflector blocks in both Japan and China. However, even these grades exhibit slight anisotropy. The final polycrystalline product contains many long ‘thin’ (and not so ‘thin’) microcracks within the crystallite structures that make up the coke particles. Similar, but much smaller, cracked structures are to be found in the ‘crushed filler flour’ used in the binder, and in well-graphitized parts of the binder itself. It is these microcracks that are responsible for the excellent thermal shock resistance of artificial polycrystalline graphite. They also provide ‘accommodation,’ which further modifies the response of bulk properties to the crystal behavior in both the unirradiated and irradiated polycrystalline graphite. Typical properties of several nuclear graphite grades are given in Table 2. One can see that polycrystalline graphite has about 20% porosity by comparing the bulk density with the theoretical density for graphite crystals (2.26 g cm3). About 10% of this is open porosity, the other 10% being closed.
4.11.4 Graphite Core Fast Neutron Fluence, Energy Deposition, and Temperatures Since the late 1940s, many journal papers, conference papers, and reports have been published on the change in properties in graphite due to fast neutron damage. Many different units have been used to define graphite damage dose (or fluence). It is important to understand the basis of these units because historic data are still being used to justify models
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Graphite in Gas-Cooled Reactors
Table 2
Typical properties of several well-known grades of nuclear graphite
Property
PGA
CSF
Gilsocarbon
IG-110
H451
Production method Direction Density (g cm3) Thermal conductivity (W m1 K) CTE, 20–120 C (106 K1) CTE, 350–450 C (106 K1) CTE, 500 C (106 K1) Young’s modulus (GPa) Poisson’s ratio Strength, tensile (MPa) Strength, flexural (MPa) Strength, compressive (MPa)
Extruded WG AG 1.74 200 109
Extruded WG AG 1.66 155 97
Press-molded WG AG 1.81 131
Iso-molded WG AG 1.77 116
Extruded WG AG 1.76 158 137
0.9
1.2
3.1
1.5 8.0
3.5 4.8
2.8
11.7 5.4 0.07 17 11 19 12 27 27
used in assessments for component behavior in reactors. Indeed, some of these historic data, for example, stored energy and strength, will also be used to support decommissioning safety assessments. Early estimations of ‘graphite damage’ were based on the activation of metallic foils such as cobalt, cadmium, and nickel. Later, to account for damage in different reactors, equivalent units, such as BEPO or DIDO equivalent dose, were used where the damage is referred to damage at a standard position in the BEPO, Calder Hall, or DIDO reactors. The designers of plutonium production reactors preferred to use a more practical unit related to fuel burnup (megawatts per adjacent tonne of uranium, MW/Atu). Researchers also found that the calculation of a flux unit, based on an integral of energies above a certain value, was relatively invariant to the reactor system and used the unit En > 0.18 MeV and other variants of this. Today, the favored option is to calculate the fluence using a reactor physics code to calculate the displacements per atom (dpa). However, in the field of nuclear graphite technology historic units are still widely used in the literature. For example, reactor operators have access to individual channel burnup which, with the aid of axial ‘form factors,’ can be used to give a measure of average damage along the individual channel length. Fortunately, most, but not all, of these units can be related by simple conversion factors. However, care must be taken; for example, the unit of megawatt days per tonne of uranium (MWd t1) is not necessarily equivalent in different reactor systems. When assessing the analysis of a particular component in a reactor, one must be aware that a single detailed calculation of a peak rated component in the
4.3 10.9 0.21 17.5 23.0 70.0
3.6
4.5
9.8 0.14 24.5 39.2 78.5
4.0
4.4 8.51 15.2 55.3
0.15
5.1 7.38 13.7 52.7
center of the core may have been carried out to give spatial, and maybe temporal, distribution of that component’s fluence (and possibly temperature and weight loss). These profiles may have then been extrapolated to all of the other components in the core using the core axial and radial ‘form factors.’ In doing this, some uncertainty will be introduced and clearly, some checks and balances will be required to check the validity of such an approach.
4.11.5 Dosimetry (Graphite Damage Dose or Fluence) In a nuclear reactor, high energy, fast neutron flux leads to the displacement of carbon atoms in the graphite crystallites via a ‘cascade.’ Many of these atoms will find vacant positions, while others will form small interstitial clusters that may diffuse to form larger clusters (loops in the case of graphite) depending upon the irradiation temperature. Conversely, vacancy loops will be formed causing the lattice structure to collapse. These vacancy loops will only become mobile at relatively high temperatures. The production of transmutation gas from impurities is not an issue for highly pure nuclear graphite, as the quantities of gas involved will be negligible and the graphitic structure is porous. The change in graphite properties is a function of the displacement of carbon atoms. The nature and amount of damage to graphite depends on the particular reactor flux spectrum, which is dependent on the reactor design and position, as illustrated in Figure 6. It is impractical to relate a spectrum of neutron energies to a dimensional or property change at a
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Graphite in Gas-Cooled Reactors
1800 TE rig in BEPO Hollow fuel element in BEPO Empty fuel channel in BEPO Empty lattice position in PLUTO
1600
Flux per unit lethargy f (v)
1400
Hollow fuel element in PLUTO 1200 1000 800 600 400 200 0 10
100
1000 Energy (keV)
10000
100000
Figure 6 Flux spectrums for various reactor positions used in graphite irradiation programs. Modified from Simmons, J. Radiation Damage in Graphite; Pergamon: London, 1965.
single point in a material such as graphite. Therefore, an ‘integrated flux’ is used and is discussed later.
Table 3 Relationship for BEPO equivalent flux (thermal) at a central lattice position to other positions in BEPO and other irradiation facilities
4.11.5.1 Early Activation Measurements on Foils
Position
Factor
BEPO lattice BEPO hollow slug BEPO empty fuel channel Windscale Piles Windscale Piles thermostats NRX fast neutron plug MWHs American data MWd/CT
1 2.27 0.63 1.29 1.29 1.54 1015 5.5 1017
Although one cannot directly measure the damage to graphite itself, it is possible to measure the activation of another material, because of nuclear impacts adjacent to the position of interest. This activation may then be related to changes in graphite properties. This was done in early experiments using cobalt foils and by measuring the activation arising from the 59Co(n,g)60Co reaction. This reaction has a cross-section of 38 barns and 60Co has a half-life of 5.72 years, which need to be accounted for in the fluence calculations. Such foils were included in graphite experiments in BEPO and the Windscale Piles, and are still used today for irradiation rig validation and calibration purposes. In these early experiments, after removal from the reactor, cobalt foils were dissolved in acid, diluted, and the decay rate measured. A measure of fluence could then be calculated from knowledge of the following:
the solution concentration the time in the reactor the decay rate the activation cross-section
Source: Simmons, J. H. W. The Effects of Irradiation on Graphite; AERE R R 1954; Atomic Energy Research Establishment, 1956.
Unfortunately the 59Co(n,g)60Co reaction is mainly a measure of thermal flux and atomic displacements in graphite are due to fast neutrons. An improvement was the use of cobalt/cadmium foils, but this was not really satisfactory. Measurements made in this way are often given the unit, neutron velocity time (nvt). Table 3 gives an example of thermal flux determined from cobalt foils defined at a standard position in the center of a lattice cell in BEPO. Graphite damage at other positions in other reactors could then be related to the standard position in BEPO.
Graphite in Gas-Cooled Reactors
4.11.5.2 Reactor Design and Assessment Methodology: Fuel Burnup When designing a nuclear reactor core, a channel ‘rating’ can be related to the reactor power and weight of uranium in a particular channel. This channel ‘rating’ can be related to a rate of change in the graphite properties. The channel rating is given in MW/Atu and over time the channel burnup as megawatt days per adjacent tonne of uranium (MWd/Atu). Note that this unit is literally the (power in a particular channel) (number of days) (weight of uranium in that channel). No account is taken of refueling. However, fuel burnup is a function of reactor design and therefore, the equivalence concept was used and damage was related to a standard position. In the United Kingdom, the change in graphite property was defined at a standard position in a Calder Hall reactor to give Calder equivalent dose. This was defined as the dose at a position on the wall of a fuel channel in a Calder Hall reactor. In the Calder Hall design, the lattice pitch is 8 in. The standard position was chosen to be in a 3.55-in.-diameter fuel channel at a point on the shortest line between the centers of two fuel channels. The fuel is assumed to be 1.15-in.-diameter natural uranium metal rods. Calder equivalent dose was then used as a function to relate graphite property change to fuel burnup. Kinchin14 had measured the change in graphite electric resistivity as a function of distance into the BEPO reflector. By normalizing this change, he defined a ‘graphite damage function’; see also Bell et al.15 Thus, graphite damage at some position in a reactor core graphite component could be defined as a function of the following: source strength distance between position and source attenuation in damage with distance through the intervening graphite The damage function is a measure of the last two bullet points. The source strength is related to fuel burnup. The graphite damage function df is defined as fðRg Þ df ¼ R
½2
where f(Rg) is the damage absorption curve for an equivalent distance through BEPO graphite ‘Rg’ of density 1.6 g cm3, and R is the distance through graphite between the source and position of interest. Note that nonattenuating geometric features, that is, holes, need to be accounted for.
335
Calder equivalent rating Pe can now be defined as Pe ¼
AdfP ACalder dfCalder
½3
where ‘ACalder’ and ‘A’ are the uranium fuel crosssectional areas in Calder (1.04 in.2) and in the reactor under consideration, respectively; ‘dfCalder’ and ‘df ’ are the values of the damage at the Calder standard position (1.395) and in the reactor under consideration, respectively, and ‘P’ is the fuel rating in the reactor under consideration. Thus, a graphite property change in a reactor under assessment can be related to the equivalent graphite property change at the Calder standard position. However, in a real reactor there is more than one fuel channel. There may also be absorbers or empty interstitial holes, the fuel rating will change with burnup, and the fuel will be replaced from time to time. Therefore, a more complex, multiple source calculation is required to take account of the actual channel rating and the geometric features of the core. This is normally done by considering a 5 5 lattice array: X Bi fðRg Þ i ½4 df ¼ B R i i where Bi and B are the accumulated fuel burnup at the ith and reference source, respectively, f(Rg)i is the damage absorption function corresponding to thickness Rg for ith source, and Rg is the distance between the ith source and target. This method was successfully used to design the Magnox reactors. However, because of the higher enriched oxide fuel and more complex fuel design in the AGRs, this approach became less satisfactory and new ‘damage functions’ that accounted for the new fuel and geometry were calculated using Monte Carlo methods (made possible by the introduction of the digital computer). This method was until recently still used in industry codes such as ‘Fairy’ (National Nuclear Company) and ‘GRAFDAM’ (UKAEA). 4.11.5.2.1 Calder effective dose
When only low-dose irradiation graphite property data were available, it was assumed that irradiation damage could be obtained at one temperature and that the property change versus dose (fluence) curves could be adjusted for all other temperatures using the so called R(y) curve: Calder effective calder equivalent ¼ RðyÞ ½5 dose dose
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Graphite in Gas-Cooled Reactors
However, the use of R(y) is valid only for very low fluence and it should no longer be used, although one may come across its use in historic papers. 4.11.5.3
Equivalent Nickel Flux
Nickel foils were used to give a measure of the damage to graphite through the 58Ni(n,p)58Co reaction. This reaction has a mean cross-section of 0.107 barns and 58Co has a half-life of 71.5 days. The change in graphite thermal resistivity was measured in the TE10 experimental hole in BEPO and the nickel flux was also measured at the same position. It was assumed that the graphite displacement rate fd was equal to the nickel flux fNi at this position. For comparison, the change in graphite thermal resistivity was then measured at various other positions in BEPO, as given in Table 4 Later, the same exercise was repeated in PLUTO, the sister reactor to DIDO at Harwell, and the ratio compared to that at other positions. In this case, the ratio appeared to be largely invariant to position. Table 5 gives a few examples of the many measurements made.16 It was decided that the activity produced in nickel fNi could be related to the graphite damage rate by a
Table 4 Ratio of graphite damage to nickel flux as measured in BEPO Position
Ratio bfd/fNi
Experimental hole TE10 Hollow fuel element Empty fuel channel (at three positions) Experimental hole E2/7
1.0 (definition) 0.43 1.0 0.75
Modified from Bell, J.; Bridge, H.; Cottrell, A.; Greenough, G.; Reynolds, W.; Simmons, J. Philos. Trans. R. Soc. Lond. A Math. Phys. Sci. 1962, 254(1043), 361–395. b is a proportionality factor.
Table 5 Ratio of graphite damage to nickel flux as measured in PLUTO Position
Ratio fd/fNi
C4 – inside fuel element stainless steel thimble D3 – inside fuel element stainless steel thimble C4 – inside fuel element aluminum thimble D4 – empty fuel element
0.518 0.468 0.507 0.564
Modified from Bell, J.; Bridge, H.; Cottrell, A.; Greenough, G.; Reynolds, W.; Simmons, J. Philos. Trans. R. Soc. Lond. A Math. Phys. Sci. 1962, 254(1043), 361–395.
factor. However, care was still required with respect to the choice of reactor and irradiation location. Thus, a definition of damage based on a standard position in DIDO and a calculation route for equivalent DIDO nickel flux (EDNF) were devised. It should be noted that there are difficulties related to a standard based on measurements made with nickel foils and the 58Ni(n,p,)58Co reaction because of the short half-life of 58Co and the interfering effect of the 58Co(n,g)59Co reaction. A method by Bell et al.15 which went back to measuring activation of cobalt foils and the 59Co(n,g)60Co reaction, and then calculating the ratio fNi/fCo, was used for a short while. This method used the following relationships: 115g fuel elements fNi =fCo ¼ 0:378 0:504b 150g fuel elements fNi =fCo ¼ 0:502 0:530b where ‘b’ is the fuel burnup. However, this was not very satisfactory and it was clear that a validated calculation route was desirable, and is now becoming practicable through development in computer technology. 4.11.5.4 Integrated Flux and Displacements per Atom The rate of change of a material property can be related to displacement rate of carbon atoms (dpa s1). However, it is not possible to directly measure dpa s1 in graphite, but dpa s1 can be related to the reactor flux. The flux depends on reactor design, and varies with position in the reactor core. Neutron flux is a measure of the neutron population and speed in a reactor. In a reactor, neutrons move at a variety of speeds in randomly orientated directions. Neutron flux is defined as the product of the number of neutrons per unit volume moving at a given speed, as given by eqn [6] below. number number cm ¼ n v ½6 f cm2 s cm3 s However, as there is a spectrum of neutrons, with many velocities, this is not a useful unit for the material scientist. Therefore, integrated flux is used over a range of energies E1 to E2 as given by eqn [7]. ð E2 fðEÞdE ½7 f¼ E1
In this way, a measure of neutron damage at any position within a structural component can be defined as follows. For a material such as graphite,
Graphite in Gas-Cooled Reactors
337
flux multiplied by the nickel cross-section at the standard position in DIDO, and s0 is the average nickel cross-section for energies >1 MeV, which is equal to 0.107 barn. The value of fs at this position is 4 1013 n cm2 s1. The carbon displacement rate can be calculated using eqn [10]. ðð fðE1 ÞsðE1 ; E2 ÞnðE2 ÞdE1 dE2 ½10 fd ¼
energy E1 to produce a recoil atom with energy E2, and v(E2) is a ‘damage function’ giving the number of atoms displaced from their lattice site by recoil energy E2. The carbon displacement rate, fds, at a standard position in DIDO is 5.25 108 dpa. The derivation of the damage function (Figure 7) is on the basis of billiard ball mechanics, energy losses to the lattice due to impacts, and to forces associated with excitation of the lattice. The early Kinchin and Pease17 form of the damage function was found to underestimate damage in graphite. To give greater dpa, it was recommended that ‘Lc’ was artificially increased, but this was not satisfactory. The Thompson and Wright18 damage function was used in the official definition of EDNF. However, the Norgett et al.19 damage function is used in most modern reactor physics codes and it has been recently shown that there is little difference in the calculation of graphite damage using either of these latter two functions.20,21 It is assumed that the ratio of dpa to nickel flux (fds/fs) at the standard position, which is equal to 1313 1024 dpa (n cm2 s1)1, can be equated to the same ratio fd/fNi in the reactor of interest as given by eqn [11]: fds fd ¼ ¼ 1313 1024 dpaðn cm2 s1 Þ1 fs fNi ½11
where f(E1) is the flux of neutrons with energy E1, s(E1, E2) is the cross-section for a neutron with
This value was derived using the Thompson and Wright damage function and an early flux spectrum
the damaging power (displacement rate), fd, can be expressed as an integrated flux as given in eqn [8]. 1 ð
fd ¼
cðEÞfðEÞdE
½8
0
where f(E) is the neutron flux with energies from E to E þ dE and C(E) is a function to describe the ability of neutrons to displace carbon atoms. 4.11.5.4.1 DIDO equivalent flux
At the standard position in a hollow fuel element, the nickel flux, fs, can be defined by eqn [9]. 1 fs ¼ s where
1 Ð
1 ð
fs ðEÞsNi ðEÞdE
½9
0
fs ðEÞsNi ðEÞdE is the integral of neutron
0
10 000
Number of displaced atoms
1000
Thompson-Wright Norgett, Robinson, and Torrens Kinchin-Pease (Lc = 25 keV) Kinchin-Pease (Lc = 12 keV)
100
10
1
0.1
10
102
103
104 105 Energy (eV)
106
107
108
Figure 7 Comparison of various damage function models that describe the number of displaced atoms versus energy of primary knock-on atom.
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Graphite in Gas-Cooled Reactors
Table 6 Comparison of calculated and measured graphite damage rates using the Thompson and Wright model Location
Calculated
Measured/ standard
DIDO hollow fuel element PLUTO empty lattice position DR-3 empty lattice position BR-2, Mol, hollow fuel element HFR-Petten core BEPO TE-10 hole BEPO empty fuel channel BEPO hollow fuel channel Windscale AGR replaced fuel stringer B Windscale AGR replaced fuel stringer D Windscale AGR loop stringer Windscale AGR loop control stringer Windscale AGR fuel element – inner ring Windscale AGR fuel element – outer ring Calder x-hole Dounreay fast reactor core
1.00 0.975 0.975 1.00 1.02 2.31 2.36 0.98 2.70
1.00 1.22 0.90 0.90 1.0 2.04 2.04 0.87 2.28
2.71
2.03
2.60 2.60
2.08 2.51
1.18
1.06
1.39
1.06
2.12 0.46
2.10 0.50
Table 7 Energies, cross-sections, and mean number of displacement for various particles Particles
Energy (eV) Cross-section Mean number of (cm2) displacements per collision
1 106 2 106 3 106 4 106 Protons 1 106 5 106 10 106 20 106 Deuterons 1 106 5 106 10 106 20 106 a-Particles 1 106 5 106 10 106 20 106 Neutrons 103 104 105 106 107 Electrons
14.5 1024 15.0 1024 15.5 1024 16.0 1024 7.8 1021 1.56 1021 7.8 1021 3.9 1021 1.56 1020 3.12 1021 1.6 1021 7.8 1022 1.25 1019 2.5 1020 1.25 1020 6.25 1021 4.7 1024 4.7 1024 4.6 1024 2.5 1024 1.4 1024
1.6 1.9 2.3 2.5 4–5.5 4–5.5 4–6 4–6 4–5 4–6 4–6 4–6.5 4–5 4–6 4–6.5 4–6.5 2.83 28.3 280 480 500
Modified from Marsden, B. J. Irradiation damage in graphite due to fast neutrons in fission and fusion systems; IAEA, IAEA TECDOC1154; 2000.
Source: Simmons, J. Radiation Damage in Graphite; Pergamon: London, 1965.
for the standard position in DIDO. Hence, the EDNF or fd can be calculated at the position of interest. The equivalent DIDO nickel dose (fluence) (EDND) is derived by integrating EDNF over time, as given in eqn [12]:
Table 8 Displacements (1021) per unit fluence for energies above E1 for various systems
ðt EDND ¼ fd ðt Þdt
½12
0
Table 6 compares the calculated and measured graphite damage rates in various systems using the Thompson and Wright model. Finally, for those wishing to try and reproduce damage in graphite using ion beams, Table 7 gives the energies, cross-sections, and mean number of displacement for various particles. 4.11.5.5
Energy Above 0.18 MeV
Dahl and Yoshikawa22 noted that for energies above 0.065 MeV, eqn [13] was reasonably independent of reactor spectrum under consideration:
Spectrum
E1 ¼ 0.067 MeV
E1 ¼ 0.18 MeV
PEGGI ETR(N-8) EBR-II DFR HFR Average
0.719 0.697 0.693 0.690 0.683 0.701 2.6%
0.738 0.810 0.769 0.790 0.779 0.774 4.5%
Source: Morgan, W. Nucl. Technol. 1974, 21, 50–56.
1 Ð
fðE>E1 Þ ¼
0
fðEÞsðEÞnðEÞdE 1 Ð E1
fðEÞdE
½13
Equation [13] is the integral of graphite displacement for a position in the particular reactor of interest, divided by the integral of flux from E1 (0.065 MeV in this case) to infinity at the same position. Table 8 gives this ratio for two other values of E1.
Graphite in Gas-Cooled Reactors
4.11.5.6
Equivalent Fission Flux (IAEA)
An IAEA committee recommended the use of equivalent fission flux23 as given by eqn [14]. 1 Ð P d
ðEÞfðE; t ÞdE
0 fG ¼ 1 1 Ð P Ð ðEÞwðEÞdE= wðEÞdE 0
d
½14
0
Equation [14] is essentially graphite dpa divided by a normalized fission flux. A similar unit is defined by Simmons4 in his book. However, the use of this unit was never taken up for general use. 4.11.5.7
Fluence Conversion Factors
Table 9 gives the conversion factor from other units to EDND. The following should be noted: EDND is a definition, Calder equivalent dose and other units relating damage to fuel ratings are approximate, BEPO equivalent dose is a thermal unit and should be avoided, Energies above En are a good approximation, dpa is directly proportional to EDND. 4.11.5.8
Irradiation Annealing and EDT
The reasoning behind the use of equivalent DIDO temperature (EDT) is that if two specimens are irradiated to the same fluence over two different time periods, the specimen irradiated faster will contain the most irradiation damage. The reasoning is that the specimen irradiated at the slower rate would have a longer time available to allow for ‘annealing’ out of defects caused by fast neutron damage as outlined below. The rate of accumulation of damage dC/dt can be described by eqn [15]. Conversion factors to EDND
Table 9 Fluence unit
2
EDND (n cm ) Equivalent fission dose (n cm2) Calder equivalent dose (MWd At1) BEPO equivalent dose (n cm2) En > 0.05 MeV (n cm2) En > 0.18 MeV (n cm2) En > 1.0 MeV (n cm2) dpa (atom/atom)
Conversion factor 1.0 0.547 1.0887 1017 0.123 0.5 0.67 0.9 7.6162 1020
Modified from Marsden, B. J. Irradiation damage in graphite due to fast neutrons in fission and fusion systems; IAEA, IAEA TECDOC-1154; 2000.
dC / dt
f E exp kT
339
½15
where f is the flux, E is the activation energy, T is temperature (K), and k is Boltzmann’s constant. Equating the damage rate for two identical samples at different flux levels f1 and f2 and different temperatures T1 and T2, f f 1 ¼ 2 E E exp exp kT1 kT2
½16
Rearranging this gives the EDT relationship: 1 1 k ðf Þ ¼ ln d 2 y1 y2 E ðfd Þ1
½17
The term on the left is the difference in the reciprocal of the temperatures in the two systems (temperature has traditionally been given the symbol ‘y ’ when used in this context) and the term on the right contains the natural log of the damage flux (or fluence) in the two systems divided by each other. In practice, the activation energy E is an empirical constant. The use of EDT has recently been investigated24 at temperatures above 300 C. The authors concluded that the use of EDT was inappropriate (Figure 8). However, below 300 C, there was some evidence of the applicability,15 but at these lower temperatures there is little reliable data. Therefore, the use of the EDT concept is not recommended for modern graphite moderated reactors where the graphite is usually irradiated above 300 C. 4.11.5.9 Summary of Fast Neutron Dose (Fluence) 1. Care must be taken when interpreting graphite data because of the variety of fast neutron dose units used. Older data in particular should be treated with care. 2. ‘Graphite damage’ has been equated to activation of nickel at a standard position in DIDO. This can now be calculated and equated to dpa. 3. ‘Graphite damage’ may also be equated to channel burnup which can also be equated to dpa. 4. ‘Graphite damage’ can also be equated to En > 0.18 MeV. 5. EDT is not applicable to irradiation temperatures above 300 C; there is some evidence that it may be applicable below 300 C.
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2
Dimensional change (%)
PLUTO DFR 1
0
−1
−2
−3
0
50
100
150
200
250
300
Fluence (1020 n cm−2 EDND) Figure 8 Comparison of dimensional changes on Gilsocarbon graphite samples irradiated in DFR with similar samples irradiated in PLUTO. Reproduced from Eason, E. D.; Hall, G.; Heys, G. B.; et al. J. Nucl. Mater. 2008, 381, 106–113.
6. There are conversion factors between all these units but these are subject to various degrees of uncertainty.
4.11.6 Graphite ‘Energy Deposition’ (Nuclear Heating) The heat generated in the graphite (or energy deposition) is required for the calculation of the graphite temperatures, and in the case of CO2-cooled systems, it is required for the calculations of radiolytic weight loss. Both of these requirements are important in graphite stress analysis calculations. In the case of CO2-cooled systems it is assumed that the graphite radiolytic oxidation rate is proportional to the heat generated in the graphite. However, it is ionizing irradiation that causes the dissociation of the CO2. The energy deposition is produced by the interaction of graphite atoms with three types of particles: Neutron interactions with graphite atoms (40%). Fission g-rays (60%). Secondary g-rays caused by absorption by materials outside the moderator (e.g., steel fuel pins in AGRs) and by inelastic scattering of carbon atoms (1% in a Magnox reactor and 10% in an AGR). The main source of gammas and neutrons arises from the fuel, mainly from prompt fission, but there are some from delayed fission. The ratios given above are for a central position in the core and for initial fuel loading. The ratio may change with position in the core and with graphite
weight loss. Furthermore, in graphite material test programs, the ratio between neutron and g-heating is likely to be significantly different, because of the different materials used to construct the various reactor cores. It is therefore important that this ratio is known and the implication of a change in this ratio on material property changes, that is, the implication of the ratio between fast neutron damage versus radiolytic weight loss on graphite property changes, is understood. The gamma and neutron spectrum varies with distance from the fuel and will vary with graphite density (i.e., will change with weight loss) and fuel design. A reactor is run at constant power, and therefore, as weight loss increases, the spectrum (gamma and neutron) will change and become harsher (higher neutron and g-flux). In the graphite, charged electrons are produced because of the following: 1. Compton scattering interaction of gamma with electrons within the carbon atoms. 2. Pair product in electrostatic field associated with carbon atoms. 3. Photoelectric absorption. Compton scattering predominates, but electrons and charged carbon ions are also produced because of the displacement of carbon atoms in the moderator, and in principle this could be calculated. Energy deposition is the energy released from the first collisions of primary gamma and neutrons.25 Energy deposition is calculated in watts per gram (W/g) and the spatial distribution can be calculated using reactor physics codes such as McBend
Graphite in Gas-Cooled Reactors
(http://www.sercoassurance.com/answers/), WIMS, and WGAM. However, a crude estimation of energy deposition can be made by assuming that 5% of the reactor power is generated in the graphite. This heat can then be proportioned to the rest of the core using interpolation and form factors, and estimates of the distribution within a moderator brick. In conclusion, energy deposition is required to calculate graphite temperatures and radiolytic oxidation rates. Energy deposition can be estimated but is most accurately calculated using reactor physics codes. However, care must be taken because the ratio between neutron heating and g-heating, or more appropriately a direct measure of the ionizing irradiation, is important. 4.11.6.1 The Use of Titanium for Installed Sample Holders During the construction of the Magnox and AGR reactors, graphite specimens were placed into ‘installed sample holders,’ the intention being that these samples could be removed at a later date to give information on the condition of the graphite core. To enhance the radiolytic weight loss of the graphite in the installed sample holders, titanium was used. Although this only slightly increased the g-heating, it did increase the number of electrons produced, because of an increase in pair production and Compton scattering caused by the higher atomic number or ‘Z-value’ of titanium compared to graphite (22 and 6, respectively).
4.11.7 Radiolytic Oxidation 4.11.7.1
Introduction
In carbon dioxide (CO2)-cooled reactors, two types of oxidation can occur. The first is thermal oxidation which is purely a chemical reaction between graphite and CO2. This reaction is endothermic and is negligible below about 625 C and is not important up to 675 C. The second is radiolytic oxidation that occurs when CO2 is decomposed by ionizing radiation (radiolysis) to form CO and an active oxidizing species, which attacks the graphite. Radiolytic oxidation occurs predominantly within the graphite open porosity. 4.11.7.2
Ionizing Radiation
Ionizing irradiation can be defined as that part of a radiation field capable of ionization (charge separation)
341
in CO2 either directly or indirectly. This leads to the creation of reactive species, which may react with the carbon atoms at the surfaces (external and more importantly internal) of the graphite components. 4.11.7.2.1 Energy deposition
Historically, ‘energy deposition’ has been used for a surrogate for ionizing irradiation, most probably because it is easy to measure using calorimetry and can be estimated from the reactor power. Energy deposition, sometimes referred as ‘dose rate,’ in the units of W/g of graphite, is a measure of the total energy absorbed in the gas in unit time from the scattering of g-radiation and fast neutrons. For a typical Magnox reactor, energy deposition is composed of approximately the following components: 36% from the neutrons 58% from the gamma 6% from the interaction of graphite atoms within the moderator Of these, it is only the last two that directly contribute to ionization of the carbon dioxide gas, mainly through Compton scattering. These ratios will be slightly different in an AGR. An assumption is made that the dose rate received by the graphite is the same as that absorbed by carbon dioxide within the pores of the graphite and that a fraction k of the fission energy from the fuel causes heating in the moderator. For a typical Magnox reactor, k is 5.6% of the thermal power. The unit GC is defined as the number of carbon atoms gasified by the oxidizing species produced by the absorption of 100 eV of energy in the CO2 contained within the graphite pores; GC for pure CO2 ¼ 3. 4.11.7.3
Radiolytic Oxidation Mechanism
The exact mechanism of radiolytic oxidation in a carbon dioxide-cooled reactor is complex and has been a matter of debate for some time; the most satisfactory explanation has been given by Best et al.26 However, in its most simplistic form the mechanism can be described as follows: In the gas phase, ionizing radiation CO2 ! CO þ O
½I
CO þ O ! CO2
½II
where O* is an activated state-oxidizing species. Thus, after ionization the carbon monoxide and oxidizing species rapidly recombine back into carbon dioxide and to an uninformed observer, carbon
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Isothermal radiolytic oxidation rate (10−8 h−1 W kg−1)
342
0.8 0.7 0.6
0.4–0.6% CO 1% CO 2% CO
0.5 0.4 0.3 0.2 0.1 0 0
100
200
300
400
500
600
700
800
Calculated CH4 concentration (vpm)
Figure 9 GC as function of CO and CH4 concentration (41 bar, 673 K).
O þ C ! CO
3
5 vpm H2O 15 to 30 vpm H2O 200 to 400 vpm H2O
2 G–c
dioxide would appear to be stable in an irradiation field. However, in the presence of graphite which typically contains 20% porosity, 10% of which is initially accessible to the carbon dioxide gas, at the graphite pore surface (mainly internal) carbon atoms are oxidized. This can be simplistically described as
1
½III
The principal oxidizing species is still under debate, but the most favored candidate is the negatively charged ion, CO 3.
0 0
1
(a)
2
3
CO (%)
1.5
Inhibition
The rate of oxidation can be reduced by the addition of carbon monoxide (CO) and moisture (H2O) and can be greatly reduced by the addition of methane (CH4), as illustrated in Figures 9 and 10. As described above the radiolytic oxidation process produces CO and if CH4 is added, moisture will be one of the by-products of the reaction.
1.0 G–c
4.11.7.4
0.5
0.0 0
4.11.7.5
Internal Porosity
As supplied, graphite components contain a significant amount of both open and closed porosity in a variety of shapes and sizes, from the nm scale to the mm scale, as illustrated in Section 4.11.3. The open pore volume (OPV) is defined as the volume of pores accessible to helium, closed pore volume (CPV) is the volume of pores not accessible to helium, and the total pore volume (TPV) is the volume of open and closed pore. The effect of pore size on the radiolytic oxidation rate was investigated by Labaton et al.27 who found
(b)
200
400
600
800
CH4(vpm)
Figure 10 Inhibition in carbon dioxide due to the addition of carbon monoxide, moisture, and methane. (a) GC as function of CO and H2O concentration and (b) GC as function of methane (CH4) concentration. Reproduced from Best, J. V.; Stephen, W.; Wickham, A. Prog. Nucl. Energy 1985, 16(2), 127–178.
the maximum range to be 2.5–5 mm. Taking this into account and referring to eqns [I]–[III] above, the oxidation process will be expected to be more efficient in the smaller pores than in the larger pores.
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This is because in the smaller pores the distance to the wall is less, making it less likely, compared to the case for larger pores, that the active species would be deactivated by collision in the gas phase. To account for this difference in oxidation rate with pore size for modeling purposes, in the case of the Magnox reactors which did not have CH4 routinely added to the coolant, a pragmatic approach of defining ‘pore efficiency’ was adopted, whereas in the case of the AGRs where CH4 is routinely added, a reactive pore volume (RPV) was defined as being the volume of pores oxidizing in CH4-inhibited coolant gas. It is also clear that as the oxidation process proceeds, closed porosity will be opened and the pore size distribution will change, thereby changing the oxidation rate. 4.11.7.6 Prediction of Weight Loss in Graphite Components The methodologies used to predict the oxidation rate in Magnox reactors are based on work by Standring28 as discussed below: P 273 Weight of CO2 in the ½18 ¼ er0 pores of 1g of graphite 14:7 T where P (psi) is the gas pressure, T (K) the temperature, and r0 is the density of CO2 at standard conditions for temperature and pressure (STP) (g cm3). The dose rate to the graphite can then be given in watts as follows: P 273 W ½19 Dose rate to graphite ¼ eDr0 14:7 T where D (W g1 s1) is the ‘energy deposition rate’ or ‘dose rate’ and e is the OPV (cm3 g1). This reasoning can be taken further to give eGC DP Percentage initial ¼ 145 % per year ½20 oxidation rate; g0 T Standring and Ashton29 measured the OPV and CPV in PGA as a function of weight loss (Figure 11). In the specimens they examined, there appeared to be a small amount of pores which opened rapidly before the pore volume increased linearly as a function of weight loss over the range of the data. To account for this behavior, they modified eqn [20] by defining an effective OPV as ‘ee’: g0 ¼ 145
ee GC Dp % per year T
½21
343
Standring further developed this reasoning into a relationship for the cumulative weight loss, Ct , at a constant dose rate: h g t i 0 Ct ¼ A exp 1 ½22 A e where A ¼ ð100P 1Pe Þ and Pe is the effective initial OPV 3 3 in cm cm . However, a reactor is operated at constant power. Replacing the dose rate in eqn [21] by kPt/Wm, where Pt is the reactor thermal power, k is the fraction of the reactor power absorbed in the graphite (5%), and Wm is the weight of the moderator, gives
g0 ¼ 145ee GC
kPt P % per year Wm T
½23
From eqn [23], it is clear that the rate of oxidation will increase with loss of moderator mass. It was shown by Standring that the cumulative weight loss, Ct, for a reactor operated at constant power is given by A2 Ct A Ct ¼ g0 t log 1 þ ½24 100Pe A 100 This equation yields higher weight loss than the constant dose rate equation. This approach was used to design the early Magnox stations. However, as higher weight loss data became available from the operating Magnox stations, it was found necessary to modify the relationship to account for the pore distribution with increasing oxidation. 4.11.7.7 Weight Loss Prediction in Inhibited Coolant It had not been possible to regularly add CH4 as an inhibitor to the coolant in the Magnox reactors because of concerns regarding the metallic components in the coolant circuit. However, the higher rated AGRs were designed with this in mind by selecting denser graphite and adding CH4 gas as an oxidation inhibitor. The addition of an inhibitor causes the process of radiolytic weight loss to be more complex than that for Magnox reactors as the oxidation rate becomes a complex function of the coolant gas composition. This is because gas composition, and hence, graphite oxidation rate, is not uniform within the moderator bricks and keys as CH4 is destroyed by radiolysis and may thus be depleted in the brick interior. In addition, methane destruction gives rise to the formation of carbon
344
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Open pore volume (cm−3 per 100 cm−3)
50 45 40 35 30 25 20 15 10 5 0 0
5
10
15 20 Weight loss (%)
25
30
35
0
5
10
15 20 Weight loss (%)
25
30
35
(a)
Closed pore volume (cm−3 per 100 cm−3)
7 6 5 4 3 2 1 0
(b)
Initial density ~ 1.68 g cm−3 Initial density ~ 1.74 g cm−3
Figure 11 (a) Open and (b) closed pore volume in pile grade A as a function of radiolytic weight loss.
monoxide and moisture which may be higher in the brick interior. Graphite oxidation forms carbon monoxide, thereby further increasing CO levels in the brick interior. These destruction and formation processes are gas composition dependent and the flow rates of these gases within the porous structure are dependent upon graphite diffusivity and permeability values which change with graphite weight loss. The exact mechanism of radiolysis in a CH4inhibited coolant is complex and the radicals are disputed. However, from a practical point of view the mechanisms for oxidation and inhibition can be considered as given below: In the gas phase ionizing radiation ! CO þ O CO2
½IV
CO þ O ! CO2 CH4 ! P
½V ½VI
where O* is the activated oxidizing species formed by radiolysis of CO2 and P is a protective species formed from CH4 oxidation. At the graphite surface (mainly internal porosity), O þ C ! CO
½VII
O þ P ! OP
½VIII
where OP is the deactivated gaseous product of CH4 destruction. An altogether more satisfactory explanation and model for the effect of pore structure on corrosion in gas mixtures containing carbon monoxide, CH4, and
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water was developed by Best and Wood30 and Best et al.,26 who gave a relationship for GC with respect to a pore structure parameter F and to P, the probability of graphite gasification resulting from species which reach the pore surface: GC ¼ 2:5FP
½25
The inhibited-coolant radiolytic oxidation rate is usually referred to as the graphite attack rate. Data on initial graphite attack rate have been obtained in experiments carried out in various materials test reactors (MTRs)31 for Gilsocarbon and to some extent other types of graphite (Figure 12). From Figure 12, it can be seen that the oxidation rate does not go on exponentially increasing as predicted by earlier low-dose work, but the increasing rate saturates at about 3 times the initial oxidation rate. The approach to predicting temporal and spatial weight loss in graphite components irradiated in inhibited coolant is to use numerical analysis to solve the diffusion equations given below: Methane concentrations rT ðD10 rðC1 Þ rðn C1 ÞÞ K1 ¼ 0 Moisture concentrations rT ðD20 rðC2 Þ rðn C2 ÞÞ þ K1 STOX ¼ 0 Carbon monoxide concentrations rT ðD30 rðC3 Þ rðn C3 ÞÞ ½26
þ K1 STOX þ K2 STOX2 ¼ 0
The basic unknowns are the CH4, C1, moisture, C2, carbon monoxide, C3, and gas concentration profiles.
In the CH4 part of eqn [26], the first term is the pure diffusion contribution, and D10 is the effective diffusion coefficient in graphite of CH4 in CO2. The second term is the contribution from porous flow due to permeation, and n is the velocity vector for CO2 flow through the graphite pores, and K1 is the sink term for CH4 destruction. In the moisture part of eqn [26], the first term is again the pure diffusion contribution, and D20 is the effective diffusion coefficient in graphite of moisture in CO2. The second term is the contribution from porous flow. K1STOX is the source term for moisture formation from CH4 destruction in accordance with CH4 þ 3CO2 ! 4CO þ 2H2 O
Fractional graphite weight loss
0.20
0.15
0.10
0.05
0.00 10
20
½IX
In the carbon monoxide part of eqn [26], the first term is the pure diffusion contribution, and D30 is the effective diffusion coefficient in graphite of carbon monoxide in CO2. The second term is the contribution from porous flow. K1STOX is defined above and K2STOX2 is the source term of carbon monoxide formation from graphite oxidation. The various terms in the diffusion equations must be updated at each time-step for changes in coolant composition, dose rate, attack rate, and all parameters controlling graphite pore structure, diffusivity, and permeability which change with oxidation. These equations can be solved numerically using finite difference or finite element techniques to give point wise, temporal distributions of weight loss in a graphite component.
0.25
0
345
30
40
50
60
Fluence (MWh kg−1)
Figure 12 Typical experimental weight loss dose relationship from materials test reactor experiments.
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4.11.8 Graphite Temperatures
temperatures are compared with the few brick thermocouples that are installed in the moderator. The codes are also fine-tuned to these. In conclusion, the calculation of graphite temperatures is complex and involves the calculation of heat transfer flow to the fuel and flow calculations. Graphite temperature predictions should be compared to measurements taken from thermocouples located in most graphite cores.
Graphite component temperature depends on radiation and convection (and conduction in the case of light-water gas-cooled reactors) heat transfer from the fuel and heat generated in the graphite by neutron and g-heating, that is, energy deposition as discussed above. Therefore, a detailed knowledge of the coolant flow is important. Thermohydraulic codes such as Panther (http:// www.sercoassurance.com/answers/) are used to calculate heat generated in graphite blocks. These codes estimate the following: 1. 2. 3. 4.
4.11.9 Variation of Fluence, Temperature, and Weight Loss in a Reactor Core
The heat generated in the fuel. The coolant flow. The heat transfer to the graphite. The heat ‘energy deposition’ in the graphite.
The flux, temperature, and weight loss will vary within each individual graphite component, for example, moderator brick. In addition, the mean component flux, temperature, and weight loss will vary throughout the core. When designing a graphite core, in order to extrapolate data from one component, which has been analyzed in detail, to the other core components, ‘form factors’ are often used, as illustrated in Figure 13. In typical graphite-moderated reactors, the axial (vertical) flux varies approximately as a cosine with the maximum at center, whereas the radial flux is usually a flattened cosine as illustrated in Figure 13. The exact form of these profiles can be calculated using reactor physics codes. The mean core rating can be calculated from eqn [27]:
The calculations take account of graphite weight loss and change in thermal conductivity of the graphite due to fast neuron damage and radiolytic oxidation. The largest uncertainty is probably associated with the size of flow bypass paths and flow resistance. In an AGR, the temperature at the outside of the brick is lower than the temperature at the inside because of the interstitial flow, whereas in an Reaktor Bolshoy Moshchnosti Kanalniy (RBMK) the temperature is hotter at the brick outside. Using the brick ‘boundary conditions’ including energy deposition temperatures calculated by the thermohydraulic code, a standard finite element code such as ABAQUS can easily be used to calculate the spatial distribution of temperature with the graphite component. Thermal transient temperatures can also be calculated using a standard finite element code. Often, the temperature distribution is calculated for a central brick, and the temperatures in the bricks in the rest of the core are calculated using interpolation/extrapolation, that is, form factors as described in Section 4.11.9. The calculated
Rating ¼
(a)
½27
and at the time of interest the mean core burnup can be calculated by eqn [28]: Core reactor days at ¼ burnup power power
Reflector
Core
=
reactor weight of fuel power in reactorðMWd t1 Þ
=
weight of fuel in reactor ½28 ðMWd t1 Þ
MHA
Individual channel About 320 channels in an AGR
Core
(b)
Figure 13 Form factors. (a) Graphite moderator with reflector and (b) graphite moderator flux profile (form factors).
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Thus, a mean moderator brick burnup can be calculated by multiplying the mean core burnup by the axial and radial form factor for the particular brick of interest. 4.11.9.1
code. The radial temperature can be assumed to follow the radial flux profile. Thus, an approximate mean gas temperature for an individual moderator brick may be obtained.
Fuel End Effects
4.11.10 Distribution of Fluence Within an Individual Moderator Brick
The relatively small gap between fuel elements has a pronounced effect on the damage to the graphite moderator bricks. This is particularly noticeable in the brick dimensional changes, in both AGRs and RBMK reactors. In assessments, this detail needs to be accounted for and may require a three-dimensional reactor physics calculation. 4.11.9.2
Having obtained the component mean fluence, temperature, and weight loss, the variation of these parameters throughout the particular component of interest is required. The fluence reduces exponentially away from the fuel in the radial direction, but is influenced by surrounding fuel sources. The exact distribution is usually calculated using a reactor physics code for a 5 5 array pertinent to the area of interest. Figure 14 is an example for the Windscale Piles. Thus, the spatial and temporal fluence distribution throughout a graphite component can be calculated. The component temperature can be calculated using finite element analysis through knowledge of the surrounding gas temperature, accounting for the
Temperature and Weight Loss
By using the same ‘form factors,’ the moderator brick mean weight loss can be estimated, assuming that weight loss is proportional to burnup or fluence. The gas temperature will vary roughly linearly in the axial (vertical) direction from the inlet temperature T1 to the outlet temperature T2. A more detailed profile may be calculated using a thermohydraulics
1.94
All values ⫻ 1011 1.97 2.03 2.14 2.05 2.31 2.21
2.64 2.95
2.11 2.28
2.44 2.39
2.32 2.41
2.61 2.49
3.04
3.42
2.52
2.69 3.44
4.13
2.56
2.99 2.63
4.15
3.41
2.79 3.06
4.19
2.72
3.55 4.14
2.83
18.26 21.18 23.48
347
4.21
4.06
3.44
3.47
2.72
3.08 2.69
3.05
2.80
2.70
Figure 14 Nickel flux distribution in a quarter cell calculated for the Windscale Piles. Courtesy of A. Avery.
348
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5% of the reactor heat which is generated within the graphite. Graphite weight loss variation within a component is more complex and is calculated by various empirical industry codes. If the axial variation in fluence, temperature, and weight loss along the brick length is deemed to be important, three-dimensional physics, temperature, and weight loss calculations will be required.
4.11.11 Fast Neutron Damage in Graphite Crystal Structures Atomic displacements due to fast neutron irradiation modify the ‘crystallite’ dimensions and most of their material properties. Neutron energies of around 60 eV are required to permanently displace carbon atoms from the lattice. However, most damage in graphite is due to fast neutron energies >0.1 MeV; a typical thermal reactor has neutron energies of up to 10 MeV, with an average of 2 MeV. High-energy neutrons knock an atom out of the lattice, leading to a cascade of secondary knock-ons. This process knocks atoms into interstitial positions between the basal planes, leaving vacant positions within the lattice. Many of the interstitial atoms will immediately find and fill these vacancies. However, others may form semistable Frenkel pairs or other small clusters or ‘semistable’ clusters. With increasing fast neutron damage, the stability, size, and number of these clusters will change depending on the irradiation temperature. The higher the irradiation temperature, the larger are the interstitial clusters or ‘loops.’ This process leads to considerable expansion in the graphite crystal ‘c’ axis. Conversely, vacancy loops also form and grow in size with increased irradiation temperature. It has been postulated that this process will cause the lattice to collapse leading to the ‘a’ axis shrinkage observed on irradiating graphite crystal structures. This process is illustrated in Figure 15. Thrower32 carried out an extensive review of transmission electron microscopic (TEM) studies of defects in graphite, particularly those produced by fast neutron irradiation. He demonstrated that interstitial loops and vacancy loops could be distinguished by tilting the specimen. He was able to observe vacancy loops in graphite irradiated only at and above 650 C, whereas interstitial loops and defects were observed at all temperatures of interest to reactor graphite. It is proposed that the dimensional change in bulk polycrystalline graphite may be understood by eqn [29]33:
2 DLc Dc r0 þ ffi Lc r1 c
½29
where Lc is the crystal dimension perpendicular to the basal plane, ‘c’ is the atomic lattice parameter, and r0 and r1 are the mean defect radius and mean half separation of defects in the basal plane, respectively. However, it was noted that this does not completely explain the expansion. In order to explain basal plane contraction it is necessary to postulate that vacancy lines cause the collapse of the basal planes.34–36 More recent atomistic calculations due to Telling and Heggie37 have sought to explain the process by the ‘buckling’ of basal planes until they twist round upon themselves. This latter explanation is more satisfying as it accounts for the atomistic bonding around the edges of the interstitial loops and vacancies. However, more HRTEM (high-resolution transmission electron microscopy) observations and other techniques are required to validate these theories. Whichever mechanism is correct, empirical observations made on HOPG, and some natural crystal flakes,35 show that graphite crystal structures expand in ‘c’ axis and shrink in the ‘a’ axis, the degree of deformation being a function of fast neutron fluence and irradiation temperature. Crystal dimensional change is discussed in more detail later in this section. 4.11.11.1
Stored Energy
It would not be appropriate to continue without some discussion on stored (or Wigner) energy. The perfect crystal configuration is the lowest energy state for the graphite lattice. However, irradiation damage will considerably alter that configuration. Wigner38 predicted that the increased lattice vibration due to heating would allow carbon atoms to rearrange themselves into lower energy states, and that in doing so energy would be released in the form of heat. Early experience in operating graphitemoderated plutonium production and research reactors at low temperatures in the United States, Russia, France, and the United Kingdom proved that this assumption was correct. The highest value of stored energy measured was 2700 J g1.15 If all of this were released under adiabatic conditions, the temperature rise would be 1500 C. Fortunately, that is not the case. Furthermore, the accumulation of stored energy is insignificant above an irradiation temperature of 300 C, it is difficult to accidentally release the stored energy above an irradiation
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~0.01 µm
~500 eV
349
0.001 µm
1 keV
Vacant lattice site Interstitial atom
Displacement cascade
Interstitial defects
Vacancy defects
Layer plane
Dose increase
Interstitial loop
Single vacancy
Submicroscopic cluster of 4 ± 2 atoms
Interstitial diffusing to loop
Layer increase
Di vacancy
Dose increase
Vacancy line
Vacancy loop Figure 15 Formation of interstitial and vacancy loops in graphite crystals. Modified from Simmons, J. Radiation Damage in Graphite; Pergamon: London, 1965.
temperature of 150 C, and only limited self-sustaining energy release of stored energy can be achieved in graphite irradiated below 100 C. Thus, stored energy is now of consideration in the United Kingdom only in the decommissioning of shutdown reactors such as the Windscale Piles and BEPO and other similar overseas systems, although there are graphite ‘thermal columns’ in some research reactors that may require periodic assessment. The reason for this is the nature of the irradiation damage sites with respect to irradiation temperature. In graphite irradiated in the early facilities,
at temperatures between about ambient and 150 C, point defects associated with Frenkel pairs and small loops can diffuse only slowly through the lattice to form larger, more stable loops because of the low irradiation temperature. However, thermal annealing at temperatures above the irradiation temperature can readily release the stored energy, and under certain circumstances, this release can be self-sustaining over certain temperature changes. (A ‘rule of thumb’ temperature of 50 C above the irradiation temperature is often cited as a ‘start of release temperature.’ However, this is misleading as a heat balance needs to be
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900 25 ⬚C 30 ⬚C 80–100 ⬚C 99 ⬚C 125 ⬚C 140 ⬚C 150 ⬚C 180 ⬚C
800
Total stored energy (J g−1)
700 600 500 400 300 200 100 0 0.0
0.2
0.4
0.6
0.8 20
1.0
1.2
1.4
1.6
–2
Fluence (10 n cm nvt) Figure 16 The accumulation of stored energy as a function of fluence and temperature.
considered when assessing energy release rates. Thus, 50 C above the irradiation temperature can be considerably overconservative.) The accumulation of stored energy, measured by burning irradiated graphite samples in a bomb calorimeter, is given as a function of fluence and temperature in Figure 16. At low fluence, stored energy quickly accumulates reaching a plateau at high fluence. Many measurements were made in the Windscale Piles, BEPO, Hanford, and Magnox reactors that clearly illustrated this behavior.15 To fully understand the thermal stability of graphite containing stored energy, the most appropriate measure is the rate of release of stored energy measured using a differential scanning calorimeter (DSC) as illustrated in Figure 17. A graphite sample is heated in the DSC usually at a constant rate of 2.5 C min1. In simple terms, two runs are made and the heat capacity of the samples measured in each case. When the heat capacities from the two runs are subtracted, the energy release rate is easily obtained as a function of heating temperature. This can be compared to the specific heat of graphite as given in Figure 17. When the rate of release of energy is below the specific heat, energy needs to be added to continue the process. When the rate of release is above the specific heat, the process is self-sustaining. This behavior was used to ‘anneal’ the Windscale Piles; a ‘hit and miss’ strategy that ended in damage to the fuel cartridges and eventually a
‘metal uranium fire.’ (Contrary to ‘common folklore,’ the graphite did not burn in the Windscale incident. A limited amount of graphite was oxidized leading to enlargement of fuel and control channels but it was the metal uranium that burnt. Graphite is very difficult to burn and requires large amounts of heat and oxygen or air, applied to crushed graphite in a fluidized bed or in similar form.39) The form of this rate of release curve is a function of (1) the amount of stored energy in the sample, (2) the temperature the sample was irradiated at, (3) the fluence the sample had been irradiated to, (4) the release temperature, and (5) the heating rate. Unfortunately, there are no comprehensive datasets of these five parameters that allow a robust empirical model to be derived for assessing the stability of graphite containing stored energy. The models that usually exist take the worstcase rate of release curve and fit an Arrhenius type equation to the rate of release curve. dS EðT ; SÞ ¼ uS exp ½30 dt KT ðt Þ where S is the stored energy remaining, t is time, T(t) is temperature in (K) as a linear function of time (T ¼at in the case of the DSC test and is nonlinear in most practical cases), K is Boltzmann’s constant, and E(T, S) is the activation energy as a function of the stored energy remaining and temperature and u is a frequency factor usually taken as 7.5 1013 s1.40
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3.0 Rate of release of stored energy (J g−1 ⬚C−1)
Specific heat capacity Dowel 4 Dowel 11
2.5
Dowel 14 Dowel 24
2.0
Dowel 30 Dowel 36
1.5
1.0
0.5
0.0 0
50
100
150
200
250
300
350
400
450
500
Temperature (⬚C) Figure 17 Typical rate of release of stored energy. Modified from Bell, J. C.; Gray, B. S. Stored Energy Studies Made on Windscale Pile Graphite Since October 1957; TRG Report 84(W), UKAEA, 1961.
It can be appreciated that the exact solution of eqn [30] requires a substantial amount of information from several rate of release curves from several samples, which is seldom available. Thus, a practical approach is usually taken, the simplest of which is to assume a single activation energy. However, this is not very satisfactory and more elegant approaches using variable or discrete activation energies can be found.41–44 Having derived a satisfactory model for the rate of release using a DSC, it then can be applied to a practical situation using commercially available computer codes such as ‘user subroutine’ facilities.42 In assessing practical situations, it is important to use an energy balance that accounts for heat applied, heat generated by the release of stored energy, the heat capacity of the graphite itself, and heat lost to the surroundings. If the heat generation is intense and oxygen is available, the heat generated by graphite oxidation should be taken into account. However, the latter case should be unnecessary as a professional scientist or engineer would not design a system or process that would approach such conditions. It should be noted that irradiated graphite thermal conductivity and total stored energy are directly correlated15; see eqn [31]. Therefore, the thermal conductivity will improve as energy is released. K0 1 J g1 ½31 S ¼ 27:2 K
The data that eqn [30] is based on is derived from rate of release curves obtained using a relatively fast heating rate. In dealing with irradiated graphite waste, much slower rates of heating are often required. Graphite samples taken from the Windscale Piles 40 years after the incident showed little change in the dS/dt curves,45 indicating that diffusion of atoms at around ambient temperature is extremely slow. Nevertheless, conditions relevant to any proposed encapsulation technique and repository will need to be accounted for in determining if heat released from stored energy is an issue. The rate of release curves given in Figure 17 are only to a temperature of around 450 C. It had been observed, by comparing the energy released in the DSC with the energy released on a similar sample in a bomb calorimeter, that not all of the stored energy had been released in the samples heated to a maximum of 450 C in the DSC. It was found that on increasing the temperature to around 1600 C, a second peak could exist46,47; see Figure 18. It was observed that the ‘200 C’ peak reduced in size and moved to a slightly higher temperature with increased irradiation, presumably as the irradiation induced defects became more stable, and the plateau between the two peaks increased in height and approached the specific heat value.15 The first of these phenomena could explain why it became more and more difficult to ‘anneal’ the Windscale Piles39 and the second had the implication that
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Graphite in Gas-Cooled Reactors
500 Graphite-specific heat Rate of release
450 400
Energy (cal g−1 ⬚C−1)
350 300 250 200 150 100 50 0 0
200
400
600
800
1000
1200
1400
1600
1800
2000
Temperature (⬚C) Figure 18 Schematic of the high-temperature rate of release curve. Reproduced from Rappeneau, J.; Taupin, J. L.; Grehier, J. Carbon 1966, 4(1), 115–124.
eventually the rate of release curve would remain above the specific heat up to 1600 C with the consequent safety implications. Fortunately, the second of these phenomena proved to be incorrect. It is interesting to note that there is a correlation (eqn [32]) between the height of the plateau at 400 C and total stored energy.15 dS S J g1 C1 ¼ dT½400 1670
½32
This equation, although not exact, was often used in reactor graphite sampling programs to avoid having to measure total stored energy. Most of the discussion on stored energy above is relevant only to low temperature reactor systems, with graphite temperatures operating from ambient to 150 C. When graphite is irradiated at higher temperatures, in practice above about 100 C, the dS/dt does not exceed the graphite specific heat. One of the operating rules for the UK Magnox reactors was that the dS/dt, as measured on surveillance samples, should always be below 80% of the specific heat, which proved to be the case. 4.11.11.2 Crystal Dimensional Change As previously discussed, graphite crystal structures, in the form of HOPG, have been observed to swell in
the ‘c’ axis direction and contract in the ‘a’ axis direction for all measured fluence and irradiation temperatures. Figure 19 shows early data obtained by Kelly et al.35 It is clear from Figure 19 that the rate of swelling and shrinkage significantly changes between 200 and 250 C, indicating that the defect population is becoming more stable above this temperature range. HOPG data is an important input into multiscale models of irradiation damage in graphite.48 For the purpose of understanding irradiation damage in operating reactors, data would be required ideally from 140 to 1400 C, the maximum fluence being dependent on the irradiation temperature. Unfortunately, the dataset is far from complete. The data due to Brockelhurst and Kelly49 is the most complete set of HOPG irradiation data covering the fluence and part of the temperature range appropriate to AGRs (Figure 20). In the same paper, the authors show the effect of final heat treatment, between 2000 and 3000 C, on the crystal dimensional change rate of HOPG. The data showed that the lower the heat treatment, the faster is the dimensional change rate, indicating that the dimensional change rate of a poorly graphitized component would be expected to be greater than that of majority of the components.50 Previously, Kelly and Brocklehurst51 had shown that boron doping also significantly increased the dimensional
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353
30 150 °C 170 °C 200 °C 250 °C 350 °C 450 °C 650 °C
Dimensional change (%)
25 20 15 10 5 0 0
5
10
15
20
25
30
Fluence (1020 n cm-2 EDND) c-axis 1
Dimensional change (%)
0 -1 -2 150 °C 170 °C 200 °C 250 °C 300 °C 350 °C 450 °C 650 °C
-3 -4 -5 -6 -7 0
5
10
15
20
25
30
Fluence (1020 n cm-2 EDND) a-axis Figure 19 Dimensional changes in highly orientated pyrolytic graphite as a function of fast neutron fluence and temperature. Modified from Kelly, B. T.; Martin, W. H.; Nettley, P. T. Philos. Trans. R. Soc. Lond. A Math. Phys. Sci. 1966, 260 (1109), 37–49.
change rate in HOPG, and this was again reflected in the behavior of doped polycrystalline graphite.50 HOPG high-temperature data was mainly obtained by investigators interested in the behavior of HTR fuel coatings.52 Some of this data is for low-density pyrolytic carbons and it is not always made clear which material the data refers to. Figure 21 shows all the data known to the author, and it is clear that there is some inconsistency. 4.11.11.3 Coefficient of Thermal Expansion There are only two reasonable sets of data for the CTE in HOPG. Data for the lower temperatures (150–250 C) is given in Figure 22. This data is
associated with the dimensional change data given in Figure 19. At these low temperatures, there is a significant increase in dimensional changes with increased fluence. This could explain the increase, and the subsequent decrease, in CTE in the ‘c’ direction. It is interesting that this is reminiscent of the behavior of medium grained, semi-isotropic polycrystalline data, as discussed later. In the ‘a’ direction, there is significant scatter in the data, possibly due to the difficulty in measuring such low values of CTE. However, the behavior appears to indicate an increase to a plateau, reminiscent of the behavior of needle coke anisotropic polycrystalline data. The higher temperature HOPG CTE data given in Figure 23 appears to be invariant to the increasing
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70
c-axis, 430 ⬚C c-axis, 600 ⬚C a-axis, 430 ⬚C a-axis, 600 ⬚C
Dimensional change (%)
60 50 40 30 20 10 0 −10 −20 −30
0
20
40
60
80
100
120
140
160
180
200
Fluence (1020 n cm−2 EDND) Figure 20 High-dose highly orientated pyrolytic graphite data. Reproduced from Brocklehurst, J. E.; Kelly, B. T. Carbon 1993, 31(1), 179–183.
fluence, although the maximum fluence is limited to 30 1020 n cm2 EDND. The author is not aware of any data at higher irradiation temperatures. 4.11.11.4 Modulus Changes to C33 and C44 in HOPG and natural graphite crystals have been reported at 50, 650, and 1000 C53 and at 150 C.54 For HOPG, the 150 C data indicated that C33 slightly reduced with increasing irradiation (Figure 24) and this was attributed to the increase in ‘c’ axis lattice spacing. However, there is no clear trend at the other temperatures (Figure 24). In the case of shear, at a very low temperature of 50 C there was a significant increase in C44, but at higher temperatures the increase was less (Figure 25). The data for natural crystal showed similar trends but there was significantly more scatter. The trend in the increase in C44 at the lower temperature would go towards explaining the increase in modulus in polycrystalline data at low fluence. However, it is surprising that the increase is only modest at the higher temperatures, although the maximum fast neutron fluence is very low and data is required at the intermediate temperatures. Seldin and Nezbeda53 also measured the shear strength but unfortunately there is considerable scatter and no definite trend. 4.11.11.5 Thermal Conductivity Taylor et al.55 measured the change in thermal conductivity in HOPG with fast neutron
irradiation. The thermal conductivity along the basal planes (the ‘a’ direction) is much greater than the value perpendicular to the basal planes (the ‘c’ direction). Taylor et al. also measured the change in thermal resistivity in irradiated graphite, and when this data is normalized, the data indicated that thermal resistivity temperature dependence changed with irradiation as given in Figure 26. This is the so-called ‘d’ curve that is used in the United Kingdom to predict thermal resistivity in irradiated graphite. 4.11.11.6
Raman
Figure 27 gives Raman spectra for unirradiated and irradiated graphite as well as for baked carbon. In the spectral range shown, there is a prominent G-peak at 1580 cm1 associated with the basal plane bond stretching of ‘c’ axis sp2 atoms. The D-peak at 1350 cm1 is associated with the breathing mode of sp2 atoms and disordered carbon structure. The second D-peak at 2700 cm1 is indicative of the crystalline structure of the graphite. In Figure 28 the normalized positions of the G- and D-peaks, and the ratio of the peak intensities are compared for various graphites (unirradiated and irradiated). HOPG is obviously the most ordered structure followed by the PGA needle coke graphite and then the medium grained graphite grades. The most disordered materials are the baked carbon (NBG-18 baked) followed by irradiated BEPO (a UK test reactor) graphite.
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120 Kelly (450 ⬚C) Kelly (650 ⬚C) Kelly (900 ⬚C) Kelly (1200 ⬚C)
Dimensional change (%)
100
Bokros et al. (900 ⬚C) Bokros et al. (1075 ⬚C) Bokros et al. (1250 ⬚C)
80 60 40 20 0 0
20
40
100 60 80 Fluence (1020 n cm−2 EDND) c-axis
120
140
160
0
20
40
100 60 80 Fluence (1020 n cm−2 EDND) a-axis
120
140
160
0
Dimensional change (%)
−5 −10 −15 −20 −25 −30
Figure 21 High-temperature dimensional change data for highly orientated pyrolytic graphite.
Figure 29 qualitatively demonstrates that there is a relationship between Raman spectra and crystal structure disorder. The higher the disorder, the higher are the D- and G-peak wave number and I(D/G) ratio. In Figure 29(a), the crystal length La has been calculated from the full width at half maximum (FWHM) using the method proposed by Tuinstra and Koenig.56 Both figures demonstrate that Raman can be used to quantify the disorder in the graphite structure, either as manufactured or due to irradiation.
4.11.12 Property Changes in Irradiated Polycrystalline Graphite Fast neutron irradiation and, in the case of carbon dioxide-cooled reactors, radiolytic oxidation change many of the properties of graphite. The
properties of interest to the nuclear engineer are the following: Stored energy – a function of fast neutron damage and temperature, due to damage to the graphite crystallites, but not affected by radiolytic oxidation other than by a reduction in mass. Specific heat – a function of temperature but not affected by fast neutron irradiation or radiolytic oxidation other than by stored energy, which may be considered separately. Dimensional changes – a function of fast neutron damage and irradiation temperature. There is also some evidence of modification by radiolytic oxidation. It is also modified by stress (seeirradiation creep). CTE – a function of temperature, fast neutron damage, irradiation temperature, and stress. There is evidence that it is not modified by radiolytic weight loss.
Graphite in Gas-Cooled Reactors
Coefficient of thermal expansion (10−6 K−1)
356
30 28 26 24 22 20 18 16 14 12 10 0
2
4
6
10
8
12
14
16
Coefficient of thermal expansion (10−6 K−1)
Fluence (1020 n cm−2 EDND) c-axis 2.0 1.5 1.0 0.5 0.0 −0.5 −1.0 −1.5 −2.0 −2.5 0
2
4
6
8
10
12
14
16
18
Fluence (1020 n cm−2 EDND) a-axis 150 ⬚C
170 ⬚C
200 ⬚C
250 ⬚C
Figure 22 Low-temperature changes to coefficient of thermal expansion in highly orientated pyrolytic graphite. Modified from Kelly, B. T.; Martin, W. H.; Nettley, P. T. Philos. Trans. R. Soc. Lond. A Math. Phys. Sci. 1966, 260(1109), 37–49.
Thermal conductivity – a function of temperature, fast neutron damage, and irradiation temperature. It is significantly modified by radiolytic weight loss. Young’s modulus – a function of temperature, fast neutron damage, and irradiation temperature. It is significantly modified by radiolytic weight loss. Strength (tensile, compressive, flexural, and fracture) – a function of temperature, fast neutron damage, and irradiation temperature. It is significantly modified by radiolytic weight loss. Electrical resistivity – a function of temperature, fast neutron damage, and irradiation temperature. It is probably modified by radiolytic weight loss.
Irradiation creep – a function of fast neutron damage, irradiation temperature, and stress. These property changes are illustrated in Figure 30. These dimensional changes, property changes and creep mechanisms are correlated, some more strongly than others. This has been taken advantage of in various semiempirical models for irradiation damage in graphite.57–60 In discussing the irradiation behavior of polycrystalline graphite, it is useful to split these changes into low, medium, and high fluence effects. At low irradiation, fluence changes in polycrystalline graphite are strongly correlated with the crystallite changes discussed elsewhere; see Section 4.11.11. Typical mechanisms would
Coefficient of thermal expansion (10−6 K−1)
Graphite in Gas-Cooled Reactors
357
30 25 20 15 10 350 ⬚C 450 ⬚C 650 ⬚C
5 0 0
5
10
15
20
25
30
Coefficient of thermal expansion (10−6 K−1)
Fluence (1020 n cm−2 EDN) c-axis 0
−1
−2
−3 300 ⬚C 350 ⬚C 450 ⬚C 650 ⬚C
−4
−5
0
5
10 15 20 Fluence (1020 n cm−2 EDND) a-axis
25
30
Figure 23 High-temperature changes to coefficient of thermal expansion in highly orientated pyrolytic graphite. Reproduced from Brocklehurst, J. E.; Kelly, B. T. Carbon 1993, 31(1), 179–183.
be the accumulation of stored energy, pinning (leading to a rapid increase in Young’s modulus), and the rapid decrease in thermal conductivity. At medium fluence, several of the properties saturate, such as Young’s modulus and thermal conductivity. At high fluence, when crystallite growth in the ‘c’ direction has taken up much of the accommodation provided by Mrozowski cracks,11 and larger ‘cracks,’ the polycrystalline structure starts to become strained, thereby generating new cracking. At extremely high fluence, beyond that experienced in a modern power production reactor, the crystallite swelling becomes so large that the polycrystalline structure completely breaks down leading to a rapid decrease in modulus and strength.
Each of the property changes is discussed in more detail below. In attempting to understand the behavior of polycrystalline graphite, reference is made to the irradiation behavior of HOPG, as previously discussed in Section 4.11.11. This is because HOPG is considered to be a representative model material for the individual crystallite structures in polycrystalline graphite.
4.11.13 Averaging Relationships Before looking at each of the properties individually, it is first worth considering the methods developed to relate changes in the crystallites to the bulk
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40
Elastic modulus C33 (GN m−2)
35 30 25 20 15 10
50 ⬚C 150 ⬚C 650 ⬚C 1000 ⬚C
5 0 0.0
0.5
1.0 1.5 Fluence (1020 n cm−2 EDND)
2.0
2.5
Elastic modulus C44 (GN m−2)
Figure 24 Reduction in C33 in irradiated highly orientated pyrolytic graphite. Modified from Seldin, E. J.; Nezbeda, C. W. J. Appl. Phys. 1970, 41(8), 3389–3400; Summers, L.; Walker, D. C. B.; Kelly, B. T. Philos. Mag. 1966, 14(128), 317–323.
identical crystals with the correct graphite symmetry, and small element volumes containing only graphite could be chosen so that their internal and external stress could be considered to be uniform.61 From first principles and by applying the laws of thermodynamics for the assumptions given above, Simmons derived the following relationship for the bulk dimensional change rate:
5 4 3 2 1 0 0
10
20
30
50
40
60
70
80
Fluence (1017 n cm−2)
Elastic modulus C44 (GN m−2)
50 ⬚C 0.5
½33
where dexx/dg, dxc/dg, and dxa/dg are the linear dimensional change rate in polycrystalline graphite in some direction ‘x’ and the crystallite dimensional change rates in the ‘a’ and ‘c’ directions respectively. Simmons also derived the following relationship for the bulk CTE:
0.4 0.3 0.2 0.1 0.0 0.0
dexx dxc dxa ¼ Ax þ ð1 Ax Þ dg dg dg
650 ⬚C 1000 ⬚C 0.2
0.4
0.6
0.8
1.0
1.2
axx ¼ Ax ac þ ð1 Ax Þaa 1.4
Fluence (1020 n cm−2) 650 and 1000 ⬚C
Figure 25 Changes in C44 with irradiation in highly orientated pyrolytic graphite. Modified from Seldin, E. J.; Nezbeda, C. W. J. Appl. Phys. 1970, 41(8), 3389–3400.
properties of polycrystalline graphite. In the 1960s, Simmons derived a model based on the following assumptions: polycrystalline graphite could be considered to be a single phase, porous aggregate of
½34
where axx, ac, and aa are the linear CTE in polycrystalline graphite in some direction ‘x’ and the crystallite CTE in the ‘a’ and ‘c’ directions, respectively. The so-called structure factor ‘Ax’ is the summation rate of change in rate of the crystallite stresses with respect to the change in bulk stress, as illustrated in the equations below. X @s033;n ½35 an Ax ¼ @sxx n
Fractional change in thermal conductivity (K0 K-1)
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0.7 270 ⬚C 300 ⬚C 350 ⬚C 400 ⬚C 460 ⬚C 510 ⬚C 780 ⬚C
0.6 0.5 0.4 0.3 0.2 0.1 0.0 0
50
(a)
100 150 200 Fluence (1020 n cm−2 EDND) 5 Normalized change in thermal resistance
Change in thermal resistance (cm 0K W-1)
10.0
1.0
0.1
0.0
250
300
30 ⬚C 150 ⬚C 300 ⬚C 450 ⬚C 150 ⬚C (annealed) 300 ⬚C (annealed) 450 ⬚C (annealed)
4 3 2 1 0
0
(b)
359
100 200 300 400 500 600 700 800 Temperature (⬚K)
0
(c)
100 200 300 400 500 600 700 Temperature (⬚K)
800
Figure 26 Changes in the thermal conductivity of highly orientated pyrolytic graphite. (a) Change in thermal conductivity in the ‘a’ direction as a function of fluence, (b) Change in thermal resistance as a function of temperature, and (c) Normalized change in thermal resistivity as a function of irradiation and measurement temperature. Reproduced from Taylor, R.; Kelly, B. T.; Gilchrist, K. E. J. Phys. Chem. Solids 1969, 30, 2251–2267.
X @ 0 0 ð1 Ax Þ ¼ an ðs þ s11;n Þ @sxx 22;n n For a more detailed derivation of these equations see Hall et al.61 In addition, Simmons4 provided evidence that there is a linear relationship between unirradiated CTE and initial dimensional change rate for polycrystalline graphite (Figure 31). Brocklehurst and Bishop62 later found a similar relationship in brominated graphite. Expressions to those of Simmons have been derived by Sutton and Howard12: apar ¼ K1 gac þ K2 baa aperp ¼ K3 gac þ K4 baa
½36
where K1, K2, K3, K4, g, and b are crystal accommodation and Bacon13 crystal orientation factors. A similar but more complex relationship was also derived by Jenkins.63 In the discussion of the individual
properties, the anisotropic graphite PGA and the semi-isotropic Gilsocarbon graphite are used as examples.
4.11.14 Dimensional Change As discussed previously (see Section 4.11.11.2), when irradiated, graphite crystallites, as simulated using HOPG, expand significantly in the ‘c’ direction and shrink in the ‘a’ direction. These dimensional changes are reflected in the behavior of polycrystalline graphite, but the volumetric changes, although relatively large are much smaller than those seen in HOPG. The reason for this is attributed to the many microcracks, which range in size from the nano- to microscale; see Figure 32. While these cracks can accommodate the crystal growth in the ‘c,’ the shrinkage in the ‘a’ direction will directly be reflected in the polycrystalline behavior. However, as previously
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1000
2000
discussed, the crystallite dimensional change rate is much greater below 300 C than above that temperature.
3000
BEPO-20
HTR 1 baked
4.11.14.1
HTR 2
The blocks of PGA were manufactured by extruding needle-shaped filler particles mixed with a pitch binder. During the extrusion process, the needle-shaped filler particles tended to align with the ‘a’ axis parallel, and the ‘c’ axis perpendicular, to the extrusion direction. Thus, the final product (or block) had two orthotropic directions: parallel to the extrusion direction (WG) and perpendicular to the extrusion direction (AG). This strong orientation in direction is reflected not only in the unirradiated properties but also in the irradiation properties and dimensional changes. Dimensional
HTR 1 Gilsocarbon PGA HOPG
1000
2000
3000
Wavenumber (cm−1)
Pile Grade A
10
Normalized D peak position (cm−1)
(a)
(b)
Graphite grade
BEPO-16
Graphite grade
1.4 1.2
I(D)/(G)
1.0 0.8 0.6 0.4
(c)
BEPO-20
BEPO-16
BEPO-13a
BEPO-1
HTR 3 baked
HTR 3
HTR 3 irradiated
HTR 2 irradiated
HTR 1 baked
HTR 2
HTR 1
Gilsocarbon
PGA
0.0
HOPG
0.2
Graphite grade
Figure 28 Relative position and ratio of I(D/G) by graphite grade and condition. (a) Normalized position of G-Peak, (b) Normalized position of D-Peak, (c) Ratio of the D-peak and G-peak intensities. Courtesy of A. Jones, University of Manchester.
BEPO-20
BEPO-13a
BEPO-1
HTR 3
HTR 3 baked
HTR 3 irradiated
HTR 2 irradiated
HTR 2
HTR 1 baked
BEPO-16
0
BEPO-20
BEPO-13a
BEPO-1
HTR 3
HTR 3 baked
HTR 3 irradiated
HTR 2 irradiated
HTR 2
HTR 1 baked
HTR 1
PGA
Gilsocarbon
0
5
HTR 1
5
PGA
10
Gilsocarbon
15
HOPG
Normalized G peak position (cm−1)
Figure 27 Raman spectra for unirradiated and irradiated graphite and baked carbon. Courtesy of A. Jones, University of Manchester.
HOPG
Intensity (a.u.)
BEPO-16
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Graphite in Gas-Cooled Reactors
1.4 1.2 Irradiated
1.0 I(D/G)
Baked
0.8 HTR 3
0.6
HTR 1
HTR 2
PGA
0.4
Gilsocarbon BEPO-1
0.2 0.0 0
20
40
60
(a)
80 La (A)
100
120
140
120 Jones Knight and White
Coke
FWHM (cm−1)
100 80
Glassy carbon
Irradiated graphite
60 40
HTR 1 baked
HTR 3 baked
20 Polycrystalline graphites
HOPG
0 0
0.2
0.4
0.6
(b)
0.8
1
1.2
1.4
I(D/G)
0.00
−0.05
1
−2
20
5 4
1.5
3
1.0 2
2
10 0.5
−3
0 0
(a)
2.0 E/Eo−1
−1
6
30 esu
0
7
2.5
4
CTE (K−1 ⫻ 10−6)
0.05
8
3.0
40
50
100
150
Fluence (n cm−2 ⫻ 1020 EDND)
1 0
0.0 0
200
(b)
Ko/K−1
2
0.10 Dimensional change (%)
Dimensional change rate (% per fluence)
Figure 29 Quantitative relationship between I(D/G) ratio and crystallite length (La) and full width at half maximum. (a) I(D/G) as a function of crystal size (La) and (b) FWHM as a function of I(D/G). Courtesy of A. Jones, University of Manchester.
50
100
150
200
Fluence (n cm−2 ⫻ 1020 EDND)
Figure 30 Schematic of the irradiation-induced changes in Gilsocarbon graphite irradiated at 550 C (note that there will be a similar set of curves for each irradiation temperature). (a) Dimensional change, dimensional change rate, and coefficient of thermal expansion and (b) Factorial change in Young’s modulus (E/E01) and thermal conductivity (K0/K1), and irradiation creep (elastic strain units, esu).
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Dimensional change rate (% per 1020 n cm−2)
0.5
0.5
0.4
0.3 0.2 0.1 0.0 0
1
2
0.3
3
4
5
6
CTE (10−6 K−1) 80 ⬚C
(b) 0.08 Dimensional change rate (% per 1020 n cm−2)
Dimensional change rate (% per 1020 n cm−2)
0.6
0.4
0.2
0.1
0.0 1
0
2
3
4
CTE (10−6 K−1) 30 ⬚C
(a)
0.06 0.04 0.02 0.00 −0.02 −0.04
0 (c)
1
2
3
4
5
CTE (10−6 K−1) 180 ⬚C Parallel to extrusion
Perpendicular to extrusion
Figure 31 Correlation between initial growth rate and unirradiated coefficient of thermal expansion. Modified from Simmons, J. Radiation Damage in Graphite; Pergamon: London, 1965.
(a)
1 µm
(b)
500 µm
Figure 32 Graphite microstructure showing Mrozowski cracking and larger porosity. (a) Mrozowski cracks and (b) Polarized optical image of Gilsocarbon microstructure. Courtesy of A. Jones, University of Manchester.
change MTR data for PGA is given in Birch and Brocklehurst64 for both the parallel (WG) and perpendicular (AG) directions. In the parallel direction and below 300 C, the large dimensional
change rate in the crystallite ‘c’ axis causes the graphite to swell. Above 300 C and parallel to the extrusion direction, the graphite shrinks in the lower fluence range. This behavior can
Graphite in Gas-Cooled Reactors
6
150 ⬚C 200 ⬚C
Dimensional change (%)
5
250 ⬚C 300 ⬚C
219 ⬚C 225 ⬚C
350 ⬚C 450 ⬚C
363
650 ⬚C
4 3 2 1 0 −1 −2 0
10
20
30
40
50
60
50
60
Fluence (1020 n cm−2 EDND)
(a)
0.0
Dimensional change (%)
−0.5 −1.0 −1.5 −2.0 −2.5 −3.0 0 (b)
10
20 Fluence
30 (1020
n cm−2
40 EDND)
Figure 33 Low- to medium-fluence irradiation dimensional change in pile grade A graphite. (a) perpendicular to extrusion and (b) parallel to extrusion.
be compared to that of HOPG, as given in Figure 33. If PGA is irradiated to a higher fluence, the shrinkage rate reduces until the graphite begins to expand or ‘turns around,’ as illustrated in Figure 34. ‘Turnaround’ is associated with the closure of the Mrozowski cracks; see Figure 32. When all of the accommodation provided by the cracks has been taken up, the larger ‘c’ crystallite dimensional change rate would be expected to dominate the ‘a’ axis shrinkage rate. This behavior has been used, with some success, to model the dimensional change behavior in PGA, Gilsocarbon, and Russian GR-280 graphite.48,65,66
4.11.14.2
Gilsocarbon
As previously discussed, Gilsocarbon graphite for the AGRs was manufactured by molding (or pressing) the spherical filler particles and blocks, resulting in a semi-isotropic graphite with an anisotropy ratio of 1.01 (based on the ratio of the orthotropic CTE values). Dimensional change MTR data for Gilsocarbon over a wide range of temperatures is given in Figure 35. There are two sets of data at each temperature; one WG and one AG. This illustrates how isotropic the properties of Gilsocarbon are, even when irradiated. In Figure 35, it is also clearly illustrated that the higher the irradiation
364
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4
Dimensional change (%)
2 0 −2 −4 −6 −8
Parallel to extrusion Perpendicular to extrusion
−10 0
40
20
80
60
100
120
140
160
Fluence (1020 n cm−2 EDND) Figure 34 High-fluence irradiation dimensional change in pile grade A graphite irradiated at 600 C. Reproduced from Brocklehurst, J. E.; Kelly, B. T. Carbon 1993, 31(1), 155–178.
430 ⬚C 600 ⬚C 900 ⬚C 940 ⬚C 1240 ⬚C 1430 ⬚C
6
Dimensional change (%)
5 4 3 2 1 0 −1 −2 −3 −4
0
50
100 Fluence
(1020
150 n cm−2
200
250
EDND)
Figure 35 Dimensional change in Gilsocarbon graphite. Modified from Birch, M.; Brocklehurst, J. E. A review of the behaviour of graphite under the conditions appropriate for the protection of the first wall of a fusion reactor; UKAEA, ND-R-1434(S); 1987; Brocklehurst, J. E.; Kelly, B. T. Carbon 1993, 31(1), 155–178.
temperature, the sooner the turnaround is reached. At very low fluence, semi-isotropic graphite swells. This swelling can be quite significant as demonstrated by the irradiation of semi-isotropic NBG10 at 294 and 691 C.67 This behavior has been attributed to the annealing out of residual machining stresses or shrinkage strains, but there are no microstructural or other experimental observations to validate this reasoning. When graphite is irradiated past turnaround and reaches its original volume, sometimes referred to
as ‘critical fluence,’ the structure of the graphite begins to break down, as illustrated in Figure 36. 4.11.14.3 Effect of Radiolytic Oxidation on Dimensional Change When designing the UK AGRs, irradiation experiments in a carbon dioxide atmosphere were carried out in BR-2 at Mol, Belgium. These experiments were designed to obtain high radiolytic weight loss (35%) in a very short time, and hence, a low fast
Graphite in Gas-Cooled Reactors
(a)
365
(b)
Figure 36 Gilsocarbon specimens irradiated to very high irradiation fluences; the original specimens were 7 mm long by 5 mm diameter. (a) Gilsocarbon irradiated to 285 1020 n cm2 EDND þ0.9% DV/V0 and (b) Gilsocarbon irradiated to 271 1020 n cm2 EDND 33% DV/V0. Reproduced from Brocklehurst, J. E.; Kelly, B. T.; Brown, R. G. An examination of surface condition and pore volume of graphite specimens irradiated at high fast neutron doses into volume expansion; UKAEA, SL-CON-2; Sept 1977.
3
Dimensional change (%)
2 1 0 −1 −2 −3 −4 −5
0
50
100
150
200
250
300
Fluence (1020 n cm−2 EDND) NA
NA oxidized, x<9
NA oxidized, 9<x<12
NA oxidized, 12<x<21
NA oxidized, 21<x<22
NA oxidized, 22<x<33
PV (NA unimpregnated, x~8)
Simmons (curve A)
Figure 37 Dimensional changes in preoxidized samples. Modified from Brocklehurst, J. E.; Edwards, J. The fast neutron-induced changes in dimensions and physical properties of near-isotropic graphites irradiated in DFR; UKAEA, TRG Report 2200(S); 1971.
neutron fluence. Some of these specimens, referred to as ‘preoxidized,’ were reirradiated in an inert atmosphere in other MTRs, including DFR, with some achieving reasonably high irradiation fluence. Figure 37 gives the dimensional change behavior of some of these experiments.68
The results show a clear correlation between preoxidized weight loss and dimensional change behavior, indicating an increased dimensional change and delay in turnaround with increased preoxidized weight loss. This has clear implications for the AGRs and required further investigation. Unfortunately, the
Graphite in Gas-Cooled Reactors
20–120 C but, for increased accuracy, it is better to measure CTE over a much larger temperature range. For use in graphite component assessment, the CTE needs to be converted to an appropriate temperature range, that is, room temperature to irradiation temperature (20 Tirr). At irradiation temperatures above 300 C, crystallite CTE, as measured on HOPG,49 is assumed to be invariant to fast neutron fluence. Thus, the irradiation changes in polycrystalline CTE are assumed to be related to closure of Mrozowski-type cracks.
graphite MTR experiments designed to carry out simultaneous radiolytic oxidation and fast neutron damage under power reactor conditions were abandoned because of the closure of the UKAEA MTRs at Harwell in 1990. It is therefore unclear how significant this behavior is for the AGRs. However, there are now MTR experiments being undertaken in HFR (High Flux Reactor) at Petten, the Netherlands to try and address this. 4.11.14.4 Dimensional Change Rate The constitutive models used to predict stresses in polycrystalline components often do not use dimensional change directly but use dimensional change rate. The dimensional change rate and dimensional change of Gilsocarbon graphite irradiated at 550 C are compared in the schematic shown in Figure 38. The turnaround in dimensional change rate occurs earlier than turnaround in dimensional change. In channel-type reactors such as an AGR or Magnox reactor, it is the turnaround in rate that is associated with the peak inbore stress. Thus, when planning a nuclear graphite MTR experiment, it is important to obtain data in the low to medium fluence range, as well as at high fluence.
4.11.15.1
As previously discussed, PGA graphite is an anisotropic material, and this is also reflected in the CTE irradiation data. CTE measurements on irradiated PGA from an MTR program are given in Figure 39. Unfortunately, there is a considerable amount of scatter in this data, probably related to the difficulty in measuring CTE on irradiated graphite at the time these measurements were taken. However, there are some interesting trends illustrated in Figure 39. At the lower irradiation fluence and irradiation temperature below 300 C, there is a rapid increase in CTE, which in the case of the perpendicular (AG) direction, is followed by a rapid reduction. This is reminiscent of the irradiation CTE behavior of HOPG in the ‘c’ direction when irradiated below 300 C.35 At the higher irradiation temperatures (above 300 C), the CTE first increases and then increases again. The initial increase has been associated with closure of Mrozowski cracks due to ‘c’ axis crystal growth.
4.11.15 Coefficient of Thermal Expansion The mechanism that drives the changes to the CTE is, at present, not well understood because of the lack of microstructural studies in this area. In the United Kingdom, CTE data is usually quoted over the range
Dimensional change (%)
0.5
Pile Grade A
0.08
Dimensional change Dimensional change rate
0.0
0.06
−0.5
0.04
−1.0
0.02
−1.5
0.00
−2.0
−0.02
−2.5
−0.04
−3.0 0
50
100
150
−0.06 200
Dimensional change rate (% per 1020 n cm−2 EDND)
366
Fluence (1020 n cm−2 EDND)
Figure 38 Schematic of dimensional change and dimensional change rate in Gilsocarbon graphite irradiated at 550 C.
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367
3.0 150 °C 200 °C 225 °C 250 °C 300 °C 350 °C 450 °C 650 °C
CTE (10-6 K-1)
2.5 2.0 1.5 1.0 0.5 0.0 0
10
(a)
20
30 40 50 Fluence (10 20 n cm−2 EDND)
60
7
150 °C 200 °C 250 °C 300 °C 350 °C 450 °C 650 °C
6 CTE (10-6 K-1)
70
5
4
3
2
(b)
0
10
20
30
40
50
60
Fluence (1020 n cm−2 EDND)
Figure 39 Changes in the coefficient of thermal expansion of pile grade A graphite with irradiation. (a) parallel to extrusion and (b) perpendicular to extrusion. Modified from Birch, M.; Brocklehurst, J. E. A review of the behaviour of graphite under the conditions appropriate for the protection of the first wall of a fusion reactor; UKAEA, ND-R-1434(S); 1987.
The fall in CTE in the perpendicular direction has been attributed to the generation of new cracks caused by the large growth of the crystallites. Electron micrographs due to Sutton and Howard12 appear to support the first assumption. Unfortunately, there is no evidence for the latter assumption. 4.11.15.2 Gilsocarbon Irradiation CTE data for semi-isotropic Gilsocarbon is given in Figure 40. This later CTE data has much less scatter. As with PGA, there is an initial rise in CTE, attributed to the closure of Mrozowski-type cracks, followed by a fall with increasing fluence. Whether this is the same mechanism as the perpendicular PGA data is unclear. However, this behavior is typical of medium-grained nuclear graphite. At very high fast neutron fluence, the CTE appears to start to saturate at a lower value than in the case of
virgin Gilsocarbon, although there is insufficient data to confirm this behavior. 4.11.15.3 Methodology for Converting Between Temperature Ranges Using Simmons’s relationship,61 the instantaneous CTE at two different temperatures represented by ax and a0x can be written as ax ¼ Ax ac þ ð1 Ax Þaa a0x ¼ Ax a0c þ ð1 Ax Þa0a
½37
Rearranging leads to the expression ax aa 0 ða0c a0a Þ þ a0a ¼ Ai ax þ Bi ½38 ax ¼ ac aa or for mean CTE að20Ti Þ ¼Ai að20120Þ þBi
½39
368
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7 430 ⬚C 600 ⬚C 940 and 1240 ⬚C
6
CTE (10−6 K−1)
5 4 3 2 1 0 0
50
100 150 Fluence (1020 n cm−2 EDND)
200
250
Figure 40 Irradiation coefficient of thermal expansion data for Gilsocarbon graphite. Modified from Birch, M.; Brocklehurst, J. E. A review of the behaviour of graphite under the conditions appropriate for the protection of the first wall of a fusion reactor; UKAEA, ND-R-1434(S); 1987; Brocklehurst, J. E.; Kelly, B. T. Carbon 1993, 31(1), 155–178.
6
CTE (10−6 K−1)
5
4
3
Inert Interchange (inert and oxidizing) Continuous oxidation
2 0
10
20
30 40 50 Weight loss (%)
60
70
80
Figure 41 Coefficient of thermal expansion of Gilsocarbon graphite irradiated in inert and oxidizing atmospheres.
Experimental data for aa and ac as a function of temperature can be used to calculate the temperature dependence of A and B. The method has been compared to other empirical approaches used generally in industry and was found to give similar results.69 However, the authors have never seen the method validated for irradiated, radiolytically oxidized graphite. 4.11.15.4 Effect of Radiolytic Oxidation on CTE Figure 41 shows data for the combined fast neutron irradiation and radiolytic oxidation of Gilsocarbon graphite over a limited fluence range, which appears to show no effect of radiolytic oxidation on CTE.
This is also supported by data on thermally oxidized (up to 60% weight loss) PGA and Gilsocarbon graphite70 (see Figure 42).
4.11.16 Thermal Conductivity Thermal conductivity in nuclear graphite is usually determined by measuring thermal diffusivity using the laser flash method at 30 C. The mechanism for thermal conductivity in graphite over the temperatures of interest in nuclear reactors is lattice vibration (phonon) conductance. There is a pronounced reduction in thermal conductivity with increased temperature attributed to phonon– phonon scattering. At low irradiation fluence, there
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369
9 8
CTE (10−6 K−1)
7 6 5 4 3 2 Gilsocarbon PGA (perpendicular to extrusion) PGA (parallel to extrusion)
1 0 0
10
20
40 30 Weight loss (%)
50
60
70
Figure 42 Coefficient of thermal expansion of thermally oxidized graphite. Modified from Hacker, P. J.; Neighbour, G. B.; McEnaney, B. J. Phys. D Appl. Phys. 2000, 33, 991–998.
is a significant decrease in thermal conductivity due to fast neutron irradiation, attributed to an increase in scattering in the damaged lattice. This decrease in thermal conductivity (or increase in thermal resistivity) saturates in the medium fluence range. At very high fluence, there is a secondary decrease attributed to microcracking due to high crystallite strains. There is also evidence of change in temperature dependence with irradiation. The thermal conductivity in crystallite basal plane is much larger than that perpendicular to basal plain; thus, Ka Kc. The thermal resistivity can be described by the equation below: 1 1 1 1 ¼ þ þ Ka K B KU KD
½40
U – Umklapp scattering (German for turnover/ down) or phonon–phonon scattering (due to increase in temperature) D – scattering due to defects (caused by irradiation) B – boundary scattering (structural effects) Changes to these resistances will be reflected in the thermal conductivity of polycrystalline graphite. Thermal conductivity is significantly decreased by radiolytic oxidation. Data is usually presented as the reciprocal of conductivity, that is, thermal resistivity. 4.11.16.1 Pile Grade A Although PGA is significantly anisotropic over the range of interest to the Magnox reactors, the change in thermal resistivity can, for practical purposes, be considered as invariant to grain direction. Changes
in thermal resistivity in PGA graphite are given in Figure 43. There is a significant change in the rate of increases in thermal resistivity between 250 and 300 C, giving a similar trend to the change in crystal growth rates between these two temperatures. The increase in thermal resistivity is significant (a factor of 100) at 150 C, for a relatively low fluence. The low irradiation temperature data for PGA in Figure 43 do not reach a high enough fluence to saturate. However, at the higher temperatures, data is near saturation. 4.11.16.2
Gilsocarbon
The irradiation-induced changes in the thermal resistivity of Gilsocarbon graphite are isotropic, as seen in Figure 44. This data extends to a much higher fluence than the PGA data, and the secondary increase in thermal conductivity, associated with cracking of the microstructure by the large swelling of the crystallites is clearly seen. This behavior occurs after dimensional change turnaround. 4.11.16.3 Thermal Conductivity Temperature Dependence of Irradiated Graphite The temperature dependence of irradiated HOPG parallel to the basal plane (the ‘a’ direction) was measured by Taylor et al.55 When the data is normalized to 30 C, a single curve was produced, the socalled temperature-dependent ‘d’ curve: dðT Þ ¼
Ka;i ð30Þ Ka;i ðT Þ
½41
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Fractional change in thermal resistivity (K0/K−1)
120 150 ⬚C 200 ⬚C 225 ⬚C 250 ⬚C 300 ⬚C 350 ⬚C 450 ⬚C 650 ⬚C
100
80
60
40
20
0 0
5
10
15
20
25
30
35
40
45
50
Fluence (1020 n cm−2 EDND)
Figure 43 Changes in thermal resistivity of pile grade A graphite. Modified from Birch, M.; Brocklehurst, J. E. A review of the behaviour of graphite under the conditions appropriate for the protection of the first wall of a fusion reactor; UKAEA, ND-R-1434(S); 1987.
Fractional change in thermal resistivity (K0 / K−1)
12
10
8
6
4 430 ⬚C 600 ⬚C 940 ⬚C 1240 ⬚C
2
0 0
50
100
150
200
250
Fluence (1020 n cm−2 EDND)
Figure 44 Changes in thermal resistivity of Gilsocarbon graphite. Adapted from Birch, M.; Brocklehurst, J. E. A review of the behaviour of graphite under the conditions appropriate for the protection of the first wall of a fusion reactor; UKAEA, ND-R-1434(S); 1987.
This curve is used in the United Kingdom to predict the thermal conductivity of irradiated polycrystalline graphite at the required assessment temperature. However, above around 300 C, ‘d’ can be taken as unity for all practical purposes. 4.11.16.4 Predicting the Thermal Conductivity of Irradiated Graphite for Reactor Core Assessments A method derived by Kelly71 is often used in the prediction of irradiated graphite thermal conductivity
in the assessment of UK reactor cores. The thermal resistivity, 1/Ki(T), induced because of irradiation as a function of temperature, T, can be given by the difference between the thermal resistivity due to fast neutron damage, and the unirradiated thermal resistivity: 1 1 1 ¼ Ki ðT Þ Kirr ðT Þ K0 ðT Þ
½42
However, dðT Þ ¼
Ka;i ð30Þ Ka;i ðT Þ
½43
Graphite in Gas-Cooled Reactors
and as Ka Kc, Ka is assumed to dominate in polycrystalline graphite, Ka,i (T) ¼ Ki (T) and Ka,i (30) ¼ Ki(30) ¼ Kirr(30), and hence, 1 1 1 ¼ þ dðT Þ Kirr ðT Þ K0 ðT Þ Ki ð30Þ
½44
The irradiation-induced fractional changes in thermal resistivity measured at 30 C are available for graphite irradiated at various temperatures, Tirr. Therefore, for graphite irradiated and measured at 30 C, we can write K0 ð30Þ K0 ð30Þ ¼ 1 ¼f ½45 Ki ð30Þ Kirr ð30; 30Þ Substituting in the previous equation then gives 1 1 f K0 ðT Þ ¼ þ dðT Þ or Kirr ðT Þ K0 ðT Þ K0 ð30Þ Kirr ðT Þ ¼ 1 þ dðT Þf
K0 ðT Þ K0 ð30Þ
½46
Thus, the irradiated thermal conductivity can be predicted at any temperature for graphite irradiated in an inert atmosphere. The effect of weight loss and high fast neutron fluence is dealt with by a version of the ‘product’ rule leading to 1 1 K0 ð30Þ K0 ¼ þ dðT Þf Sk ½47 K ox Kirr ðT Þ K0 ð30Þ K0 ðT Þ
The addition in these two factors in this way appears to be arbitrary as the author does not know of any validation of this method. Other methods do exist, but they depend on many more measurements of thermal conductivity at various irradiation and measurement temperatures.72,73
4.11.17 Young’s Modulus The unirradiated stress–strain behavior of graphite is nonlinear and exhibits hysteresis and permanent set. It is different in tension to compression (Figure 45) and graphite is also much stronger in compression than in tension. Similar curves can be found for Gilsocarbon.74 On irradiation, in an inert atmosphere, there is a rapid and significant increase in modulus attributed to pinning of dislocations in the basal plane. Also, the stress–strain behavior becomes almost linear (Figure 46). This increase soon saturates, but there is a secondary increase attributed to structure tightening (or closure of porosity due to high crystal strain). Finally, at very high fluence, there is a rapid fall in modulus due to the degeneration of the graphite microstructure. Brown75 also showed that the Vickers hardness of isotropic graphite was considerably increased by irradiation. Young’s modulus is significantly reduced by radiolytic oxidation. Graphite Young’s modulus increases with increasing temperature (Figure 47), which is attributed to the tightening of the structure, presumably because of the closure of microcracks; Maruyama et al.76 tested samples in vacuum to avoid thermal oxidation.
10
25
8
20 Stress (MN m -2)
Stress (MN m -2)
where Sk is the high dose reduction in thermal conductivity and [K0/K]ox is the reduction in thermal conductivity due to radiolytic oxidation.
6
4
15
10
5
2
Loading Unloading
Loading Unloading 0 0.00 (a)
371
0.04 0.08 Strain (%)
0 0.0
0.12 (b)
0.2
0.4
0.6
Strain (%)
Figure 45 Stress–strain curves of unirradiated pile grade A graphite. (a) Tension and (b) Compression. Modified from Losty, H.; Orchard, J. In Fifth Conference on Carbon; Pergamon, 1962; pp 519–532.
372
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16
Irradiated (1.0 ⫻ 1019 n cm-2 EDND at 50 ⬚C) -2 14 Failure stress = 47 MN m
Stress (MN m-2)
12 10 Unirradiated failure stress = 23 MN m-2
8 6 4 2 0 0
0.02
0.04
0.06
0.08 Strain (%)
0.1
0.12
0.14
0.16
Figure 46 Change in the stress–strain behavior of graphite due to fast neutron irradiation. Modified from Brocklehurst, J. E. Chem. Phys. Carbon 1977, 13, 145–272.
However, note the significant difference in strength at room temperature between samples tested in air and in vacuum. Several other authors have reported similar findings, attributing the difference to adsorbed moisture.77 The increase in strength with temperature is significant above 600 C, making it of interest only for HTR reactor components. Also of interest in Figure 47 is the correlation between modulus and strength. 4.11.17.1 Relationship Between Static and Dynamic Young’s Modulus Irradiated and unirradiated Young’s modulus of nuclear graphite is usually measured either by an impulse or frequency method, giving the dynamic Young’s modulus (DYM). However, for use in component stress analysis assessments, the static Young’s modulus (SYM) is required. In addition, the UK irradiation creep modulus was originally defined using SYM, so there is a need to interconvert between the two measurements. DYM is always higher in value than SYM. Unfortunately, SYM has historically been defined in various ways, that is, as the chord between the origin at zero stress and half, or in some cases two-thirds the failure strength in tension or bending. There is limited amount of data on the SYM/ DYM ratio for UK nuclear graphite, as given in Figure 48.78 The data is for unirradiated graphite and for graphite irradiated in an inert atmosphere. The ratio for irradiated graphite is higher than for unirradiated graphite. At present, there are no
published data on this ratio for radiolytically oxidized graphite, although safety case requirements have rekindled effort in this area in the United Kingdom. It has been shown79 that for fine-grained IG-110 nuclear graphite and PGX reflector graphite, as well as ASR-ORB baked carbon, the ratio of static to dynamic Young’s modulus strongly depends on the chord length chosen to define SYM. 4.11.17.2
Pile Grade A
The orthotropic values of Young’s modulus in PGA graphite are seen in the changes due to fast neutron irradiation; see Figure 49.64 All of these measurements are DYM measured on small samples irradiated in an MTR. Again, there is a distinct difference in the behavior below 300 C compared to that above 300 C. At low temperature (below 300 C) and fluence, the initial increase attributed to pinning reaches a peak and then rapidly decreases. This rapid decrease has been attributed to large crystallite ‘c’ axis expansion, although there are no microstructural observations to verify this postulation. There is considerable scatter in the data, which may reflect the measurement technique at that time. In some temperature ranges, the trends are defined by very few points. 4.11.17.3
Gilsocarbon
Gilsocarbon graphite data extends to a much higher fluence; see Figure 50.50,64 For assessment purposes, the initial increase or pinning is normally assumed to
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373
70
Flexural strength (MN m-2)
60 50 40 30 20 10
In air In vacuum
0 0
200
400
(a)
600
800
1000
1200
1400
Temperature (⬚C)
Relative change in sf/sf0 and E/E0
1.2 1.0 0.8 0.6 0.4 0.2
Flexural strength Elastic modulus
0.0 0
200
(b)
400
600
800
1000
1200
1400
1600
Temperature (⬚C)
Figure 47 The strength and modulus of IG-11 as a function of test temperature. (a) Flexural strength and (b) Flexural strength and modulus. Reproduced from Maruyama, T.; Eto, M.; Oku, T. Carbon 1987, 25(6), 723–726.
Static Young⬘s modulus (GN m−2)
30 25 20 15 10 5
Unirradiated Irradiated
0 0
5
10 15 20 Dynamic Young⬘s modulus (GN m−2)
25
30
Figure 48 Ratio between static and dynamic Young’s modulus. Modified from Brocklehurst, J. E.; Brown, R. G. The relation between strength, modulus, and structural changes of isotropic graphites irradiated to high fast neutron doses in DFR; UKAEA, TRG-M-5985 (AB 7/22015); 1972.
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Fractional change in Young’s modulus (E/E0−1)
374
1.8 1.6 1.4 1.2 1.0 0.8 150 ºC 200 ºC 250 ºC 300 ºC 350 ºC 450 ºC 650 ºC
0.6 0.4 0.2 0.0 0
5
10
Fractional change in Young’s modulus (E/E0−1)
(a)
15 20 25 Fluence (1020 n cm-2 EDND)
30
35
40
2.0 1.8 1.6 1.4 1.2 1.0 0.8 0.6 150 ºC 200 ºC 250 ºC 450 ºC 650 ºC
0.4 0.2 0.0 0
5
(b)
10 15 20 Fluence (1020 n cm-2 EDND)
25
30
Figure 49 Changes in Young’s modulus of pile grade A due to fast neutron irradiation. (a) Parallel to extrusion and (b) Perpendicular to extrusion. Modified from Birch, M.; Brocklehurst, J. E. A review of the behaviour of graphite under the conditions appropriate for the protection of the first wall of a fusion reactor; UKAEA, ND-R-1434(S); 1987.
occur instantaneously on irradiation. After saturation, there is a significant increase in modulus with increase in fluence followed by a reduction. This behavior is often referred to as a ‘structural’ effect. At very high fluence, the decrease in modulus due to degradation of the microstructure is clearly illustrated (Figure 50). 4.11.17.4 Separation of Structure and Pinning Terms Various proposed graphite irradiation models, including the UKAEA irradiation creep models, require the pinning and structural terms to be separated. Originally, the pinning and structure terms for Gilsocarbon were separated by ‘eye,’ which is subjective. Recently, more sophisticated statistical, pattern recognition and curve fitting methods have been used.60
Having separated out the pinning term, the structure term is a function of irradiation temperature and dose, whereas the pinning term is only a function of temperature. The pinning term is assumed to not affect creep rate, whereas the structure term affects it. In some models, the structure term is also assumed to be a function of radiolytic weight loss and to be correlated to dimensional change. 4.11.17.5 Effect of Radiolytic Weight Loss on Dimensional Change and Young’s Modulus Dimensional change and modulus MTR data on small preoxidized samples, see Figure 51, appear to indicate that radiolytic weight loss would be expected, not only to increase dimensional change shrinkage and delay turnaround, but also to delay the structural
Fractional change in Young’s modulus (E/E0−1)
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3.5
375
430 ⬚C 600 ⬚C 940 ⬚C 1240 ⬚C 1430 ⬚C
3.0 2.5 2.0 1.5 1.0 0.5 0.0 −0.5 −1.0 0
50
100 Fluence
150 (1020
n cm−2
200
250
EDND)
Figure 50 Changes in Young’s modulus of Gilsocarbon due to fast neutron irradiation. Modified from Birch, M.; Brocklehurst, J. E. A review of the behaviour of graphite under the conditions appropriate for the protection of the first wall of a fusion reactor; UKAEA, ND-R-1434(S); 1987; Brocklehurst, J. E.; Kelly, B. T. Carbon 1993, 31(1), 155–178.
increase in Young’s modulus. This correlation has been used as the basis of models to predict dimensional change in radiolytically oxidized graphite.51 4.11.17.6 Small Specimen Strength Graphite is stronger in bend than in tension, and stronger in compression than in bend. Irradiated strength tends to be correlated with Young’s modulus. Graphite strength is significantly reduced by radiolytic oxidation. Losty and Orchard80 used the Griffith theory to try and demonstrate that the change in strength can be related to the square root of modulus as follows: Failure stress s is proportional to the square root of the product of strain energy release rate g and Young’s modulus divided by the crack length c. 2gE ½48 s2 / pc Thus, assuming that strain energy release rate, g, is not changed by irradiation and the critical crack, 1=2 s E ¼ ½49 s0 E0 Figure 52 was purported to support this relationship. However, statistical scrutiny of the data given in this figure revealed that there is not enough data to support the argument in favor of the square root law, or even a relationship to another power. There is a significant amount of data indicating that the relationship may be more appropriate as a direct relationship.81
4.11.18 Effect of Radiolytic Oxidation on Thermal Conductivity, Young’s Modulus, and Strength Figure 53 illustrates that thermal conductivity, strength, and Young’s modulus are all significantly reduced by radiolytic oxidation.82 The data is usually fitted to a simple exponential decay83 of the form [P/P0] ¼ exp(lx). However, there must be a practical ‘percolation limit’ to this law when all the porosity joins together and properties reduce to zero. This was recently highlighted in the statistical analysis of high weight loss PGA data by McNally et al.84
4.11.19 The Use of the Product Rule The product rule has been used for many years in the United Kingdom to combine changes due to fast neutron irradiation and radiolytic oxidation for strength, modulus, and thermal conductivity. Examples are given below: Strength s s s ¼ ½50 s0 s0 irr s0 oxidation Young’s modulus E E E E ¼ E0 E0 P E0 S E0 oxidation Thermal conductivity
½51
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1
Dimensional change (%)
0 −1 −2 −3 −4
0% weight loss 5.8% weight loss 13.4% weight loss 23.1% weight loss 38.4% weight loss Trend curve A Trend curve B
−5 −6 −7
0
20
40
(a)
100 60 80 Fluence (1020 n cm−2 EDND)
120
140
160
Young’s modulus structure term
2.0 1.8 1.6 1.4 0% weight loss 5.8% weight loss 13.4% weight loss 23.1% weight loss 38.4% weight loss
1.2 1.0 0.8 0
(b)
20
40
60
80
100
120
140
160
Fluence (1020 n cm−2 EDND)
Figure 51 Correlations between dimensional change and Young’s modulus structure term in Gilsocarbon. (a) Dimensional changes in pre-oxidized Gilsocarbon and (b) Young’s modulus structure terms in pre-oxidized Gilsocarbon. Modified from Schofield, P.; Brown, R. G.; Daniels, P. R. C.; Brocklehurst, J. E. Fast neutron damage in Heysham II/Torness moderator graphites (final report on 3-temperature zone rig); UKAEA, NRL-M-2176(S); 1991.
1 1 K0 ð30Þ K0 ¼ þ dðT Þf Sk ½52 K ox Kirr ðT Þ K0 ð30Þ K0 ðT Þ In recent years, it has been realized that the use of the product rule is simplistic, and most probably, only applicable for low irradiation data, up to a fluence not far beyond dimensional change turnaround and only for relatively low weight loss. Therefore, there has been a recent trend to use empirical fits to reactor or MTR data where available.
4.11.20 Irradiation Creep in Nuclear Graphite By the late 1940s, it was known that graphite components, when subjected to fast neutron irradiation, suffered significant dimensional change. It was thought that, because of the flux gradient across the brick section, these dimensional changes would generate significant stresses in hollow graphite moderator blocks and that this would lead to significant
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4.5
Fractional change in s/s0
4.0 3.5 3.0 2.5 2.0 1.5 1.0 0.5 0.0 0
50
100
150
200
250
300
350
Fluence (1020 n cm−2 EDND) Gilsocarbon (300–400 ⬚C) Gilsocarbon (560 ⬚C) Petroleum coke (300–440 ⬚C)
Petroleum coke (560 ⬚C) Pitch coke (300–440 ⬚C) Unidentified coke (300–440 ⬚C)
Unidentified coke (560 ⬚C) E/E0 Ö(E/E0)
Figure 52 Change in strength of irradiated Gilsocarbon graphite compared to change in modulus. Modified from Brocklehurst, J. E. Chem. Phys. Carbon 1977, 13, 145–272.
Thermal conductivity (k/ko) Young’s modulus (E/Eo)
Relative change in property
1
Tensile strength (s/so) Compressive strength (s/so)
0.5 e−2.8x
e−3.8x
e-5.1x
0 0.0
0.1
0.2 Porosity
0.3
0.4
Figure 53 Change in thermal conductivity, Young’s modulus, and strength due to radiolytic oxidation. Modified from Adam, R. W.; Brocklehurst, J. E. Mechanical tests on graphite with simulated radiolytic oxidation gradients; UKAEA, ND-R-853(S) (AB 7/26300); 1983.
component failures within a few years of reactor operation. Therefore, the fuel channels in the early reactors, such as the Windscale Piles, were designed to avoid the buildup of stress. By the 1950s, it was realized that there was an irradiation-induced mechanism that was relieving
stresses generated by dimensional change and the term ‘irradiation induced plasticity’85,86 was coined to describe this mechanism. Later, around 1960, the term ‘irradiation creep’87 started to be used for the difference between dimensional change in loaded and unloaded graphite specimens irradiated to the same dose.
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4.11.20.1 Dimensional Change and Irradiation Creep Under Load Compressive stress increases, and tensile stress decreases, the irradiation-induced dimensional change of graphite as illustrated in Figure 54. In these experiments, two matching graphite samples, a loaded specimen and an unloaded ‘control’ specimen, were irradiated adjacent to each other in an MTR, and dimensional change in the direction of load was measured. However, as well as change in dimension in the load direction, there are also dimensional changes perpendicular to the load direction as shown in Figure 55. An irradiation creep curve can be simply obtained by subtraction of the unloaded dimensional change curve from the crept dimensional change curve, as illustrated in Figure 56. However, for practical use in assessments, this would require data for a range of temperatures and fast neutron fluences covering all of the expected operating conditions. Also, in the case of carbon dioxide-cooled reactors, the effect of
Dimensional change (%)
0 −1 −2 −3 Reference Loaded
−4 0
50
100
150
200
250
Fluence (1020 n cm−2 EDND)
(a)
Dimensional change (%)
1 0 −1 −2
4.11.20.2 Types of Irradiation Creep Experiments There are three categories of graphite creep experiments. The first type was the restrained creep experiments. In these experiments a graphite specimen, usually dumbbell in shape, is restrained from shrinkage by a tube or split collar manufactured from graphite that shrinks less than the specimen of interest. In the case of anisotropic graphite, the tube or split collar could be manufactured with its longitudinal axis aligned with the more dimensionally stable grain direction; the specimen would be manufactured with its axis perpendicular to this. These types of experiments are relatively easy to deploy but are difficult to assess as the load is not directly measured and a ‘creep law’ has to be assumed in the assessment of the results.88 A variation on these experiments was the graphite spring tests used in Calder Hall to define primary creep.86 A second important type of experiment was the ‘out-of-pile measurements’ technique.89,90 These experiments, importantly, give ‘real-time results.’ However, this type of experiment is difficult to install in a reactor and results are obtained only for one specimen. The final type of experiment is the in-pile rig loading using a string of samples and usually taking advantage of the MTR flux gradient to obtain data on samples to various levels of fluence. There have been various designs of simple strings of specimens loaded either in tension or in compression. Tensile creep tests are vulnerable, that is, if one specimen fails, results for the whole string of specimens could be lost. However, there have been various rig designs aimed at overcoming this problem.
−3 Reference Loaded
−4 0
(b)
the rate of radiolytic oxidation on creep rates would have to be quantified and understood. In addition, changes to the CTE and Young’s modulus with irradiation creep have been observed, which further complicate assessment technology.
50
100
150
200
250
Fluence (1020 n cm−2 EDND)
Figure 54 Dimensional changes of loaded ATR-2E graphite. (a) compression and (b) tension. Modified from Haag, G. Properties of ATR-2E Graphite and Property Changes due to Fast Neutron Irradiation; FZJ, Ju¨l-4183; 2005.
4.11.20.3
The UKAEA Creep Law
Irradiated creep experiments were carried out between 350 and 650 C on both PGA and Gilsocarbon graphite.91 Some low fluence experiments were also carried out in Calder Hall86 which were used to define the so-called ‘primary creep.’ Creep strain data ecr was normalized to elastic strain units (esu) by dividing by the applied stress (s) and multiplying by the
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379
3
Dimensional change (%)
2
1
0
−1 Reference Loaded
−2 0
50
100 Fluence
150 (1020
n cm−2
250
200
EDND)
Figure 55 Lateral dimensional changes of loaded ATR-2E graphite. Modified from Haag, G. Properties of ATR-2E Graphite and Property Changes due to Fast Neutron Irradiation; FZJ, Ju¨l-4183; 2005.
2.0
Creep strain (%)
1.5
1.0
0.5
Compressive Tensile
0.0 0
50
100
150 20
Fluence (10
−2
n cm
200
250
EDND)
Figure 56 Irradiation creep curves for ATR-2E at 500 C. Modified from Haag, G. Properties of ATR-2E Graphite and Property Changes due to Fast Neutron Irradiation; FZJ, Ju¨l-4183; 2005.
unirradiated SYM (E0) as given below: esu ¼
E0 ecr s
½53
Surprisingly, they found that by doing this, the creep data for these two types of graphite, with very different microstructures, could be fitted to a simple ‘creep equation’ of the form s s ½54 ecr ¼ ½expð4gÞ þ 0:23 g E0 E0
where g is the fast neutron dose. This is illustrated in Figure 57. The primary creep strain is assumed to be recoverable on removal of the load while still under irradiation. Some evidence for this came from out-of-pile measurement experiments such as the FLACH experiments.90 However, if the specimens had been left unloaded for longer duration, more than 1 esu may have been recovered. In addition to this, an experiment carried out on precrept samples, that is,
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1.2
Creep strain (%)
1.0 0.8 0.6 BR–2 (300 – 650 ⬚C, 900 psi)
0.4
Isotropic graphites (compressive) Isotropic graphites (tensile) PGA (perpendicular) PGA (parallel) Pyrolytic graphites
0.2 0.0 0
10
20
30
40
50
60
70
Fluence (1020 n cm−2)
(a)
16 14
Elastic strain units
12 10 8 6
Isotropic graphites (BR-2, 300 – 650 ⬚C) PGA (BR-2 (300 – 650 ⬚C, perpendicular) PGA (BR-2 (300 – 650 ⬚C, parallel) Pyrolytic graphites (BR-2, 300 - 650 ⬚C) PGA graphites (PLUTO, 300 ⬚C) Calder hall (140 – 350 ⬚C)
4 2 0 0
10
20
(b)
30
40
50
60
70
Fluence (1020 n cm−2)
Figure 57 UK creep data used to define the UK creep law. (a) creep strain and (b) elastic strain units. Modified from Brocklehurst, J. E. Irradiation Damage in CAGR Moderator Graphite; Northern Division, UKAEA, ND-R-1117(S); 1984.
samples of PGA and Gilsocarbon irradiated in a creep experiment in the BR-2 reactor and then irradiated with the load removed in DIDO and DFR respectively, exhibited more than 1 esu (a recovery in the region of 6 esu in the case of Gilsocarbon in DFR). 4.11.20.4 Observed Changes to Other Properties 4.11.20.4.1 Coefficient of thermal expansion
Significant differences have been observed between the unstressed CTE and stressed CTE, as illustrated in Figure 58. Compressive creep strain was found
to increase the CTE, and tensile creep strain to decrease the CTE. The changes in CTE caused by irradiation creep have similarities to those caused by the application of stress on unirradiated graphite. Figure 59(a) shows the changes in CTE in irradiated, crept specimens plotted as a function of creep strain92 and Figure 59(b) gives the changes in CTE in unirradiated graphite due to stress.93 At room temperature, the average CTE of an isotropic graphite with no porosity should be the average of the crystallite CTEs, that is, the crystallite CTEs are 27.0 106 K1 and 1.0 106 K1 in the ‘c’ and ‘a’ directions respectively, giving an average
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381
6
CTE (10−6 K−1)
5
4
3
Unstressed Compressive Tensile
2
0
10
20
30
40
50
60
Fluence (1020 n cm−2 EDN)
(a)
Change in CTE (10−6 K−1)
6
2
1
IM1-24 (parallel) SM2-24 (parallel) SM2-24 (perpendicular) VNMC (parallel) VNMC (perpendicular)
2 0 (b)
10
20
30 40 Creep strain(%)
50
60
Figure 58 Additional changes to the coefficient of thermal expansion in loaded materials test reactor specimens. (a) Modified from Brocklehurst, J. E.; Brown, R. G. Carbon 1969, 7(4), 487–497 and (b) Modified from Brocklehurst, J. E.; Kelly, B. T. A review of irradiation induced creep in graphite under CAGR conditions; UKAEA, ND-R-1406(S); 1989.
of 8.0 106 K1. From Figure 59(a), which is for Gilsocarbon with an unirradiated CTE of 4.0 106 K1, it is interesting to note that the increase in CTE in compression is approaching that value. 4.11.20.4.2 Young’s modulus
The early BR-2 experiments showed little evidence of change in modulus with irradiation creep. However, other authors94 did find evidence of a change (Figure 60). Later, higher temperature creep experiments carried out in the United Kingdom92 also showed a change in Young’s modulus. Figure 61(a) shows the fractional change in Young’s modulus E/E0, as a function of different compressive stresses. In Figure 61(b), the data has been normalized by dividing the crept Young’s modulus Ec by the irradiated value of the unstressed specimen Ei. 4.11.20.5 Lateral Changes Creep in metals is purported to be related to plastic flow occurring at constant volume. However, artificial
polycrystalline graphite is a porous material with a crystal structure considerably different from that of metals. It is therefore not surprising that in irradiation creep, graphite does not deform with a constant volume. As with much of the irradiation creep data, the quality of the transverse creep data is poor and inconsistent. However, from the data that does exist, the irradiation creep ratio does appear not to be constant with creep strain (Figure 62). 4.11.20.6
Creep Models and Theories
It is unfortunate that a validated set of graphite irradiation creep data covering the range of temperatures and fluences of interest for power producing reactors, as well as radiolytic oxidation in the case of carbon dioxide-cooled reactors, does not exist. In addition, there are no microstructural studies available to give an insight into the mechanism involved in irradiation creep in graphite. This has lead to much speculation and several model proposals.
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3
Change in CTE (10−6 K−1)
2
1
−5
−4
−3
−2
0
−1
0
1
2
−1
Compression (PLUTO) Compression (BR-2) Tension (BR-2)
−2 Creep strain (%)
(a)
3
Change in CTE (10−6 K−1)
2
1
0
Longitudinal stain gauge results
−1
Longitudinal stain gauge results Transverse strain gauge results
−60
−50
−40
−30
−20
−10
−2 0
10
20
Stress (MN m−2)
(b)
Figure 59 Synergy between changes in coefficient of thermal expansion in irradiation creep specimens and change in unirradiated, stressed graphite. (a) additional change in CTE as a function of creep strain and (b) change in CTE in unirradiated stressed samples.
Change in Young’s modulus (E/E0)
2.0 1.6 1.2 SM1-24 (axial), irradiation temperature = 850- 920 ⬚C
0.8 0.4
0.0 MPa (Capsule 76M-18A)
0.0 MPa (Capsule 77M-10A)
3.3 MPa (Capsule 76M-18A)
3.3 MPa (Capsule 77M-10A)
4.5 MPa (Capsule 76M-18A) 6.5 MPa (Capsule 76M-18A)
4.5 MPa (Capsule 77M-10A) 6.5 MPa (Capsule 77M-10A)
0.0 0
2
4
6
8 24
Fluence (10
10
12
14
16
−2
n cm , E > 29 fJ)
Figure 60 Changes in Young’s modulus in tensile crept and uncrept specimens. Reproduced from Oku, T.; Fujisaki, K.; Eto, M. J. Nucl. Mater. 1988, 152(2–3), 225–234.
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Change in Young’s modulus (E/E0)
0.5
383
0.9 ⫻ 1021 n cm−2 (1050 ⬚C) 1.8 ⫻ 1021 n cm−2 (1050 ⬚C) 1 ⫻ 1021 n cm−2 (850 ⬚C) 2 ⫻ 1021 n cm−2 (850 ⬚C)
0.4 0.3 0.2 Fail
0.1 0.0 0
5
10
15 20 25 30 35 -2 Compressive stress (MN m )
(a)
40
45
50
Creep modulus/elastic modulus
1.1
1.0
1+3(creep modulus)
0.9 1050 ⬚C 850 ⬚C
0.8
-6
-5
-4
(b)
-3 -2 -1 0 Compressive creep strain (%)
1
2
Figure 61 Changes in Young’s modulus in irradiated-creep experiments. (a) changes to Young’s modulus as a function of stress and fast neutron fluence and (b) normalized Young’s modulus as a function of creep strain. Modified from Brocklehurst, J. E.; Kelly, B. T. A review of irradiation induced creep in graphite under CAGR conditions; UKAEA, ND-R-1406(S); 1989.
4.11.20.6.1 UKAEA creep law
In the United Kingdom, to extend the creep law to higher fluence and to account for radiolytic oxidation, the following approach was taken by the UKAEA. Creep strain, decr/dg, was assumed to be defined by decr s s ¼ aðT Þ ½expðbgÞ þ bðT Þ dg Ec Ec
½55
where a(T), b(T ), and b are temperature-dependent functions equal to 1.0, 0.23, and 4.0, respectively in the AGR and Magnox temperature ranges, and s and g are stress and irradiation fluence, respectively. The need for the temperature dependence outside this range was defined by data for HTRs obtained in the United States and Russia (Figure 63).
The ‘creep modulus’ in the UKAEA model was defined as Ec ¼ E0 SE½ox
½56
where E0 is the unirradiated SYM and S is the irradiation temperature- and fluence-dependent structure term derived from the irradiated modulus data (Section 4.11.17.4). To account for radiolytic weight loss, E[ox] is a modulus weight loss term defined as E/E0 ¼ exp(lx) where l is an empirical constant equal to about 4.0 and x is the fractional weight loss. There are no rigorous observational data to underpin this model other than a few data points for preoxidized graphite given in Figure 64. It should be noted that in Figure 64, the only two
384
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0.6
Poisson’s ratio
0.5 0.4 0.3 0.2 SM2-24 VNMC IM1-24
0.1 0.0 0
1
2
(a)
3
4
6
5
Longitudinal creep strain (%)
Transverse creep strain (%)
0.4
0.0 −0.4 −0.8 −1.2 −1.6 0.0
ATR-2E (tension, 500 ⬚C) ATR-2E (compression, 500 ⬚C) H-337 (compression, 550 ⬚C) H-337 (compression, 800 ⬚C) AGOT (compression, 550 ⬚C) AGOT (compression, 800 ⬚C) H-451 (compression, 900 ⬚C)
0.5
1.0
(b)
1.5
2.0
2.5
3.0
3.5
Longitudinal creep strain (%)
Creep coefficient ( kg cm−2)−1 (neutron cm−2)−1
Figure 62 Various transverse irradiation creep data. (a) UKAEA data and (b) US data. Modified from Brocklehurst, J. E.; Kelly, B. T. A review of irradiation induced creep in graphite under CAGR conditions; UKAEA, ND-R-1406(S); 1989.
1E−24
1E−25
1E−26
Russian graphite American graphite (EGCR) American graphite (CGB)
1E−27 0
100
200
300
400 500 600 Temperature (⬚C)
700
800
900
1000
Figure 63 Temperature dependence of the secondary creep coefficient b(T) from US and Russian data.
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12 Compressive Tensile Low density graphite (equivalent to 25% weight loss) UK creep law
Elastic strain units (esu)
10
9.5%
8
6 9.8% 27.8% 4
25.5%
6%
2
0 0
5
10
15
20
25
30
35
40
45
50
Fluence (1020 n cm-2 EDND) Figure 64 Preoxidized irradiation creep data. Modified from Brocklehurst, J. E.; Kelly, B. T. A review of irradiation induced creep in graphite under CAGR conditions; UKAEA, ND-R-1406(S); 1989.
samples with a significant amount of weight loss are irradiated to a relatively low fluence. Similarly, there are some, even less convincing, data on creep samples initially irradiated to high fast neuron fluence before loading.92 The main criticisms of the UKAEA creep model, for inert conditions, is that it gives a very poor fit to the high fluence creep data obtained in Germany and the United States, as discussed in the next section. 4.11.20.6.2 German and US creep model
This model was devised for helium-cooled HTR applications where radiolytic oxidation was of no concern. The form of some of the US and German data is given below (Figure 65). There appears to be a difference between tension and compression at high fluence. However, this is the only data that shows this and it is not clear if it is a real effect. It was assumed that microstructural changes at medium to high fluence would modify the creep rate and account for the shape of these curves. It was assumed that this could be accounted for by modifying the secondary creep coefficient in the UK creep law by the following expression: s ecrðsecondaryÞ ¼ K E0 DV =V 0 0 ½57 K ¼K 1m ðDV =V0 Þm where K0 is the UK creep coefficient, DV/V0 is the change in volume (which is a function of fluence and
irradiation temperature), (DV/V0)m is the volume change at volumetric ‘turnaround,’ and m is a graphite grade and temperature-dependent variable. 4.11.20.6.3 Further modifications to the UKAEA creep law: interaction strain
The theory originally developed by Simmons in the 1960s reported in detail by Hall et al.61 relating the polycrystalline dimensional change rate and CTE with crystallite dimensional change rate, and CTE has been further developed95 in an attempt to explain the shape of the graphite irradiation creep behavior at high dose. The proposed theory argues that if the dimensional change rate in polycrystalline graphite can be related to the CTE, and because irradiation creep has been observed to modify CTE of the loaded specimen differently to that seen in unloaded specimens, changing the CTE by creep would be expected to change the dimensional change rate and hence, the dimensional change in the loaded specimen. This leads to the introduction of the so-called ‘interaction strain.’ The theory behind this methodology is described below. Considering two specimens (a crept specimen and an unloaded control) being irradiated under identical conditions; in the unloaded control specimen, by applying the Simmons equations, the bulk dimensional change rate gx and bulk CTE ax can be defined by gx ¼ ð1 Ax Þga þ Ax gc ax ¼ ð1 Ax Þaa þ Ax ac
½58
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3.0 ATR-2E (stress = 5 Mpa) Tension (900 ⬚C) Tension (300 ⬚C) Tension (500 ⬚C) Compression (500 ⬚C)
Creep strain (%)
2.5 2.0 1.5 1.0 0.5 0.0 0
50
150
100 20
Fluence (10
–2
n cm
200
EDND)
Figure 65 High-fluence German and US data.
where ga and gc are crystal dimensional change rate in the a and c directions, respectively, and aa and ac are the crystal CTE in the a and c directions, respectively. Ax is referred to as the structure factor and by rearrangement ax aa Ax ¼ ½59 ac aa Thus,
gx ¼ ga þ gT
ax aa ac aa
½60
where gT is the crystal shape rate factor and is equal to gc ga. Similarly for the loaded specimen, 0 a aa ½61 gx0 ¼ ga þ gT x ac aa Therefore, the difference between the dimensional change rates of the unloaded and loaded specimen is Da ½62 gx0 gx ¼ gT ac aa or gx0
¼ gx þ gT
Da ac aa
½63
where Da is the change in CTE under load (a0x ax). This leads to the following definitions: The true dimensional change in the loaded specimen ¼ the dimensional change in the control þ the interaction term True creep ¼ dimensional change in loaded specimen true dimensional change in loaded specimen
Apparent creep ¼ dimensional change in loaded specimen dimensional change in control Thus, the interaction term gT acDa aa is included in the finite element analysis of graphite components. The limited data that exists on irradiated HOPG indicates that the dimensional change rate of graphite crystallites increases with increasing fluence in the ‘c’ direction and decreases in the ‘a’ direction for all measured irradiation temperatures and dose range. For irradiation temperatures of 450 and 600 C, the data indicates that ac and aa remain invariant to fluence. However, below 300 C the crystal CTE appears to change. There are no crystal CTE data for higher temperatures. It should also be noted that Simmons equations imply that Ax ¼
ax aa g x g a ¼ ac aa g c g a
½64
Close examination of typical graphite irradiation data, say for Gilsocarbon irradiated in the temperature range where crystal data are available (450 and 600 C), shows that the relationship given above does not hold. In fact, the Simmons relationship and measured data diverge at low dose. This is attributed to Simmons assuming that polycrystalline graphite can be considered as a loose collection of crystallites with no mechanical interactions. Others95 have added an extra ‘pore generation’ term to the Simmons dimensional change relationship to try and reconcile these issues, but again there is no real validation of these models.
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The use of this interaction term did not gain wide (international) acceptance as it appeared to be using the Simmons relationship beyond its applicability and did not explain the difference between compressive and tensile loading at high fluence. 4.11.20.6.4 Recent nuclear industry model
Recently Davies and Bradford96 have developed a far more complex creep model as given below: ðg ak1 k1 g s 0 expk1 g dg0 exp ec ¼ E0 Sðg; T ÞW ðxÞ g0 ¼0
ðg b s 0 dg0 þ 1 expk2 g E0 Sðg; T ÞW ðxÞ g0 ¼0
ok3 k3 g þ exp E0
ðg g0 ¼0
s 0 expk3 g dg0 Sðg; T ÞW ðxÞ
½65
where a, 1 esu (where this is defined by s/E0); k1, 0.0857e(1630.4/T); b, 0.15 esu per 1020 n cm2 EDN; k2, 0.0128e(1270.8/T); o, 5 esu; k3, 0.4066e(1335.9/T); s, stress (P); E0, unirradiated SYM (Pa) appropriate to the stress applied (0.84 DYM); S(g, T), structure term representing structural induced changes to creep modulus (a function of fluence and temperature); W(x), oxidation term representing oxidation-induced changes to creep modulus (a function of weight loss, x, which is a function of fluence). The lateral strain ratio for the primary and recoverable terms is assumed to be equal to the elastic Poisson’s ratio. The lateral strain ratio for secondary creep, nsc, is assumed to follow the relationship nsc ¼ 0:5½1 3Sc ðgÞ Sc is a structural connectivity term that the authors have used in model fits for other graphite property changes.57 This model certainly fits the available inert data better than the previous models, although it cannot be tested against radiolytically oxidizedgraphite data as there is none. 4.11.20.7 Final Thoughts on Irradiation Creep Mechanisms Two main models for the mechanism of irradiation creep have been put forward but neither has any microstructural observations to support them. The first suggestion is that a model by Roberts and Cottrell97 for a-uranium may be appropriate. This model proposes that the graphite crystallite structures will yield and shear because of the generation
387
of stresses caused by dimensional change. However, it is difficult to envisage such a yield and shear mechanism in crystalline graphite. The second model98 suggests that under load, the crystallite basal planes will slide because of a pinning and unpinning mechanism during irradiation. Such a mechanism is described in detail by Was99 with relation to metals and could explain primary creep and secondary linear creep. However, if irradiation creep in graphite is associated with basal plane slip due to pinning–unpinning, it is surprising that in PGA, irradiation creep is less in the WG or parallel to the basal plane direction than it is in the AG or perpendicular to the basal plane direction (Figure 57). Another possibility is that stress modifies the crystal dimensional change rate itself. In support of this are X-ray diffraction measurements100 that showed that the lattice spacing in compressive crept specimens is less than that in the unstressed control specimens (Figure 66). Such a mechanism would explain the PGA data and could be related to the change in CTE and the observed annealing behavior. However, the data and experimental fluence and creep range given are very limited. It is clear that changes to the lattice spacing in crept graphite would be an area worthy of further investigation in future irradiation creep programs. Irradiation creep in the graphite crystallite will be reflected in the bulk deformations observed in creep specimens and in reactor components. Changes to the bulk microstructure due to radiolytic oxidation would be expected to influence this bulk behavior, as would large crystal dimensional changes at very high fluence (past dimensional change turnaround). It would be expected that at very high fluence the behavior of graphites with differing microstructures would diverge; this appears to be the case from the limited high fluence data available.
4.11.21 Concluding Remark Nuclear grade graphite has been used, and is still used, in many reactor systems. Furthermore, it provides an essential moderator and reflector material for the next-generation high-temperature gas-cooled nuclear reactors that will be capable of supplying high-temperature process heat for the hydrogen economy. Hence, nuclear graphite technology remains an important topic. Although there is a wealth of data, knowledge, and experience on the design and operation of graphite-moderated reactors,
388
Graphite in Gas-Cooled Reactors
3.56 Unrestrained Restrained
3.54
Interlayer spacing (Å)
3.52 3.50 3.48 3.46 3.44 3.42 3.40 3.38 3.36 0
2
4
6
8
10
12
14
16
18
20
Fluence (GWd te−1) Figure 66 The effect of stress and irradiation on the interlayer spacing of graphite. Modified from Francis, E. L. Progress Report for the JNPC-Materials Working Party: Graphite Physics Study Group; UKAEA, TRG-M-2854 (AB 7/17604); 1965.
the need for present plants to extrapolate beyond current data and to predict the behavior of new graphite grades operating for longer lifetimes at higher temperatures than before means there is still a substantial amount of work for the graphite specialist. Future understanding and validation of property/ microstructural change relationships that enable the prediction and interpolation of existing databases and the development of new graphite grades is now possible using new characterization, modeling, and computation techniques. These allow the investigation of mechanisms and graphite behavior that were previously impossible or impractical to conduct. Areas of particular interest are obtaining a better understanding of the mechanism of dimensional change and irradiation creep, and the development of a validated graphite failure model.
References 1. 2. 3. 4. 5. 6.
Delle, W.; Koizlik, K.; Nickel, H. Graphitic Materials for Use in Nuclear Reactors: Parts 1 and 2; Karl Thiemig: Munich, 1978. Nightingale, R. E. Nuclear Graphite; Academic Press: New York and London, 1962. Reynolds, W. Physical Properties of Graphite; Elsevier: Amsterdam, 1968. Simmons, J. Radiation Damage in Graphite; Pergamon: London, 1965. Pacault, A. Les Carbones; Masson et Cie: Paris, 1965. Kelly, B. T. Physics of Graphite; Applied Science: London, 1981.
7. Kelly, B. T. In Materials Science and Technology; Kramer, E. J., Ed.; Wiley: New York, 1994; pp 365–417. 8. Burchell, T. D. In Carbon Materials for Advanced Technologies; Timothy, D. B., Ed.; Elsevier Science: Oxford, 1999; pp 429–484. 9. Kretchman, H. F. The Story of Gilsonite; American Gilsonite: Salt Lake City, UT, 1957. 10. Amelinckx, S.; Delavignette, P.; Heerschap, M. Chem. Phys. Carbon 1965, 1, 1–71. 11. Mrozowski, S. Mechanical strength, thermal expansion and structure of cokes and carbons. In 1st and 2nd Conferences on Carbon, University of Buffalo, Buffalo, New York, 1956; pp 31–45. 12. Sutton, A. L.; Howard, V. C. J. Nucl. Mater. 1962, 7(1), 58–71. 13. Bacon, G. E. J. Appl. Chem. 1956, 6, 447–481. 14. Kinchin, G. Variation of graphite damage in the reflector of BEPO and the calculation of damage in graphite moderated piles; AERE, AERE N/PC 20; 1951. 15. Bell, J.; Bridge, H.; Cottrell, A.; Greenough, G.; Reynolds, W.; Simmons, J. Philos. Trans. R. Soc. Lond. A Math. Phys. Sci. 1962, 254(1043), 361–395. 16. Marsden, B. J. Irradiation damage in graphite due to fast neutrons in fission and fusion systems; IAEA, IAEA TECDOC-1154; 2000. 17. Kinchin, G. H.; Pease, R. S. Rep. Prog. Phys. 1955, 18(1), 1–51. 18. Thompson, M. W.; Wright, S. B. J. Nucl. Mater. 1965, 16(2), 146–154. 19. Norgett, M. J.; Robinson, M. T.; Torrens, I. M. Nucl. Eng. Des. 1975, 33(1), 50–54. 20. Allen, D.; Sterbentz, J. The validity of the use of equivalent DIDO nickel dose. In International Nuclear Graphite Specialists Meeting 5 (INGSM-5), Plas Tan-YBwlch, Wales, 2004. 21. Allen, D.; Thornton, D.; Harris, A.; Sterbentz, J. J. ASTM Int. 2006, 3(9), 365–417. 22. Dahl, R.; Yoshikawa, H. Nucl. Sci. Eng. 1963, 17, 55. 23. Morgan, W. Nucl. Technol. 1974, 21, 50–56.
Graphite in Gas-Cooled Reactors 24. 25. 26. 27. 28. 29. 30. 31. 32. 33. 34. 35. 36. 37. 38. 39. 40. 41. 42.
43. 44. 45.
46. 47. 48. 49. 50. 51. 52. 53. 54. 55. 56. 57.
Eason, E. D.; Hall, G.; Heys, G. B.; et al. J. Nucl. Mater. 2008, 381, 106–113. Sidebotham, E. Nuclear heating in the moderator of Calder Hall type reactors; UKAEA, TRG 295; 1962. Best, J.; Stephen, W.; Wickham, A. J. Prog. Nucl. Energy 1985, 16(2), 127–178. Labaton, V.; Ashton, B.; Lind, R.; Tait, J. Carbon 1969, 7(1), 59–75. Standring, J. J. Nucl. Energy A/B 1966, 20, 201–217. Standring, J.; Ashton, B. W. Carbon 1965, 3, 157–165. Best, J.; Wood, C. Carbon 1975, 13(6), 481–488. Kelly, B. T.; Johnson, P.; Schofield, P.; Brocklehurst, J. E.; Birch, M. Carbon 1983, 21(4), 441–449. Thrower, P. Chem. Phys. Carbon 1969, 5, 217–319. Reynolds, W. N.; Thrower, P. A. Philos. Mag. 1965, 12(117), 573–593. Horner, P.; Williamson, G. K. Carbon 1966, 4(3), 353–363. Kelly, B. T.; Martin, W. H.; Nettley, P. T. Philos. Trans. R. Soc. Lond. A Math. Phys. Sci. 1966, 260(1109), 37–49. Kelly, B. T.; Martin, W. H.; Price, A. M.; Bland, J. T. Philos. Mag. 1966, 14(28), 343–356. Telling, R. H.; Heggie, M. I. Philos. Mag. 2007, 87(31), 4797–4846. Wigner, E. P. J. Appl. Phys. 1946, 17(11), 857–863. Arnold, L. Windscale 1957: Anatomy of a Nuclear Accident, 2nd ed.; Macmillan: London, 1995. Nightingale, R. E. Thermal annealing kinetics of interlayer spacing damage in irradiated graphite. HW-37406, 1955. Lasithiotakis, M.; Marsden, B.; Marrow, J.; Willets, A. J. Nucl. Mater. 2008, 381(1–2), 83–91. Minshall, P.; Wickham, A. J. The description of Wigner energy and its release from Windscale Pile graphite for application to waste packaging and disposal. In Proceedings of the Technical Committee Meeting ‘‘Nuclear Graphite Waste Management’’, Manchester, 1999. Primak, W. Phys. Rev. 1955, 100, 1677. Vand, V. Proc. Phys. Soc. 1943, 55(3), 222–246. Marsden, B. J.; Preston, S. D.; Wichham, A. J.; Tyson, A. Evaluation of Graphite Safety Issues for the British Production Piles at Windscale: Graphite Sampling in Preparation for the Dismantling of Pile 1 and the Further Safe Storage of Pile 2; IAEA–TECDOC–1043; IAEA, 1998. Bell, J.; Bridge, H. Stored Energy in Reactor Graphite; IGR-TN/W522 (AB 7/6077); 1957. Rappeneau, J.; Taupin, J. L.; Grehier, J. Carbon 1966, 4(1), 115–124. Hall, G.; Marsden, B. J.; Fok, S. L. J. Nucl. Mater. 2006, 353, 12–18. Brocklehurst, J. E.; Kelly, B. T. Carbon 1993, 31(1), 179–183. Brocklehurst, J. E.; Kelly, B. T. Carbon 1993, 31(1), 155–178. Kelly, B. T.; Brocklehurst, J. E. Carbon 1971, 9, 783–789. Bokros, J. C.; Guthrie, G. L.; Stevens, D. W. Carbon 1971, 9, 349–353. Seldin, E. J.; Nezbeda, C. W. J. Appl. Phys. 1970, 41(8), 3389–3400. Summers, L.; Walker, D. C. B.; Kelly, B. T. Philos. Mag. 1966, 14(128), 317–323. Taylor, R.; Kelly, B. T.; Gilchrist, K. E. J. Phys. Chem. Solids 1969, 30, 2251–2267. Tuinstra, F.; Koenig, J. L. J. Chem. Phys. 1970, 53(3), 1126–1130. Bradford, M. R.; Steer, A. G. J. Nucl. Mater. 2008, 381(1–2), 137–144.
58.
59.
60. 61. 62. 63. 64.
65. 66. 67. 68.
69. 70. 71. 72. 73. 74. 75. 76. 77. 78.
79. 80.
81.
82. 83. 84.
389
Cords, H.; Zimmermann, R. In Model for irradiation induced changes in graphite materials properties, In Fifth International Conference on Carbon and Graphite, Imperial College, London, 1978. Eason, E. D.; Hall, G.; Marsden, B. J. Development of a model of dimensional change in AGR graphites irradiated in inert environments. In Conference on Ageing Management of Graphite Reactor Cores, University of Cardiff, Wales, 2005; Neighbour, G. B., Ed.; pp 43–50. Eason, E. D.; Hall, G. N.; Marsden, B. J.; Heys, G. B. J. Nucl. Mater. 2008, 381(1–2), 145–151. Hall, G.; Marsden, B. J.; Fok, S. L.; Smart, J. Nucl. Eng. Des. 2003, 222, 319–330. Brocklehurst, J. E.; Bishop, R. Carbon 1964, 2, 27–39. Jenkins, G. J. Nucl. Mater. 1964, 13(1), 33–39. Birch, M.; Brocklehurst, J. E. A review of the behaviour of graphite under the conditions appropriate for the protection of the first wall of a fusion reactor; UKAEA, ND-R-1434(S); 1987. Hall, G. EngD thesis, University of Manchester, 2004. Hall, G.; Marsden, B. J.; Smart, J.; Fok, A. Nucl. Energy 2002, 41(1), 53–62. Burchell, T. D.; Snead, L. L. J. Nucl. Mater. 2007, 371, 18–27. Brocklehurst, J. E.; Edwards, J. The fast neutroninduced changes in dimensions and physical properties of near-isotropic graphites irradiated in DFR; UKAEA, TRG Report 2200(S); 1971. Tsang, D. K. L.; Marsden, B. J.; Fok, S. L.; Hall, G. Carbon 2005, 43, 2902–2906. Hacker, P. J.; Neighbour, G. B.; McEnaney, B. J. Phys. D Appl. Phys. 2000, 33, 991–998. Kelly, B. T. An improved method of estimating the thermal conductivity of irradiated graphite at any temperature; UKAEA, TRG-M-4180(C) (AB 7/19346); 1967. Binkele, L. High Temp. High Press. 1972, 4, 401–409. Binkele, L. J. Non Equilibrium Thermodyn. 1978, 3, 257–266. Brown, R. G. A comparison of the stress/strain curves of isotropic graphite before and after irradiation; UKAEA, TRG-M-5232(C) (AB 7/21009); 1970. Brown, R. G. Carbon 1968, 6(2), 27–30. Maruyama, T.; Eto, M.; Oku, T. Carbon 1987, 25(6), 723–726. Logsdail, D. The effect of gaseous environment on the flexural strength of graphite; AERE, AERE-R5721; 1968. Brocklehurst, J. E.; Brown, R. G. The relation between strength, modulus, and structural changes of isotropic graphites irradiated to high fast neutron doses in DFR; UKAEA, TRG-M-5985 (AB 7/22015); 1972. Oku, T.; Eto, M. Nucl. Eng. Des. 1993, 143, 239–243. Losty, H.; Orchard, J. The strength of graphite. In Fifth Conference on Carbon, Pennsylvania State University, University Park, Pennsylvania; Pergamon, 1962; pp 519–532. Marsden, B. J.; Fok, S. L.; Marrow, T. J.; Mummery, P. M. The relationship between strength and modulus in nuclear graphite. In HTR-2004, 2nd International Topical Meeting on High Temperature Reactor Technology, Beijing, China, 2004. Adam, R. W.; Brocklehurst, J. E. Mechanical tests on graphite with simulated radiolytic oxidation gradients; UKAEA, ND-R-853(S) (AB 7/26300); 1983. Knudsen, F. J. Am. Ceram. Soc. 1959, 42, 376–387. McNally, K.; Tan, E.; Warren, N.; Heys, G. B.; Marsden, B. J. A statistical analysis of the mechanical properties of PGA graphite samples taken from Magnox nuclear
390
85. 86.
87. 88. 89. 90. 91.
Graphite in Gas-Cooled Reactors reactors. In Securing the Safe Performance of Graphite Reactor Cores; The University of Nottingham: Nottingham, 2008; pp 95–102. Gray, B.; Brocklehurst, J. E.; McFarlane, A. Carbon 1967, 5, 173–180. Losty, H.; Bell, I.; Jenkins, G. The irradiation induced plasticity of graphite. In Fifth Conference on Carbon, Pennsylvania State University, University Park, Pennsylvania; Pergamon, 1962; pp 266–273. Perks, A.; Simmons, J. Carbon 1964, 1, 441–449. Tsang, D. K. L.; Marsden, B. J.; Vreeling, J. A.; van der Laan, J. Nucl. Eng. Des. 2008, 238(11), 3026–3030. Hausen, H.; Lo¨lgen, R.; Cundy, M. J. Nucl. Mater. 1977, 65, 148–156. Jouquet, G.; Kleist, G.; Veringa, H. J. Nucl. Mater. 1977, 65, 86–95. Kelly, B. T.; Brocklehurst, J. E. J. Nucl. Mater. 1977, 65, 79–85.
92. 93. 94. 95. 96. 97. 98. 99. 100.
Brocklehurst, J. E.; Kelly, B. T. A review of irradiation induced creep in graphite under CAGR conditions; UKAEA, ND-R-1406(S); 1989. Preston, S. D.; Marsden, B. J. Carbon 2006, 44(7), 1250–1257. Oku, T.; Fujisaki, K.; Eto, M. J. Nucl. Mater. 1988, 152(2–3), 225–234. Kelly, B. T.; Burchell, T. D. Carbon 1994, 32(3), 499–505. Davies, M. A.; Bradford, M. J. Nucl. Mater. 2008, 381(1–2), 39–45. Roberts, A. C.; Cottrell, A. H. Philos. Mag. 1956, 1(8), 711–717. Kelly, B. T.; Foreman, A. J. Carbon 1974, 12, 151–158. Was, G. S. Fundamentals of Radiation Materials Science: Metals and Alloys; Springer: Berlin, 2007. Francis, E. L. Progress Report for the JNPC-Materials Working Party: Graphite Physics Study Group; UKAEA, TRG-M-2854 (AB 7/17604); 1965.
4.12
Vanadium for Nuclear Systems
T. Muroga National Institute for Fusion Science, Oroshi, Toki, Gifu, Japan
ß 2012 Elsevier Ltd. All rights reserved.
4.12.1
Introduction
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4.12.2 4.12.3 4.12.4 4.12.5 4.12.6 4.12.7 4.12.8 4.12.9 4.12.10 4.12.11 4.12.12 4.12.13 References
Vanadium Alloys for Fusion Reactors Compositional Optimization Fabrication Technology Fundamental Study on Impurity Effects Thermal Creep Corrosion, Compatibility, and Hydrogen Effects Radiation Effects Tritium-Related Issues Development of Advanced Alloys Critical Issues Vanadium Alloy Development for Fusion Blankets Summary
391 392 393 396 396 398 400 401 402 403 403 404 405
Abbreviations DBTT dpa flibe GTA HFIR HIP IFMIF
Ductile–brittle transition temperature Displacement per atom Molten LiF-BeF2 salt mixture Gas tungsten arc High Flux Isotope Reactor Hot isostatic pressing International Fusion Materials Irradiation Facility IP Imaging plate ITER International Thermonuclear Experimental Reactor LMFBR Liquid Metal Fast Breeder Reactor MA Mechanical alloying PWHT Postweld heat treatment RAFM Reduced activation ferritic/martensitic REDOX Reduction–oxidation reaction TBM Test Blanket Module TBR Tritium breeding ratio TEM Transmission electron microscope
4.12.1 Introduction Vanadium alloys were candidates for cladding materials of Liquid Metal Fast Breeder Reactors (LMFBR) in the 1970s.1 However, the development was suspended mainly because of an unresolved issue
of corrosion with liquid sodium. Vanadium alloys attracted attention in the 1980s again for use in fusion reactors because of their ‘low activation’ properties. At present, vanadium alloys are considered as one of the three promising candidate low activation structural materials for fusion reactors with reduced activation ferritic/martensitic (RAFM) steels and SiC/SiC composites. Overviews of vanadium alloys for fusion reactor applications are available in the recent proceedings papers of ICFRM (International Conference on Fusion Reactor Materials).2–6 This chapter highlights the recent progress in the development of vanadium alloys mainly for application in fusion nuclear systems.
4.12.2 Vanadium Alloys for Fusion Reactors Various tritium breeding fusion blanket concepts have been studied with different combinations of structural materials, tritium breeding materials, and cooling materials. Vanadium alloys have been used in most cases with liquid lithium as the breeding and cooling materials (self-cooled V/Li blankets) for advanced concepts of DEMO (fusion demonstration power plant) and commercial fusion reactors.7,8 Because of high atomic density of Li atoms in liquid Li relative to Li-ceramics, Li–Pb, and molten-salt 391
392
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Flibe, V/Li systems can obtain high tritium breeding ratio (TBR) without using the neutron multiplier Be. A neutronics calculation showed that ‘tritium self sufficiency’ can be satisfied without Be both in Tokamak and Helical reactor systems.9 Without the necessity of using beryllium as a neutron multiplier, the replacement frequency of the blanket will be reduced because the blanket system is free from the periodic replacement due to the lifetime of Be, which can lead to enhanced plant efficiency. V/Li blankets can be designed with a simple structure as schematically shown in Figure 1. The blanket is composed of Li cooling channels made of vanadium alloys, reflectors, and a shielding area, which is in contrast to more complex solid breeder blankets that need a solid breeder zone, a neutron multiplier beryllium zone, cooling channels using gas or water, and tritium recovery gas flow channels in addition to reflectors and shielding. A self-cooled Li blanket using neutron multiplier beryllium was also designed in the Russian program.10 This concept can downsize the blanket area because of efficient tritium generation per zone. However, the blanket structure must be more
complex than V/Li and new issues need to be solved such as Li/Be compatibility. General requirements for structural materials of fusion blankets include dimensional stability, compatibility with breeder and coolants, high-temperature strength and low-temperature ductility during irradiation. For vanadium alloys, issues concerning industrial maturity such as developing large-scale manufacturing technology need to be resolved. Vanadium alloys could be a candidate structural material for molten-salt Flibe (LiF–BeF2) blankets. For this application, a concept was proposed to dissolve WF6 or MoF6 into Flibe for corrosion protection of the wall surfaces by precipitation of W or Mo and reduction of the tritium inventory in vanadium alloys by enhancing reaction from T2 to TF, which is more highly soluble in Flibe than T2.11 The TBR of Flibe/V blankets may be marginal, but the neutron shielding capability for the superconductor magnet systems may be superior relative to V/Li according to neutronics investigation.12 In this system, precipitates of Wor Mo formed as a result of reaction from T2 to TF needs to be recovered from the flowing Flibe. Table 1 summarizes the blanket concepts using vanadium alloys with the advantages and critical issues.
Flowing liquid lithium
4.12.3 Compositional Optimization
Superconducting magnet
Neutron D-T plasma
Shield
Coating with W, Be, or C
Reflector Vanadium alloy structures
Blanket Figure 1 Illustration of self-cooled Li blanket with V–4Cr–4Ti structural material. Table 1
Vanadium alloys potentially have low-induced activation characteristics, high-temperature strength, and high thermal stress factors. For the optimization of the composition, both major alloying elements and minor impurities need to be controlled. For maintaining the low activation properties, use of Nb and Mo, which used to be the candidate alloying elements for application to LMFBR, need to be avoided. Cr was known to increase the strength of vanadium at high temperature and Ti was known to enhance ductility of vanadium by absorbing interstitial impurities, mostly oxygen. However, excess Cr or Ti can
Breeding blanket concepts using vanadium alloys
Concept
V/Li
V/Be/Li
V/Flibe
Breeder and coolant materials Use of neutron multiplier Be Advantages Critical issues
Liquid Li
Liquid Li
Molten-salt Flibe
No Simple structure MHD coating, T recovery from Li
Yes High TBR MHD coating, Li/Be compatibility, T recovery from Li
No Small MHD pressure drop REDOX control, recovery of W or Mo, increase in TBR
Vanadium for Nuclear Systems
lead to loss of ductility. Hence, optimization of Cr and Ti levels for V–xCr–yTi has been investigated. It was known that with x þ y > 10%, the alloys became brittle6 as shown in Figure 2. With systematic efforts, V–4Cr–4Ti has been regarded as the leading candidate. For low activation purposes, the level of Nb, Mo, Ag, and Al needs to be strictly controlled. Large and medium heats of V–4Cr–4Ti have been made in the United States, Japan, and Russia.
100 1150 ⬚C
50
±20 ⬚C
DBTT (⬚C)
0 1100 ⬚C
950 ⬚C
–50 1000 ⬚C
–100
Annealing temperature 950 ⬚C 1000 ⬚C 1050 ⬚C 1100 ⬚C 1150 ⬚C
–150 –200 –250
5
10
15 Cr + Ti (Wt %)
20
25
Figure 2 DBTT as a function of Cr þ Ti (wt%) of V–Cr–Ti alloy for various annealing temperatures. Reproduced from Zinkle, S. J.; Matsui, H.; Smith, D. L.; Rowcliffe, A. L.; van Osch, E.; Abe, K.; Kazakov, V. A. J. Nucl. Mater. 1998, 258–263, 205–214, with permission from Elsevier.
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An especially high-purity V–4Cr–4Ti ingot produced by the National Institute for Fusion Science (NIFS) in collaboration with Japanese Universities (NIFS-HEAT-1 and 2) showed superior properties in manufacturing due to their reduced level of oxygen impurities.4 Figure 3 compares the contact dose rate after use in the first wall of a fusion commercial reactor for four reference alloys. The full-remote and full-hands-on recycle limits are shown to indicate the guideline for recycling and reuse.13 SS316LN-IG (the reference ITER structural material) will not reach the remoterecycling limit after cooling and hence the recycling is not feasible. F82H (reference RAFM steel) and NIFS-HEAT-2 behave similarly, but NIFS-HEAT-2 shows significantly lower dose rate before the 100year cooling. The dose rate of F82H and NIFSHEAT-2 reached a level almost two orders lower than the remote-recycle limit by cooling for 100 and 50 years, respectively. The dose rate of SiC/SiC composites (assumed to be free from impurities because of lack of reference composition) is much lower at <1 year cooling, but slightly higher at >100 year cooling relative to F82H and NIFS-HEAT-2.
4.12.4 Fabrication Technology Figure 4 summarizes the microstructural evolution during the breakdown process of NIFS-HEAT-2
105 104
Contact dose rate (Sv h–1)
103
Reduced activation ferritics (F82H)
102 101 100 10–1
V–4Cr–4Ti (NIFS-HEAT)
FFHR Li blanket first wall neutron 1.5 MW m–2 operation
SS316 for ITER (SS316LN-IG)
Pure SiC/SiC
Full-remote recycling
10–2 10–3 Full-hands-on recycling
10–4 10–5 10–2
10–1
100 101 102 103 Cooling time after shutdown (years)
104
Figure 3 Contact dose after use in first wall of a fusion commercial reactor for four reference alloys. SS316LN-IG: the reference ITER structural material F82H: reference reduced activation ferritic/martensitic steel NIFS-HEAT-2: reference V–4Cr–4Ti alloy SiC/SiC: assumed to be impurity-free.
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Ingot
Hot forging 1423 K Formation
Heat treatment Hot/cold roll 1373 K/RT 973 K 1273 K 1373 K 1573 K Ti-rich blocky precipitates (with N, O, C) Elongation, band structure Dissolution Ti–O–C thin precipitates Formation Coarsening Dissolution
V–C on GB
50 mm
50 mm
25 mm
1 mm
1 mm
1 mm
50 mm
Figure 4 Microstructural evolution during the breakdown process of V–4Cr–4Ti ingots. Reproduced from Muroga, T.; Nagasaka, T.; Abe, K.; Chernov, V. M.; Matsui, H.; Smith, D. L.; Xu, Z. Y.; Zinkle, S. J. J. Nucl. Mater. 2002, 307–311, 547–554.
260 240 Vickers hardness (Hv)
ingots.4 Bands of small grains aligned along the rolling direction were observed at the annealing temperature below 1223 K. The grains became homogeneous at 1223 K. The examination showed that optimization of size and distribution of Ti-CON precipitates are crucial for good mechanical properties of the V–4Cr– 4Ti products. Two types of precipitates were observed, that is, the blocky and the thin precipitates. The blocky precipitates formed during the initial fabrication process. The precipitates aligned along the working direction during the forging and the rolling processes forming band structures, and were stable to 1373 K. Since clustered structures of the precipitates result in low impact properties, rolling to high reduction ratio is necessary for making a thin band structure or homogenized distribution of the precipitates. The thin precipitates were formed at 973 K and disappeared at 1273–1373 K. At 1373 K, new precipitates, which were composed of V and C, were observed at grain boundaries. They seem to be formed as a result of redistribution of C induced by the dissolution of the thin precipitates. The impact of the inhomogeneous microstructure can influence the fracture properties.14 Figure 5 shows the hardness as a function of final heat treatment temperature for three V–4Cr–4Ti materials: NIFS-HEAT-1, NIFS-HEAT-2, and USDOE-832665 (US reference alloy).15 The hardness has a minimum at 1073–1273 K, which corresponds to the temperature range where formation of the thin precipitates is maximized. With the heat treatment higher than this temperature range, the hardness increases and the ductility decreases because the
NIFS-HEAT-1 NIFS-HEAT-2 US-DOE 832665
220 200 180 160 140 V–4Cr–4Ti 120 200 400 600 800 1000 1200 1400 1600 Annealing temperature (K)
Figure 5 Vickers hardness as a function of annealing temperature for NIFS-HEAT-1, NIFS-HEAT-2, and US-DOE 832665. Reproduced from Heo, N. J.; Nagasaka, T.; Muroga, T. J. Nucl. Mater. 2004, 325, 53–60.
precipitates dissolve enhancing the level of C, N, and O in the matrix. Based on the evaluation of various properties in addition to the hardness as a function of heat treatment conditions, the optimum heat treatment temperature of 1173–1273 K was suggested. Plates, sheets, rods, and wires were fabricated minimizing the impurity pickup and maintaining grain and precipitate sizes in Japanese, US, and Russian programs. Thin pipes, including those of pressurized creep tube specimens, were also successfully fabricated
Vanadium for Nuclear Systems
in Japan maintaining the impurity level, fine grain size, and straight band precipitate distribution by maintaining a constant reduction ratio between the intermediate heat treatments.16 The fine-scale electron beam welding technology was enhanced as a result of the efforts for fabricating the creep tubes, including plugging of end caps.17 In the United States, optimum vacuum level was found for eliminating the oxygen pick-up during intermediate annealing to fabricate thin-walled tubing of V–4Cr–4Ti.18 In Russia, fabrication technology is in progress for construction of a Test Blanket Module (TBM) for ITER (International Thermonuclear Experimental Reactor).19 Joining of V–4Cr–4Ti by gas tungsten arc (GTA) and laser welding methods was demonstrated. GTA
395
is a suitable technique for joining large structural components. GTA welding technology for vanadium alloys provided a significant progress by improving the atmospheric control. The results are summarized in Figure 6. Oxygen level in the weld metal was controlled by combined use of plates of NIFSHEAT-1 (181 wppm O) or US-8332665 (310 wppm O) and filler wire of NIFS-HEAT-1, US-8332665, or a high-purity model alloy (36 wppm O). As demonstrated in Figure 6, ductile–brittle transition temperature (DBTT) of the joint and the oxygen level in the weld metal had a clear positive relation. This motivated further purification of the alloys for improvement of the weld properties.20 Only limited data on irradiation effects on the weld joint are available at present.
15
Absorbed energy (J)
EU = 13 J
10
Plate/filler NH1/HP 128 K
US/HP 183 K
NH1/NH1
US/US
188 K
320 K
5
0 50
100
150
200 250 Test temperature (K)
350
400
US/US
300
DBTT (K)
300
US/HP
200
NH1/NH1
100
NH1/HP DBTT = +60 K/100 wppm O
0
0
50
100
150
200
250
300
350
400
Oxygen in weld metal (wppm) Figure 6 Upper: Absorbed energy of Charpy impact tests of V–4Cr–4Ti weld joints as a function of test temperature for various combinations of plates and fillers. Lower: DBTT of V–4Cr–4Ti weld joints as a function of oxygen level in the weld metal. NH1, NIFS-HEAT-2 (O: 181 wppm); US, US-DOE 832665 (O: 310 wppm); HP, high-purity model V–4Cr–4Ti alloy (O: 36 wppm). Reproduced from Nagasaka, T.; Grossbeck, M. L.; Muroga T.; King, J. F. Fusion Technol. 2001, 39, 664–668.
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The welding results in complete dissolution of TiCON precipitates and thus results in significant increase in the level of C, O, and N in the matrix. In such conditions, radiation could cause embrittlement. Some TEM observations showed enhanced defect cluster density at the weld metals. However, the overall evaluation of the radiation effects remains to be performed. Especially, elimination of radiation-induced degradation byapplying appropriate conditions of postweld heat treatment (PWHT) is the key issue. For the use of vanadium alloys as the blanket of fusion reactors, the plasma-facing surfaces need to be protected by armor materials such as W layers. Limited efforts are, however, available for developing the coating technology. A low pressure plasma-spraying method was used for coating W on V–4Cr–4Ti for use at the plasma-facing surfaces. The major issue for the fabrication is the degradation of the vanadium alloy substrates by oxidation during the coating processes. Figure 7 shows the result of bending tests of the coated samples. The crack was initiated within the W layer propagating parallel to the interface and followed by cracking across the interface. Thus, in this case, the quality of W coating layer is the issue rather than the property of the V–4Cr–4Ti substrate or the interface. Hardening of substrate V–4Cr–4Ti by the coating occurred but was shown to be in acceptable range.21 Figure 8 is a collection of the products from NIFS-HEAT-2.
Research with model V–4Cr–4Ti alloys doped with O and N provided information on the partitioning of O and N into the precipitates and matrix. The density of the blocky precipitates and thin precipitates increased with N and O levels, respectively. Figure 9 shows hardness as a function of N and O levels in V–4Cr–4Ti after melting and annealing at 1373 K for 1 h.22 Hardness after annealing at 1373 K, where only the blocky precipitates were observed in the matrix, increased to a certain extent with O level (4.5 Hv/100 wppm O), but only very weakly with N level (0.9 Hv/100 wppm N). These data suggest that, after the annealing, most of the N is included in the blocky precipitates and stable to 1373 K. On the other hand, O exists in the matrix, the blocky and the thin precipitates, and the partitioning changes with the heat treatment. Thus, for the purpose of the property control of V–4Cr–4Ti, the level of N before the heat treatment is not so important but that of O is crucial. It is to be noted, however, that N contamination during the operation can influence the properties of vanadium alloys seriously. Fundamental information on the impurity distribution and interaction with solutes and dislocations is obtained by serrated flow in tensile deformation as shown in Figure 10. Temperature and stain rate dependence of the flow showed that the serrated flow above 673 K is related to C and O and above 773 K to N. Small serration height at 673 K for NIFS-HEAT-1 (1–3 MPa) relative to that of US-832665 (9 MPa) was observed and attributed to the difference in O level.23
4.12.5 Fundamental Study on Impurity Effects 4.12.6 Thermal Creep Effects of C, O, and N on the property of vanadium are a long-standing research subject. However, research into the effects of C, O, and N on V–4Cr–4Ti is limited.
Thermal creep is a potential factor which can determine the upper operation limit of vanadium alloys.
Crack W V–4Cr–4Ti Intergranular fracture 500 µm
50 µm
10 µm Figure 7 Cross-section of W coating on V–4Cr–4Ti after bending tests. Fracture started in the W coating layer.
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397
(mm) f 4.57 ⫻ 0. 25 t ⫻ 400 mm 6.6 t 0.5 t
26 t 1.9 t
1.0 t
4.0 t 0.25 t
2d 8d f 10 ⫻ 0. 5 t ⫻ 100 m m Plates, sheets, wires, and rods
Thin pipes
20 mm Creep tubes
W coating
NIFS-HEAT-2 (V–4Cr–4Ti)
0.5 mm
W coating by plasma spraying
Laser weld joint
5 mm
Figure 8 Collection of the V–4Cr–4Ti products manufactured by the Japanese program.
V–4Cr–4Ti, as-melted V–4Cr–4Ti, 1373 K Pure V, as-melted Pure V, 1373 K
Vickers hardness (Hv)
300
250
200
150
100
50
0
200 400 600 800 1000 1200 0 Oxygen level (wppm)
200 400 600 800 1000 1200 Nitrogen level (wppm)
Figure 9 Vicker’s hardness as a function of O and N levels for V–4Cr–4Ti after melting and annealing at 1373 K for 1 h. Reproduced from Heo, N. J.; Nagasaka, T.; Muroga, T.; Matsui, H. J. Nucl. Mater. 2002, 307–311, 620–624.
Previously, uniaxial tensile creep tests and biaxial pressurized creep tube tests were carried out in vacuum for evaluation of the creep deformation characteristics. Figure 11 shows summary of the creep deformation rate as a function of applied stress.3 In this plot, the creep data were described by a powerlaw equation24: de=dt ¼ AðDGb=kT Þðs=GÞn
where de/dt is the creep rate, s is the applied stress, D is the lattice diffusion coefficient, G is the shear modulus, b is the Burgers vector, k is the Boltzmann constant, T is the absolute temperature, and A is a constant. The exponent of the function (n) changed from <1 to >10 with the increase in the stress. A new apparatus for biaxial creep testing in Li provided opportunities for examining creep
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Stress (MPa)
200
0
10
30
20 Strain (%)
Figure 10 Tensile deformation curves of V–4Cr–4Ti at various temperatures.
10-5 Uniaxial tests 310 wppm O
10-6
(de/dt)kT/DGb
n = 3.7
10-9 10-10
n = 4.3 n = 0.84
10-11 10-12
10-7
10-8
In vacuum 50 MPa 70 MPa 90 Mpa
10-9
10-10
0
2
4 6 8 Creep strain (%)
In lithium
10
12
Figure 12 Creep strain rate as a function of creep strain for the same batch of NIFS-HEAT-2 creep tubes in vacuum and Li environments. Modified from Li, M.; Nagasaka, T.; Hoelzer, D. T.; et al. J. Nucl. Mater. 2007, 367–370, 788–793; Fukumoto, K.; Nagasaka, T.; Muroga, T.; Nita, N.; Matsui, H. J. Nucl. Mater. 2007, 367–370, 834–838.
n = 13
10-7 10-8
Creep strain rate (1 s–1)
10-6 1073 K 973 K 873 K 773 K 673 K RT
surface hardening during exposure to Li. Further investigation is necessary for understanding the environmental effects on impurity redistribution and creep performance. Microstructural observations of the creep tube specimens tested at 1123 K showed free dislocations and dislocation cell at 100 and 150 MPa, respectively. This change of dislocation structure can cause the change in power-law creep behavior.27
Biaxial tests 699 wppm O 10-3
10-2
4.12.7 Corrosion, Compatibility, and Hydrogen Effects
s/G Figure 11 Thermal creep deformation rate of V–4Cr–4Ti as a function of applied stress for uniaxial and biaxial tests. The definition of the terms and the function from which n is extracted are indicated in the text. Reproduced from Kurtz, R. J.; Abe, K.; Chernov, V. M.; Hoelzer, D. T.; Matsui, H.; Muroga, T.; Odette, G. R. J. Nucl. Mater. 2004, 329–333, 47–55.
deformation in Li with that in vacuum.25 However, the correlation of creep data is subject to the alloy heat and manufacturing processes as well as test methods and environments. Figure 12 shows the comparison of the NIFS-HEAT-2 creep strain rate versus creep strain data for tests in vacuum and Li environments at 1073 K, for the same batch of NIFSHEAT-2 creep tubes.25,26 The figure clearly shows reduced strain rate in Li environments. A possible factor could be N pick-up from Li and the resulting
In a Li/V blanket, it is believed that the interior of the wall needs to be coated with insulator ceramics for mitigating the pressure drop caused by magnetohydrodynamic effects (see also Chapter 4.21, Ceramic Coatings as Electrical Insulators in Fusion Blankets). Corrosion of vanadium alloys in liquid Li might not be a concern if the entire inner wall is covered with the insulating ceramic coating. However, since the idea to cover the insulator ceramic coating again with a thin vanadium or vanadium alloy layer was presented for the purpose of preventing liquid lithium from intruding into the cracks in the ceramics coating, the corrosion of vanadium alloys in liquid lithium again attracted attention. It is known that the corrosion of vanadium alloys in liquid lithium is highly dependent on the alloy composition and lithium chemistry. Especially, the N level influences the corrosion in complex manners.28,29 Figure 13 shows a summary of the weight
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0.1
Ti 50
I
40
30 +11.9 20 +11
Ti:Cr = 2:1
+2.5
+6.7
+5.8
Weight gain (mg mm–2)
0.08
+1.2
–19.0
–22.0
V–4Cr–4Ti V–4Cr–4Ti–0.5Si V–4Cr–4Ti–0.5Al V–4Cr–4Ti–0.5Y
0.06
0.04
+7.5
V
+0.4
10
–8.2
–2.1
–2.4
–19.7 –21.0
20
–47.4
0
–52.5 –26.4
30
Cr 40
50
Figure 13 The compositions of V–Ti–Cr alloys (wt%) with increase (area I) and decrease (area II) of mass (g cm2) after holding of samples in Li at 973 K, 500 h. Reproduced from Eliseeva, O. I.; Fedirko, V. N.; Chernov, V. M.; Zavialsky, L. P. J. Nucl. Mater. 2000, 283–287, 1282–1286.
gain and loss in V–xCr–yTi systems in Li.30 High Ti alloys showed a weight increase by forming a TiN layer and high Cr alloys exhibited a weight loss as a result of the dissolution of Cr–N complexes. As the boundary of the two contradictory changes, Ti:Cr2:1 was observed. Recently, a corrosion test using monometallic thermal convection Li loop made of V–4Cr–4Ti was conducted at 973 K for 2355 h. Because of the temperature gradient, weight loss and weight gain of V–4Cr–4Ti samples occurred at the hot leg and cold leg, respectively. However, the loss rate corresponded to only <1 mm year1 and the degradation of the mechanical properties were shown to be small.31 V–4Cr–4Ti alloys have been developed mainly for use in Li environments, which are extremely reducing conditions. For the use of vanadium alloys in oxidizing conditions, a different alloy optimization may be necessary. The corrosion of vanadium alloys in oxidizing environments is of interest both for the performance of the pipe exterior out of the breeding blanket and application in non-Li coolant systems such as gas and water systems. Oxidation kinetics of vanadium alloys were studied and showed either parabolic or linear kinetics.32,33 As the surface oxide layer is not formed or, if formed, not protective to the internal oxidation, alloying with other oxide-formers is necessary for improvement. The addition of Si, Al, or Y was shown to significantly suppress the weight gain during exposure to air above 873 K as shown in Figure 14.34
773
873 973 Oxidation temperature (K)
1023
Figure 14 Weight gain of V–4Cr–4Ti with Si, Al, and Y exposed to air for 1 h. At 1023 K, the weight gain was not measured for V–4Cr–4Ti because the surface oxidized layer melted. Reproduced from Fujiwara, M.; Natesan, K.; Satou, M.; Hasegawa, A.; Abe, K. J. Nucl. Mater. 2002, 307–311, 601–604.
Natesan (BL-71 O:670 wppm)
50
DiStefano (US-832665 O:310 wppm) DiStefano (Preoxidized US-832665 O:800 wppm) Chen (SWIP-Heat O:900 wppm)
40 Total elongation (%)
10
Melted
0.02
II
Chen (NIFS-HEAT-2 O:158 wppm)
30
20
10
0
0
100 200 300 400 500 600 Hydrogen concentration (wppm)
700
Figure 15 Total elongation as a function of hydrogen concentration for V–4Cr–4Ti alloys with different O levels. Modified from DiStefano, J. R.; Pint, B. A.; DeVan, J. H. J. Nucl. Mater. 2000, 283–287, 841–846; Chen, J. M.; Muroga, T.; Qiu, S.; Xu, Y.; Den, Y.; Xu, Z. Y. J. Nucl. Mater. 2004, 325, 79–86.
However, the addition of these elements was not effective in suppressing corrosion in water. Increase in Cr level was shown to be effective, instead. The effects of oxygen level on hydrogen embrittlement have been investigated. Figure 15 compares elongation versus hydrogen concentration for V–4Cr–4Ti
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alloys with various O levels. The loss of ductility by hydrogen charging was shown to be enhanced by impurity oxygen.35,36
4.12.8 Radiation Effects A fair amount of data is available for radiation response of vanadium alloys partly because they were candidates of cladding materials of LMFBR. For example, void swelling is known to be quite small if the alloy contains Ti. However, data are limited for V–4Cr–4Ti because this composition was decided as the reference one for fusion only recently. For this alloy, the feasibility issues of radiation effects are considered to be loss of ductility at lower temperature, embrittlement enhanced by transmutant helium at high temperature, and irradiation creep at intermediate to high temperature. The mechanism of the loss of uniform elongation of vanadium alloys at relatively low temperature (<673 K) and low dose (0.1 dpa) has been a longterm research subject. Microstructural observation after tensile tests showed that radiation-induced defect clusters were lost in layer structures and the defect-free zones were accompanied by dislocation channels as shown in Figure 16.37 This fact implies flow localization during deformation. Although the mechanism of the flow localization needs further investigation, it is inferred that interaction of dislocations with
radiation-induced defect clusters, precipitates, or complexes of the two species is responsible. If the precipitates, most likely Ti-CON, play the role in this process, reduction of impurities in the matrix can improve the properties. Figure 17 compares the uniform elongation after irradiation for V–(4–5)Cr–(4–5)Ti alloys and those with doping of Al, Si, and Y. The significant increase in uniform elongation by the addition of Al, Si, and Y, which are known as getters of interstitial impurities such as O, N, and C in the matrix, suggests that the reduction of the interstitial impurities in solution enhances the radiation resistance.38 The effects of interstitial impurities on the formation of dislocation loops and precipitates were investigated by ion irradiations. Figure 18 shows temperature dependence of the densities of loops and precipitates.39 The loop density was not influenced by O level, but the precipitate density increased with O level below 973 K. Helium embrittlement is a critical issue, which is thought to determine the upper temperature limit for vanadium alloys. Past experimental evaluations of the helium effects involved uncertainties because controlled generation of helium during irradiation in a similar manner to that in fusion condition has been quite difficult. As a result, the past evaluation of the helium effects varied from weak to very strong.3 The Dynamic Helium Charging Experiment (DHCE) using fission reactors40 is one of the few potential neutron irradiation experiments with controlled variation of He/dpa ratio including typical fusion
Load-elongation curves for V–4Cr–4Ti irradiated in HFBR to 0.5 dpa Engineering stress (MPa)
700 600 500
383 K
543 K
Ttest ~ Tirr 598 K 693 K
400 300 200 100 0 0
g = 011
0.05 0.1 0.15 0.2 0.25 0.3 Normalized crosshead displacement (mm mm–1) Ttest = Tirr = 543 K
100 nm Figure 16 Tensile test curves for V–4Cr–4Ti irradiated in HFBR to 0.5 dpa and microstructure after the tensile test. Reproduced from Rice, P. M.; Zinkle, S. J. J. Nucl. Mater. 1998, 258–263, 1414–1419.
Vanadium for Nuclear Systems Annealing conditions: 900–1125 ⬚C, 3.6 or 7.2 ks
Uniform elongation (%)
12 10 8
ATR, 0.7–4.7 dpa FFTF, EBR-II, 10–54 dpa Loomis et al.,51 FFTF, 13–33dpa Zinkle et al.,52 EBR-II, 4 dpa Tsai et al.,53 ATR, 4.1–4.3 dpa Tsai et al.,54 BOR-60, 17–19 dpa Snead et al.,56 HFBR, 0.5 dpa Chung et al.,55 HFIR, 10 dpa
Creep strain (%)
V–4.8Ti–4.0Cr–Si, Al, Y V–3.8Ti–5.9Cr–Si, Al, Y V–(4–5)Cr–(4–5)Ti
Test temperature–Irradiation temperature
US-832665 698 K-Li (HFIR) NIFS-HEAT-2 698 K-Li (HFIR) NIFS-HEAT-2 731 K-Na (JOYO)
0.2
0.1
6 0 4 2 0
0
100
200 300 400 500 Irradiation temperature (⬚C)
600
700
Figure 17 Uniform elongations as a function of irradiation temperature for V–(4–5)Cr–(4–5)Ti alloys and those with addition of Si, Al, and Y. Reproduced from Satou, M.; Chuto T.; Abe, K. J. Nucl. Mater. 2000, 283–287, 367–371.
1025 O:389, N:14 wppm (loop) O:28, N:27 wppm (loop) O:389, N:14 wppm (precipitates) O.28, N:27 wppm (precipitates)
1024 Defect density (m–3)
401
50
150 100 Applied stress (MPa)
200
Figure 19 Creep strain as a function of applied stress for V–4Cr–4Ti (US-832665 and NIFS-HEAT-2) irradiated in Li (HFIR) and Na (JOYO) environments. The creep strain was normalized as that at two displacements per atom. Reproduced from Fukumoto, K.; Narui, M.; Matsui, H.; et al. J. Nucl. Mater. 2009, 386–388, 575–578.
environments. The data also compare the performances of US and Japanese reference alloys.41 It was found that the creep strain rate exhibited a linear relationship with the effective stress up to 150 MPa at 700 K and the differences with the environments and the heats are small.
4.12.9 Tritium-Related Issues
1023
1022
1021
1020 400
0
500
600
700 800 900 Temperature (K)
1000 1100
Figure 18 Densities of dislocation loops and precipitates as a function of irradiation temperature for two V–4Cr–4Ti alloys with different O and N levels (0.1 dpa by Cu ion irradiation). Reproduced from Watanabe, H.; Suda, M.; Muroga, T.; Yoshida, N. J. Nucl. Mater. 2002, 307–311, 408–411.
conditions. DHCE is highly anticipated as a potential method to extend our understanding of the helium effects. However, for conclusive evaluation, a 14 MeV neutron source is certainly necessary. The irradiation creep tests have made progress recently, partly because of the progress in fabricating high quality pressurized creep tube specimens with reduced impurity levels. Figure 19 shows the normalized creep strain as a function of applied stress by irradiation in HFIR and JOYO in Li and Na
In the blanket, the tritium inventory is not considered to be the issue once liquid Li is used as the breeding and cooling materials owing to the high hydrogen solubility of Li. The behavior of hydrogen and its isotopes in vanadium alloys is a concern for tritium retention in the first wall. Deuterium retention of V–4Cr–4Ti was investigated by deuterium ion implantation followed by thermal desorption, in comparison with other candidate first wall materials. The study showed that the retained amount at 380 K was one and two orders of magnitude larger than graphite and tungsten, respectively. For the irradiation at 773 K, the retained amount was almost the same as that of graphite and one order larger than tungsten.42 Surface composition was also known to influence the hydrogen transport. For example, the rate of absorption was highly influenced by prior heat treatment, inducing Ti surface segregation.43 Recent progress in detecting tritium by means of imaging plate (IP) enhanced the understanding of the tritium behavior in vanadium alloys. Figure 20 compares IP images of cold rolled V–4Cr–4Ti and pure V after tritium charging. Tritium is preferentially absorbed in Ti-rich precipitates that have a band structure to the rolling direction.44
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Pure Vanadium
Rolling direction
V–4Cr–4Ti (NIFS-HEAT-2)
Low
High
2 mm
Low
High
2 mm
Figure 20 Distribution of tritium measured by imaging plates for cold-rolled V–4Cr–4Ti and pure vanadium. The tritium was charged by gas absorption. After Homma, H.; Hatano, Y.; Daifuku, H.; et al. J. Nucl. Mater. 2007, 367–370, 887–891.
4.12.10 Development of Advanced Alloys The performance of structural materials can strongly influence the blanket design. Especially, the operation temperature window and expected lifetime are the key parameters. Increase in the upper operation temperature limit can enhance the blanket operation temperature and thus plant efficiency. Therefore, enhancing the high-temperature strength is the key issue for improving the performance of the blanket and thus the attractiveness of the fusion power systems. For this purpose, efforts have been made to develop advanced vanadium alloys with potential use at higher temperature. One of the relatively simple ways to enhance the strength of the alloy is to change the thermal and mechanical treatment of the alloys. Especially, formation of a high density of precipitates can strengthen the alloy. Figure 21 shows microstructure and hardness of V–4Cr–4Ti as a function of the temperature of reheating for 1 h after annealing at 1373 K for 1 h. The annealing at 1373 K dissolves most of the thin precipitates and the reheating can form new precipitates. By choosing an appropriate reheating temperature (873–973 K), the materials can be strengthened by the high density of fine precipitates. However, the strengthening by this treatment will be lost at >973 K because of the coarsening of the precipitates. To prevent the coarsening, cold work was
applied to the specimens. Figure 22 shows the minimum creep rate for standard V–4Cr–4Ti and solution annealed, aged, and cold-worked V–4Cr–4Ti. Suppression of the creep rate occurred at 1073 K but only with relatively high stresses.45 Microstructural analysis showed that the suppressive role of coldwork-induced dislocations was lost during the creep deformation by the change in the nature of the dislocations from sessile ah100i type to gliding a/2h111i type.46 Further efforts are being made, for example, to cold-work followed by aging (strain-aging-induced strengthening). High-temperature strength of V–Cr–Ti alloys can be enhanced by increasing the Cr level. However, high Cr alloys have low ductility and fabricability issues. Recent detailed survey in V–xCr–4Ti alloys showed that the strength at high temperature increases with a small change in the DBTT with the Cr level at 7%.47 High-strength vanadium alloys were made by addition of Y, O, and N to vanadium followed by mechanical alloying (MA) and hot isostatic pressing (HIP). The addition of Y, O, and N was intended to enhance mechanical properties by dispersion of Y2O3 and YN and scavenging O and N from the matrix. Alloys produced by optimization of the processes had small grains and homogeneously dispersed particles and showed higher tensile strength than those of NIFS-HEATs with moderate uniform elongation, both at room temperature and 1073 K as shown in Figure 23.48 Fine
Vanadium for Nuclear Systems
450 Larger grain Yield stress (MPa)
873 K
403
Perfectly dissolved 1373 K
400
Solution
350 Precipitation 300
200 nm 250
As solution 873 973 1073 1173 1273 1373 heat treated Reheat temperature (K)
973 K
1073 K
1273 K
1173 K
Start to dissolve Coarse precipitates (low density)
Fine precipitates
Figure 21 Hardness and microstructure of V–4Cr–4Ti as a function of reheating temperature for 1 h after annealing at 1373 K for 1 h.
4.12.11 Critical Issues
Minimum creep rate (s–1)
100–4
NIFS-HEAT-2 10–5
800 ⬚C 10–6
750 ⬚C
10–7
700 ⬚C
10–8
10–9 100
120
140
160
SAACW STD 180
200 220 240 260 280 300
Stress (MPa) Figure 22 Minimum creep rates as a function of applied stress for V–4Cr–4Ti with standard heat treatment (1273 K for 1 h: STD) and precipitate-hardening heat treatment (1373 K for 1 h, 873 K for 20 h, and cold rolled: SAACW). Reproduced from Chen, J. M.; Nagasaka, T.; Muroga, T.; Qiu, S. Y.; Li, C.; Nita, N. J. Nucl. Mater. 2008, 374, 298–303.
grain and oxide dispersion increased high-temperature strength and inhibited formation of interstitial loops in the matrix by neutron irradiation because of the enhanced defect sinks. Thus, mechanically alloyed vanadium alloys have the potential to extend both low- and high-temperature operation limits. Other efforts to improve high-temperature strength of vanadium alloys include strengthening by internal oxidation.49
With the recent progress in the fabrication technology, the number of critical issues for the development of vanadium alloys for fusion reactors has been reduced. The remaining critical issues are thermal and irradiation creep, transmutant helium effects on high temperature mechanical properties, and radiation effects on fracture properties. The effect of helium, particularly, is still uncertain and can be evaluated precisely only with the use of 14 MeV neutrons. This fact highly motivates the construction of a 14 MeV neutron source. With the progress of the properties of vanadium alloys, the blanket concepts using the alloy become more attractive. Extension of the operation temperature window and lifetime of vanadium alloys contribute to the improvement of the quality of the blanket. Therefore, exploration of advanced vanadium alloys from the current reference alloy is a valuable challenge for enhancing the expected performance, and then attractiveness, of fusion reactors.
4.12.12 Vanadium Alloy Development for Fusion Blankets In the fusion materials development strategy, the candidate structural materials are categorized into reference and advanced materials. As the reference
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800
TTest = 298 K
V–1.0 vol% Y2O3–0.7 vol% YN NIFS-HEAT-1
Stress (MPa)
600
400
TTest = 1073 K
200
0
Strain (%)
20%
400 nm
Figure 23 Tensile test curves and microstructure of V–Y2O3–YN produced by mechanical alloying (MA) in comparison with V–4Cr–4Ti (NIFS-HEAT-1). After Kuwabara, T.; Kurishita, H.; Hasegawa, M. J. Nucl. Mater. 2000, 283–287, 611.
Flexible support
Be Li WC
First wall
Figure 24 A layout of the V/Be/Li test blanket module for International Thermonuclear Experimental Reactor proposed by Russia. After Kirillov, I. R.; Shatalov, G. E.; Strebkov, Y. S.; the RF TBM Team. Fusion Eng. Des. 2006, 81, 425–432.
the advanced materials, which will contribute to increasing attractiveness of the fusion system in terms of cost of electricity and environmental benignness. It is recognized that the development of the advanced materials must also be enhanced now due to the long lead time necessary for their development. It should also be noted that vanadium alloys are the only nonferromagnetic and ductile materials of the three candidates. If the impact of the ferromagnetism of the RAFM on plasma operation should be unavoidable and the brittleness of SiC/SiC should be determined unaccepted by design studies, vanadium alloys could be the only candidate of low activation structural materials for fusion reactors. As shown in the summary of critical issues, a 14 MeV neutron source is highly necessary for the qualification of vanadium alloys. IFMIF (International Fusion Materials Irradiation Test Facility, a 14 MeV neutron source) is under design and is recognized to be essential for developing structural materials for fusion reactors. The TBM to be installed in ITER is also considered to be an important milestone for technological integration. Figure 24 shows the design of the V/Li TBM in ITER proposed by Russia.50 The development of vanadium alloys is planned to proceed with IFMIF for qualification of the alloy and ITER-TBM for technology integration, in addition to fundamental studies using fission reactors, etc.
4.12.13 Summary materials, RAFM steels were selected because they have the most matured industrial infrastructure. Development of the reference materials is crucial for the realization of DEMO (fusion demonstration power plant) in a timely manner. On the other hand, vanadium alloys and SiC/SiC were nominated as
As to the application in nuclear systems, vanadium alloys were once candidate cladding materials for LMFBR, but, at present, are considered mostly as candidate low activation structural materials for fusion reactors.
Vanadium for Nuclear Systems
Vanadium alloys, with the reference composition of V–4Cr–4Ti, are one of the three candidate low activation structural materials with RAFM and SiC/SiC. They are the only nonferromagnetic and ductile materials of the three candidates and thus are promising for advanced structural materials of fusion reactors. The self-cooled liquid lithium blanket using structural materials of vanadium alloys is an attractive concept because of the high heat transfer capability, high-temperature operation, simple structure, high tritium breeding capability, and low tritium leakage. Recent progress, especially in the fabrication technologies, has successfully reduced the number of critical issues enhancing the feasibility of the alloys for fusion application. Major remaining issues of vanadium alloys are thermal and irradiation creep, transmutant helium effects on mechanical properties, and radiation effects on fracture properties. For conclusive characterization of the helium effects, the use of IFMIF is essential. Efforts are also being made to develop advanced vanadium alloys to extend the temperature window and lifetime of vanadium alloys in fusion reactor environments.
17. 18. 19. 20. 21. 22. 23. 24. 25. 26. 27. 28. 29. 30. 31. 32. 33. 34. 35.
References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16.
Suzuki, T.; Noda, T.; Iwao, N.; Kainuma, T.; Watanabe, R. J. Nucl. Mater. 1976, 62, 205–212. Kurtz, R. J.; Abe, K.; Chernov, V. M.; et al. J. Nucl. Mater. 2000, 283–287, 70–78. Kurtz, R. J.; Abe, K.; Chernov, V. M.; et al. J. Nucl. Mater. 2004, 329–333, 47–55. Muroga, T.; Nagasaka, T.; Abe, K.; et al. J. Nucl. Mater. 2002, 307–311, 547–554. Muroga, T.; Chen, J. M.; Chernov, V. M.; et al. J. Nucl. Mater. 2007, 367–370, 780–787. Zinkle, S. J.; Matsui, H.; Smith, D. L.; et al. J. Nucl. Mater. 1998, 258–263, 205–214. Najmabadi, F.; et al. Fusion Eng. Des. 1997, 38, 3–25. Sokolov, Y. A.; et al. Fusion Eng. Des. 1998, 41, 525–529. Tanaka, T.; Muroga, T.; Sagara, A. Fusion Sci. Technol. 2005, 47, 530–534. Kolbasov, B. N.; Belyakov, V. A.; Bondarchuk, E. N.; et al. Fusion Eng. Des. 2008, 83, 870–876. Muroga, T.; Tanaka, T.; Li, Z.; Sagara, A.; Sze, D. K. Fusion Sci. Technol. 2007, 52, 682–686. Sze, D. K.; McKarthy, K. A.; Sawan, M. E.; Tillack, M. S.; Ying, A. Y.; Zinkle, S. J. Fusion Technol. 2001, 39, 746–750. Dolan, T. J.; Butterworth, G. J. Fusion Technol. 1994, 26, 1014–1018. Donahue, E. G.; Odette, G. R.; Lucas, G. E. J. Nucl. Mater. 2000, 283–287, 518–522. Heo, N. J.; Nagasaka, T.; Muroga, T. J. Nucl. Mater. 2004, 325, 53–60. Nagasaka, T.; Muroga, T.; Iikubo, T. Fusion Sci. Technol. 2003, 44, 465–469.
36. 37 38. 39. 40. 41. 42. 43. 44. 45. 46. 47. 48. 49. 50. 51.
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Fukumoto, K.; Matsui, H.; Narui, M.; Nagasaka, T.; Muroga, T. J. Nucl. Mater. 2004, 335, 103–107. Rowcliffe, A. F.; Hoelzer, D. T.; Kurtz, R. J.; Young, M. J. Nucl. Mater. 2007, 367–370, 839–843. Chernov, V. M.; et al. Nucl. Fusion 2007, 47, 839–848. Nagasaka, T.; Grossbeck, M. L.; Muroga, T.; King, J. F. Fusion Technol. 2001, 39, 664–668. Nagasaka, T.; Muroga, T.; Noda, N.; Kawamura, M.; Ise, H. Fusion Sci. Technol. 2005, 47, 876–880. Heo, N. J.; Nagasaka, T.; Muroga, T.; Matsui, H. J. Nucl. Mater. 2002, 307–311, 620–624. Fukumoto, K.; Yamamoto, T.; Nakao, N.; Takahashi, S.; Matsui, H. J. Nucl. Mater. 2002, 307–311, 610–614. Li, M.; Hoelzer, D. T.; Grossbeck, M. L.; Rowcliffe, A. F.; Zinkle, S. J.; Kurtz, R. J. J. Nucl. Mater. 2009, 386–388, 618–621. Li, M.; Nagasaka, T.; Hoelzer, D. T.; et al. J. Nucl. Mater. 2007, 367–370, 788–793. Fukumoto, K.; Nagasaka, T.; Muroga, T.; Nita, N.; Matsui, H. J. Nucl. Mater. 2007, 367–370, 834–838. Gelles, D. S.; Toloczko, M. B.; Kurtz, R. J. J. Nucl. Mater. 2007, 367–370, 869–875. Chopra, O. K.; Smith, D. L. J. Nucl. Mater. 1988, 155–157, 683. Evtikhin, V. A.; Lyublinski, I. E.; Vertkov, A. V. J. Nucl. Mater. 1998, 258–263, 1487–1491. Eliseeva, O. I.; Fedirko, V. N.; Chernov, V. M.; Zavialsky, L. P. J. Nucl. Mater. 2000, 283–287, 1282–1286. Pint, B. A.; Pawel, S. J.; Howell, M.; et al. J. Nucl. Mater. 2009, 386–388, 712–715. DiStefano, J. R.; DeVan, J. H. J. Nucl. Mater. 1997, 249, 150–158. Natesan, K.; Uz, M. Fusion Eng. Des. 2000, 51–52, 145–152. Fujiwara, M.; Natesan, K.; Satou, M.; Hasegawa, A.; Abe, K. J. Nucl. Mater. 2002, 307–311, 601–604. Chen, J. M.; Muroga, T.; Qiu, S.; Xu, Y.; Den, Y.; Xu, Z. Y. J. Nucl. Mater. 2004, 325, 79–86. DiStefano, J. R.; Pint, B. A.; DeVan, J. H. J. Nucl. Mater. 2000, 283–287, 841–846. Rice, P. M.; Zinkle, S. J. J. Nucl. Mater. 258–263, 1414–1419. Satou, M.; Chuto, T.; Abe, K. J. Nucl. Mater. 2000, 283–287, 367–371. Watanabe, H.; Suda, M.; Muroga, T.; Yoshida, N. J. Nucl. Mater. 2002, 307–311, 408–411. Smith, D. L.; Matsui, H.; Greenwood, L.; Loomis, B. A. J. Nucl. Mater. 1988, 155–157, 1359–1363. Fukumoto, K.; Narui, M.; Matsui, H.; et al. J. Nucl. Mater. 2009, 386–388, 575–578. Yamauchi, Y.; Yamada, T.; Hirohata, Y.; Hino, T.; Muroga, T. J. Nucl. Mater. 2004, 329–333, 397–400. Hayakawa, R.; Hatano, Y.; Fukumoto, K.; Matsui, H.; Watanabe, K. J. Nucl. Mater. 2004, 329–333, 411–415. Homma, H.; Hatano, Y.; Daifuku, H.; et al. J. Nucl. Mater. 2007, 367–370, 887–891. Chen, J. M.; Nagasaka, T.; Muroga, T.; Qiu, S. Y.; Li, C.; Nita, N. J. Nucl. Mater. 2008, 374, 298–303. Muroga, T.; Nagasaka, T.; Chen, J. M.; Li, Y. F.; Watanabe, H. J. Nucl. Mater. 2009, 386–388, 606–609. Sakai, K.; Satou, M.; Fujiwara, M.; Takahashi, K.; Hasegawa, A.; Abe, K. J. Nucl. Mater. 2004, 329–333, 457–461. Kuwabara, T.; Kurishita, H.; Hasegawa, M. J. Nucl. Mater. 2000, 283–287, 611. Tyumentsev, A. N.; Korotaev, A. D.; Pinzhin, Y. P.; et al. J. Nucl. Mater. 2007, 367–370, 853–857. Kirillov, I. R.; Shatalov, G. E.; Strebkov, Y. S. The RF TBM Team. Fusion Eng. Des. 2006, 81, 425–432. Loomis, B. A.; Nowicki, L. J.; Smith, D. L. J. Nucl. Mater. 1994, 212–215, 790.
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Vanadium for Nuclear Systems
52. Zinkle, S. J.; Alexander, D. J.; Robertson, J. P.; et al. Eatherly. Fusion Mater. 1996, 73. DOE/ER-0313/21. 53. Tsai, H.; Nowicki, L. J.; Billone, M. C.; Chung, H. M. Smith, D. L. Fusion Mater. 1997, 70. DOE/ER-0313/23. 54. Tsai, H.; Gazda, J.; Nowicki, L. J.; Billone, M. C.; Smith, D. L. Fusion Mater. 1998, 15. DOE/ER-0313/24.
55. 56.
Chung, H. M.; Nowicki, L. J.; Smith, D. L. Fusion Mater. 1997, 29. DOE/ER-0313/22. Snead, L. L.; Zinkle, S. J.; Alexander, D. J.; Rowcliffe, A. F.; Robertson, J. P.; Eatherly, W. S. DOE/ER-0313/23 (Dec. 31, 1997) 81. Available at http://www.ms.ornl.gov/ fusionreactor/dec97.shtml
4.13
Concrete*
D. J. Naus Oak Ridge National Laboratory, Oak Ridge, TN, USA
ß 2012 Published by Elsevier Ltd.
4.13.1
Introduction
4.13.2
Concrete Longevity, NPP Safety-Related Concrete Structures, Testing and In-Service Inspection Requirements, and Operating Experience Historical Perspective on Concrete Longevity NPP Safety-Related Concrete Structures Boiling water reactors Pressurized water reactors Testing and In-Service Inspection Requirements Operating Experience Aging and Long-Term Durability Considerations Design, Construction, Material Selection, and Maintenance Considerations Materials of Construction, Degradation Mechanisms, Damage Modeling, and Long-Term Performance of Concrete Materials Materials of construction Degradation mechanisms Damage modeling Long-term performance of concrete materials Assessment and Repair Component selection In-service inspections Nondestructive examinations Remedial methods Structural Reliability Theory Summary and Potential Research Topics
4.13.2.1 4.13.2.2 4.13.2.2.1 4.13.2.2.2 4.13.2.3 4.13.2.4 4.13.3 4.13.3.1 4.13.3.2 4.13.3.2.1 4.13.3.2.2 4.13.3.2.3 4.13.3.2.4 4.13.3.3 4.13.3.3.1 4.13.3.3.2 4.13.3.3.3 4.13.3.3.4 4.13.4 4.13.5 References
Abbreviations ACI ANS ASME ASTM BWR C3A C2S C3S C4AF CEB-FIP CSNI
American Concrete Institute American Nuclear Society American Society of Mechanical Engineers ASTM International Boiling water reactor Tricalcium aluminate Dicalcium silicate Tricalcium silicate Tetracalcium aluminoferrite International Federation for Structural Concrete Committee on the Safety of Nuclear Installations
*Prepared for the Oak Ridge National Laboratory under Contract No. DE-AC05-00OR22725
408 409 409 409 412 413 415 416 417 419 419 419 422 422 422 424 424 425 425 426 427 428 428
CFR C-S-H GDC GGBFS
Code Federal Regulations Calcium silicate hydrate General Design Criteria Ground granulated blast furnace slag IAGE WG Integrity of Components and Structures Working Group LWR Light-water reactor NEA Nuclear Energy Agency NPP Nuclear power plant NSSS Nuclear steam supply system PCA Portland Cement Association PWR Pressurized water reactor RG Regulatory guide RILEM International Union of Laboratories and Experts in Construction Materials, Systems and Structures RPV Reactor pressure vessel
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USNRC WIS
United States Nuclear Regulatory Commission University of Wisconsin
4.13.1 Introduction As concrete ages, changes in its properties will occur as a result of continuing microstructural changes (i.e., slow hydration, crystallization of amorphous constituents, and reactions between cement paste and aggregates) as well as environmental influences. These changes do not have to be detrimental to the point where concrete will not be able to meet its functional and performance requirements; however, concrete can suffer undesirable changes with time because of improper specifications, violation of specifications, or adverse performance of its cement paste matrix or aggregate constituents under physical or chemical attack. Additional information related to environmental effects on concrete is provided in molten core concrete interaction (Chapter 2.25, Core Concrete Interaction). Portland cement concrete durability is defined as its ability to resist weathering action, chemical attack, abrasion, or any other process or deterioration.1 A durable concrete is one that retains its original form, quality, and serviceability in the working environment during its anticipated service life. The materials and mix proportions specified and used should be such as to maintain concrete’s integrity and, if applicable, to protect embedded metal from corrosion.2 The degree of exposure anticipated for the concrete during its service life, together with other relevant factors related to mix composition, workmanship, and design, should be considered.3 Guidelines for production of durable concrete are available in national consensus codes and standards, such as American Concrete Institute (ACI) 3184, which have been developed over the years through knowledge acquired in testing laboratories and supplemented by field experience. Serviceability of concrete has been incorporated into the codes through strength requirements and limitations on service load conditions in the structure (e.g., allowable crack widths, limitations on midspan deflections of beams, and maximum service level stresses in prestressed members). Durability generally has been included through items such as specifications for maximum water–cement ratios, minimum cementitious materials contents, type cementitious material, requirements for entrained air, and minimum concrete cover over reinforcement. Requirements are frequently specified in terms of environmental exposure
classes (e.g., chloride and aggressive ground environments). Specifications in terms of service life requirements (e.g., short <30 years, normal 30–100 years, and long >100 years) have only recently been developed, primarily through European standards.5 Water is the single most important factor controlling the degradation processes of concrete (i.e., the process of deterioration of concrete with time is generally dependent on the transport of a fluid through concrete), apart from mechanical deterioration. The rate, extent, and effect of fluid transport are largely dependent on the concrete pore structure (i.e., size and distribution), presence of cracks, and microclimate at the concrete surface. The primary mode of transport in uncracked concrete is through the cement paste pore structure (i.e., its permeability). The dominant mechanism controlling the rates of water penetration into unsaturated or partially saturated concrete is absorption caused by capillary action of the concrete’s pore structure. To improve the durability of concrete, generally the capillary and pore size within the concrete matrix should be reduced to a minimum. Although the coefficient of permeability for concrete depends primarily on the water–cement ratio and maximum aggregate size, it is influenced by the curing temperature, drying, cementitious materials content, and addition of chemical or mineral admixtures as well as the tortuosity of the path of flow. Concrete compressive strength has traditionally been utilized as an acceptance test for concrete, but it typically is not a good indicator of durability. Many structures have been fabricated with concretes having adequate 28-day compressive strength only to lose their functionality because they were facing an environment for which they had not been designed or because the concrete had not been placed or cured correctly.6 The safety-related concrete structures in nuclear power plants (NPPs) are designed to withstand loadings from a number of low-probability external and internal events, such as earthquake, tornado, and lossof-coolant accident. Consequently, they are robust and not subjected to high enough stresses during normal operation to cause appreciable degradation. In general, this has been the case, as the performance of reinforced concrete structures in NPPs has been good. (Operating experience is discussed in the next section.) However, as the NPPs age, degradation incidences start to occur at an increasing rate, primarily due to environmental-related factors. Onefourth of all containments in the United States have experienced corrosion, and nearly half of the concrete containments have reported degradation related to either the reinforced concrete or post-tensioning
Concrete
system.7 Although the vast majority of these structures will continue to meet their functional and performance requirements during their initial licensing period (i.e., nominally 40 years), it is reasonable to assume that with the increasing age of the operating reactors there will be isolated examples where the structures may not exhibit the desired durability without some form of intervention. Currently, the United States has 104 NPP units licensed for commercial operation, which provide about 20% of the electricity supply. As all but one of the construction permits for existing NPPs in the United States were issued prior to 1978, the focus for the existing plants has shifted from design to condition assessment. Here, the aim is to demonstrate that structural margins of the plants have not eroded or will not erode during the desired service life due to aging or environmental effects. One of the key factors to maintaining adequate structural margins to protect public health and safety in the unlikely event of an accident is implementation of effective inspection and maintenance programs. An inspection program is important for identifying and characterizing any degradation that may be present in a timely manner. Once degradation has been identified, or its potential to occur established, a maintenance program is implemented to repair the degradation and arrest (as far as possible) the mechanism(s) causing the degradation. Proper maintenance is essential to the safety of NPP structures, and a clear link exists between effective maintenance and safety. Uncertainty in condition assessment can be assessed using probabilistic methods, which are also an essential ingredient of risk-informed management decisions concerning continued service of the NPP structures.
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mortars were used by the Egyptians to fabricate the Great Pyramid at Giza about 2500 BC. The Romans were the first to use hydraulic limes and discovered the benefits of pozzolans. The survival of several ancient concrete structures (e.g., Pantheon in Rome, Figure 1) attests to the durability that concrete can attain. A detailed study involving an examination of samples obtained from several ancient concrete structures utilizing physical and chemical techniques concluded that these structures survived primarily because of careful selection of materials and construction, mild climatic conditions, and the lack of steel reinforcement.9 These structures, however, were not fabricated using current ‘hydraulic Portland cement,’ as it did not exist until about 1824. Some information, however, was presented in Mallinson and Davies9 relative to samples that were obtained for testing from several structures fabricated in the mid- to late 1800s. It was concluded that the durability of these structures was not only due to high cement contents but also due to the relatively slow cement-setting times and high construction quality. These Portland cements differ somewhat from the Portland cements used to fabricate NPP concrete structures in that the formulations have changed as well as the fineness of the cement. Also, modern concretes have incorporated admixtures to improve workability, modify hardening or setting characteristics, aid in curing, and enhance the performance or durability. 4.13.2.2 NPP Safety-Related Concrete Structures All commercial NPPs in the United States contain structures whose performance and function are necessary for the protection of the safety of
4.13.2 Concrete Longevity, NPP Safety-Related Concrete Structures, Testing and In-Service Inspection Requirements, and Operating Experience 4.13.2.1 Historical Perspective on Concrete Longevity Concrete, originally based on lime that hardened by atmospheric carbonation, has been utilized as a construction material for several thousand years. Cement has been around for at least 12 My when reactions occurred between limestone and oil shale during spontaneous combustion in Israel to form a natural deposit of cement compounds.8 The oldest known concrete is from Yugoslavia and is about 7600 years old.9 Gypsum
Figure 1 Pantheon, built 119–128 AD. From Http://en. wikipedia.org/wiki/file:Pantheon_rome_2005may.jpg
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plant-operating personnel and the general public, as well as the environment. The basic laws that regulate the design (and construction) of NPPs are contained in Title 10 of the Code of Federal Regulations (CFR),10 which is clarified by Regulatory Guides (e.g., R.G. 1.29),11 NUREG reports, Standard Review Plans (e.g., Concrete and Steel Internal Structures of Steel or Concrete Containments),12 etc. In addition, R.G. 1.29 and Part 100 to Title 10 of the CFR state that NPP structures important to safety must be designed to withstand the effects of earthquakes without the loss of function or threat to public safety. These ‘safety-related’ structures are designed as Seismic Category I. Seismic Category I structures typically include those classified by the American Society of Mechanical Engineers (ASME) and the American Nuclear Society (ANS) as Classes 1, 2, and 3 (i.e., safety related). Initially, existing building codes such as the ACI Standard 318, Building Code Requirements for Reinforced Concrete, were used in the nuclear industry as the basis for the design and construction of concrete structural members. However, because the existing building codes did not cover the entire spectrum of design requirements and because they were not always considered adequate, the United States Nuclear Regulatory Commission (USNRC) developed its own criteria for the design of Category I structures (e.g., definitions of load combinations for both operating and accident conditions). Current requirements for nuclear safety-related concrete structures, other than concrete reactor vessels and concrete containments, are also based on ACI 318, but have incorporated modifications to accommodate the unique performance requirements of NPPs. These requirements were developed by ACI Committee 349 and first published in October 1976.13 This Code has been endorsed by the USNRC as providing an adequate basis for complying with the general design criteria for structures other than reactor vessels and containments.14 USNRC15 provides additional information on the design of seismic Category I structures that are required to remain functional if the Safe Shutdown Earthquake (SSE) occurs. Current requirements for concrete reactor vessels and concrete containments were developed by ACI Committee 359 and first published in 1977.16 Supplemental load combination criteria are presented in Section 3.8.1 of the USNRC Regulatory Standard Review Plan.17 However, since all but one of the construction permits for existing NPPs have been issued prior to 1978, it is unlikely that endorsed versions of either ACI 349 or ACI 359 were used in the design of many of the concrete
structures at these plants. Older plants that used early ACI codes, however, have been reviewed by the USNRC through the Systematic Evaluation Program to determine if there were any safety concerns.18 Each boiling water reactor (BWR) or pressurized water reactor (PWR) unit in the United States is located within a much larger metal or concrete containment that also houses or supports the primary coolant system components. Although the shapes and configurations of the containment can vary significantly from plant to plant, leak tightness is ensured by a continuous pressure boundary consisting of nonmetallic seals and gaskets and metallic components that are either welded or bolted together. There are several CFR General Design Criteria (GDC) and ASME Code sections that establish minimum requirements for the design, fabrication, construction, testing, and performance of the light-water reactor (LWR) containment structures. The GDC serve as fundamental underpinnings for many of the most important safety commitments in licensee design and licensing bases. General Design Criterion 2, Design Bases for Protection Against Natural Phenomena, requires the containment to remain functional under the effects of postulated natural phenomena such as earthquakes, tornadoes, hurricanes, floods, tsunami, and seiches. General Design Criterion 16, Containment Design, requires the provision of reactor containment and associated systems to establish an essentially leaktight barrier against the uncontrolled release of radioactivity into the environment and to ensure that the containment design conditions important to safety are not exceeded for as long as required for postulated accident conditions. Criterion 53, Provisions for Containment Testing and Inspection, requires that the reactor containment be designed to permit (1) appropriate periodic inspection of all important areas, such as penetrations; (2) an appropriate surveillance program; and (3) periodic testing at containment design pressure of leak tightness of penetrations that have resilient seals and expansion bellows. Current LWR containments are considered as a significant element of the USNRC’s safety policy, which employs a defensein-depth approach (i.e., successive compensatory measures are exercised to prevent accidents or mitigate damage if a malfunction, accident, or naturally caused event occurs). The defense-in-depth philosophy ensures that safety will not be wholly dependent on any single element of the design, construction, maintenance, or operation at a nuclear facility (e.g., the facility in question tends to be more tolerant of failures and external challenges).
Concrete
From a safety standpoint, the containment is one of the most important components of an NPP because, independent of the fuel barrier and reactor coolant pressure boundary barrier, it serves as the final barrier to the release of fission products to the outside environment under postulated accident conditions. During normal operating conditions, the containment is subject to various operational and environmental stressors (e.g., ambient pressure fluctuations, temperature variations, earthquakes, and wind storms). In some containment designs, the principal leak-tight barrier is surrounded by another structure (e.g., reactor or shield building) that protects the containment from external events. Ensuring Table 1
411
that the structural capacity and leak-tight integrity of the containment has not deteriorated unacceptably because of aging or environmental stressor effects is essential to reliable continued service evaluations and informed aging management decisions. More detailed information on containments is available.19 In addition to the containment, a myriad of concrete-based structures are contained as a part of an LWR plant to provide foundation, support, shielding, and containment functions. Table 1 presents a listing of typical safety-related concrete structures that may be included as part of an LWR plant.20 Relative to general civil engineering reinforced concrete structures, NPP concrete structures tend to be
Typical safety-related concrete structures in LWR plants and their accessibility for visual examination
Concrete structure Primary containment Containment dome/roof Containment foundation/basemat Slabs and walls Containment internal structures Slabs and walls Reactor vessel support structure (or pedestal) Crane support structures Reactor shield wall (biological) Ice condenser dividing wall (ice condenser plants) NSSS equipment supports/vault structures Weir and vent walls (Mark III) Pool structures (Mark III) Diaphragm floor (Mark II) Drywell/wetwell slabs and walls (Mark III) Secondary containment/reactor buildings Slabs, columns, and walls Foundation Sacrificial shield wall (metallic containments) Fuel/equipment storage pools Walls, slabs, and canals Auxiliary building Fuel storage building Control room (or building) Diesel generator building Piping or electrical cable ducts or tunnels Radioactive waste storage building Stacks Intake structures (including concrete water intake piping and canal embankments) Pumping stations Cooling towers Plant discharge structures Emergency cooling water structures Dams Water wells Turbine building
Accessibility Internal liner/complete external Internal liner (not embedded) or top surface Internal liner/external above grade Generally accessible Typically lined or hard to access Generally accessible Typically lined Lined or hard to access Generally accessible Lined with limited access Lined Lined with limited access Internal liner/partial external access Accessible on multiple surfaces Top surface Internal lined/external accessible Internal lined/partial external Generally accessible Generally accessible Generally accessible Generally accessible Limited accessibility Generally accessible Partial internal/external above grade Internal accessible/external above grade and waterline Partially accessible Accessible above grade Internal accessible/external above grade and waterline Limited accessibility External surfaces above waterline Limited accessibility Generally accessible
Source: Hookham, C. J. In-Service Inspection Guidelines for Concrete Structures in Nuclear Power Plants, ORNL/NRC/LTR-95/14; Lockheed Martin Energy Systems, Oak Ridge National Laboratory: Oak Ridge, TN, 1995.
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more massive and have increased steel reinforcement densities with more complex detailing. Information pertaining to a particular structure at a plant of interest can be obtained from sources such as the plant’s safety analysis report or docket file. Concrete structures that are considered to be ‘plant specific’ or unique have not been addressed in the discussion later, but some information provided for similar structures may be applicable. Additionally, the names of certain structures may vary from plant to plant depending on the nuclear steam supply system (NSSS) vendor, architect engineering firm, and owner preference. Typical safety-related concrete structures contained in LWR plants may be grouped into four categories: primary containments, containment internal structures, secondary containment/ reactor buildings, and other structures. 4.13.2.2.1 Boiling water reactors
Of the BWR plants that have been licensed for commercial operation in the United States, 30% utilize either reinforced or prestressed concrete primary containments. Leak tightness of each of these containments is provided by a steel liner attached to the containment inside surface by studs (e.g., Nelson studs) or by structural steel members. Exposed surfaces of the carbon steel liner are typically painted to protect against corrosion and to facilitate decontamination should it be required. A portion of the liner toward the bottom of the containment and over the basemat is typically embedded in concrete to protect it from damage, abrasion, etc. due to corrosive fluids and impact. A seal to prevent the ingress of fluids is provided at the interface around the circumference of the containment where the vertical portion of the liner becomes embedded in the concrete. BWR containments, because of provisions for pressure suppression, typically have ‘normally dry’ sections (drywell) and ‘flooded’ sections (wetwell) that are interconnected via piping or vents. Requirements for BWR containments include the following: 1. Provide an ‘essentially’ leak-tight barrier against the uncontrolled release of radioactivity to the environment for all postulated design basis accident conditions. 2. Accommodate the calculated pressure and temperature conditions resulting from a loss-ofcoolant accident. 3. Withstand periodic integrated leak-rate testing at the peak-calculated accident pressure that may be
at levels up to and including the containment design pressure. 4. Permit appropriate periodic inspection of all important components and surfaces, and the periodic testing of the leak tightness of containment penetrations. The containment vessel can also provide structural support for the NSSS and other internal equipment. The containment foundation, typically a basemat, provides the primary support and transfer of load to the earth below. Figure 2 presents a cross-section of a BWR Mark I reinforced concrete containment. Each of the three BWR primary plant types (Mark I, Mark II, and Mark III) incorporates a number of reinforced concrete containment internal structures. These structures may perform singular or several functions, including the following: 1. Radiation shielding; 2. human accessibility provisions; 3. NSSS and other equipment anchorage/support/ protection; 4. resistance to jet, pipe whip, and other loadings produced by emergency conditions; 5. boundary of wetwells and pool structures, and allow communication between drywell and wetwell (Mark II and III); 6. lateral stability for containment; 7. transfer of containment loads to underlying foundation; and 8. transfer of fuel to reactor (Mark III). As many of these functions are interrelated with the required containment functions, these structures are considered to be safety-related. Of the BWR plants that utilize steel primary containments, all but the pre-Mark plant type have reinforced concrete structures that serve as secondary containments or reactor buildings and provide support and shielding functions for the primary containment. Although the design parameters for the secondary containments of the Mark I and Mark II plants vary somewhat, the secondary containments are typically composed of beam, floor, and wall structural elements. These structures typically are safetyrelated because they provide additional radiation shielding; provide resistance to environmental/operational loadings; and house safety-related mechanical equipment, spent fuel, and the primary metal containment. Although these structures may be massive in cross-section to meet shielding or load-bearing requirements, they generally have smaller elemental
Concrete
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Polar crane
Equipment pool
Fuel storage pool
Reactor building
Drywell head Concrete containment vessel
Reactor pressure vessel
Metal liner
Biological shield Core Steel liner
Vent Grade
Reactor pressure vessel support skirt
Drywell Reactor pressure vessel pedestal
Carbon steel and stainless steel (below waterline liner)
Downcomer pipe Pressure suppression chamber Basemat Figure 2 Boiling water reactor Mark I reinforced concrete containment and reactor building.
thicknesses than primary containments because of reduced exposure under postulated accident loadings. These structures may be maintained at a slight negative pressure for collection and treatment of any airborne radioactive material that might escape during operating conditions. Other structures include such things as foundations, walls, slabs, and fuel/equipment storage pools. The spent- and new-fuel storage pools, and the pools for reactor internals storage, typically have a fourwall-with-bottom-slab configuration. The walls and slab are composed of reinforced concrete members lined on the interior surface with stainless steel. Cross-sections of these members are generally large because they must support a large pool of water
and heavy fuel/component loads produced by high-density fuel storage considerations. The fuel storage pool in Mark III plants is located within the primary containment. 4.13.2.2.2 Pressurized water reactors
Of the PWR plants that have been licensed for commercial operation in the United States, 80% utilize either reinforced or prestressed concrete primary containments. In meeting the same basic functional and performance requirements as noted for BWR containments, the concrete containments in PWR plants are of three different functional designs: subatmospheric (reinforced concrete), ice condenser (reinforced concrete), and large/dry
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Concrete
(reinforced and prestressed concrete). The primary differences between these containment designs relate to volume requirements, provisions for accident loadings/pressures, and containment internal structures layout. The PWR containment structure generally consists of a concrete basemat foundation, vertical cylindrical walls, and dome. The basemat may consist of a simple mat foundation on fill, natural cut, or bedrock or may be a pile/pile cap arrangement. Most of the plants have utilized the simple mat on fill or bedrock design. Interior containment surfaces are lined with a thin carbon steel liner to prevent leakage. Exposed surfaces of the carbon steel liner are typically painted to protect against corrosion and to facilitate decontamination should it be required. Depending on the functional design, the concrete containments can be on the order of 40–50 m in diameter and 60–70 m high, with wall and dome thicknesses from 0.9 to 1.4 m and base slab
thicknesses from 2.7 to 4.1 m. Two of the PWR plants (Bellefonte and Ginna) have rock anchor systems to which the post-tensioning tendons are attached. Figure 3 presents a cross-section for a prestressed concrete, large, dry containment. The containment internal structures in PWR plants are typically constructed of conventionally reinforced concrete and tend to be more massive in nature than the internal structures in BWR plants, because they typically support the reactor pressure vessel, steam generators, and other large equipment and tanks. In addition, these structures provide shielding of radiation emitted by the NSSS. Some of the specific functions that these structures (typically floor slabs, walls, and columns) are required to perform include the following: 1. provision of human accessibility; 2. support and separation of various plant equipment;
Dome Ring girder (tendon anchorage)
Steel liner
42.7 m Polar crane Steam generators
64.7 m
Prestressed reinforcing
Seal table Grade
Basemat 8m In-core instrument guide tubes Reactor vessel
2.5 m Reactor cavity
Figure 3 Pressurized water reactor large dry prestressed concrete containment.
Concrete
3. resistance to emergency loading conditions; 4. transfer of containment loads to containment foundation; 5. missile protection; and 6. channeling/routing steam and air through ice condensers (PWR ice condenser containments). PWR plants that utilize a metallic primary containment (large dry and ice condenser designs) are usually contained in reinforced concrete ‘enclosures’ or ‘shield’ buildings that, in addition to withstanding environmental effects, provide radiation shielding and particulate collection and ensure that the freestanding metallic primary containment is protected from the natural environment. The secondary containment consists of a vertical cylinder wall with shallow dome and is often supported by the containment basemat. Except for differences in the spent- and new-fuel storage pools, structures that fall into the other structures category are essentially the same at the PWR and BWR plants. The spent- and new-fuel storage pools for PWR plants are typically located in an auxiliary building proximate to the containment. These reinforced concrete wall and slab structures are generally massive in cross-section to support a large pool of water and the fuel elements and are lined on the water side with stainless steel. The pools are connected to the reactor/refueling cavity (inside containment) via a transfer channel that is also a safety-related structure since it must provide radiation shielding and support for the fuel transport mechanism and fuel. 4.13.2.3 Testing and In-Service Inspection Requirements One of the conditions of all operating licenses for water-cooled power reactors in the United States is that the primary reactor containments shall meet the containment leakage test requirements set forth in Appendix J, Primary Reactor Containment Leakage Testing for Water-Cooled Power Reactors, to 10 CFR Part 50.21 These test requirements provide for preoperational and periodic verification by tests of the leak-tight integrity of the primary reactor containment as well as systems and components that penetrate containment of water-cooled power reactors and establish the acceptance criteria for such tests. The purpose of these tests is to ensure that (1) leakage through the primary reactor containment and the systems and components penetrating primary reactor containment shall not exceed allowable leakage-rate values
415
as specified in the technical specifications or associated bases and (2) periodic surveillance of reactor containment penetrations and isolation valves is performed so that proper maintenance and repairs are made during the service life of the containment as well as systems and components that penetrate primary containment. Contained in this regulation are requirements pertaining to Type A, B, and C leakagerate tests that must be performed by each licensee as a condition of their operating license. Type A tests are intended to measure the primary reactor containment overall integrated leakage rate (1) after the containment has been completed and is ready for operation and (2) at periodic intervals thereafter. Type B tests are intended to detect local leaks and to measure leakage across each pressure-containing or leakage-limiting boundary for primary reactor containment penetrations (e.g., penetrations that incorporate resilient seals, gaskets, or sealant compounds and air lock door seals). Type C tests are intended to measure containment isolation valve leakage rates. Requirements for system pressure testing and criteria for establishing inspection programs and pressure-test schedules are contained in Appendix J. Appendix J to 10 CFR Part 50 also requires a general inspection of the accessible interior and exterior surfaces of the containment structures and components to uncover any evidence of structural deterioration that may affect either the containment structural integrity or leak tightness. Subsection IWL of ASME Section XI addresses reinforced and posttensioned concrete containments (Class CC). Two examination categories are provided in Subsection IWL. Examination Category L-A addresses accessible concrete surfaces and examines them for evidence of damage or degradation, such as cracks. The concrete is examined at 1, 3, and 5 years following the containment structural integrity test and every 5 years thereafter. The primary inspection method of Category L-A is visual examination (general or detailed). Examination Category L-B addresses the unbonded post-tensioning system. The unbonded post-tensioning system examination schedule is the same as for the concrete. For post-tensioned concrete containments, tendon wires are tested for yield strength, ultimate tensile strength, and elongation. Tendon corrosion protection medium is analyzed for alkalinity, water content, and soluble ion concentrations. Prestressing forces are measured for selected sample tendons. Subsection IWL specifies acceptance criteria, corrective actions, and expansion of the inspection scope when degradation exceeding
416
Concrete
the acceptance criteria is found. Additional requirements for inaccessible areas are specified in 10 CFR 50.55a(b)(2)(viii). The acceptability of concrete in inaccessible areas is to be evaluated when conditions that could indicate the presence or result in degradation to such inaccessible areas exist in accessible areas. Information on aging management programs for masonry walls22,23 and water-control structures24 is available. Inspection requirements for steel containments and liners of concrete containments are contained in Subsection IWE of ASME Section XI. Editions and addenda of the ASME Code acceptable to the USNRC are identified in 10 CFR 50.55a. 4.13.2.4
that have been observed by one organization during condition surveys of various concrete structures at both United States and foreign NPPs located in areas having several different climatic conditions.25 Some general observations derived from these results were that virtually all NPPs have experienced cracking of the concrete structures that exceeds typical acceptance criteria for width and length; numerous NPPs had groundwater intrusion occurring through the power block or other subsurface structures; and aging concerns exist for subsurface concrete structures, as their physical condition cannot be easily verified. Collectively, it was concluded in this study that the general performance of the NPP concrete structures, has been quite favorable and proper evaluation and treatment of observed degradation at an early stage is both a costeffective and necessary approach to long-term plant operations. Initially, degradation of NPP concrete structures in the United States occurred early in their life and has been corrected.26–28 Causes were primarily related either to improper material selection and construction/design deficiencies or environmental
Operating Experience
In general, the performance of NPP safety-related concrete structures has been very good. However, there have been several isolated incidences that, if not remedied, could challenge the capacity of the containment and other safety-related structures to meet future functional and performance requirements. Table 2 presents a summary of local degradation mechanisms Table 2
Condition survey results from several plants for NPP structures
Local degradation mechanism Concrete Chemical attack Efflorescence and leaching Alkali–aggregate reaction Freeze–thaw cycling Thermal exposure Abrasion/erosion Fatigue/vibration Cracking Conventional reinforcement Corrosion Prestressing system Corrosion Block walls Excessive cracking Structural steel and liners Corrosion Soil/structure issues Differential settlement Soil erosion (scour)
Plant A
B
C
D
E
F
G
H
I
J
b,c
c b,c,d
b b,c
c
c b,d
c b,d
c d
c b,d,f
a,b,c,d
d
f c
b,f a
b,c,d,f
b,c,d,f
b,f,g
b,d
b
b,f
e
g
d
c c
c,d,f,g
c a,b,c,d
c,d,g
b,d
b,d
b
a,d c
c
c,d
a,b,c,d,g
b,c,d,f,g
b,d
b
c,d b,f
e1 c d
e
d c,d
c c,e
a
c d
a – external structure (power block); b – subgrade structure (power block); c – internal structure (power block); d – water control structure (Intake, discharge, etc.); e – Containment vessel (power block); f – Other site structure (power block); g – Equipment supports (power block). Source: Gregor, F. E.; Hookham, C. J. Remant life preservation of LWR plant structures. In Transactions of the 12th International Conference on Structural Mechanics in Reactor Technology, Stuttgart, Germany, Aug 15–20; Elsevier Science: Amsterdam, The Netherlands, 1993; Paper DH06/2. 1 Corrosion limited to exposed grease can and bearing plate surface.
Concrete
effects. Examples of some of the problems attributed to these deficiencies include low 28-day concrete compressive strengths; voids under the post-tensioning tendon-bearing plates resulting from improper concrete placement; cracking of post-tensioning tendon anchor heads due to stress corrosion or embrittlement; and containment dome delaminations due to lowquality aggregate materials and absence of radial steel reinforcement or unbalanced prestressing forces.29–31 Other construction-related problems included occurrence of excessive voids or honeycomb in the concrete, contaminated concrete, cold joints, cadweld (steel reinforcement connector) deficiencies, materials out of specification, higher than code allowable concrete temperatures, misplaced steel reinforcement, posttensioning system button-head deficiencies, and water-contaminated corrosion inhibitors.26 Although continuing the service of a NPP past the initial operating license period is not expected to be limited by the concrete structures, several incidences of agerelated degradation have been reported.28–33 Examples of some of these problems include corrosion of steel reinforcement in water intake structures, corrosion of post-tensioning tendon wires, leaching of tendon gallery concrete, low prestressing forces, and leakage of corrosion inhibitors from tendon sheaths. Other related problems include cracking and spalling of containment dome concrete due to freeze–thaw damage,
Concrete cracking outside containment wall
Crease leakage outside containment wall
Exterior concrete wallcracks and spalling
Anchor head failure
417
low strengths of tendon wires, contamination of corrosion inhibitors by chlorides, and corrosion of concrete containment liners. As the plants age, the incidences of degradation are expected to increase, primarily due to environmental effects. A listing of documented concrete problem areas by plant, type reactor, and degradation is available.34 Documented information on problem areas experienced with NPP concrete structures in other countries has also been assembled.35 Figure 4 presents examples of occurrences of degradation that have been observed at NPPs. Anchor head failure and containment dome delamination shown in the figure represent occurrences related to materials selection and design, respectively, with the remainder representing aging-related occurrences.
4.13.3 Aging and Long-Term Durability Considerations In the United States, the Atomic Energy Act and regulations of the USNRC limit commercial power reactor licenses to an initial 40-year period, but permits such licenses to be renewed. (Other countries may not have a limit set on the plant operating license period, but the utility must obtain a permanent renewal of its operating license subject to numerous and continuous justifications (e.g., periodic safety
Concrete wall water infiltration
Containment dome delamination repair
Figure 4 Examples of degradation related to nuclear power plant concrete structures.
Corrosion of grease cap
Water intake structure rebar corrosion
418
Concrete
Performance in durability terms
reevaluations).) This original 40-year term for reactor licenses was based on economic and antitrust considerations – not on limitations of nuclear technology. Due to this selected period, however, some structures and components may have been engineered on the basis of an expected 40-year service life. Several nuclear power units in the United States have reached the end of their initial operating license period. To help ensure an adequate energy supply, the USNRC has established a timely license renewal process and clear requirements that are needed to ensure safe plant operation for an extended plant life. These requirements are codified in Parts 51 and 54 of Title 10, Energy, of the CFR that provides for a renewal of an operating license for an additional 20 years. In order to ensure the safe operation of NPPs, it is essential that the effects of age-related degradation of plant structures, as well as systems and components, be assessed and managed during both the current operating license period as well as subsequent license renewal periods. As these plants mature, environmental factors are going to become increasingly important. Demonstration of continued safe and reliable operation of the plants will involve implementation of a program that effectively manages aging to ensure the availability of design safety functions throughout the plant service life. Examples of considerations to be addressed by such a program for the safety-related concrete structures are identified in Figure 5 and include the following:
Actual performance
1. What environmental stressors or aging factors are most important with respect to impacting structural reliability? 2. What in-service inspection or condition assessment programs are most effective in demonstrating structural reliability, and how often should they be applied? 3. What material sampling and testing programs should be required, if any? 4. How effective are remedial measures in enhancing the reliability of the structures and extending their usable life? 5. How have the material properties changed under the influence of aging and environmental stressors? 6. What is the residual life of the structure and how might it respond to something like a design basis event? General guidance on developing an aging management program for concrete containment buildings has been developed.35 Included in this reference is information related to practices and techniques that have been utilized by various countries for assessing the fitness for service as well as inspection, monitoring, and mitigation of aging degradation of concrete containment buildings. The International Union of Laboratories and Experts in Construction Materials, Systems and Structures (RILEM) has held an international conference, prepared a report, and sponsored two workshops related to aging management of concrete structures.36–39 Finally, the Nuclear Energy
Performance spread
Example: not NPP
Minimum required performance
I0
Repair implemented
Imin Ir
Time
Current and future condition assessments
Environmental stressors/ aging factors
Example: not NPP
Characterization of materials – database
Remedial measures
Material sampling/testing
In-service inspection
Figure 5 Components of an aging management program for nuclear power plant concrete structures.
Concrete
Agency Committee on the Safety of Nuclear Installations (NEA/CSNI) under its Integrity of Components and Structures Working Group (IAGE WG) has prepared several reports and held a series of workshops that addressed various aspects of aging of NPP concrete structures.40–52 Also, there are a number of other documents that address aging of NPP concrete structures,32,53–56 as well as national programs. 4.13.3.1 Design, Construction, Material Selection, and Maintenance Considerations Design errors that can lead to subsequent deterioration of concrete structures can be placed into two categories: inadequate structural design and lack of attention to details.57 Inadequate structural design occurs when the structure is exposed to a load greater than it is capable of carrying or if it sustains greater strain than its strain capacity. Inadequate considerations of temperature change or concrete creep and accidental impact can also result in damage. Typical symptoms of inadequate design include spalling and cracking of concrete. Poor detailing of a structure may result in localized concentration of stresses that result in cracking, which in turn can permit water or chemicals to access the concrete or ponding of water to produce saturated concrete. Poor detailing does not generally lead directly to concrete failure but can contribute to the action of one of the other specific causes of concrete failure.57 Examples of inadequate structural design include insufficient concrete cover over steel reinforcement, improper sizing and placement of steel reinforcement, inadequate section geometry, inadequate provision for drainage, abrupt changes in section, material incompatibility, and inadequate provision for deflection. Poor construction practices and negligence can result from not following specified procedures or from carelessness. Poor construction practices do not lead directly to failure or deterioration of concrete but can cause defects that lead to concrete cracking. Examples of concrete cracks that can result from poor construction practices include plastic shrinkage, plastic settlement, early thermal contraction, crazing, and long-term drying shrinkage. The resulting concrete cracking then can enhance the adverse impacts of mechanisms (such as described in the next section) and lead to concrete degradation. Poor construction practices and negligence are best addressed through adequate quality assurance/quality control in conjunction with an aggressive inspection program. Examples of poor construction practice include adding additional
419
water to concrete to facilitate placement or finishing, improper mixing and curing, improper consolidation, and improper location of steel reinforcement. Lack of knowledge about the importance of careful selection and specification of materials and use of admixtures can also result in durability issues. This can include improper cement contents, use of poor quality or contaminated aggregates, incorporation of additives that can produce corrosion such as calcium chloride accelerators, and incorrect water–cement ratios. Improper or inadequate maintenance also can contribute to the deterioration of concrete structures. Examples of inadequate maintenance include moisture exposure and penetration caused by unrepaired cracks, improper application of coatings, damaged waterstops, and failure to clean drains and drain pathways. 4.13.3.2 Materials of Construction, Degradation Mechanisms, Damage Modeling, and Long-Term Performance of Concrete Materials 4.13.3.2.1 Materials of construction
Nuclear safety-related concrete structures are composed of several constituents that, in concert, perform multiple functions (e.g., load-carrying capacity, radiation shielding, and leak tightness). Primarily, these constituents can include the following material systems: concrete, conventional steel reinforcement, prestressing steel, and steel liner plate. The quality of these materials is established through regulations, qualification tests, and certification, followed by checking throughout construction. More detailed information on materials of construction than provided later is available elsewhere.35,58–60 Concrete is a composite material consisting of a binder (cement paste) and a filler of fine or fine and coarse aggregate particles that combine to form a synthetic conglomerate. Cement is a mixture of compounds made by grinding crushed limestone, clay, sand, and iron ore together to form a homogeneous powder that is then heated at very high temperatures ranging from 1400 to 1600 C to form a clinker.59 After the clinker cools, it is ground and mixed with a small amount of gypsum to regulate setting and facilitate placement. This produces the general-purpose Portland cement, which is mixed with water to produce cement paste that binds the aggregate particles together. (Current generation cements have higher tricalcium silicate (C3S) contents and are ground finer than previous cements. Current
420
Concrete
cements attain most of their compressive strength within a 28-day period, whereas the previous cements continued to gain strength after 28 days.6,61) Portland cements are composed primarily of four chemical compounds: (C3S), dicalcium silicate (C2S), tricalcium aluminate (C3A), and tetracalcium aluminoferrite (C4AF). The type of Portland cement produced (e.g., general purpose, moderate sulfate resistance and heat of hydration, high early strength, low heat of hydration, and sulfate resistant) depends on the relative amounts of the four basic chemical compounds and fineness (high early strength). The calcium silicate hydrates (C–S–H) constitute about 75% of the mass. The C–S–H gel structure is made up of three types of groups that contribute to bonds across surfaces or in the interlayer of partly crystallized tobermorite material: calcium ions, siloxanes, and water molecules. Bonding of the water within the layers (gel water) with other groups via hydrogen bonds determines the strength, stiffness, and creep properties of the cement paste. There are also a number of alternative or supplementary cementing agents that have been used in conjunction with Portland cement, and these are pulverized fly ash, ground granulated blast furnace slag (GGBFS), and silica fume. Fly ash is collected from the exhaust flow of furnaces burning finely ground coal and reacts with calcium hydroxide in the presence of water to form cement compounds consisting of calcium silicate hydrate. GGBFS is a by-product of the iron-making process and is formed by taking the hot slag, rapidly chilling or quenching it, and grinding into a powder. When mixed with water in the presence of an alkaline environment provided by the Portland cement, GGBFS hydrates to form cementing compounds consisting of calcium silicate hydrate. Silica fume is the condensed vapor by-product of the ferrosilicon smelting process. Silica fume reacts with calcium hydroxide in the presence of water to form cementing compounds consisting of calcium silicate hydrate. High alumina cement, consisting mainly of calcium aluminates, has been utilized as a cementitious material because of its rapid set and rapid strength gain characteristics and resistance to acidic environments, sea water, and sulfates. However, owing to certain conditions of temperature and humidity, the cement converts over time to a different hydrate having reduced volume (i.e., increased porosity and reduced strength), it is recommended that calcium aluminate cements not be used for structural applications (particularly in wet or humid conditions above 27 C).62
Selection of the proper water content of concrete is critical, since too much water reduces the concrete strength and insufficient water makes the concrete unworkable. Hardening of concrete occurs as a result of hydration, which is a chemical reaction in which the major compounds in the cement form chemical bonds with water molecules and become hydrates. The hardened cement paste consists mainly of calcium silicate hydrates, calcium hydroxide, and lower proportions of calcium sulfoaluminate hydrate either as ettringite or monosulfate. About 20% of the hardened cement paste volume is calcium hydroxide. The pore solution is normally a saturated solution of calcium hydroxide within which high concentrations of potassium and sodium hydroxides are present. Proper curing of the concrete during this stage is essential, as it affects the concrete’s durability, strength, water-tightness, abrasion resistance, volume stability, and resistance to freezing and thawing. Since cement is the most expensive ingredient in concrete, it is desirable to utilize the minimum amount necessary to produce the desired properties and characteristics. Aggregate typically occupies 60–75% of the volume of concrete and therefore its characteristics strongly influence the chemical, physical, and thermal properties of concrete, its mix proportions, and economy. (The balance of the concrete mix generally consists of 10–15% cement, 15–20% water, and air (5–8% if entrained).) Aggregates thus are important with respect to the concrete durability. The aggregates come in various shapes, sizes, and material types ranging from fine sand particles to large coarse rocks. Selection of the aggregate material is determined in part by the desired characteristics of the concrete. Aggregate materials are available ranging from ultra-lightweight (e.g., vermiculite and perite) to lightweight (e.g., expanded clay shale or slate-crushed brick), normal weight (e.g., crushed limestone or river gravel), and heavyweight (e.g., steel or iron shot). Sometimes chemical or mineral admixtures are added during the mixing process to enhance durability (air entrainment), improve workability (enhanced placement and compaction), modify hardening and setting characteristics, aid in curing, reduce heat evolution, or provide other property improvements.63 The concrete typically used in nuclear safetyrelated structures consists of Type II Portland cement,59 fine aggregates (e.g., sand), water, various minerals, or chemical admixtures for improving properties or performance of the concrete and either normal-weight or heavy-weight coarse aggregate.
Concrete
American Society of Testing and Materials (ASTM) C 150,64 Type II Portland cement, typically has been used because of its improved sulfate resistance and reduced heat of hydration relative to the generalpurpose Type I Portland cement. Both the water and fine and coarse aggregates are normally acquired from local sources and subjected to material characterization testing prior to use. Coarse aggregate can consist of gravel, crushed gravel, or crushed stone. Chemical (e.g., air-entraining or water-reducing) or mineral (e.g., fly ash or ground granulated blast furnace slag) admixtures have been utilized in many of the mixes to impart improved characteristics or performance. For those concrete structures in NPPs that provide primary (biological) radiation shielding, heavy-weight or dense aggregate materials, such as barites, limonites, magnetites, and ilmenites, may have been used to reduce the section thickness and meet attenuation requirements. The constituents are proportioned and mixed to develop Portland cement concrete that has specific properties. Depending on the characteristics of the specific structure, the concrete mix may be adjusted to provide increased strength, higher durability, or better workability for placement. The hardened concrete typically provides the compressive loadcarrying capacity for the structure. Specified concrete unconfined compressive strengths typically have ranged from 13 to 55 MPa, with 35 MPa being a typical value achieved at 28 days age. Concrete tensile strength is about one-tenth to one-fifth of its compressive strength, so concrete cannot be relied upon to withstand very high tensile stresses. This limitation is overcome by embedding steel reinforcement in the concrete so that the concrete and steel work in concert. In addition to resisting tensile loads, the bonded steel reinforcement is used to control the extent and width of cracks, especially where it is desirable to reduce member cross-sections. Steel reinforcement is also used in compression members to safeguard against the effects of unanticipated bending moments that could crack or even fail the member. The effectiveness of reinforced concrete as a structural material depends on the interfacial bond between the steel and concrete so that it acts as a composite material, the passivating effect of the highly alkaline concrete environment to inhibit steel corrosion (see next section), and the similar coefficients of thermal expansion of the concrete and steel. Most of the mild, or conventional, reinforcing steels used in NPPs to provide primary tensile and shear load resistance/transfer consist of
421
plain carbon steel bar stock with deformations (lugs or protrusions) on the surface. These bars typically conform to ASTM A 61565 or A 70666 specifications. The minimum yield strength for the steel reinforcement ranges from 280 to 520 MPa, with the 420 MPa strength material being most common. Post-tensioning is a method of reinforcing (or strengthening) concrete with high-strength tendons to resist tensile loadings and to apply compressive forces to the concrete to provide increased resistance to concrete cracking. A number of NPP concrete containment structures utilize post-tensioned steel tendons that are designed to have (1) consistently high strength and strain at failure, (2) serviceability throughout their lifetime, (3) reliable and safe prestressing procedures, and (4) ability to be retensioned and replaced (nongrouted systems). The tendons are installed within preplaced ducts in the containment structure and post-tensioned from one or both ends after the concrete has achieved sufficient strength. After tensioning, the tendons are anchored by button-heads, wedges, or nuts. Corrosion protection is provided by filling the ducts with wax or corrosioninhibiting grease (unbonded) or portland cement grout (bonded). (Although bonded post-tensioning tendons are less vulnerable to local damage than ungrouted tendons, ungrouted tendons have been primarily used in the United States because the grouted tendon systems cannot be visually inspected, mechanically tested, or retensioned in the event of a larger than anticipated loss of prestressing force.) Supplemental conventional reinforcing is also used to minimize shrinkage or temperature effects and to provide local load-carrying capacity or load transfer. Three major categories of post-tensioning system exist depending on the type of material utilized to fabricate the tendons: wire, strand, or bar that conform to ASTM specifications A 421,67 A 416,68 and A 722,69 respectively. Minimum tensile strengths range from 1620 to 1725 MPa for the A 421 material and 1725 to 1860 MPa for the A 416 material. The A 722 material has a minimum tensile strength of 1035 MPa. Typical NPP tendon systems group sufficient numbers of wires, strands, or bars to have minimum ultimate strengths ranging from 2000 to 10 000 kN. The trend has been to increase the strength of the tendons to reduce the total number (e.g., in the early 1970s, the typical tendon had a capacity of 3000 kN and since then has progressed to capacities of 10 300 and 15 300 kN).19 With the exception of Robinsion 2 (bar tendons) and Three Mile Island 2 (strand tendons), plants that have post-tensioned
422
Concrete
containments utilize unbonded tendons so that the tendons can be inspected and replaced (if necessary). Bellefonte and Ginna each has grouted tendons (rock anchors) to which tendons are attached. Leak tightness of reinforced and post-tensioned concrete containment vessels is provided by a steel liner plate. A typical liner is composed of steel plate stock <13 mm thick, joined by welding, and anchored to the concrete by studs (Nelson studs or similar conforming to ASTM A 10870), structural steel shapes, or other steel products. PWR containments and the drywell portions of BWR containments are typically lined with carbon steel (ASTM A 3671 or A 51672). The liners of LWR fuel pool structures typically consist of stainless steel (ASTM A 27673 or A 30474). The liners of wetwells also have used carbon steel materials such as ASTM A 285,75 A 516, and A 537.76 Certain LWR facilities also have used carbon steel clad with stainless steel weld metal for liner members. Although the liner’s primary function is to provide a leak-tight barrier, it acts as part of the formwork during concrete placement and may be used in the support of internal piping/equipment. The liner is not considered to contribute to the strength of the structure. 4.13.3.2.2 Degradation mechanisms
Reinforced concrete structures almost from the time of construction can start to deteriorate in one form or the other as a result of exposure to the environment (e.g., temperature, moisture, and cyclic loadings).77 The rate of deterioration is dependent on the component’s structural design, materials selection, quality of construction, curing, and aggressiveness of environmental exposure. Termination of a component’s service life occurs when it no longer can meet its functional and performance requirements, it becomes obsolete, or the maintenance costs become excessive. Primary mechanisms (factors) that, under unfavorable conditions, can produce premature deterioration of reinforced concrete structures include those that impact either the concrete or steel reinforcing materials (i.e., mild steel reinforcement or post-tensioning systems). Degradation of the concrete can be caused by adverse performance of either its cement-paste matrix or aggregate materials under chemical or physical attack. In practice, these processes may occur concurrently to reinforce each other. In nearly all physical and chemical processes influencing the durability of concrete structures, dominant factors involved include the transport mechanisms within the pores and cracks and the presence of water. Chemical attack
may occur in several forms: efflorescence or leaching; attack by sulfate, acids, or bases; delayed ettringite formation; and alkali–aggregate reactions. Physical attack involves the degradation of concrete due to external influences and generally involves cracking due to exceeding the tensile strength of concrete, or loss of surface material. Physical attack mechanisms for concrete include salt crystallization, freezing and thawing, thermal exposure/thermal cycling, abrasion/ erosion/cavitation, irradiation, fatigue or vibration, biological attack, and settlement. Degradation of mild steel reinforcing materials can occur as a result of corrosion, irradiation, elevated temperature, or fatigue effects, with corrosion being the most likely form of attack. Post-tensioning systems are susceptible to the same degradation mechanisms as mild steel reinforcement plus loss of prestressing force primarily due to tendon relaxation and concrete creep and shrinkage. Of these, corrosion and loss of prestressing force are the most pertinent from the perspective of NPP durability. Additional information on durability of NPP reinforced concrete structures is available.34 4.13.3.2.3 Damage modeling
Extensive research and studies have been carried out to determine the durability of concrete under various service conditions, and thus information on the progressive changes in the physical and chemical nature of concrete under such conditions is available through technical committees and publications of organizations such as the ACI, International Union of Laboratories and Experts in Construction Materials, Systems and Structures (RILEM), and International Federation for Structural Concrete (CEB-FIP). Available damage models for reinforced concrete in large measure have addressed corrosion of steel reinforcement in concrete and new construction. Improved damage models and guidelines for their use are desired to predict failure probability of a degraded concrete structure, either at present or at some future point in time. Additional investigations also are desired with respect to synergistic effects involving more than one degradation factor and the interaction of loading and environment. 4.13.3.2.4 Long-term performance of concrete materials
Data on the long-term performance of the reinforced concrete materials are of importance for demonstrating the durability of the NPP concrete structures and in predicting their performance under the influence of pertinent aging factors and environmental
Concrete
stressors. This information also has application to establishing limits on hostile environmental exposure for these structures and to developing inspection and maintenance programs that will prolong component service life and improve the probability of the component surviving an extreme event such as a loss-of-coolant accident. Prior reviews of research conducted on concrete materials and structures indicate that only limited data are available on the long-term (40–80 years) properties of reinforced concrete materials.26 Where concrete properties have been reported for conditions that have been well documented, the results were generally for concretes having ages <5 years or for specimens that had been subjected to extreme, nonrepresentative environmental conditions such as accelerated corrosion or aging. Relatively few investigations have been reported providing results on examinations of structures that had been in service for the time period of interest, 20–100 years, and they did not generally provide the ‘high quality’ information (e.g., baseline material characteristics and changes in material properties with time) that is desired for meaningful assessments to indicate how the structures have changed under the influence of aging factors and environmental stressors. Limited data on the long-term performance of reinforced concrete materials reported in the literature, results from concrete cores removed from NPPs, and specimens cast in conjunction with NPP facilities have been reported.78 As shown in Figure 6, these results generally show an increase in compressive strength (relative to the 28-day reference
Relative compressive strength
PCA-0.36-moist PCA-0.36-air PCA-0.62-moist PCA-0.62-air
Wis-0.51-inside Wis-0.51-outside Wis-0.67-inside Wis-0.67-outside
2.5
423
strength) at a decreasing rate with age, but the data obtained from the literature were for concrete ages 50 years and the nuclear plant data for ages 27 years. (Using results obtained from concrete cores removed from residential buildings and bridges, one reference indicates that although the concrete strength and modulus exhibit an increase with age, the ability of concrete to resist shear and torsion may decrease.79) Long-term laboratory results in the figure for the Portland Cement Association (PCA) and University of Wisconsin (WIS) studies that attained the largest increases in strength were generally for concrete mixes having high water–cement ratios (lower reference compressive strength) and that had access to moisture for continued cement hydration. NPP concretes had higher reference compressive strengths and were essentially maintained under sealed conditions. With the availability of decommissioned NPPs and plant modifications requiring removal of materials, opportunities exist to obtain samples for use in providing an improved understanding of the effects of extended exposure under the unique conditions found in NPPs. In addition to aging, areas of interest would be the effects of longterm thermal loadings at moderate temperature levels and effects of irradiation on load-bearing concrete structures operating for >40 years. Additional applications of a concrete material sampling activity would be for assessment of construction quality, development of improved damage models, assessment and validation of nondestructive testing methods, and evaluation of the performance of repair activities.
M-CC(640) M-CC(659) M-AUX(634.5) TROJAN Mix E-2 WYLFA
ANL-FC EBRII-SB EBRII-BS GETR-HW GETR-BS ANL-HD
HEY SHAM I HEY SHAM II Hartlepool Torness Sizewell
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Figure 6 Normalized concrete compressive strength data obtained from the literature and by testing nuclear power plantrelated concrete samples.
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With respect to post-tensioning systems, current examination programs such as ASME Section XI Subsection IWL80 are adequate for determining the condition of the post-tensioning system materials and evaluating the effects of conventional degradation. Isolated incidences of wire failure due to corrosion have occurred. Leakage of tendon sheathing filler (ungrouted tendons) has occurred at a few plants but, except for the potential loss of corrosion protection, the problem appears to be primarily aesthetic.81 Tendon forces at most plants are acceptable by a significant margin, but larger than anticipated loss of force has occurred at a few older plants. The hypothetical effect of reduced prestressing force and degradation of prestressing tendons (e.g., broken wires) has been investigated for a typical PWR posttensioned concrete containment during a loss-ofcoolant accident using finite-element analysis.56 (Results for the scenario investigated indicated that loss of prestressing force leads to increased concrete cracking at lower pressurization levels, complete failure of selected hoop tendons can have a significant impact on the containment ultimate capacity, and failure of selected vertical tendons does not have a significant impact on ultimate capacity.) With the potential use of grouted tendon systems in some of the new reactor designs proposed for construction in the United States, improved guidance on in-service inspection of grouted tendons is desired. Other potential research topics related to post-tensioning systems include development of an improved relationship between the end-anchorage force measured by the lift-off test and change in mean force along the tendon length for unbonded tendons, as well as an assessment of the validity of using estimates of time-dependent loss of prestressing force based on limited-duration relaxation tests (e.g., 1000 h) and concrete creep results (e.g., 6 months): at a plant 60-year old, this involves application of time factors of 500 and 120, respectively. 4.13.3.3
Assessment and Repair
Operating experience has demonstrated that periodic inspection, maintenance, and repair are the essential elements of an overall program to maintain an acceptable level of reliability for structures over their service life. Assessment and management of aging in NPP concrete structures requires a more systematic approach than simple reliance on existing code margins of safety.82 What is required is the integration of structural component function, potential degradation mechanisms, and appropriate control
programs into a quantitative evaluation procedure. A methodology for demonstrating the continued reliable and safe performance of these structures should include (1) identification of structures important to public health and safety; (2) identification of environmental stressors, aging mechanisms and their significance, and likely sites for occurrence; (3) a monitoring- or in-service-inspection-based methodology that includes criteria for resolution of existing conditions; and (4) a remedial measures program. 4.13.3.3.1 Component selection
The most effective structural condition assessment programs are those that focus on the components most important to safety and at risk due to environmental stressor effects. Aging assessment methodologies have been developed to provide a logical basis for identifying the critical concrete structural elements and degradation factors that can potentially impact the performance of these structures.83 An evaluation of the impact on plant risk due to structural aging can also be used in the selection of structural components for evaluation.84 Probabilistic risk assessments conducted to date indicate that the structural systems generally play a passive role in mitigating design basis (or larger) internal initiating events: a notable exception being the pressure-retaining function of the containment following a degraded core incident involving failure of the reactor pressure vessel. The structural components play an essential role in mitigating extreme events initiated by earthquake, wind, and other extreme influences, and their failure probabilities due to external events can be higher. Moreover, failure of major structural components may impact the operation of a number of mechanical and electrical systems and lead to so-called common cause failures. Thus, deterioration of structural components and systems due to aging and other aggressive environmental influences may be more serious in terms of overall plant risk than might be evident from a cursory examination of their role in accident mitigation. The significance of structural aging and deterioration to plant risk can be evaluated by considering the impact that they have on risk associated with external initiating events, especially earthquakes. It is in mitigating the effects of strong ground motion due to earthquakes that structural systems play a particularly significant role. The apparent impact of structural aging can be investigated using a margins analysis to assess suitability for continued service. Sensitivity analysis can help to identify the structures of importance that should warrant particular attention.
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A third approach involves the combination of finite-element analysis and nondestructive testing methods for evaluation of aging and degradation of concrete containments.85 The CONMOD Project objective was to find a practical means to determine the condition of a containment structure as well as how this condition can be expected to change with time under the influence of various loading conditions, including aging. Applications of the approach developed include (1) identification of the critical parts of a structure for nondestructive evaluation including critical parameters, (2) updated structural analyses using input from nondestructive evaluations, (3) prediction of nondestructive responses for a known condition at a given time using finite-element method modeling techniques, and (4) prediction of the nondestructive evaluation responses using finiteelement modeling techniques based on a known condition and how this will change because of aging processes. One of the conclusions of this study was that development of new containment designs should focus on establishing rules, designs, and novel ideas on how to significantly improve the accessibility of the concrete structures for diagnostic investigations.
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information on instances that can result from or lead to structural distress (e.g., crack mapping); (3) determination of the need for additional surveys or application of destructive or nondestructive testing methods; (4) analysis of results; and (5) preparation of a report presenting conclusions and recommendations. More detailed information on guidelines on conduct of surveys of existing civil engineering buildings is available.88–91 Some general guidance on assessment of NPP degradation is also available.92–95 However, NPP reinforced concrete structures present special challenges for development of acceptance criteria because of their massive size, limited accessibility in certain areas, stochastic nature of past and future loads, randomness in strength, uncertainty in material changes due to aging and possibly degradation, and somewhat qualitative nature of some nondestructive evaluation techniques. Improved guidelines and criteria to aid in the interpretation of condition assessment results, including development of probability-based degradation acceptance limits, are required. (Some information on probability-based crack acceptance limits for beams and shear walls considering loss of steel area and concrete spalling is available.96)
4.13.3.3.2 In-service inspections
In-service inspection programs have the primary goal of ensuring that the NPP structures have sufficient structural margins to continue to perform in a reliable and safe manner.86,87 A secondary goal is to identify environmental stressors or aging factor effects before they reach sufficient intensity to potentially degrade structural components. Routine observation, general visual inspections, leakage-rate tests, and destructive and nondestructive examinations are techniques available to identify areas of NPPs that have experienced degradation. Determination of the existing performance characteristics and extent and causes of any observed distress is accomplished through a structural condition assessment that routinely initiates with a general visual inspection to identify suspect areas followed by application of destructive or nondestructive examinations to quantify the extent and significance of any observed degradation. Basic components of a condition assessment include (1) a review of ‘as-built’ drawings and other information pertaining to the original design and construction so that information, such as accessibility and position and orientation of embedded steel reinforcing and plates in concrete, is known prior to the site visit; (2) detailed visual examination of structure to document easily obtained
4.13.3.3.3 Nondestructive examinations
Application of nondestructive examination methods to NPP reinforced concrete structures presents challenges: wall thicknesses can be in excess of 1 m; structures often have increased steel reinforcement density with complex detailing; there can be a number of penetrations or cast-in-place items; accessibility may be limited because of the presence of liners or other components, harsh environments, or structures located below ground; experience with nondestructive examinations of NPP concrete structures is somewhat limited; and methods utilized for the NPP structures are often based on equipment developed for other materials or technologies. Available methods are relatively good at identifying cracking, voids, and delaminations as well as indicating the relative quality of concrete. Methods for determining concrete properties, however, generally are somewhat more qualitative than quantitative because they tend to be indirect in that they often require the development of correlation curves for relating a measured parameter (e.g., ultrasonic velocity or rebound number) to a property (e.g., concrete compressive strength). Information on identification and description of methods for determining the strength of concrete and evaluation of concrete structures is available.97–100 A practical
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guide related to nondestructive examination of concrete, which not only identifies and describes the capabilities, limitations, and applications of the various methods that are available but also presents results from a number of examples, has been developed.101 The status of nondestructive examination methods and priorities for its development with respect to examination and instrumentation and monitoring of concrete structures in nuclear plants was addressed by prior NEA/CSNI IAGE workshops.41,45,51 It was noted that although nondestructive examination techniques have been successfully used on a variety of reinforced and post-tensioned concrete structures, there has been somewhat limited experience in their use to evaluate typical NPP safety-related structures. With respect to these structures, three conditions exist where performing inspections or conduct of nondestructive examinations is not straightforward and requires development – inspection of thickwalled, heavily reinforced concrete sections, basemats or foundations, and inaccessible portions of a containment metallic pressure boundary. Information summarizing the activities conducted addressing these conditions has been documented.102 Noninvasive techniques for characterization, inspection, and monitoring of thick-walled, heavily reinforced concrete sections to provide additional assurances of their continued structural integrity are desirable (e.g., as-built or current structural features determination, flaw detection and characterization, identification of honeycomb areas and embedded items, and location of voids adjacent to the liner). Methods that can be used to inspect the basemat without the requirement for removal of material and techniques that can detect and assess corrosion are of particular interest. Acoustic (e.g., ultrasonic pulse velocity, spectral analysis of surface waves, impact echo, and acoustic tomography), radar, and radiography appear to have potential for application to thick-walled, heavily reinforced concrete structures in NPPs; however, additional development is required. The most commonly used type of foundation for both concrete and steel NPP containments is a mat foundation, which is a flat, thick slab supporting the containment, its interior structures, and any shield building surrounding the containment.103 As such, the concrete foundation elements of NPPs are typically either partially or totally inaccessible for inspection unless adjacent soil, coatings, waterproof materials, or portions of neighboring components or structures are removed. As a result, indirect methods related to environmental qualification are often
utilized to indicate the potential for degradation of the NPP concrete foundations.20 This is generally done through an evaluation of the surrounding medium (e.g., air, soil, humidity, groundwater, or cooling water). Methods employed are based primarily on chemical evaluations to assess the presence and concentration of potentially aggressive ions (e.g., sulfates or chlorides). In addition to an assessment of the aggressiveness of the surrounding environment, the CFR requires a complete description of the effects of groundwater levels and other hydrodynamic effects on the design bases of the plant foundations and other structures, systems, and components important to safety.104 Inspection of inaccessible portions of metallic pressure boundary components of NPP containments (e.g., fully embedded or inaccessible containment shell or liner portions, the sand pocket region in Mark I and II drywells, and portions of the shell obscured by obstacles such as platforms or floors) requires special attention. Embedded metallic portions of the containment pressure boundary may be subjected to corrosion resulting from groundwater permeation through the concrete; a breakdown of the sealant at the concrete–containment shell interface that permits entry of corrosive fluids from spills, leakage, or condensation; or in areas adjacent to floors where the gap contains a filler material that can retain fluids. NPP inspections have identified corrosion of the steel containment shell in the drywell sand cushion region, shell corrosion in ice condenser plants, corrosion of the torus of the steel containment shell, and concrete containment liner corrosion. Corrosion incidences such as these may challenge the containment integrity and, if throughwall, can provide a leak path to the outside environment. Several techniques have been investigated that exhibit potential for performing inspections of inaccessible portions of NPP metallic pressure boundaries (i.e., ultrasonics, electromagnetic acoustic transducers, half-cell potential measurements, magnetostrictive sensors, and multimode guide waves).102 However, these techniques tend to be time consuming and costly because they tend to examine only a small area at a time. A technique that can be applied remotely to perform global inspections and determines the overall condition of the containment metallic pressure boundary in a cost- and performance-effective manner is desired. 4.13.3.3.4 Remedial methods
Deterioration of reinforced concrete generally will result in cracking, spalling, or delamination of the
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cover concrete. Whenever damage is detected, corrective actions are taken to identify and eliminate the source of the problem, thereby halting the degradation process. The first step in any repair activity is a thorough assessment of the damaged structure or component including evaluation of (1) cause of deterioration, (2) extent of deterioration, and (3) effect of deterioration on the functional and performance requirements of the structure or component. From this information, a remedial measures strategy is formulated based on the consequence of damage (e.g., effect of degradation on structural safety), time requirements for implementation (e.g., shutdown requirements, immediate or future safety concern), economic aspects (e.g., partial or complete repair), and residual service life requirements (e.g., desired residual service life will influence action taken).105 Basic remedial measures options include (1) no active intervention; (2) more frequent inspections or conducting specific studies; (3) if safety margins are presently acceptable, taking action to prevent deterioration from getting worse; (4) carrying out repairs to restore deteriorated or damaged part of the structure to a satisfactory condition; and (5) demolishing and rebuilding all or part of the structure. Basic guidance on the repair of degraded structures is available,27,106 and a workshop has been held addressing repair of NPP concrete structures.46 Results of the workshop indicate that improved guidance is required on the assessment of defects (e.g., cracks), and information is desired on the performance and effectiveness of subsequent repairs to concrete structures in NPPs (e.g., durability of repair materials). Information on past performance and current practices for repair materials and systems for general civil engineering structures is being assembled (http://projects.bre.co. uk/conrepnet/pages/contents.htm).
4.13.4 Structural Reliability Theory If properly designed and constructed, the concrete structures in NPPs generally have substantial safety margins; however, additional information for quantifying the available margins of degraded structures is desired. In addition, how age-related degradation may affect dynamic properties (e.g., stiffness, frequency, and dampening), structural response, structural resistance/ capacity, failure mode, and location of failure initiation is not well understood. A better knowledge of the effects of aging degradation on structures and passive components is necessary to help ensure that the
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current licensing basis is maintained under all loading conditions.96 Decisions as to whether to invest in maintenance and rehabilitation of structures, systems, and components as a condition for continued service and risk mitigation, and the appropriate level of investment, should consider the nature and level of uncertainties in their current condition and in future demands.107,108 Recent advances in structural reliability analysis, uncertainty quantification, and probabilistic risk assessment make it possible to perform such evaluations and to devise uniform risk-based criteria by which existing facilities can be evaluated to achieve a desired performance level when subjected to uncertain demands.109,110 Consideration of in situ conditions, redundancy, and uncertainties in important engineering parameters often can lead to significant economic benefits when assessing the condition of an existing structure in a (possibly) degraded condition, and the maintenance or rehabilitation strategies that might be required as a condition for future service. Reliability-based approaches have been applied to the NPP concrete structures111,112 and in evaluation of the prestress level in concrete containments with unbonded tendons.113 Degradation effects can be quantified with fragility curves developed for both undegraded and degraded components.114 Fragility analysis is a technique for assessing, in probabilistic terms in the presence of uncertainties, the capability of an engineered system to withstand a specified event. Fragility modeling requires a focus on the behavior of the system as a whole and specifically, on things that can go wrong with the system. The fragility modeling process leads to a median-centered (or likely) estimate of system performance, coupled with an estimate of the variability or uncertainty in performance. The fragility concept has found widespread usage in the nuclear industry, where it has been used in seismic probabilistic safety and/or margin assessments of safetyrelated plant systems.115 The fragility modeling procedures applied to degraded concrete members can be used to assess the effects of degradation on plant risk and can lead to the development of probability-based degradation acceptance limits. This approach has been applied to a limited extent to degraded flexural members and shear walls.96 Additional work is desired in this area for the purpose of refining and applying the time-dependent reliability methodology for optimizing in-service inspection/maintenance strategies and for developing and evaluating improved quantitative models for
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predicting future performance (or failure probability) of a degraded concrete structure, either at present or at some future point in time.
4.13.5 Summary and Potential Research Topics As concrete ages, changes in its properties will occur as a result of continuing microstructural changes (i.e., slow hydration, crystallization of amorphous constituents, and reactions between cement paste and aggregates), as well as environmental influences. These changes do not have to be detrimental to the point that concrete will not be able to meet its performance requirements. Concrete, however, can suffer undesirable changes with time because of improper specifications, a violation of specifications, or adverse performance of its cement paste matrix or aggregate constituents under either physical or chemical attack. Concrete durability and the relationship between durability and performance, a review of the historical perspective related to concrete and longevity, a description of the basic materials that comprise reinforced concrete, and information on the environmental factors that can affect the performance of NPP concrete structures have been provided. Primary aspects related to management of aging of NPP concrete structures have been noted: degradation mechanisms, damage models, and material performance; assessment and repair (i.e., component selection, in-service inspection, nondestructive examinations, and remedial actions); and estimation of performance at present or some future point in time (i.e., application of structural reliability theory to the design and optimization of inservice inspection/maintenance strategies, and determination of the effects of degradation on plant risk). Several areas have been identified where additional research would be of benefit to aging management of NPP concrete structures: (1) compilation of material property data for long-term performance and trending, evaluation of environmental effects, and assessment and validation of nondestructive evaluation methods; (2) evaluation of long-term effects of elevated temperature and radiation on concrete behavior; (3) improved damage models and acceptance criteria for use in assessments of the current condition as well as estimation of the future condition of the structures; (4) improved constitutive models and analytical methods for use in determination of nonlinear structural response (e.g., accident
conditions); (5) nonintrusive methods for inspection of thick-walled, heavily reinforced concrete structures and basemats; (6) global inspection methods for metallic pressure boundary components (i.e., steel containments and liners of concrete containments) including inaccessible areas and backside of liner; (7) data on application and performance (e.g., durability) of repair materials and techniques; (8) utilization of structural reliability theory incorporating uncertainties to address time-dependent changes to structures to ensure that minimum accepted performance requirements are exceeded and to estimate on-going component degradation to estimate end of life; and (9) application of probabilistic modeling of component performance to provide risk-based criteria to evaluate how aging affects structural capacity.
References 1. ACI. J. Am. Concr. Inst. 1977, 74(12), 573–609. 2. BSI. Concrete – Guide to Specifying Concrete, BS 5328; British Standards Institution: London, 1997. 3. BSI. Structural Use of Concrete – Code of Practice for Design and Construction, BS 8110; British Standards Institution: London, 1985. 4. ACI. Building Code Requirements for Structural Concrete and Commentary; ACI Standard 318–05; American Concrete Institute: Farmington Hills, MI, 2005. 5. BSI. Concrete – Complementary British Standard to BS EN 206–1, BS 8500; British Standards Institution: London, 2002. 6. Aı¨tcin, P.-C. Cement Concr. Res. 2000, 30, 1349–1359. 7. USNRC. Issuance of Final Amendment to 10 CFR }50.55a to Incorporate by Reference the ASME Boiler and Pressure Vessel Code (ASME Code), Section XI, Subsection IWE and Subsection IWL; SECY-96–080; US Nuclear Regulatory Commission: Washington, DC, 1996. 8. University of Illinois. Department of Materials Science and Engineering, Urbana, IL, http://matse1.mse.uiuc. edu/concrete/hist.htm. 9. Mallinson, L. G.; Davies, I. Ll. A Historical Examination of Concrete, EUR 10937 EN; Commission of European Communities: Luxembourg, 1987. 10. OFR. Title 10 – Energy, Code of Federal Regulations; Office of the Federal Register: Washington, DC, 2006. 11. USNRC. (2007). Seismic Design Classification, Regulatory Guide 1.29 (Rev. 4); US Nuclear Regulatory Commission: Washington, DC, 2006. 12. USNRC. Regulatory Standard Review Plan, NUREG0800 Section 3.8.3; US Nuclear Regulatory Commission: Washington, DC, 1981. 13. ACI. Code Requirements for Nuclear Safety Related Concrete Structures and Commentary, ACI Standard 349; American Concrete Institute: Farmington Hills, MI, 1976. 14. USNRC. Safety-Related Concrete Structures for Nuclear Power Plants (Other Than Reactor Vessels and Containments) (for Comment Issue), Regulatory Guide 1.142; US Nuclear Regulatory Commission: Washington, DC, 1978.
Concrete 15.
16. 17. 18.
19.
20.
21. 22. 23.
24.
25.
26.
27.
28.
29. 30. 31.
32.
33.
USNRC. Standard Format and Content of Safety Analysis Reports for Nuclear Power Plants (LWR Edition), Regulatory Guide 1.70, Rev. 3; US Nuclear Regulatory Commission: Washington, DC, 1978. ACI. Code for Concrete Reactor Vessels and Containments, ACI Standard 359; American Concrete Institute: Farmington Hills, MI, 1977. USNRC. Regulatory Standard Review Plan, NUREG0800 Section 3.8.1; US Nuclear Regulatory Commission: Washington, DC, 1981. Lo, T.; Nelson, T. A.; Chen, P. Y.; Persinko, P.; Grimes, C. In Proceedings of the American Society of Civil Engineers Conference on Structural Engineering in Nuclear Facilities, Raleigh, NC, Sept 10–12, 1984; Ucciferro, J., Ed.; ASCE: Reston, VA, pp. 847–861. FIB. Nuclear Containments – State-of-the-Art Report, Bulletin 13; International Federation for Structural Concrete (fib), Federal Institute of Technology: Lausanne, Switzerland, 2001. Hookham, C. J. In-Service Inspection Guidelines for Concrete Structures in Nuclear Power Plants, ORNL/ NRC/LTR-95/14; Lockheed Martin Energy Systems, Oak Ridge National Laboratory: Oak Ridge, TN, 1995. USNRC. Federal Register 1995, 60(186), 49495–49505. USNRC. Masonry Wall Design, Inspection Enforcement Bulletin 80–11; US Nuclear Regulatory Commission: Washington, DC, 1980. USNRC. Lessons Learned from Regional Inspections of Licensee Actions in Response to IE Bulletin 80–11, Inspection Notification 87–67; US Nuclear Regulatory Commission: Washington, DC, 1987. USNRC. Inspection of Water-Control Structures Associated with Nuclear Power Plants, Regulatory Guide 1.127, Rev. 1; US Nuclear Regulatory Commission: Washington, DC, 1978. Gregor, F. E.; Hookham, C. J. In Transactions of the 12th International Conference on Structural Mechanics in Reactor Technology, Stuttgart, Germany, Aug 15–20; Elsevier Science: Amsterdam, 1993; Paper DH06/2. Naus, D. J. Concrete Component Aging and Its Significance Relative to Life Extension of Nuclear Power Plants, NUREG/CR-4652; US Nuclear Regulatory Commission: Washington, DC, 1986. Krauss, P. D. Repair Materials and Techniques for Concrete Structures in Nuclear Power Plants, ORNL/NRC/LTR-93/28; Martin Marietta Energy Systems, Inc., Oak Ridge National Laboratory: Oak Ridge, TN, 1994. Ashar, H.; Bagchi, G. Assessment of Inservice Conditions of Safety-Related Nuclear Plant Structures, NUREG 1522 US Nuclear Regulatory Commission: Washington, DC, 1995. Ashar, H.; Naus, D. J.; Tan, C. P. Concr. Int. 1994, 16(5), 30–34. Ashar, H.; Tan, C. P.; Naus, D. J. Concr. Int. 1994, 16(6), 58–61. Ashar, H.; Jeng, D. Effectiveness of in-service inspection requirements of prestressed concrete containments – U.S. experience. In Proceedings of Second International Conference on Containment Design and Operation, Toronto, ON, Oct 14–17, 1990. Braverman, J. I.; Hofmayer, C. H.; Morante, R. J.; Shteyngart, S.; Bezler, P. Assessment of Age-Related Degradation of Structures and Passive Components for US Nuclear Power Plants, NUREG/CR-6679 US Nuclear Regulatory Commission: Washington, DC, 2000. USNRC. Degradation of Prestressing Tendon Systems in Prestressed Concrete Containments, NRC Information
34.
35.
36.
37.
38.
39.
40.
41.
42. 43. 44.
45. 46.
47.
48.
49.
429
Notice 99–10, Rev. 1, US Nuclear Regulatory Commission: Washington, DC, 1999. Naus, D. J. Primer on Durability of Nuclear Power Plant Reinforced Concrete Structures – A Review of Pertinent Factors, NUREG/CR-6927 US Nuclear Regulatory Commission: Washington, DC, 2007. IAEA. Assessment and Management of Major Nuclear Power Plant Components Important to Safety: Concrete Containment Buildings, IAEA-TECDOC-1025; International Atomic Energy Agency: Vienna, 1998. EXPERTCENTRUM. In Proceedings of the International Conference Life Prediction and Aging Management of Concrete Structures; Ja´vor, T., Ed.; EXPERTCENTRUM: Bratislava, Czech Republic, 1999; Proceedings 8. RILEM. In Considerations for Use in Managing the Aging of Nuclear Power Plant Concrete Structures; Naus, D. J., Ed.; RILEM Publications S.A.R.L.: Cachan Cedex, France, 1999; Report 19. RILEM. In Proceedings of International Workshop on Life Prediction and Aging Management of Concrete Structures, PRO 16, Cannes, France, Oct 16–17, 2000; Naus, D. J., Ed.; RILEM Publications S.A.R.L.: Cachan Cedex, France. RILEM. In Proceedings of 2nd RILEM International Workshop on Life Prediction and Aging Management of Nuclear Power Plant Concrete Structures, Paris, France, May 5–6; Naus, D. J., Ed.; RILEM Publications S.A.R.L.: Cachan Cedex, France, 2003; Report 29. NEA. Report of Task Group Reviewing National and International Activities in the Area of Ageing of Nuclear Power Plant Concrete Structures, NEA/CSNI/R(95)19: OECD Nuclear Energy Agency: Iss-les-Moulineaux, France, 1996. NEA. Development Priorities for Non-destructive Examination of Concrete Structures in Nuclear Plant, NEA/CSNI/R(98)6; OECD Nuclear Energy Agency: Issles-Moulineaux, France, 1998. NEA. Finite Element Analysis of Degraded Concrete Structures, NEA/CSNI/R(99)1: OECD Nuclear Energy Agency: Issy-les-Moulineaux, France, 1999. NEA. NPP Containment Prestress Loss Summary Report, NEA/CSNI/R(99)11: OECD Nuclear Energy Agency: Issy-les-Moulineaux, France, 1999. NEA. Joint WANO/OECD-NEA Workshop: Prestress Loss in NPP Containments, OCDE/GD(97)9; OECD Nuclear Energy Agency: Issy-les-Moulineaux, France, 1999. NEA. Workshop on the Instrumentation and Monitoring of Concrete Structures, NEA/CSNI/R(2000)15; OECD Nuclear Energy Agency: Issy-les-Moulineaux, France, 2001. NEA. OECD-NEA Workshop on the Evaluation of Defects, Repair Criteria and Methods of Repair for Concrete Structures of Nuclear Power Plants, NEA/CSNI/R(2002)7; OECD Nuclear Energy Agency: Issy-les-Moulineaux, France, 2002; Vols 1 and 2. NEA. Finite Element Analysis of Ageing Reinforced and Prestressed Concrete Structures in Nuclear Plants, NEA/CSNI/R(2002)13; OECD Nuclear Energy Agency: Issy-les-Moulineaux, France, 2002. NEA. Report of Task Group Reviewing Activities in the Area of Ageing of Concrete Structures Used to Construct Nuclear Power Plant Fuel-Cycle Facilities, NEA/CSNI/R(2002)14; OECD Nuclear Energy Agency: Issy-les-Moulineaux, France, 2002. NEA. Electrochemical Techniques to Detect Corrosion in Concrete Structures in Nuclear Installations, NEA/CSNI/R(2002)21; OECD Nuclear Energy Agency: Issy-les-Moulineaux, France, 2002.
430 50.
51.
52.
53. 54.
55.
56.
57.
58. 59. 60. 61. 62. 63. 64. 65.
66.
67.
68.
69.
Concrete NEA. In Proceedings of the CSNI/RILEM Workshop on Use and Performance of Concrete in NPP Fuel Facilities, NEA/CSNI/R(2004)8 OECD Nuclear Energy Agency: Issy-les-Moulineaux, France, 2004. NEA. CSNI Workshop on Ageing Management of Thick Walled Concrete Structures, ISI, Maintenance and Repair, Instrumentation Methods and Safety Assessment in Terms of LTO, NEA/SEN/SIN/IAGE(2008)7; OECD Nuclear Energy Agency: Issy-les-Moulineaux, France, 2008. NUCPERF. In Proceedings of Workshop on Corrosion and Long Term Performance of Reinforced Concrete in Nuclear Power Plants and Waste Facilities, Cadarache, France, Mar 27–30; Foct, F., Fe´ron, D., L’Hostis, V., Eds.; EDP Sciences: France, 2006. EPRI. Class I Structures License Renewal Industry Report: Revision 1, EPRI TR-103842 (NUMARC 90–06); Electric Power Research Institute: Palo Alto, CA, 1994. EPRI. PWR Containment Structures License Renewal Industry Report: Revision 1, EPRI TR-103835 (NUMARC 90–01); Electric Power Research Institute: Palo Alto, CA, 1994. Naus, D. J.; Oland, C. B.; Ellingwood, B. R. Report on Aging of Nuclear Power Plant Reinforced Concrete Structures, NUREG/CR-6424; US Nuclear Regulatory Commission: Washington, DC, 1996. Smith, J. A. Capacity of Prestressed Concrete Containment Vessels with Prestressing Loss, SAND2001–1762; Sandia National Laboratories: Albuquerque, NM, 2001. USACE. US Army Corps of Engineers, Evaluation and Repair of Concrete Structures – Engineering Manual, EM 1110–2–2–2002; Waterways Experiment Station: Vicksburg, MS, 1995. ACI. Guide for Use of Normal Weight and Heavyweight Aggregates in Concrete, ACI 221R-96; American Concrete Institute: Farmington Hills, MI, 2001. ACI. Cementitious Materials for Concrete, ACI Education Bulletin E3–01; American Concrete Institute: Farmington Hills, MI, 2001. ACI. Chemical Admixtures for Concrete, ACI 212.3R-04; American Concrete Institute: Farmington Hills, MI, 2004. PCA. Concr. Technol. Today 1996, 17(2), 1–4. BSI. The Structural Use of Reinforced Concrete in Buildings, CP 114, Amendment 1; British Standards Institution: London, 1965. Troxell, G.; Davis, H.; Kelly, J. Composition and Properties of Concrete; McGraw-Hill: New York, 1968. ASTM. Standard Specification for Portland Cement, ASTM C 150-05; ASTM International: West Conshohocken, PA, 2005. ASTM. Standard Specification for Deformed and Plain Carbon Steel Bars for Concrete Reinforcement, ASTM A 615/A 615M-06; ASTM International: West Conshohocken, PA, 2006. ASTM. Standard Specification for Low-Alloy Steel Deformed and Plain Bars for Concrete Reinforcement, ASTM A 706/A 706M-06; ASTM International: West Conshohocken, PA, 2006. ASTM. Standard Specification for Uncoated StressRelieved Steel Wire for Prestressed Concrete, ASTM A 421/A 421M-05; ASTM International: West Conshohocken, PA, 2005. ASTM. Standard Specification for Steel Strand, Uncoated Seven-Wire for Prestressed Concrete, ASTM A 416/A416M-05; ASTM International: West Conshohocken, PA, 2005. ASTM. Standard Specification for Uncoated HighStrength Steel Bars for Prestressing Concrete, ASTM
70. 71. 72.
73. 74.
75.
76.
77. 78.
79.
80.
81.
82.
83.
84. 85.
86.
A 722/A 722M-05; ASTM International: West Conshohocken, PA, 2005. ASTM. Standard Specification for Steel Bar, Carbon and Alloy, Cold-Finished, ASTM A 108-03; ASTM International: West Conshohocken, PA, 2003. ASTM. Standard Specification for Carbon Structural Steel, ASTM A 36/A 36M-05; ASTM International: West Conshohocken, PA, 2006. ASTM. Standard Specification for Pressure Vessel Plates, Carbon Steel, for Moderate- and Lower-Temperature Service, ASTM A 516/A 516M-06; ASTM International: West Conshohocken, PA, 2006. ASTM. Standard Specification for Stainless Steel Bars and Shapes, ASTM A 276-06; ASTM International: West Conshohocken, PA, 2006. ASTM. Standard Specification for Carbon and Alloy Steel Bars Subject to End-Quench Hardenability Requirements, ASTM A 304-05; ASTM International: West Conshohocken, PA, 2005. ASTM. Standard Specification for Pressure Vessel Plates, Carbon Steel, Low- and Intermediate-Tensile Strength, ASTM A 285/A 285M-03; ASTM International: West Conshohocken, PA, 2003. ASTM. Standard Specification for Pressure Vessel Plates, Heat-Treated, Carbon–Managanese–Silicon Steel, ASTM A 537/A 537M-06; ASTM International: West Conshohocken, PA, 2006. Browne, R. D. Durability of Reinforced Concrete Structures, New Zeal Concr Construct, Parts 1 and 2, Sept and Oct 1989. Oland, C. B.; Naus, D. J. Summary of Materials Contained in the Structural Materials Information Center, ORNL/ NRC/LTR-94/22; Martin Marietta Energy Systems, Oak Ridge National Laboratory: Oak Ridge, TN, 1994. Levtchitch, V.; Kvasha, V.; Boussalis, H.; Chassiakos, A.; Kosmatopoulos, E. Seismic performance capacities of old concrete. In Proceedings of 13th World Conference on Earthquake Engineering, Vancouver, Canada, Aug 1–6, 2004; Paper no. 2182. ASME. Code for Concrete Reactor Vessels and Containments, Section XI, Subsection IWL, ASME Boiler and Pressure Vessel Code; American Society of Mechanical Engineers: New York, 2005. Naus, D. J.; Oland, C. B. An Investigation of Tendon Sheathing Filler Migration in Concrete, NUREG/CR-6598; US Nuclear Regulatory Commission: Washington, DC, 1998. Christensen, J. A. NPAR approach to controlling aging in nuclear power plants. In Proceedings of the 17th Water Reactor Safety Information Meeting, NUREG/CP-0105, Washington, DC 1990; Vol. 3, pp 509–529. Hookham, C. J. Structural Aging Assessment Methodology for Concrete Structures in Nuclear Power Plants, ORNL/NRC/LTR-90/17; Martin Marietta Energy Systems, Oak Ridge National Laboratory: Oak Ridge, TN, 1991. Ellingwood, B. R. Aging Effects on Probabilistic Risk Assessment, NUREG/CR-6425; US Nuclear Regulatory Commission: Washington, DC, 1996. Jovall, O.; Larsson, J.-A.; Shaw, P.; Touret, J.-P.; Karlberg, G. Concrete containment modeling and management, conmod. In Transactions of the 17th International Conference on Structural Mechanics in Reactor Technology, Prague, Czech Republic, Aug 17–22, 2003, Paper H04-4. OFR. Appendix J – Primary Reactor Containment Leakage Testing for Water-Cooled Reactors; Code of Federal Regulations, 10 CFR Part 50; Office of Federal Register: Washington, DC, 1995; pp 748–753.
Concrete 87. 88. 89. 90.
91. 92. 93.
94.
95. 96.
97. 98. 99. 100.
ASME. ASME Boiler and Pressure Vessel Code; American Society of Mechanical Engineers: New York, 2007. ASCE. Guidelines for Structural Condition Assessment of Existing Buildings, SEI/ASCE 11-99; American Society of Civil Engineers: Reston, VA, 2000. FIB. Inspection and Maintenance of Reinforced and Prestressed Concrete; Fe´de´ration Internationale du Be´ton: Lausanne, Switzerland, 1986. FIB. Monitoring and Safety Evaluation of Existing Concrete Structures – State-of-the-Art Report, Bulletin 22; International Federation for Structural Concrete (fib), Federal Institute of Technology: Lausanne, Switzerland, 2003. ACI. Guide for Conducting a Visual Inspection of Concrete in Service, ACI 201.1R; American Concrete Institute: Farmington Hills, MI, 2008. ACI. Evaluation of Existing Nuclear Safety Related Concrete Structures, ACI 349R.3-02; American Concrete Institute: Farmington Hills, MI, 2002. ASME. Section XI, Rules for Inservice Inspection of Nuclear Power Plant Components; ASME Boiler and Pressure Vessel Code; American Society of Mechanical Engineers: New York, 2007. Concrete Society. Diagnosis of Deterioration in Concrete Structures – Identification of Defects, Evaluation, and Development of Remedial Action, Technical Report No. 54; Concrete Society: Berkshire. RILEM. Mater. Struct. 1994, 27, 362–369. Braverman, J. I.; Miller, C. A.; Ellingwood, B. R.; et al. Probability-Based Evaluation of Degraded Reinforced Concrete Components in Nuclear Power Plants, NUREG/ CR-6715; US Nuclear Regulatory Commission: Washington, DC, 2001. ACI. Nondestructive Test Methods for Evaluation of Concrete in Structures, ACI 228.2R-98; American Concrete Institute: Farmington Hills, MI, 1998. ACI. In-Place Methods for Estimating Concrete Strength, ACI 228.1R-03; American Concrete Institute: Farmington Hills, MI, 2003. Malhotra, V. M.; Carino, N. J. Handbook of Nondestructive Testing of Concrete; CRC Press: Boca Raton, FL, 1991. Bungey, J. H. Testing Concrete in Structures – A Guide to Equipment for Testing Concrete in Structures, Technical Note 143; Construction Industry Research and Information Association: London, 1992.
101.
102.
103. 104. 105.
106. 107. 108. 109. 110. 111.
112.
113. 114. 115.
431
Force Technology. A PRACTICAL Guide to Nondestructive Examination of Concrete; Nordic Innovation Center Report; Force Technology: Helsingborg, Sweden, 2004. Naus, D. J. Inspection of Nuclear Power Plant Structures – Overview of Methods and Related Applications, ORNL/TM-2007/191; Oak Ridge National Laboratory: Oak Ridge, TN, 2009. USNRC. ‘Foundations,’ Standard Review Plan, NUREG-0800, Revision 2, Section 3.8.5; US Nuclear Regulatory Commission: Washington, DC, 2007. USNRC. ‘Groundwater,’ Standard Review Plan, NUREG-0800, Revision 2, Section 2.4.12; US Nuclear Regulatory Commission: Washington, DC, 2007. Price, W. F.; Bamforth, P. B.; Glass, G. K. Review of European Repair for Corrosion Damaged Reinforced Concrete, Report No. 1303/91/5823; Taywood Engineering: R&D Division: London, 1993. Emmons, P. H. Concrete Repair and Maintenance Illustrated; R.S. Means: Kingston, MA, 1993. Elms, D. G. Risk assessment. In Engineering Safety; Blockley, D., Ed.; McGraw-Hill: Berkshire, 1992; 28–46. Ellingwood, B. R. Progr. Struct. Eng. Mater. 2001, 3(2), 170–179. Ellingwood, B. R.; Wen, Y. K. Progr. Struct. Eng. Mater. 2005, 7(2), 56–70. Wen, Y. K.; Ellingwood, B. R. Earthquake Spectra, EERI 2005, 21(3), 861–877. Mori, Y.; Ellingwood, B. Methodology for ReliabilityBased Condition Assessment – Application to Concrete Structures in Nuclear Plants, NUREG/CR-6052; US Nuclear Regulatory Commission: Washington, DC, 1993. Ellingwood, B. R.; Naus, D. J. In Modeling Complex Engineering Structures; Melchers, R. E., Hough, R., Eds.; American Society of Civil Engineers: Reston, VA, 2007; Chapter 6, 137–170. Anderson, P.; Hansson, M.; Thelandersson, S. Struct. Saf. 2008, 30(1), 78–89. Braverman, J. I.; Miller, C. A.; Hofmayer, C. H.; Ellingwood, B. R.; Naus, D. J.; Chang, T. Y. Int. J. Nucl. Eng. Des 2004, 228(1–3), 283–304. Kennedy, R. P.; Ravindra, M. Int. J. Nucl. Eng. Des. 1984, 79(1), 47–68.
4.14
Fracture Toughness Master Curve of bcc Steels
T. Planman VTT, Espoo, Finland
W. L. Server ATI Consulting, Pinehurst, NC, USA
ß 2012 Elsevier Ltd. All rights reserved.
4.14.1
Introduction
434
4.14.1.1 4.14.1.2
Historical Review of Fracture Toughness Determination for Ferritic Steels Theoretical Background Leading to Use of EPFM and Data Distributions for Ferritic Steels Master Curve Methodology as Developed by Wallin History Basics Methodology Analysis of Fracture Toughness Test Data for Master Curve Application Standard Test and Analysis Procedure (ASTM E 1921) Test specimens used Determination of reference temperature T0 Data qualification Determination of lower bound curves Limits of applicability Aspects of Applying Small and Miniature Specimens Effect of Constraint Effect of Ductile Crack Growth Inhomogeneous Materials Application to Integrity Assessments Transferability of Test Data Accomplished Analyses for Specific RPV Integrity Assessments PTS test by Framatome PTS test by ORNL Application to Lifetime Assessment Assessment of Irradiation Embrittlement Changes Definition of Reference Curves and Their Use On Correlations Between T0 and Other Related Parameters Crack Arrest Reference Temperature TKIa Dynamic Versus Static Fracture Toughness T0 Versus Charpy V-notch Transition Temperatures Summary and Conclusions
434
4.14.1.3 4.14.1.3.1 4.14.1.3.2 4.14.1.3.3 4.14.2 4.14.2.1 4.14.2.1.1 4.14.2.1.2 4.14.2.1.3 4.14.2.1.4 4.14.2.1.5 4.14.2.2 4.14.2.3 4.14.2.4 4.14.2.5 4.14.3 4.14.3.1 4.14.3.2 4.14.3.2.1 4.14.3.2.2 4.14.4 4.14.4.1 4.14.4.2 4.14.5 4.14.5.1 4.14.5.2 4.14.5.3 4.14.6 References
Abbreviations bcc COD C(T) DC(T) EPFM LEFM NDT
Body-centered cubic Crack opening displacement Compact tension specimen Disk-shaped compact tension specimen Elastic–plastic fracture mechanics Linear-elastic fracture mechanics Nil-ductility transition temperature
NPP RPV SE(B)
436 438 438 439 440 442 442 442 443 443 444 444 445 447 449 450 451 451 455 455 455 456 456 457 458 459 459 460 460 460
Nuclear power plant Reactor pressure vessel Single-edge bend specimen
Symbols a a0
Surface crack depth Original (physical) crack depth/length
433
434
Fracture Toughness Master Curve of bcc Steels
b0 B B0 Bnet c d d¯ dc dN E J Jc Je Jp K, KI K0 K0Tref K0F KIc KI eff KI F KJc KJc(0.xx) KJc(limit) KJc(med) Kmin M m N P Pf Pfr Pmax PQ Pr{A/I} Pr{I} Pr{I/O} Pr{O} Pr{P/I} Pr{V} Pr{V/O} r rpl RTNDT RTT0
Specimen initial ligament (W a0) Specimen thickness Reference thickness Net specimen thickness Half length of surface crack Particle diameter Mean particle diameter Critical particle diameter Particle diameter scale factor Modulus of elasticity J-integral J at fracture instability Elastic component of J Plastic component of J Stress-intensity factor Fracture toughness scaling factor corresponding to 63.2% probability K0 corresponding to specific reference temperature Local K0 value at crack tip Linear-elastic fracture toughness Effective stress-intensity factor Local stress intensity at crack tip Elastic–plastic fracture toughness Fracture toughness for specific probability KJc capacity limit for specimen Average median KJc of test data Theoretical minimum fracture toughness Constraint coefficient Weibull exponent or material-dependent constant Number of initiators in volume element Specimen load Cumulative failure probability Probability of particle fracture Maximum load in KIc test Limit load in KIc test Conditional probability of arrest Probability of cleavage initiation Conditional probability of cleavage initiation Probability of ‘no event’ Conditional probability of propagation Probability of void initiation Conditional probability of void initiation Number of valid test data Radius of crack tip plastic zone Reference temperature by ASME Code Reference temperature by ASME Code Case
s T T0 T0 deep T0(margin) T28 J T41 J Ti TKIa Tref Tstress W Z b gp di Da G n s s0 spart sys syy V
Distance dimension along crack front Temperature Master Curve reference temperature 100 MPa √m T0 measured for deep crack case T0 including margin for uncertainty 28 J Charpy-V transition temperature 41 J Charpy-V transition temperature Single test temperature Crack arrest reference temperature 100 MPa √m Reference temperature T-stress parameter Specimen width Coefficient for selected confidence level Sample size uncertainty factor Particle surface energy Censoring parameter Stable crack growth Gamma function Particle size distribution shape factor or Poisson’s ratio Stress Stress scaling factor Particle stress Yield stress Normal stress Material constant
4.14.1 Introduction 4.14.1.1 Historical Review of Fracture Toughness Determination for Ferritic Steels Fracture mechanics is an engineering discipline which concerns the behavior of crack-like defects in structures or components and their effect on integrity. Initially conceived by Griffith during World War I, early applications were limited to the study of fracture of highly brittle materials (e.g., glass).1 Interest in the discipline languished until World War II, when 25% of the all-welded US Liberty ships experienced brittle fracture, exposing the urgent need to understand failure in ferritic structural steels and weldments. The earliest development in fracture mechanics of metals was focused on linear-elastic theory for understanding the fracture behavior of primarily high-strength steels and aluminum alloys. Application to brittle cleavage fracture in structures
Fracture Toughness Master Curve of bcc Steels
made of welded ferritic steels of low-moderate strength evolved later. One of the basic quantities of fracture mechanics is the stress-intensity factor, K, which is used to describe the loading condition of a cracked structure as a function of crack depth, a, and the applied stress, s. In the simplest case, a wide plate containing a central crack (2a), the loading condition can be expressed using the stress-intensity factor in the form: pffiffiffiffiffi K ¼ s pa ½1 Increasing the stress eventually results in a situation where the crack starts to propagate. Depending on the material and loading condition, this can occur by a ductile, cleavage, intergranular, or some mixedmode mechanism. Once the crack starts to propagate, the critical stress intensity (i.e., the fracture toughness) of the material has been exceeded. The fracture toughness is thus a material property, but it may be strongly affected by many environmental factors like temperature and air humidity. Using fatigue precracked test specimens, the fracture toughness of a material can be determined experimentally. This fracture toughness quantity can be used to evaluate the integrity of real structures with real or postulated flaws. The fracture mechanism can be different depending on the material, and for ferritic steels, different fracture modes are possible at different temperatures due to the ductile–brittle transition. Therefore, different approaches for the characterization of fracture mechanics are needed. The fracture toughness in body-centered cubic (bcc) ferritic steels exhibits a temperature dependence characterized by: (1) a low toughness cleavage initiation shelf at low temperatures; (2) an increasing transitional rise in toughness (going from cleavage to a mixture of cleavage and ductile-tearing fracture) with increasing temperature defined arbitrarily as a ductile–brittle transition temperature; and (3) an upper shelf characterized by fully ductile initiation fracture (Figure 1). The lower shelf and the region around the ductile–brittle transition temperature can be characterized using linear-elastic fracture mechanics (LEFM). LEFM considers material defects (flaws and cracks) and the effects of those defects on brittle cleavage crack behavior. LEFM is based on elastic stress analysis of the stress–strain field in the vicinity of the crack tip and a singularity called the stress-intensity factor, K. The linear-elastic theory was soon followed by elastic–plastic fracture mechanics (EPFM), which involved a different type of singularity parameter called the J-integral. Determination of
435
Ductile tearing J–R curve
KJc KIc
Transitional region Statistical size effect No statistical size effect KJc KIc
KJc
KJc
Ductile fracture
K KIc Jc Lower shelf Temperature
Figure 1 Schematic description of the fracture toughness transition region and parameters used to characterize fracture toughness in the lower shelf, over the transition region, and in and near the upper shelf where ductile cracking gradually becomes the predominant fracture mode.
material J-integral fracture resistance (J–R) curves expanded the scope of application to also include stable crack growth characterized by ductile tearing. Regardless of which theory is being used, it is necessary to know the material resistance to fracture, that is, the fracture toughness of the material being evaluated. Standardized test methods for determining material fracture toughness properties have been developed. In LEFM, fracture toughness is characterized by the parameter KIc; in EPFM, the initiation toughness parameter Jc (often converted to an approximate equivalent K value termed KJc) is used to characterize the onset of unstable crack growth under significant crack-tip plastic deformation conditions. (The statistical size effect and the elastic– plastic parameter KJc are associated with the Master Curve methodology discussed in Section 4.14.1.3.3. The linear-elastic parameter KIc is not presently recommended to characterize the transition region, as shown in Figure 1, due to the inherently large scatter of data in this region.) The J–R curve determination and parameters for the onset of stable crack growth are described in separate standards or sections of standards (not discussed here in detail).2 Historically, LEFM concepts for determining the fracture toughness of ferritic, bcc steels have been used, often together with conventional Charpy V-notch impact tests, to characterize the lower shelf and the transition fracture toughness region. There have been few alternatives to the LEFM methodology when combined with Charpy V-notch transition temperature results, as this is the current
436
Fracture Toughness Master Curve of bcc Steels
methodology applied for irradiated reactor pressure vessel (RPV) integrity following the ASME Boiler and Pressure Vessel Code. The LEFM approach itself is simple, because only the load record and specimen dimensions are needed for KIc determination, that is, due to the qualification requirements, the test tends to be invalid if there is any significant plastic area under the load versus displacement record. This restriction imposes a major disadvantage in that the amount of test material needed is often large, even if only one large specimen is tested. In this respect, the J-integral EPFM concepts using the parameter JIc are more applicable, as they make possible testing with smaller specimens due to less severe size requirements. Although current KIc and crack opening displacement (COD) testing standards better correspond to the latest fracture mechanics understanding (e.g., size restrictions relative to test specimen ligament and thickness dimensions), these standards generally are no longer recommended for characterizing the transition behavior of ferritic steels; they are more applicable to cases where the fracture mode is known to be ductile or possibly quasicleavage, and the material shows predominantly elastic behavior. The reason is that these older standards do not account for the statistical nature of the brittle fracture process in ferritic steels. More recently, a statistical assessment methodology, called the Master Curve procedure, has been developed as an improved method for characterizing the material fracture toughness (both LEFM and EPFM) of ferritic bcc materials, and for characterizing the temperature dependence of the transition temperature fracture toughness curve. It is the purpose of this chapter to provide a summary review of the Master Curve methodology. The following summary review of the Master Curve fracture toughness approach provides the basis and general framework for the methodology, but it also focuses on some key technical details that are often misunderstood. The reader should consult several of the references for greater detail regarding the various considerations needed in applying the Master Curve methodology for structural integrity assessments. In this chapter, the discussion is first devoted to LEFM involving the standard methods for experimentally determining the value of KIc; then, the more advanced approach based on EPFM and the approximate equivalent KJc is reviewed. Finally, the Master Curve procedure is discussed in depth and is the primary focus of this chapter.
4.14.1.2 Theoretical Background Leading to Use of EPFM and Data Distributions for Ferritic Steels The fracture toughness of ferritic steels has been characterized by numerous different parameters. It is not the purpose here to discuss history, so the main parameter is linear-elastic KIc and its use with respect to the elastic–plastic KJc. The KIc parameter has, in the past, been one of the most commonly used parameters, also for structural steels, but its limitations for describing the transition behavior controlled by both cleavage and ductile cracking are widely recognized (discussed later in detail) today. Many high-alloyed quenched and tempered steels, which exhibit practically no plastic deformation, still have moderate fracture toughness, KIc, and can be used and no additional benefit is achieved by using an elastic–plastic parameter like KJc. For these steels, the measurement of fracture toughness at one or a few temperatures is all that is necessary. For lowalloyed structural steels, which typically exhibit a pronounced ductile–brittle transition and may be loaded in a wide temperature range, the situation is different. In this case, an elastic–plastic parameter is needed. An example of such an application is an irradiated RPV where safety and performance have to be demonstrated in accident and abnormal service conditions which are more severe (i.e., at lower temperatures) than in normal operation. The integrity analyses must be based on the material fracture toughness covering the entire transition temperature range. In the mid and lower temperature portions of the transition curve, cleavage fracture has to be explicitly considered, and the parameter KJc and Master Curve analysis of data are the preferred methods to determine fracture toughness. It should be noted that application of the methodology is not just limited to ferritic steels, but the fracture is generally cleavage or a stress-controlled type of mechanism. When a cracked material is loaded, a plastic zone will develop at the crack tip. The size of this plastic zone depends on the crack tip loading (stress intensity) and the material yield strength. The radius of the plastic zone (rpl) can be expressed for mode I loading (tension loading perpendicular to crack plane) in a simple form as follows: rpl ¼
1 KI 2 2p sys
½2
Fracture Toughness Master Curve of bcc Steels
where KI is the stress-intensity factor and sys is material yield strength. The plastic zone size is thus a measure of plasticity at the crack tip and can be used to assess the applicability of different fracture toughness parameters. In predominantly elastic cases, the plastic zone size can be very small, and the material may be analyzed using LEFM. The parameter KIc is generally determined without correction term for plasticity, when the plastic zone size is small in relation to the specimen dimensions. Otherwise, an elastic–plastic parameter such as the J-integral should be used. For a compact specimen [C(T)] geometry and loading, the value of KI can be calculated directly from the load (Pi) and the original crack length (a0/W): a Pi 0 KðiÞ ¼ pffiffiffiffiffiffiffiffiffiffiffiffiffiffi f ½3 WBBnet W where W is specimen width, B is thickness, Bnet is net thickness, and function f (a0/W ) is defined dependent on specimen crack depth. For elastic–plastic conditions where a large plastic zone has developed and some stable crack growth can occur, one normally cannot use the LEFM parameter KIc and eqn [3] without at least some correction term(s). The use of EPFM is more practical to determine directly the J-integral which takes into account the plastic component of the work done to the specimen or component. The fracture toughness equivalence KI, denoted as KJc, can then be converted from the J-integral, which is first divided into elastic and plastic components: J c ¼ Je þ Jp
½4
where the elastic component of J is calculated from the elastic K (Ke) as follows: ðKe Þ2 ð1 v2 Þ ½5 E where E is the elastic modulus and n is Poisson’s ratio. The plastic component of J (Jp) is calculated from the total and elastic work using the measured load versus load line displacement or crack opening data. The value of Jc is converted to KJc as follows: rffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi E ½6 KJc ¼ Jc 1 n2 Je ¼
Despite the limitations of LEFM as discussed earlier, one of the basic standards applied in the past in fracture toughness testing has been American Society for Testing and Materials (ASTM) E 399 which
437
defines the methodology for KIc determination. This standard has recently been revised and issued as part of ASTM E 1820-08 (Annex 5),2 but the basic approach is essentially the same as in the previous versions of ASTM E 399. Because these standards, as well as a corresponding linear-elastic European standard, BS 7448, are still being used to assess some ferritic steels, it is important to know how they differ from the Master Curve approach used in ASTM E 1921-08.3 In particular, the scope of the application requires discussion since it affects how KIc data should be analyzed with respect to KJc. A common feature of LEFM KIc standards (such as the previous ASTM E 399) is that they express the fracture toughness as a single value, which should be a material property characterizing the resistance of a material to fracture. The value of KIc should be insensitive to specimen size, if the measured value fulfills the specified size criteria. When these conditions are met, the stress–strain condition at the crack tip has been thought to be predominantly plane strain, thus ensuring sufficient constraint to produce a minimum fracture toughness value for the material. The qualification to obtain a valid KIc measurement requires that relatively large test specimens have to be tested. The value of KIc is supposed to represent a lower limiting value of fracture toughness, but the method (ASTM E 1820, Annex 5) does not adequately cover ferritic steels where fracture by a cleavage mechanism in the transition or lower shelf region has known statistical characteristics different from ductile initiation fracture toughness.2 Based on the current knowledge of fracture mechanics, confirmed by numerical finite element model analyses of local stress–strain fields, many of the arguments for the LEFM parameter KIc are not valid, especially for ferritic steels.4 First, for crack tip constraint, the KIc size requirement has been shown to be overly conservative, leading to highly oversized specimens. Also, it has been shown that KIc is not a size-insensitive parameter; a size adjustment, similar to that made for Master Curve specimens (discussed later in Section 4.14.1.3), should also be made to the values of KIc to correct for size effects and to make the data comparable with KJc. Justification of this size effect argument has been demonstrated in many comparisons made between the KJc and the older KIc data measured with different size specimens.4 It is important to note that increasing the specimen size (both ligament and thickness) gradually diminishes the effect of the size adjustment, which means a reduced,
Fracture Toughness Master Curve of bcc Steels
but still existing, dependence on specimen size even with large dimensions (discussed later in Section 4.14.1.3.3). One should also remember that the Master Curve size adjustment is valid only in the transition region where cleavage fracture initiation is expected to occur, even after some stable crack extension before cleavage fracture. Due to the size effect, the fracture toughness decreases in the transition region with increasing specimen size whenever cleavage fracture is encountered. Examples of applying the statistical size adjustment are given in Section 4.14.3. As mentioned previously, ASTM E 1820 (Annex 5) actually invalidates the KIc determination if the value exhibits transition behavior indicative of some cleavage fracture. If KIc values characterize only the upper shelf behavior of ferritic steels, no size effect typical of the transition behavior should exist even if later cleavage initiation occurs. Thus, KIc determinations performed as per the previous ASTM standard versions do not necessarily fulfill the requirement or follow the recommendations now in ASTM E 1820. This issue is significant, because it concerns an essential qualification criterion of KIc determination. Another related issue is that ASTM E 1820-08 does not specify selection of the test temperature to make sure that there is no cleavage fracture in the test. No posttest measures are required or recommended to confirm that the test data are not affected by cleavage initiation. The question arises: which of the reported KIc values should be size-adjusted consistently with KJc data and which should not be size corrected? Another important aspect is the 95% secant requirement (Pmax/PQ 1.1), which limits stable crack growth to very small amounts for typical structural steels exhibiting a rising tearing resistance curve.5 The methodology for KIc determination according to the ASTM E 1820 standard is thus not applicable for high-tearing resistance or high-toughness steels, which is also mentioned in Annex 5. The size requirement for KIc is essential, since it means that the specimen (ligament) size is the only factor which can be affected in pursuing a valid test result, if the test has to be performed at a specific temperature. If the secant requirement cannot be met, it is possible that no value of KIc can be determined at that temperature. For all of these reasons, fracture toughness testing in the transition region is recommended to be made following standard ASTM E 1921, which takes into account the statistical nature of cleavage fracture. Direct fracture toughness determination for the reactor vessel surveillance programs of nuclear power plants (NPPs) was one of the first applications of the
Master Curve methodology. Compared to the conventional Charpy V-notch methods, the Master Curve concept represents a new approach which makes possible direct fracture toughness determination with only a few relatively small specimens, which is a more efficient use of limited test material. Using the Master Curve method allows statistical confidence to be applied to the directly measured data. The traditional practice of estimating fracture toughness from Charpy data (using correlations) is more unreliable due to large uncertainties associated with the correlations and the subsequent safety margins to meet regulatory requirements. However, the Charpy-based methods are still in use and will continue to be used until a large amount of surveillance capsule Master Curve data is available. Existing data from correlations is discussed in Section 4.14.5, and material reference curves are discussed in Section 4.14.4.2. Use of these Charpy-based methods may remain to be the only way of estimating static fracture toughness in some cases, but is not the preferred approach for future surveillance programs. 4.14.1.3 Master Curve Methodology as Developed by Wallin 4.14.1.3.1 History
Master Curve methodology is based on the observation that the fracture toughness transition curve for any ferritic steel has the same shape, no matter the steel (see Figure 2). Thus, a single ‘Master Curve’ can be used for all ferritic steels; the curve is simply shifted along the temperature axis to match a mean fracture toughness value, which is established from measured fracture toughness data for the particular steel being evaluated. The primary material fracture
200
KJc = 30 + 70 exp [0.019 (T-T0)] Master Curve
KJc (MPa Öm)
438
150
100
50
0 -80
-60
-40
-20
0
20
40
60
80
T-T0 (⬚C) Figure 2 Definition of Master Curve which is defined as a representative mean fracture toughness curve for most structural bcc steels with moderate strength.
Fracture Toughness Master Curve of bcc Steels
toughness parameter is the transition temperature, T0, which characterizes the Master Curve position, and is defined as the temperature where the median fracture toughness is 100 MPa √m (Figure 2). The advantages of Master Curve technology over past methods for estimating the fracture toughness of materials (particularly irradiated materials) are (1) it is based on direct measurements of the property of interest (e.g., fracture toughness); (2) it provides a direct method of establishing the transition curve for irradiated materials (instead of inferring a shift in an assumed baseline bounding curve using Charpy data); and (3) it can be used for materials even with a limited availability of archival materials. The development of Master Curve methodology was started in the 1980s by Wallin and his coworkers by the introduction of a mathematical model to describe the probability of cleavage fracture initiation in a material containing a distribution of potential fracture initiators (flaws). The model was completed by including the temperature dependence of KJc, which was estimated empirically from a dataset including various ferritic structural steels. The scatter definition, the size adjustment, and the definition of the temperature dependence are the basic elements of the Master Curve methodology described in ASTM E 1921.6–8 The approach has been verified in several round-robin and research programs.9 The first version of the Master Curve standard comprised a procedure for analyzing only singletemperature test data; in later versions, the approach was extended to consider multitemperature test data. The multitemperature approach requires finding a maximum likelihood solution for the value of the transition temperature, T0, from data measured over a range of temperatures, rather than at a single temperature. The first version of ASTM E 1921 was approved in 1997 and issued in 1998 (ASTM E 192197). The multitemperature approach was included in the second revision, after which several other revisions with some minor changes have been released. The present revision, ASTM E 1921-08, describes procedures for the experimental determination of the elastic–plastic fracture toughness, KJc, estimation of the reference temperature, T0, and principles for the lower bound curve definition of fracture toughness (Figure 3). Further detailed descriptions on the methodology and applications are given in McCabe et al.10 and Sattari-Far and Wallin11. The model12 has also been validated numerically to more accurately describe the true fracture behavior and the stress–strain distribution of bcc steels on a
439
Measurement of elastic– plastic KJc data
Statistical estimation to determine reference temperature T0
Construction of KJc–T reference curves Figure 3 Master Curve fracture toughness determination according to ASTM E 1921.
micromechanical scale.13 The advanced numerical assessment capabilities presently available for multiscale modeling of materials have made it possible to validate the main elements of the model. Despite several further developments primarily related to material inhomogeneity (in the basic model material macroscopic homogeneity is assumed), the basic approach being applied today is essentially the same as that developed over 20 years ago. No other deficiencies or assumptions requiring readjustment have been identified. Further verification of the approach and the validity of the empirically determined temperature dependence have been conducted. Some aspects, such as the lower shelf definition (Kmin), are practically impossible to verify only using experimental methods, so that numerical modeling studies have been very instructive. The Master Curve methodology is currently being used in both structural integrity and lifetime assessments. Typical areas of application are pressure vessels and piping, nuclear RPV surveillance programs, other energy production structures, off-shore structures, and various welded components and bimetallic joints. 4.14.1.3.2 Basics
The applied fracture model12 defines a conditional probability of fracture assuming that the distribution of flaws in the material follows a statistical distribution. The conditional probability term takes into account the possibility of void formation (blunting of a crack initiator) and the crack arrest–propagation event (Figure 4). The fracture event is controlled in the model by assessing the criticality of a single crack initiator from the weakest link principle. The basic elements of the methodology – the scatter definition and the specimen size correction6,7 – are based on this cleavage fracture model. It is assumed that the material has uniform macroscopic properties and
440
Fracture Toughness Master Curve of bcc Steels
Stress applied to material element Pr{V/O}
No initiation
Void initiation Pr{A/I}
Arrest
V = Volume
Pr{I/O} Cleavage initiation
Pr{P/I} Propagation Failure
Pr{I}, N
{
Pr{O}
s = Stress
s
Cleavage initiator distribution
Pr{I} = Probability of cleavage initiation Pr{V} = Probability of void initiation Pr{O} = Probability of ‘no event’ Pr{I/O} = Conditional probability of cleavage initiation (no prior void initiation) Pr{V/O} = Conditional probability of void initiation (no prior cleavage initiation) Pr{P/I} = Conditional probability of propagation (in the event of cleavage initiation) Pr{A/I} = Conditional probability of arrest (in the event of cleavage initiation) Figure 4 Definition of the conditional probability of cleavage fracture. Reproduced from Wallin, K.; Laukkanen, A. Eng. Fract. Mech. 2008, 75(11), 3367–3377.
that no global interaction exists between the crack initiators. An overview (basic equations) of the cleavage fracture model (also called the Wallin, Saario, To¨rro¨nen (WST) model) is described next. The conditional probability of cleavage initiation, Pr{I/O}, is expressed as a product of the probability of having a cleavage initiation and that of not having a void initiation as follows: PrfI=Og ¼ PrfIgð1 PrfV=OgÞ
½7
Equation [7] can be approximated as: PrfI=Og PrfIgð1 PrfIgÞ
½8
The WST model expresses Pr{I/O} as the product of the particle fracture and nonfracture probabilities so as to take into account the previously broken particles which do not contribute to the cleavage process as follows: 1 ð ½9 PrfI=Og Pfr ð1 Pfr ÞPfd g@d dc
where the critical particle size dc is defined by a Griffith type expression as follows: dc ¼
2pE 0 gp s2yy
½10
where syy is the tensile stress ahead of the crack tip, E 0 is the plane strain modulus of elasticity, and gp is the particle surface energy.
The probability of particle fracture is described by a Weibull-type dependence accounting for particle size (d) and particle stress (spart) as follows: 3 ! spart m d Pfr ¼ 1 exp ½11 s0 dN where dN and s0 are scaling factors. The particle size distribution, P{d}, is given by eqn [12], which describes the size distribution with two parameters, the average particle size (d ) and the shape factor n. For pressure vessel steels, the shape factor appears generally to be in the range 4–6. ðn 2Þðn1Þ d n n2 ½12 exp Pfd g ¼ Gðn 1Þ d d =d A detailed description of the WST model is given in Wallin et al.12 and Wallin and Laukkanen.13 The recent numerical validation of the model is presented in Wallin and Laukkanen.13 4.14.1.3.3 Methodology
The methodology of determining the Master Curve value of T0 is described next and follows that given in ASTM E 1921-08.3 ASTM E 1921-08 also covers the testing procedure and specimen preparation, which are not described here. The reference temperature, T0, is defined as the temperature at which the mean fracture toughness for a 1-in. (25.4 mm) thick fracture toughness specimen equals 100 MPa √m. It is the main parameter used to
Fracture Toughness Master Curve of bcc Steels
define the curves for the mean and lower bound fracture toughness. If the dataset covers several test temperatures, T0 is found as a solution of equations giving the maximum likelihood estimate to this value. Other than the LEFM testing standards for fracture characterization, there is no specified limit for the minimum specimen size. However, the number of specimens to be tested has to be at least six and generally increases when the specimen size is decreased in order to produce a statistically acceptable confidence level for the estimate. The method also includes a censoring procedure that allows the use of adjusted invalid test data, which contain statistically useful information. The probability of cleavage fracture initiation is described as a three-parameter Weibull distribution, which defines the relationship between the cumulative failure probability (Pf) and the fracture toughness level before or at KI as follows: " # KI Kmin 4 ½13 Pf ðKJc KI Þ ¼ 1 exp K0 Kmin where Kmin is the theoretical lower bound fracture toughness, set normally to 20 MPa √m for steels with yield strength from 275 to 825 MPa, and K0 is the scale parameter corresponding to Pf ¼ 63.2%. The Weibull exponent is assumed to have a constant value equal to 4 based on theoretical and experimental arguments.7 The measuring capacity (maximum KJc) of a specimen depends on its dimensions (ligament size) and the material yield strength as follows: sffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi Eb0 sys ½14 KJcðlimitÞ ¼ Mð1 n2 Þ where E is the modulus of elasticity, n is Poisson’s coefficient, b0 is the initial specimen ligament length (specimen width minus the initial crack depth, W a0), M is the constraint value (usually set equal 95% No size adjustment
to 30), and sys is the material yield strength at the test temperature. Toughness values exceeding this measuring capacity require censoring and are set at the level of maximum KJc. The expression for predicting the specimen size effect is based on the cleavage fracture model. The fracture toughness (KJc(x)) corresponding to the desired specimen thickness or crack front length (Bx) is obtained from the values of KJc(0) and B0, respectively, as follows: 1=4 B0 ½15 KJcðxÞ ¼ Kmin þ ðKJcð0Þ Kmin Þ Bx When analyzing measured data according to ASTM E 1921, the size adjustment is normally made to 1 in. (Bx ¼ 1 in. or 25.4 mm). Note that the influence of side grooves on the specimen thickness is ignored. After the size adjustment is made for each KJc measurement, the data for different size specimens can be described as one population following the same Master Curve form and scatter as shown in Figure 5. The aforementioned formula applies in the transition range, and it is not necessary to perform the size conversion at temperatures below (T0 50 C), because the size effect diminishes (see the fracture toughness curve 50 MPa √m shown in Figure 6). The procedure for estimating the maximum likelihood solution for T0 from data measured at various temperatures was published in 199514 and added to the second and later revisions of ASTM E 1921. The value of T0 is solved iteratively from the following equation, which includes a factor d for data censoring: n X di expf0:019½Ti T0 g i¼1
11 þ 77expf0:019½Ti T0 g n X ðK JcðiÞ Kmin Þ4 expf0:019½Ti T0 g i¼1
95%
ð11 þ 77expf0:019½Ti T0 gÞ5
¼0
95%
Statistical size adjustment KJC
KJC
5% Small specimens
441
5%
5%
Small specimens
Large specimens T Figure 5 Schematic presentation of Master Curve size adjustment.
Large specimens T
½16
442
Fracture Toughness Master Curve of bcc Steels KJc(med) (MPa Öm, B = 25 mm)
350
250
0.27
KJc = 200 MPa Öm
80
100 120 140 160 180 200 220
0.24
KJc = 150 MPa Öm
0.21
200
min. 6 specimens
0.18
KJc = 100 MPa Öm
150
60
0.15
n
B = 25 (MPa Öm) KJc
300
KBJc= 25 = (KJc-Kmin)*(B/25 mm)0.25 + Kmin
100
0.12
KJc = 50 MPa Öm
0.09
50
min. 7 specimens min. 8 specimens
0.06 0
20
40
60
80 100 120 140 160 180 200 B (mm)
Figure 6 Effect of the statistical size adjustment at different levels of KJc.
where KJc(i) is the size-adjusted fracture toughness measured at temperature Ti and di is the censoring coefficient: di ¼ 1 if the KJc(i) datum is valid (less than the limit determined from eqn [14]) or di ¼ 0 if the KJc(i) datum is not valid but may be included as censored data. When the value of T0 has been solved, the fracture toughness curves for specific levels of fracture probability are obtained from the following equations for probability level 0.xx: 1=4 1 KJcð0:xxÞ ¼ 20 þ ln 1 0:xx f11 þ 77 exp½0:019ðT T0 Þg
½17
The mean curve (50% failure probability) for 25.4 mm specimen thickness is obtained as a function of temperature from KJcðmeanÞ1T ¼ 30 þ 70 exp½0:019ðT T0 Þ
½18
4.14.2 Analysis of Fracture Toughness Test Data for Master Curve Application 4.14.2.1 Standard Test and Analysis Procedure (ASTM E 1921) 4.14.2.1.1 Test specimens used
ASTM E 1921 specifies three specimen types for testing: standard compact tension (C(T)), diskshaped compact tension (DC(T)), and single-edge notched bend (SE(B)). The type shall be selected based on the form and dimensions of the end-product or component (plate, forging, bar, etc.). The C(T)
0.03 0.00 -60 -50 -40 -30 -20 -10 0 10 20 30 40 50 60 T-T0 (⬚C) Figure 7 The number of specimens required by ASTM E 1921 for T0 determination (n ¼ 1/number of specimens) and how it depends on temperature and median KJc.
specimen is commonly used, however, in the nuclear power industry, the SE(B) specimen with a square cross-section is preferred since that is a form that can be used directly from the Charpy V-notch test specimens contained in most current RPV surveillance programs. The standard test method does not directly specify the minimum specimen size, but only an upper limit for the measuring capacity (KJc(limit)) depending on the size of the ligament (b0 ¼ W a0) (see eqn [14]). A value exceeding this limit may be included in the analysis as censored data, even if the test did not occur by cleavage fracture initiation. Those values shall be lowered to the limit value and included as censored (with reduced weight as indicated in eqn [16]) data. It should be noted that censored values shall not be included in determining the minimum number of specimens required for the T0 determination (discussed in Section 4.14.2.1.3). When very small-size specimens are used, more specimens need to be tested for a valid analysis. The effect of test temperature (T T0) or the corresponding median KJc on the required minimum number of valid test values is given in Figure 7, where parameter n is a weight factor (ASTM E 1921) representing the inverse of the number of specimens needed for overall validity. When using subsize specimens, the expected number of specimens required for a valid analysis is, however, larger. The estimated numbers of specimens needed with different size single-edge notched bend specimens are given in Table 1.15 Two SE(B) specimen configurations are specified: square (W ¼ B) or rectangular (W ¼ 2B) cross-sections.
Fracture Toughness Master Curve of bcc Steels Table 1 Expected number of specimens needed for a valid Master Curve analysis with different size single-edge notched bend specimens Specimen type
Expected number of specimens
10 10 5 10 55 34 3.3 3.3
7 7 12 28 40
Source: Wallin, K.; Planman, T.; Valo, M.; Rintamaa, R. Eng. Fract. Mech. 2001, 68, 1265–1296.
55
1
W = 10
60⬚
b=5 a0 = 5
0.5
B=5
45⬚
4.5
Figure 8 Example of a 5 10 mm single-edge notched bend specimen design for fracture toughness testing according to ASTM E 1921.
Due to the same ligament size (b0), the W ¼ 2B geometry has the same measuring capacity as the C(T) specimen with the same thickness. Commonly used configurations are, for example, in RPV surveillance programs, the Charpy-size (10 mm square crosssection) and the half Charpy-size (5 10 mm rectangular cross-section), which also have the same measuring capacity (equal b0). For typical medium strength (quenched and tempered) structural steels, these specimens have been found to give almost identical results.15 Side grooving of specimens is optional (total side-grooved depth may not exceed 0.25B), but recommended to increase stress triaxiality near the specimen surfaces.15 The side-grooved half Charpy-size design, with dimensions, is shown in Figure 8. It should be noted that the total specimen thickness (not net thickness) shall always be used in the Master Curve analysis independent of the side-grooving. Testing with miniature specimens are discussed in Section 4.14.2.2.
443
4.14.2.1.2 Determination of reference temperature T0
There are two methods of determining the value of T0: the single-temperature method, which is used when the tests have been conducted at a single temperature; and the multitemperature method, used if testing is performed at more than one temperature. In the first case, the analysis is simpler and can be performed in analytical form, but the results do not provide further insight on temperature dependence or the lower shelf behavior. The multitemperature analysis can be performed only iteratively, but it gives more comparative information on temperature dependence, scatter, and the lower shelf behavior. As for their effects on T0, both methods give statistically equal confidence levels if the numbers of the valid test data are equal and the measurements have been made close enough to the final T0 (see Figure 7). If sufficient specimens are available, the multitemperature method is generally preferred since there is less risk in exceeding the limits of T0 50 C (discussed next) and the maximum KJc from eqn [14]. 4.14.2.1.3 Data qualification
Data qualification is described in detail in ASTM E 1921 and is not repeated in detail here. In addition to actions associated with the test procedure and equipment, there are measures which have to be undertaken after each test or test series to ensure that only valid data will be included in the final T0 analysis. The optimum test temperature range for the T0 determination is generally selected iteratively by taking into account the already measured data. For example, three or four tests may be conducted at a selected temperature and then a preliminary T0 is determined from the data; subsequent test temperatures are then based on that preliminary T0. As discussed previously, the Master Curve cleavage fracture model is accurate only in the transition area, where stress state and cleavage crack initiation are the main controlling factors for the fracture event. The data used for T0 determination should therefore be populated in the mid-transition area rather than near the lower or upper shelf. The shape of the transition curve causes the uncertainty of the T0 determination to increase when data measured near the lower shelf are used. There can be an optimum range within the ASTM E 1921 validity range of T0 50 C (Figure 9) depending upon test specimen size. The resultant data outside of the validity range are excluded from the analysis but still can be compared to the Master Curve determined for the
444
Fracture Toughness Master Curve of bcc Steels
The lower and upper tolerance bound (KJc(0.xx)) for the estimated fracture toughness in KJc ¼ f (T) is calculated from a revised T0 (T0(margin)) as follows:
200 95% KJC (MPa Öm)
150
KJc(limit)
T0 ðmarginÞ ¼ T0 þ DT0
5%
1=4 1 1 0:xx
11 þ 77 exp 0:019 T T0 ðmarginÞ
KJcð0:xxÞ ¼ 20 þ ln
100
50
0 -75
½20
-50
-25
0 T-T0 (⬚C)
25
50
75
Figure 9 Example of a Master Curve analysis showing the validity window for material with yield strength 455 MPa and specimen size 10 10 mm SE(B).
½21
where 0.xx is the selected cumulative failure probability. For 1, 2, and 5% cumulative failure probability, the bounds are as follows:
KJcð0:01Þ ¼ 23:5 þ 24:4 exp 0:019 T T0ðmarginÞ ½22
KJcð0:02Þ ¼ 24:1 þ 29:0 exp 0:019 T T0ðmarginÞ
valid temperature range data. The closer the test temperature is to T0, the fewer the number of test specimens that are needed. However, when testing small specimens, the maximum KJc limit is closer to the 100 MPa √m level at T0, and the test temperature generally has to be moved to temperatures below T0 with more specimens than the minimum number needed for a valid analysis.
KJcð0:05Þ ¼ 25:2 þ 36:6 exp 0:019 T T0ðmarginÞ ½24 When the dataset consists of several test temperatures, the median KJc is obtained from the basic relationship (eqn [18]) as follows: eq
KJcðmedÞ ¼
4.14.2.1.4 Determination of lower bound curves
A lower bound curve can be constructed to correspond to a lower limiting fracture probability, which normally is set to 5% or 2%, taking into account the uncertainty of the T0 determination. The uncertainty of determining T0 depends on the number of specimens used to establish T0. The uncertainty (DT0) is defined from a normal distribution with two variables, the test temperature, and the number of specimens used for the T0 determination, as follows: b DT0 ¼ pffiffiffi Z r
½23
½19
where b ¼ 18–20 C, depending on the value of T T0, r is the number of valid test data used to determine T0, and Z is the confidence level (e.g., Z85% ¼ 1.44). The median KJc is used to determine the value of b and the uncertainty of T0 according to ASTM E 1921. When KJc(med) is equal to or greater than 83 MPa √m, b ¼ 18 C. Alternatively, b ¼ 20 can be used for all values of KJc(med) not less than the minimum of 58 MPa √m.
r 1X f30 þ 70 exp½0:019ðTi T0 Þg ½25 r i¼1
where r is the number of valid test data. The 2% lower bound curve is sometimes used as a criterion to determine if the material should be analyzed by taking into account possible material inhomogeneity. This further analysis can be done using the SINTAP procedure (discussed in Section 4.14.2.5), which ensures that a conservative lower bound definition is obtained regardless of possible low fracture toughness values.16,17 If there are values below the 2% curve, the data preferably should be analyzed using the not yet standardized multimodal procedure.17 An example of the 2% curve and the effect of the basic SINTAP analysis are shown in Figure 10. 4.14.2.1.5 Limits of applicability
The applied cleavage fracture model strictly applies only to the transition region of the material, although the model also includes a term to take into account the conditional probability of crack propagation. This term is needed because in and near the lower shelf, the fracture event is mainly controlled by crack
445
Fracture Toughness Master Curve of bcc Steels
propagation and therefore cannot be correctly described by only the crack initiation term. The situation is complicated by the fact that the lower shelf is not always close to the assumed, theoretically estimated, constant value of 20 MPa √m (although it usually is). It is also assumed (Section 4.14.1.3.2) that the crack initiators are randomly distributed and that no global interaction exists between the crack initiators. The material is also assumed to be macroscopically homogeneous. The method applies for transgranular, cleavage fracture events, although it may be used, with caution, for intergranular-type fracture especially when the fracture event is predominantly stress controlled (as is typical in the lower transition area). Figure 11 shows scanning electron microscope views of both transgranular cleavage and intergranular fracture mode surfaces. The scope of applicability and the limitations of the method are also discussed in Sections 4.14.2.5 and 4.14.3.1 and the application to structures in Section 4.14.3.2.
4.14.2.2 Aspects of Applying Small and Miniature Specimens An advantage of the Master Curve method is that it makes possible the use of Charpy-size and even smaller specimens for a valid determination of fracture toughness and T0. The number of tests always has to be determined so that the minimum required confidence level is achieved for the estimation. ASTM E 1921 describes a special weighting system to ensure a sufficient confidence level for the analysis. The final check can be made only after the value of T0 is rather well known, when the adequacy of tests can be determined from the condition P ni ri 1, where ri is the number of valid data in the valid temperature range i and ni is the weight factor of this range. Normally, the required number can be tentatively determined only after some tests have been conducted, because the optimal test temperature range is not known beforehand. With very small specimens, the final number of tests needed for a
250
250
KJC (MPa Öm)
150
95%
200 KJC (MPa Öm)
Cleavage Excessive DCG T0 = 69 ⬚C B0 = 25 mm
200
M = 30
100
5% 2%
50
150
Cleavage Excessive DCG T0 = 74 ⬚C B0 = 25 mm
95%
M = 30
100 5% 50 SINTAP
0 (a)
0
20
40
60 80 T (⬚C)
100
120
0 0
140 (b)
20
40
60 80 T (⬚C)
100
120
140
Figure 10 Example of the 2% lower bound definition for a dataset with DT0 ¼ 9 C, assuming b ¼ 18 C (a), and the same dataset analyzed with the SINTAP procedure showing in this case no inhomogeneity (b). Material: irradiated A508 Cl. 2 steel, 10 10 mm single-edge bend specimens (data with excessive ductile crack growth are indicated).
Figure 11 Examples of fracture surfaces of ferritic steels: the left one appearing as pure cleavage (crack initiation shown) and the right one as mostly intergranular mode.
446
Fracture Toughness Master Curve of bcc Steels
10
Charpy size specimen 55
10
10
5 ⫻ 10 Charpy specimen 55
5 ⫻ 5 specimen 5
4
3 ⫻ 4 (KLST) specimen
5
27
3
27
5
Figure 12 Possible single-edge bend specimen geometries for testing using ASTM E 1921; dimensions are in millimeter.
valid estimate will likely be larger than the given minimum of six because of the smaller validity window, which is reduced with decreasing specimen size and the material yield strength (see Figure 7). Examples of possible small and miniature SE(B) specimen geometries for fracture toughness testing are compared in Figure 12. As mentioned previously, a commonly used specimen configuration for irradiated RPV steels is the full Charpy-size geometry with 10 10 mm crosssection. For most applications, this geometry provides a sufficiently large validity window (Figure 9), because as few as six specimens may be sufficient for a valid estimate. From the present experience for irradiated steels with different size specimens, Charpy square or half Charpy rectangular SE(B) or 0.4T or 0.5T C(T) geometries are generally optimum when the amount or form of the test material is limited. In some cases, even smaller test specimens may be required due to very small amounts of test material. A comparison made between T0 and scatter estimates from test results measured with miniature specimens (i.e., smaller than the Charpy-size) shows that the definitions of scatter and the measuring capacity (specimen size) criterion (eqn [14]) apply even to miniature specimens that are of 3 4 mm and 3.3 3.3 mm cross-section SE(B).15 The results indicate no bias between the T0 estimates measured with the miniature specimens compared to the overall
mean values of T0, which is shown in Figure 13 for the Charpy-size specimen and three subsize SE(B) specimens. The comparison demonstrates, in all respects, applicability of the miniature size specimens for the fracture toughness estimation using the Master Curve approach. In some datasets (e.g., on grade 15 Kh 2 MFA and on the International Atomic Energy Agency (IAEA) reference material JRQ), the smallest specimens (3 4 and 5 5 mm) produced some low T0 values, which most likely were caused by macroscopic inhomogeneity encountered with such small specimen dimensions. An example of miniature specimen results for the A508 Cl. 3 steel FFA (a French steel grade) is presented in Figure 14, demonstrating nearly consistent fracture toughness data independent of the specimen size. Another benefit of using the Master Curve approach is that the uncertainty associated with the T0 estimation can be determined and taken into account for assessing a conservative, realistic estimate for the lower bound fracture toughness. As discussed previously, ASTM E 1921 does not set limits for the specimen size or configuration; however, the minimum number of test results, dependent on test temperature relative to T0, is predefined to ensure an acceptable minimum confidence level for the estimate. If a larger uncertainty for the T0 estimation is accepted, the minimum number of specimens required can be reduced from that given in the standard. Correspondingly, having less than the minimum
447
Fracture Toughness Master Curve of bcc Steels
100
100
ASTM E 1921
ASTM E 1921 50 T0 5 ⫻ 10 (⬚C)
T0 10 ⫻ 10 (⬚C)
50 0 -50 -100
0 -50 -100
10 ⫻ 10 mm
5 ⫻ 10 mm
-150
-150 -150
-100
0 -50 T0 All (⬚C)
50
-150
100
-50
0
50
100
T0 All (⬚C)
100
100 ASTM E 1921
ASTM E 1921
50
50 T0 3 ⫻ 4 (⬚C)
T0 5 ⫻ 5 (⬚C)
-100
0 -50 -100
0 -50 -100
5⫻ 5 mm
3⫻ 4 mm
-150
-150 -150
-100
-50
0
50
100
-150
T0 All (⬚C)
-100
-50
0
50
100
T0 All (⬚C)
Figure 13 Comparison of the values of T0 measured with normal and subsize Charpy specimens relative to the mean of all specimens (10 10, 5 10, 5 5, and 3 4 mm single-edge bend specimens). Reproduced from Wallin, K.; Planman, T.; Valo, M.; Rintamaa, R. Eng. Fract. Mech. 2001, 68, 1265–1296.
number of valid data in the test series does not necessarily invalidate the dataset, but it does result in a lower confidence level of the estimate. It is essential that the most realistic confidence level is estimated and taken into account in final integrity assessments. Very small specimens (around 3 3 mm) tend to give slightly (1–3 C) higher values of T0 compared to 10 10 mm specimens.15 This trend is likely due to the censoring procedure, which screens out proportionally more data from the upper tail of the dataset than from the lower tail. This screening affects both the scatter caused by possible material inhomogeneity and that from statistical outliers. The optimal test temperature range for miniature specimens has been proposed to be 50 C T – T0 20 C. Even though more specimens are needed when smaller specimens are tested, the consumption of test material becomes smaller, even if more than the minimum number of specimens were tested
for the T0 estimate. In this respect, the 5 5 mm specimen is the least material consuming SE(B) size (12 specimens are needed for the standard estimate if sys ¼ 500 MPa).15 Using the 5 5 or 3 4 mm specimen geometry, it is possible to prepare up to 8 or 12 subsize specimens from the halves of one tested 10 10 mm specimen (see Figure 15). When selecting specimen configuration, it should be noted that with deeply cracked specimens having the ligament size equal to or less than the thickness, the ligament is the primary dimension limiting the specimen-measuring capacity, not thickness. Also, with slim (reduced thickness) and very small specimens (like the 3 4 mm cross-section), it is recommended to side-groove the specimens to increase stress triaxiality near the surfaces. 4.14.2.3
Effect of Constraint
Loading a cracked structure creates a local stress concentration ahead of the crack tip. A situation
448
Fracture Toughness Master Curve of bcc Steels
A508 Cl. 3 (FFA) sY = 434 MPa B = 3–10 mm W = 4–10 mm
350
KJC (MPa Öm)
300 250 200
10⫻10 5⫻10 5⫻5 3⫻4 T0 = –114 ⬚C B0 = 25 mm M = 30 censoring
150 100
95% 5%
50 0
-160 -140 -120 -100 T (⬚C)
-80
-60
-40
Figure 14 Master Curve analysis of steel FFA (A508 Cl. 3) KJc data measured with Charpy-size and different subsize or miniature specimens (10 10, 5 10, 5 5, and 3 4 mm single-edge bend specimens). Reproduced from Wallin, K.; Planman, T.; Valo, M.; Rintamaa, R. Eng. Fract. Mech. 2001, 68, 1265–1296.
Charpy-size specimen KLST-size specimen 10 4 3 27
10
55 Figure 15 Preparation of 3 4 mm miniature specimens from a 10 10 mm specimen.
where brittle fracture initiation is likely only in this restricted area around the maximum stress corresponds to small-scale yielding conditions. The high stress area is localized ahead of the crack tip extending to the border of the crack tip plastic zone or three to five times the distance from the crack tip to the stress maximum.18 In such a situation, the stress distribution ahead of the crack can be described correctly with the J-integral. When loading exceeds small-scale yielding, large-scale yielding is involved and the J-integral can no longer describe the crack tip stress distribution correctly. At that point, the measuring capacity of the specimen has been exceeded. The limit for measuring capacity is normally given as a function of the material yield strength and specimen ligament in the form of eqn [14]. The value of M has been proposed from a variety of finite element calculations to be from 50 up to 200.18 Based on present understanding and examination of many
experimental studies, a value of M ¼ 30 can be regarded as a realistic estimate for cleavage fracture with most structural steels. Exceeding the load level given by eqn [14] will lead to gradually increasing amounts of ductile tearing before cleavage fracture initiation. The basic analysis procedure described in ASTM E 1921 is intended for specimens and structures for which at least a moderate level of constraint (triaxial stress state) is achieved. With deep-cracked specimens or thick-wall structures including internal cracks, the constraint is typically high enough for the standard T0 analysis. The situation is different with low-constraint geometries like surface cracks and thin-sheet structures where the small-scale yielding condition may not prevail (note that the limit for this condition also depends on material strength properties). In principle, the basic Master Curve approach can also be applied for such lowconstraint conditions, but the estimation may be overly conservative due to high plastic deformation which is outside the applied fracture model assumptions. In ASTM E 1921 test conditions, sufficient constraint is assured by defining valid tests only as those that exhibit brittle fracture initiation below or at the capacity limit value (eqn [14]) and by limiting the ductile crack extension prior to brittle fracture initiation. The probability of cleavage initiation is controlled by the narrow zone ahead of the crack tip where smallscale yielding condition prevails. Several approaches like the Tstress, the Q-parameter, and small-scale yielding corrections have been developed to account for the effect of plastic deformation due to low constraint.19 The Tstress analysis is relatively straightforward since a simple elastic stress analysis can be used instead of a numerical large-scale yielding model. Another advantage of using the Tstress is that it may be performed assuming no change in temperature dependence which allows the constraint effect to be described only as a shift in T0. Consequently, the Tstress yields conservative estimates compared to the more complex Q-parameter approach. On the other hand, the Q-parameter is more accurate for very low-constraint situations. On the basis of a simplified linear-elastic analysis, the correction to T0 due to low-constraint SE(B) geometry (a0/W 0.1) with a negative Tstress has been expressed in the form19: T0 T0 deep þ Tstress < 0
Tstress for 10MPa C1 ðfor SEðBÞ specimens onlyÞ
½26
Fracture Toughness Master Curve of bcc Steels
20
T0-T0 deep (⬚C)
0
290–350 Mpa 490–680 Mpa 720–1380 Mpa
Best fit Best fit Best fit
-20 -40 -60 -80
-100 -800
T0 ≈ T0 deep+Tstress/10 MPa / ⬚C-1 Tstress < 0
-600
-400
-200
0
200
Tstress (MPa) Figure 16 Effect of Tstress on the value of T0 based on tests made with single-edge bend specimens having different crack depths. Reproduced from Wallin, K. Eng. Fract. Mech. 2001, 68(3), 303–328.
where T0 is the corrected value and T0 deep is the value measured for a deep crack case. Equation [26] is an empirical result from data consisting of only SE(B) specimens (Figure 16). Additional work has been conducted to refine the expression by including test data from C(T) specimens and by a comparison of solutions based on Tstress and the Q-parameter. Thus, a more accurate formula for estimating the T0 from Tstress and the T0 of a deep crack case has been proposed in the form20: T0 T0 deep þ
Tstress 12MPa C1
for Tstress <300 MPa
ðfor SEðBÞ and CðTÞ specimensÞ
½27
The main difference between eqns [26] and [27] is that the latter also covers positive values of Tstress up to 300 MPa, whereas in eqn [26], it is assumed that only a negative Tstress has a marked effect on the value of T0. 4.14.2.4
Effect of Ductile Crack Growth
The effect of ductile tearing on the measured J-integral is addressed in ASTM E 1921 by limiting the amount of ductile crack growth prior to brittle fracture to avoid inaccuracies from excessive plastic deformation. In some standards, specific formulas for correcting the effect of excessive ductile crack growth have also been presented. It is possible to correct in this way for small crack growth relative to the excessive plastic deformation that occurs, but the statistical effect associated with the probability of cleavage fracture initiation cannot be corrected. Ductile tearing preceding brittle fracture affects the measured fracture toughness by increasing the
449
volume ahead of the crack tip where brittle fracture initiation can occur, which increases the cumulative probability of failure. This statistical crack growth effect is comparable to the statistical size effect, which also is due to the increased volume of material under stress, increasing the probability of brittle fracture initiation as the crack length increases. In addition to increasing the volume of potential cleavage initiators, ductile tearing also tends to change the crack tip stress distribution. On the other hand, a small amount of ductile tearing can be regarded as beneficial since it can increase the stress triaxiality at the crack tip. This triaxiality effect is of minor importance when compared to the statistical effect that is discussed next. The statistical effect is due to the preceding ductile crack growth and is dependent on the amount of the crack growth.18,21 A simplified expression for the fracture probability (Pf) at stress intensity (KI), considering a small amount of ductile tearing (Da 1 mm), is given in the form18,22: " # B KI Kmin 4 2O2 Da Pf ¼ 1 exp 1þ 2 B0 K0 Kmin KI ð2m þ 1Þ ½28 where m and O are material-dependent constants, B is specimen thickness, B0 and K0 are scale factors, and Kmin ¼ 20 MPa √m. The value of O has been estimated to be 5500 MPa for medium strength structural steels (sy ¼ 300–550 MPa). Equation [28] can be simplified by setting (2m þ 1) ¼ 1 when O becomes 4700 MPa. Parameter m is defined from the ductiletearing power law function of the form: f ðDaÞ ¼ J1mm Dam
½29
where J1 mm is the value of J-integral at 1 mm crack growth and Da is crack growth. By substituting the expression for failure probability (eqn [13]) into eqn [28], one can derive a more practical formula giving a corrected value for KI (KI eff) due to ductile crack growth as follows: 1=4 2O2 Da ½30 KIeff ¼ Kmin þ ðKI Kmin Þ 1 þ 2 KI ð2m þ 1Þ Equations [29] and [30] can be used to take into account the increased fracture probability due to small amounts of ductile tearing, in addition to a possible standard correction for plastic deformation. Consideration of the statistical ductile-tearing effect may become relevant for high-strength steels with a
450
Fracture Toughness Master Curve of bcc Steels
low-strain hardening capacity or for steels exhibiting low ductile-tearing resistance. An example of analyses corrected for ductile crack growth is shown in Figure 17. The material is a thermally embrittled pressure vessel steel (A508 Cl. 3), which therefore has a high T0 (þ69 C). In this case, the correction lowers the fracture toughness in the upper transition region, but has a negligible effect on the value of T0 and the behavior near the lower shelf. 4.14.2.5
Inhomogeneous Materials
In the Master Curve brittle fracture model, it is assumed that the steel is macroscopically homogeneous. In real wrought steels, the macro- and microstructures are seldom fully homogeneous due to various effects occurring during production and cooling, such as segregation of elements and development of inclusions. The resultant inhomogeneity may manifest itself as excessive data scatter. However, in welds or dissimilar metal joints, the material in the associated heat-affected zone (HAZ) may exhibit strongly distributed ‘inhomogeneity’ so that two or more subpopulations can be distinguished in the test data. Often in the final application of the fracture toughness data, it is only essential to determine a statistically sound conservative estimate for the lower bound fracture toughness, which can then be used to construct a conservative reference curve for the material. In most cases, even those where materials show slight inhomogeneity, the standard Master Curve 250
KJC (MPa Öm)
200 150
Cleavage Ductile T0 = 69 ⬚C B0 = 25 mm DCG corrected
95%
100 5% 50 0 -100
-50
0
50 T (⬚C)
100
150
Figure 17 Effect of ductile-tearing correction on a pressure vessel steel (embrittled A508 Cl. 3 with sy ¼ 676 MPa, and specimen thickness 10–20 mm). The vertical lines indicate the ASTM E 1921 temperature validity area. From Wallin, K.; Planman, T., Eds. In Use and Applications of the Master Curve for Determining Fracture Toughness, Proceedings of the Workshop MASC 2002, Helsinki/Stockholm, June 12–14, 2002; VTT Industrial Systems: Espoo.
estimation is sufficient, without further assessments. For testing the data population for possible inhomogeneity, several approaches have been developed. A common feature for all of these approaches is that they are based on the standard Master Curve, but are extended to properly take into account the inhomogeneity. The following extended approaches have been proposed for analyzing inhomogeneous materials: 1. Simplified SINTAP procedure for determining a conservative lower bound estimate of fracture toughness when the inhomogeneity is of randomly distributed type.16,20 2. Bimodal or multimodal models for analyzing datasets showing distributed inhomogeneity (typically welds and HAZ materials).17,20 3. Model for analyzing randomly inhomogeneous materials, including, for example, macroscopic segregations.17,20 Of the three methods, SINTAP is the simplest and the one recommended for initial assessment of the quality of data; it can generally be applied in conservative structural integrity analysis when material inhomogeneity is suspected. An advantage of the SINTAP procedure is that it can be performed for a small dataset, unlike the bimodal or multimodal analyses, which require 15 or more data points for a valid assessment. It should be noted that the SINTAP analysis does not produce a statistically representative description of the whole dataset, but its primary purpose is to determine whether the dataset is homogeneous and secondly, to develop a generally conservative lower bound when possible material inhomogeneity is present. The model for analyzing randomly distributed inhomogeneities gives a statistically correct description of the fracture toughness data, but requires a large dataset and is more complicated to perform than the SINTAP analysis. When there are a sufficient number of fracture toughness data points, the bimodal/multimodal and the randomly distributed inhomogeneous models should be used to better identify a more statistically correct description for a modified Master Curve. The simplified SINTAP procedure is described next. The SINTAP analysis is based on the maximum likelihood method like the basic Master Curve procedure and it consists of three steps as follows. Step 1 is the standard estimation giving the first estimate of T0 and the median fracture toughness according to ASTM E 1921. The resulting T0 is used as an input value for Step 2.
Fracture Toughness Master Curve of bcc Steels
Step 2 is a lower tail maximum likelihood estimation and is performed so that all data exceeding the median fracture toughness are censored by substituting d ¼ 0 and by reducing the corresponding KJc values to the median curve. If the resulting new T0 (T0–2) is lower than the T0 from Step 1 (T0–1), then T0-SINTAP ¼ T0–1; otherwise, Step 2 shall be repeated using the last T0 estimate (T0–2) as a new input value until a constant T0 (T0–2 adj) is obtained. If the number of valid data is ten or more, T0-SINTAP ¼ T0–2 adj, otherwise Step 3 shall be performed. Step 3 gives the minimum value estimate of T0 and is performed only if the number of valid test data is less than 10. In Step 3, an additional safety factor is incorporated for cases where the number of tests is small. Here, the value of T0 is estimated for each single data point to find the maximum value of T0 (T0 max). If the resulting T0 max>T0–2 adj þ 8 C, T0-SINTAP ¼ T0 max, otherwise T0-SINTAP ¼ T0–2 adj.
4.14.3 Application to Integrity Assessments 4.14.3.1
Transferability of Test Data
The statistical size adjustment enables one to extend the fracture toughness estimation to specimens and structures with different crack front lengths. A longer crack front means that a larger volume of material is exposed to tensile stress ahead of the crack tip, which increases the probability that a crack exceeding the critical size will exist in this volume leading, according to the weakest link theory, to brittle fracture. This size effect is addressed in a conservative way by the size adjustment formula (eqn [15]). Generally, the size adjustment is made for all test specimens so that a single reference thickness (normally 1 in.) is used for the Master Curve. As described earlier, when approaching the lower shelf, the size effect diminishes to zero for both EPFM KJc data and LEFM KIc data; note that KIc data are often characterized as being size independent, but as previously described in the transition region, a size correction appears to be applicable. An example is presented in Figure 20, which shows LEFM fracture toughness values (including both KIc and KQ) from a
The steps and the data censoring in Steps 1 and 2 are described schematically in Figure 18. The flow chart of the iterative procedure is shown in Figure 19.
P = 50% Censoring KJc (MPa Öm)
Specimen measuring capacity limit
KJc (limit) = Data used for MML estimate of T0
Eb0sys 30(1−ν2)
T (⬚C) Lower tail MML estimate P = 50%
Step 2 Censoring
Minimum value estimate Step 3 used only when n < 10
KJc
P = 50%
KJc Single data used to estimate T0
Data used for MML estimate of T0
T (⬚C)
451
T (⬚C)
Figure 18 Data censoring in SINTAP Steps 1 and 2 and the T0 determination in Step 3.
452
Fracture Toughness Master Curve of bcc Steels
Step 1: ASTM E 1921 MML estimation T0–1
T0–2
Step 2: Lower tail MML estimation T0–2 (adj.) T0–2 > T0–1 No
n < 10
T0 = T0–2
No
T0 = T0–1
Yes
Yes
Repeat Step 2 until constant T0–2 using the last T0–2 in estimation
Step 3: Minimum value estimation T0 max Yes
T0 max – T0–2 > 8 ⬚C No
T0 = T0 max
T0 = T0–2 (n = total number of specimens)
Figure 19 Flow chart for the SINTAP procedure.
A533B Cl. 1 (HSST 02) sY = 480 MPa center
100
KIC KQ 1T 2T 4T 6T
200
4&6T
KIC (MPa Öm)
KIC (MPa Öm)
150
250
150
A533B Cl. 1 (HSST 02) sY = 480 MPa center KIC KQ 1T 2T 4T 6T B0 = 25 mm
100
50 50 1&2T 1&2&4&6T 0
−150
−50
−100
0
T (⬚C)
0
−150
−100
−50
0
T (⬚C)
Figure 20 Fracture toughness data (KIc and KQ) from the heavy-section steel technology program measured with different size specimens showing the data before (left) and after (right) the size adjustment. Material: A533B Cl.1 steel of HSST 02 (sys ¼ 480 MPa). Modified from Sattari-Far, I.; Wallin, K. Application of Master Curve Methodology for Structural Integrity Assessments of Nuclear Components; SKI Report 2005:55; Swedish Nuclear Power Inspectorate: Stockholm, 2005; p 178.
large dataset measured in the heavy-section steel technology (HSST) program with different size specimens ranging from 25.4 up to 152 mm (6 in.) thickness.23 The data from different specimen sizes follow the same Master Curve prediction after the size adjustment as shown in the second plot in Figure 20. Another example of applying the size adjustment is presented in Figure 21, showing LEFM KIc data measured by MPA (Materialpru¨fungsanstalt Universita¨t Stuttgart, Germany) with different size
specimens, including two very large ones (B ¼ 500 mm).24 A distinct size effect is visible in the uncorrected data (Figure 21, plot on the left), but not in the data size adjusted to the 25.4 mm specimen thickness (plot on the right in Figure 21). The above transferability principle also applies for a structure when the size adjustment is made for a real or postulated crack front length, if the material and load conditions remain constant over the whole crack length. If they do not, their variation has to
Fracture Toughness Master Curve of bcc Steels
20 MnMoNi 5 5 (KS 15) sY = 602 MPa B = 25–500 mm 350
20 MnMoNi 5 5 (KS 15) sY = 602 MPa B = 25–500 mm 200 B = 25 mm B = 50 mm B = 500 mm
300 250 KIC (MPa √m)
KIC (MPa √m)
150
100
200
B = 25 mm B = 50 mm B = 500 mm T0 = – 19 ⬚C B0 = 25 mm 95%
150 100
50
5%
50 0
453
−100
−50
0
50
0
−100
−50
0
T (⬚C)
50
T (⬚C)
Figure 21 Fracture toughness data measured on steel 20MnMoNi55 (KS 15, sys ¼ 602 MPa) by MPA without size adjustment (left) and size adjusted to specimen thickness 25 mm (right). Modified from Sattari-Far, I.; Wallin, K. Application of Master Curve Methodology for Structural Integrity Assessments of Nuclear Components; SKI Report 2005:55; Swedish Nuclear Power Inspectorate: Stockholm, 2005; p 178.
be taken into account. This situation is typical in a thick-wall structure, where the crack front is not usually straight and the temperature and loading may significantly vary along the crack length. By combining the three-parameter Weibull distribution (eqn [13]) and the formula for the statistical size adjustment (eqn [15]), one can derive a formula for the cumulative failure probability of the form: ( ) B KI Kmin 4 ½31 Pf ¼ 1 exp B0 K0 Kmin where B is the specimen thickness, B0 is the normalization thickness (usually 1 in. or 25 mm), Kmin is the minimum fracture toughness (20 MPa √m), KI is the stressintensity factor, and K0 is the scaling fracture toughness corresponding to 63.2% failure probability. In this formula, KI represents the crack driving force, whereas K0 is specific for the material, and B for the specimen or geometry of the structure or component. In test configurations with simple and small specimen geometry, both KI and K0 can be taken to be constants over the whole crack length with sufficient accuracy. In a real structure, where a real three-dimensional surface crack is possible, both KI and K0 may vary depending on the location (F) along the crack front and should be treated as variables (Figure 22). Note in Figure 22 the designation of the surface crack depth (a), the crack length on the surface (2c), the total surface crack front length (s), the angle f along the crack front, and the wall thickness (t). A more
S t f a c Figure 22 Definition of quantities along the surface crack front.
general expression for the cumulative failure probability can thus be expressed in the form25: 9 8 s < ð K K 4 ds = IF min ½32 Pf ¼ 1 exp : K0F Kmin B0 ; 0
By defining an effective stress-intensity factor (KI eff) corresponding to a specific reference temperature (Tref), which can be the minimum temperature along the crack front, eqns [31] and [32] can be combined to determine the effective stress intensity corresponding to the same failure probability as eqn [32]: KIeff Tref ¼
8s <ð K :
0
IF Kmin K0F Kmin
ðK0Tref Kmin Þ þ Kmin
4
91=4 ds = B0 ; ½33
KI F is obtained from a stress analysis as a function of location (F). K0Tref is the standard, high constraint
454
Fracture Toughness Master Curve of bcc Steels
Master Curve K0, corresponding to the reference temperature (Tref) along the crack front and has the form for 63.2% failure probability of: ½34
K0 F is the local K0 value, based on local temperature and constraint and can be expressed in the form: K0 F ¼ K0T ;Tstress ¼ 31 þ 77 exp 0:019 T T0 deep
Tstress 10 MPa C1
½35
Equation [35] is directly applicable with the ASME Code Case N-629 fracture toughness reference curve,26 since it is written in terms of the standard deep specimen T0.11(ASME Code Case N-629 and N-631 allow the determination of RTT0 when T0 is measured, see Section 4.14.4.2) Equations [33]–[35] give the effective crack driving force, normalized to represent a standard Master Curve 25.4 mm crack front (B0) and the minimum temperature along the crack front. It should be noted that the area of applicability of the constraint correction based on the Tstress has not yet been fully established.19 The Tstress actually is a LEFM concept that does not work when excessive plasticity is present. In this case, a more advanced concept should be used, such as the Q-parameter or a local approach.27 The Tstress equation (eqn [26] as applied in eqn [35]) works well for the negative values of Tstress, but the effect saturates at higher values. However, in actual components, the Tstress is generally negative. The new Tstress equation [27] is expected to be valid up to the Tstress value of 300 MPa.25 The fracture toughness can be expressed either with the 5% lower bound Master Curve, which can be expressed in the form:
K5%;Tref ¼ 25:2 þ 36:6 exp 0:019 Tref T0 deep ½36 or by using the fracture toughness reference curve from ASME Code Case N-629 or N-631.26,28 Details on these Code Cases are presented in Section 4.14.4.2. The following expressions are derived: KICASME;Tref ¼ 36:5 þ 11:4
exp 0:036 Tref T0 deep
½37
or KICASME;Tref ¼ 36:5 þ 3:083 expf0:036ðTref RTT0 þ 56 CÞg
½38
KlC, N-629 200 KI eff, KIC (MPa Öm)
K0Tref ¼ 31 þ 77 expf0:019ðTref T0 Þg
250
150
100
Kl eff KIC 5% MC
50
0 -50
0
50 100 Tmin-T0 deep (⬚C)
150
200
Figure 23 Comparison of Master Curve analysis of real flaws (5% curve) and the ASME Code Case N-629 – curve. Reproduced from Sattari-Far, I.; Wallin, K. Application of Master Curve Methodology for Structural Integrity Assessments of Nuclear Components; SKI Report 2005:55; Swedish Nuclear Power Inspectorate: Stockholm, 2005; p 178.
The curves are compared in Figure 23. Note that the fracture toughness curve is not directly compared to the crack driving force estimated from stress analysis. Instead, the fracture toughness is compared to an effective driving force, which accounts for the local stress and constraint state and temperature along the crack front, as well as the crack front length. In this way, it is possible to combine the classical fracture mechanics and Master Curve analyses, and to present the comparison in a conventional format. One should remember, however, that postulated flaws often contain unrealistically long crack fronts. An assumed quarter thickness elliptical surface flaw (a/t ¼ 1/4; c/a ¼ 3), as used in the ASME Code for pressure–temperature operating curves for RPVs, may be used from a conservative deterministic driving force perspective (as was the intention), but from a statistical size adjustment point of view, this assumption is too conservative. If such postulated flaws are analyzed using KI eff, an additional size adjustment to the Master Curve is recommended. Note that s for a 200-mm thick RPV would correspond to an a of 50 mm, 2c of 300 mm, and an s of about 400 mm. A more realistic maximum crack front length (s) is 150 mm or less. This value of s is also consistent with much of the original KIc data for the ASME Code KIc curve and therefore justifiable in terms of the functional equivalence principle. The form of KI eff for an excessively large postulated flaw (s 150 mm) becomes11:
455
Fracture Toughness Master Curve of bcc Steels
:
0
IF Kmin K0F Kmin
4
91=4 150ds = B0 s ;
ðK0Tref Kmin Þ þ Kmin
200
150
½39
where s is the postulated crack front length. Applied nondestructive evaluation (NDE) techniques and other evidence suggest that a value of s less than 150 mm generally can be justified. If warm prestress (WPS) transients need to be analyzed, they can be assessed based on the maximum KI along the crack front and eqns [36]–[38]. 4.14.3.2 Accomplished Analyses for Specific RPV Integrity Assessments A typical pressurized thermal shock (PTS) experiment can be conducted by loading a thick-wall pressure vessel with an embedded crack by an external load and a severe thermal transient. The aim of a PTS experiment is to load the test vessel so that unstable crack initiation occurs. In connection with these large-scale PTS tests, a material test program is typically performed to produce the necessary material property data. If the thermal transient includes large temperature changes, any crack initiation after the peak temperature may be affected by the WPS effect. WPS involves the material loaded at the peak temperature during the thermal transient to a stress-intensity level that is higher than the critical fracture toughness at lower temperatures during the transient. Because the material was preloaded to a higher toughness level, the critical fracture toughness at a lower temperature of the transient is higher than without the preloading. The WPS effect is complex and is related to the crack tip stress state and is expected to be most pronounced in short transients where strain ageing will not diminish the effect. 4.14.3.2.1 PTS test by Framatome
In the 1980s, Framatome performed a thermal shock pressure test for a thick-wall pressure vessel, including a very long but shallow crack.29 The crack length (2c) in the 230-mm-thick cylinder was 1000 mm and the depth (a) was 17 mm. The vessel material (A508 Cl. 3) was characterized by KIc tests conducted with large, 75 and 100-mm thick, C(T) specimens. The aim was to characterize the material directly ahead of the crack front. The data from these characterization tests have been reanalyzed using the Master Curve
KIC (MPa Öm)
KIeff Tref ¼
8s <ð K
3T-CT B = 75 mm 4T-CT B = 100 mm TSE 2c = 1000 mm Filled point 1st initiation T0 = -35 ⬚C (25 mm) B0 = 1000 mm
95%
5%
100
50
0 -100
-50
0 T (⬚C)
50
100
Figure 24 Data from the thermal shock test performed by Framatome29 showing the first initiation (filled symbol) and two other initiations (open symbols) together with fracture toughness data (B0 ¼ 1000 mm). The Master Curve analysis was performed by Wallin.30
method.30 The analysis was performed by making the size adjustment of the material test data to the 1000 mm crack length, corresponding to that of the vessel. After this correction, the material characterization data and the crack tip load data for the test vessel for the observed first, second, and third initiations should fall on the same Master Curve. The comparison in Figure 24 shows that the vessel initiations occurred within the 5% and 95% fracture probability bounds estimated for the material, that is, two initiations fell almost on the mean Master Curve and the third close to the 95% probability level. In this case, the result indicates no WPS effect between the initiations, although some effect can be expected due to some warm prestressing during the decreasing temperature of the transient. The C(T) fracture toughness and the PTS test data generally coincide for all three initiations. 4.14.3.2.2 PTS test by ORNL
Oak Ridge National Laboratory (ORNL) performed several PTS tests as part of the HSST program. One test (TSE-7) was performed in the 1980s on a pressure vessel (steel A508 Cl. 2) with 152-mm wall thickness.31 The initial crack shape was semielliptic, but the form changed in the test to something very irregular, which complicated the analysis of this test. The material was characterized by KJc tests made in the transition region with 25.4 mm C(T) specimens. In the PTS test, neither KI nor temperature was constant along the crack length. Due to subsequent inaccuracies in the original KI data, the Master Curve reanalysis30 was performed using the apparent maximum
456
Fracture Toughness Master Curve of bcc Steels
KI and the local temperature of the crack tip area. The crack length 2c ¼ 37 mm was used as the effective crack length. Due to the crack shape changes in the test, only the first initiation was included in the Master Curve analysis. The estimated stress intensity in the first initiation is shown together with the fracture toughness data adjusted to a crack length of 37 mm in Figure 25. Despite the inaccuracies, the result of the test is also consistent with the measured fracture toughness data.
4.14.4 Application to Lifetime Assessment 4.14.4.1 Assessment of Irradiation Embrittlement Changes The Master Curve methodology was originally developed for applications like RPV surveillance programs, for which the conventional methods are generally less accurate (based on Charpy V-notch energy shifts) or not suitable due to specimen size requirements (LEFM KIc and KIa tests). As a direct measurement approach, the Master Curve approach is preferred over the correlative and indirect methods, based mostly on the Charpy test, used in the past to assess irradiated RPV integrity. It is therefore reasonable to expect that the future determination of plant-operating limits will be based on the Master Curve and related methods rather than on the 350
KJC, KIC (MPa Öm)
300 250
1T-CT B = 25 mm TSE-7 2c = 37 mm T0 = –32 ⬚C (25 mm) B0 = 37 mm
M = 30 95%
200 150 100
5%
50 0
−100
−50
0
50
T (⬚C)
Figure 25 Data from the heavy-section steel technology (TSE-7) test showing the first initiation (filled symbol) and the fracture toughness results (B0 ¼ 37 mm). Reproduced from Cheverton, R. D.; Ball, D. G.; Bolt, S. E.; Iskander, S. K.; Nanstad, R. K. Pressure vessel fracture studies pertaining to the PWR thermal-shock issue: Experiment TSE-7; NUREG: CR-4304 (ORNL-6177); Oak Ridge National Laboratory: Oak Ridge, TN, 1985, p 133. The Master Curve analysis was performed by Wallin.30
indirect methods or the trend curves based on the chemical composition of materials and their expected neutron fluence. The advantages achievable with the ASTM E 1921 methodology especially for RPV applications are as follows: Direct fracture toughness estimation of the RPV using irradiated small SE(B) or C(T) type specimens and determination of a statistically correct mean behavior for ageing assessment and a realistic lower bound curve definition for integrity assessments.15,32–34 Expansion of the analysis to cover issues related to material inhomogeneity and the quality of measured data are possible utilizing the proposed statistical methods. Expansion of the maximum likelihood estimation to take into account specimen-specific fluence data is possible when needed. Different data and those measured with different size and type specimens can be included in the analysis (including LEFM KIc with caution). (Note that C(T) and SE(B) specimens may show a (usually about 8 oC) bias due to different geometry so that C(T) specimens yield higher T0.) Utilization of correlations between the parameters characterizing different loading conditions like crack arrest and dynamic loading. A typical situation especially for older RPVs is that the existing material data consists of very miscellaneous information on the properties of materials, such as test results measured with numerous different material conditions, test standards, specimen types, equipment, etc. In such a situation, a method for handling and analyzing all available data in a synergistic way can provide significant savings especially if there are no archive test materials available for additional testing. These aspects of characterization are described schematically in Figure 26. It is important to note that, based on present knowledge, the shift in Charpy transition temperature (e.g., DT41 J) due to neutron irradiation on average is close to or less than the transition temperature shift in fracture toughness (DT0 from the Master Curve method); that is, the shift in Charpy data is generally unconservative in respect of the corresponding shift in T0. However, there is large scatter in the relationship between these two shifts, and caution is needed when assessing equivalence. Although there are no specific requirements for Master Curve testing in RPV surveillance programs
Fracture Toughness Master Curve of bcc Steels
Surveillance specimens
CVN
3-PB
Neutron dosimetry Prefatiguing
457
Fluence E (T) USE LE (T)
Instrumented CVN-tests Correlations
ASTM E 1921 test
KIc KIR KIa
& Reconstitution
KJc (T)
Prefatiguing
KJc (T)
ASTM E 1921 test Dynamic KJc test
KJd (T)
Data evaluation using the Master Curve approach
Statistically defined fracture toughness of irradiated RPV steels Figure 26 Characterization of irradiated reactor pressure vessel materials and related parameters.
in the current Codes and Regulations, the methodology has been applied in national RPV surveillance programs, and numerous retroactive analyses have been made using data measured in the past in the surveillance and material characterization programs of NPPs. Today, there are also many publications and guidelines which can be used as guidance for using the Master Curve and the related methodologies effectively and in a proper manner. Applications for nuclear grade pressure vessel materials and irradiated materials are addressed in the IAEA publication Technical Report Series No. 429.33 Chapter 4.05, Radiation Damage of Reactor Pressure Vessel Steels, provides additional information and mechanistic details of ‘Radiation Damage of Reactor Pressure Vessel Steels.’ 4.14.4.2 Definition of Reference Curves and Their Use The use of a Master fracture toughness Curve is not a new concept. The ASME Boiler and Pressure Vessel Code, Section III, Appendix G, has used a lower bound static fracture toughness curve that is a ‘Master Curve’ indexed using the reference temperature
RTNDT.35 Currently, the ASME KIc and KIR curves, indexed to the RTNDT of the material, describe the fracture toughness of the RPV and its lower bound variance with temperature. These curves were adopted in the early 1970s as a lower bound representation to a relatively small set of linear-elastic fracture toughness (KIc) and linear-elastic arrest toughness (KIa) values for 11 heats of RPV steel.36 The use of RTNDT to normalize temperature was intended to account for the heat-to-heat differences in fracture toughness transition temperature, thereby collapsing the fracture toughness data onto a single curve. However, RTNDT is not always successful in this regard, often providing a very conservative characterization of fracture toughness. RTNDT is the material/heat-specific ASME Code-defined temperature per Section III, NB-2300,37 based on a combination of drop-weight nil-ductility transition temperature (NDT) and Charpy V-notch tests (>68 J) for nonirradiated materials; or for irradiated materials, RTNDT is the nonirradiated RTNDT (IRT) plus the shift in the 41 J Charpy V-notch temperature to account for irradiation (assumed to be DRTNDT). Appendix A to Section XI of the ASME Code38 uses the same lower bound Master Curve as Section III,
458
Fracture Toughness Master Curve of bcc Steels
Appendix G, for crack arrest toughness, and another Master Curve (again indexed using RTNDT) for static initiation fracture toughness. This approach has been used now for over 30 years in the US nuclear industry and has been shown to be very reliable in that there have been no vessel failures, albeit very conservative for most materials. Kim Wallin’s direct fracture toughness Master Curve provides a more complete representation of the material fracture toughness. The Master Curve reference temperature T0 can be used in an analogous manner as RTNDT to index the position of the Master Curve. The obvious advantages of this approach are: The index temperature itself is based on measured fracture toughness rather than Charpy V-notch and drop-weight tests. The Wallin Master Curve has a well-described statistical shape that allows for better-defined direct use in either deterministic or probabilistic analyses. Direct measurement of irradiated fracture toughness eliminates the need to add a shift to an initial value for many applications and it provides, to some extent, the possibility to extrapolate outside the already characterized fluence area. ASME Code Cases N-62926 and N-63128 were published in 1998 and utilize the ASTM E 1921 test method for determining T0. These Code Cases permit the use of a Master Curve -based index temperature (RTT0 ¼ T0 þ 19.4 C) as an alternative to RTNDT. Code Case N-629 is for Section XI applications for both irradiated and nonirradiated RPV steels; Code Case N-631 is essentially the same Code Case, but it is for Section III design applications for only nonirradiated RPV steels. These Code Cases allow the determination of RTT0 when T0 is measured. Application to RPV integrity requires the knowledge of uncertainties associated with the use of a measured RTT0 in place of RTNDT. The use of Master Curve in the United States has been limited to a few examples for which the Nuclear Regulatory Commission (NRC) has written a safety evaluation (SE): The first use was indirect in that Master Curve data were used to justify a lower value of nonirradiated RTNDT for some Linde 80 welds rather than defining a value of RTT0 39; this application for the Zion RPVs was actually submitted and approved before ASTM E 1921-97 or the ASME Code Cases were finalized. The key approved use of Master Curve was for the Kewaunee RPV. The NRC did not accept the
utility submittal approach,40 but modified it to reflect their interpretation of approximating the current Charpy shift-based regulatory approach.41 This interpretation resulted in the use of a deterministic Margin term that was larger than that used for the Charpy data application. However, there was still enough beneficial gain using the Master Curve approach for determining an irradiated RTT0 (over the current regulatory approach) to allow the utility to move forward to replace steam generators and pursue license renewal. The most recent SE was issued for the plants that have vessels containing Linde 80 welds. The Babcock and Wilcox (B&W) fabricated welds were used in all of the B&W design vessels and some Westinghouse design vessels fabricated by B&W. The initial RTNDT for these weld metals has always been uncertain since these weld metals tend to have low upper shelf levels that often fall below 68 J after irradiation. Since 68 J is part of the transition temperature definition for RTNDT, these welds may be unduly affected by the Charpy 68 J temperature requirements. Therefore, the B&W Owners Group developed a program to better define the initial nonirradiated RTNDT using the Master Curve and the RTT0 approach in Code Cases N-629 and N-631. Their approach utilized Charpy V-notch testing to get the 41 J transition temperature change for assessing the effects of radiation embrittlement in the same manner as currently used for RTNDT. Irradiated surveillance program materials were evaluated using the Master Curve to compare with the predictive method of initial RTT0 þ DT41 J. The methodology was accepted by the NRC but requires explicit margins to be applied.42
4.14.5 On Correlations Between T0 and Other Related Parameters The Master Curve transition temperature T0 for quasistatic loading conditions is statistically precise and measurable following testing standards such as ASTM E 1921. Extension of this same type of approach to dynamic loading (KId and KJd) and crack arrest (KIa) situations seems logical, and empirical studies with a large variety of structural steels have confirmed a similar relationship. In both cases, the material property (KJd and KIa) can generally be
Fracture Toughness Master Curve of bcc Steels
described with the same temperature dependence as the quasistatic initiation fracture toughness, but with an increased value of T0. The correlations between T0 and other related toughness parameters are discussed next. 4.14.5.1 Crack Arrest Reference Temperature TKIa For an integrity assessment of real structures, it is often necessary to have information not only on the initiation fracture toughness but also on the crack arrest toughness, KIa. A definition of the reference curve for crack arrest toughness is given in the ASME Code, Appendix A of Section XI. The parameters describing crack initiation, including the Master Curve T0 and the related parameter RTT0 , cannot be used to directly describe the crack arrest toughness. Associated with the development of the Master Curve concept, studies have concluded that it is possible to develop correlations describing the relationship between the crack initiation and arrest toughness.43 These studies have focused on clarifying which elements of the Master Curve approach should be modified for assessing crack arrest; and finding possible correlations between initiation and arrest parameters. Due to different mechanisms and differences in factors controlling fracture initiation and arrest events (e.g., the local properties are crucial for crack initiation, but not so critical for crack arrest), the weakest link theory applied in the Master Curve approach is not directly suitable for crack arrest. The analyses of nine well-defined crack arrest datasets, consisting of various pressure vessel base and weld metals (including those used to construct the ASME reference curve), confirm that43: No statistical size adjustment should be made to crack arrest data. The scatter seems to be material independent, but lower than the scatter for crack initiation. KIa data follow the same Master Curve temperature dependence as KJc. Crack arrest data can thus, in general, be described following the same Master Curve approach, using the reference temperature TKIa to characterize the temperature corresponding to the crack arrest toughness at a level of 100 MPa √m, consistent with the crack initiation transition temperature T0. To clarify if a reasonable correlation exists between T0 and TKIa , a total of 54 datasets, for
459
which both crack initiation and arrest data were available, have been analyzed with the Master Curve concept.43 The result shows an exponentially decreasing trend for TKIa T0 as a function of T0, but the standard deviation of this correlation is high. Taking into account the observed scatter, a rough estimate for the maximum TKIa can be obtained from the value of T0 at a confidence level of 85% from: TKIa ¼ T0 þ DT þ 19 C
½40
where DT is obtained from the quasistatic T0 as follows:
sys T0 þ 273 þ DT ¼ exp 5 136:3 C 683:3MPa
½41
where sys is the material yield strength. Equation [41] is recommended only for steels where the nickel content is less than 1%. The crack arrest toughness (TKIa ) can also be assessed from instrumented Charpy V-notch data using the correlation developed between TKIa and the temperature corresponding to the crack arrest force of 4 kN. The method has been used for assessing the crack arrest toughness of irradiated RPV steels from the existing Charpy V-notch data. This correlation and its application are described in Wallin.43 4.14.5.2 Dynamic Versus Static Fracture Toughness The Master Curve concept, developed originally for quasistatic loading condition, has proven to be useful for dynamic KJd tests conducted at a high loading rate. Assuming that only the value of T0 is rising (along with the material yield strength) with increasing loading rate, dK/dt, the dynamic fracture toughness, KJd and associated T0 should, in principle, be estimated from the value of T0 measured by static tests. This dependence was empirically evaluated in 1997 by Wallin using a large dataset consisting of dynamic and static fracture toughness data measured for various structural steels (yield strength ranging from about 200 to nearly 1000 MPa). The Master Curve method was applied to both the static and dynamic loading rates.44 Based on these results, as well as an IAEA round-robin exercise (report to be published in the IAEA Report Series within the framework of the Technical Working Group on Life Management of Nuclear Power Plants), the Master Curve approach
460
Fracture Toughness Master Curve of bcc Steels
appears to be fully applicable to dynamic fracture toughness measurements conducted in the ductileto-brittle transition region. 4.14.5.3 T0 Versus Charpy V-notch Transition Temperatures In the case of dynamic loading with notched specimens, the correspondence with the fracture toughness test and T0 is complicated due to several uncertainties associated with the Charpy V-notch impact test. First, the loading situation is very different in the dynamic loading of a notched specimen compared to the quasistatic loading of a fatigue precracked specimen. Due to the differences in the loading conditions, the measured Charpy energy includes a significant proportion of both crack initiation and propagation, and often some energy associated with crack arrest; whereas, the quasistatic SE(B) test in the transition region characterizes mainly crack initiation conditions. Additionally, the inherent data scatter and curve fitting required to obtain Charpy V-notch parameters increase the uncertainty of estimating and correlating the transition temperatures. Correlations between the Charpy V-notch temperatures T28 J and T41 J versus T0 are presented in Sattari-Far and Wallin.11 The correlations, which are based on data from over 200 pressure vessel steels, are currently being reassessed in more detail, but in applications where the direct estimation of T0 is not possible, the correlations can be used as indicated below (s is standard deviation):
T0 ¼ T28J 19 C ðs ¼ 22 CÞ
½42
T0 ¼ T41J 26 C ðs ¼ 25 CÞ
½43
probably will be used in parallel into the foreseeable future. Once a sufficient amount of reference data have been measured using the Master Curve method, it will gain even further acceptance. Also, further understanding of the limits of applicability for different steels will be obtained. Over this transfer period, the correlations developed between the different methods should play a significant role, providing support for properly analyzing data and encouraging the use of the more advanced methods. Note that some correlations, like those proposed for estimating crack arrest toughness from Charpy V-notch tests have brought new applications for the instrumented Charpy test. The overall trend in fracture mechanics testing is toward characterization and methods which allow the use of small and/or moderate-size specimens simulating the true loading conditions and accounting for the expected micromechanisms of fracture. The Master Curve approach and its implementation in the ASTM E 1921 methodology have proven to be a valuable and powerful analysis tool for a wide variety of applications involving ferritic steels.
References 1. 2. 3.
4. 5.
4.14.6 Summary and Conclusions The Master Curve methodology has been described as an advanced, direct technique of determining the fracture toughness of ferritic structural steels. The application of the methodology has increased during the last decade and spread worldwide extending beyond the initial applications associated with NPP surveillance and integrity assessment programs. Today, the methodology is well known and increasingly accepted by safety authorities as a standardized method for application in safety assessments. Methods based on conventional approaches, such as the Charpy V-notch test, are still widely used, and
6. 7. 8.
9.
10. 11.
Griffith, A. A. Philos. Trans. Royal Soc. London 1921, 221, 163–198. ASTM E 1820 REV A. Standard Test Method for Measurement of Fracture Toughness; ASTM International, Dec 2008. ASTM E 1921-08. Test Method for Determination of Reference Temperature, T0, for Ferritic Steels in the Transition Range; Annual Book of ASTM standards, July 2008; Vol. 03.01. Wallin, K. J. ASTM Int. 2005, 2(4), 17–37. Wallin, K. The misconception(s) of ASTM E 399, plane-strain fracture toughness KIc. In Life Management and Maintenance for Power Plants, VTT Symposium 234; BALTICA VI, Helsinki/Stockholm, June 8–10, 2004; Veivo, J., Auerkari, P., Eds.; Vol. 2, pp 379–392. Wallin, K. Eng. Fract. Mech. 1985, 22(1), 149–163. Wallin, K. Eng. Fract. Mech. 1984, 19(6), 1085–1093. Wallin, K. Fracture toughness transition curve shape for ferritic structural steels. In Joint FEFG/ICF International Conference on Fracture of Engineering Materials and Structures, Singapore, Aug 6–8, 1991. Wallin, K. Master Curve analysis of ductile to brittle transition region fracture toughness round robin data. The ‘EURO’ fracture toughness curve; VTT Publications 367; Technical Research Centre of Finland: Espoo, 1998. McCabe, D. E.; Merkle, J. G.; Wallin, K. An introduction to the development and use of the master curve method; ASTM manual series, MNL 52; 2005, p 67. Sattari-Far, I.; Wallin, K. Application of Master Curve Methodology for Structural Integrity Assessments of Nuclear Components; SKI Report 2005:55; Swedish Nuclear Power Inspectorate: Stockholm, 2005; p 178.
Fracture Toughness Master Curve of bcc Steels 12. Wallin, K.; Saario, T.; To¨rro¨nen, K. Metal Sci. 1984, 18(1), 13–16. 13. Wallin, K.; Laukkanen, A. Eng. Fract. Mech. 2008, 75(11), 3367–3377. 14. Wallin, K. In Fracture: Theory and Applications; Kirk, M., Bakker, A., Eds.; American Society for Testing and Materials: Philadelphia, PA, 1995; pp 519–537, ASTM STP 1244. 15. Wallin, K.; Planman, T.; Valo, M.; Rintamaa, R. Eng. Fract. Mech. 2001, 68, 1265–1296. 16. SINTAP – Structural Integrity Assessment Procedures for European Industry. Project BE95–1426, Final Procedure, British Steel Report, Rotherham 1999. 17. Wallin, K.; Nevasmaa, P.; Laukkanen, A.; Planman, T. Eng. Fract. Mech. 2004, 71(16–17), 2329–2346. 18. Wallin, K. 1993. In Constraint Effects in Fracture; Hacket, E. M., Schwalbe, K.-H., Dodds, R. H., Jr., Eds.; ASTM: Philadelphia, PA, 1991; pp 264–288. 19. Wallin, K. Eng. Fract. Mech. 2001, 68(3), 303–328. 20. Wallin, K. Eng. Fract. Mech. 2010, 77(2), 285–292. 21. Wallin, K. Eng. Fract. Mech. 1989, 32, 523–531. 22. Wallin, K.; Planman, T., Eds. In Use and Applications of the Master Curve for Determining Fracture Toughness, Proceedings of the Workshop MASC 2002, Helsinki/ Stockholm, June 12–14, 2002; VTT Industrial Systems: Espoo. 23. Marston, T. U. Flaw Evaluation Procedures – Background and Application of ASME, Section XI, Appendix A, EPRI NP-719-SR: Electric Power Research Institute: Palo Alto, CA, 1978. 24. Kussmaul, K.; Fo¨hl, J.; Roos, E. Some conclusions with regard to the safety assessment of cracked components drawn from the research program Integrity of Components (FKS II) at the present state. In 12th MPA Seminar on Safety and Reliability of Pressure Components, Oct 9–10, 1986, Stuttgart. 25. Wallin, K. In Proceedings of the 17th European Conference on Fracture, Brno, Czech Republic, Sept 2–5, 2008; pp 138–148. 26. ASME Boiler and Pressure Vessel Code Case N-629, Use of Fracture Toughness Test Data to Establish Reference Temperature for Pressure Retaining Materials, Section XI, Division 1, 1999. 27. O’Dowd, N.; Shih, C. J. Mech. Phys. Solids 1992, 40(5), pp. 939–963. 28. ASME Boiler and Pressure Vessel Code Case N-631, Use of Fracture Toughness Test Data to Establish Reference Temperature for Pressure Retaining Materials Other Than Bolting for Class 1, Vessels Section III, Division 1, 1999. 29. Pellissier-Tanon, A.; Devaux, J. C.; Houssin, B. In 6th European Conference on Fracture (ECF 6): Fracture Control of Engineering Structures; Elst, H. C.,
30. 31.
32.
33.
34.
35. 36.
37. 38. 39. 40. 41.
42. 43. 44.
461
Bakker, A., Eds.; Engineering Materials Advisory Services: Warley, UK, 1986; Vol. II, pp 903–916. Wallin, K. Nucl. Eng. Des. 1997, 174(3), 205–218. Cheverton, R. D.; Ball, D. G.; Bolt, S. E.; Iskander, S. K.; Nanstad, R. K. Pressure vessel fracture studies pertaining to the PWR thermal-shock issue: Experiment TSE-7; NUREG:CR-4304 (ORNL-6177); Oak Ridge National Laboratory: Oak Ridge, TN, 1985, p 133. Sokolov, M.; Wallin, K.; McCabe, D. In Fatigue and Fracture Mechanics, 28th National Symposium of Fracture Mechanics, Saratoga Springs, New York, June 1996; Underwood, J. H., MacDonald, B. D., Mitchell, M. R., Eds.; ASTM: Philadelphia, PA, 1997; Vol. 28, ASTM STP 1321. International Atomic Energy Agency. Guidelines for Application of the Master Curve Approach to Reactor Pressure Vessel Integrity in Nuclear Power Plants; Technical Reports Series No. 429; IAEA: Vienna, 2005; p 105. Wallin, K. In Advances in Fracture Research, 9th International Conference on Fracture, Sydney, Australia, Apr 1–5, 1997; Pergamon: New York, 1997; Vol. 5, pp 2333–2344. ASME Boiler and Pressure Vessel Code, Section III, Appendix G, 2002. PVRC Ad Hoc Group on Toughness Requirements. PVRC Recommendations on Toughness Requirements for Ferritic Materials; Welding Research Council Bulletin No. 175; Aug 1972. ASME NB-2331. ASME Boiler and Pressure Vessel Code, Rules for Construction of Nuclear Power Plants, Division 1, Section III, SubSection NB, Class 1 Components, 2002. ASME Boiler and Pressure Vessel Code. Section XI, Appendix G, 2002. Yoon, K. K J. Press. Vess. Technol. 1995, 117, 378–382. Lott, R. G.; Kirk, M. T.; Kim, C. C. Master Curve Strategies for RPV Assessment; WCAP-15075; Nov 1998. Nuclear Regulatory Commission. Safety evaluation by the Office of Nuclear Reactor Regulation to include the use of a Master Curve-based methodology for reactor pressure vessel integrity assessment; Docket No. 50–305; Kewaunee Nuclear Power Plant, May 2001. Nuclear Regulatory Commission. Initial RTNDT of Linde 80 Weld Materials; Safety Evaluation for Topical Report BAW2308, Revision 1; 2005. Wallin, K. In Fatigue and Fracture Mechanics; Chona, R., Ed.; ASTM: Philadelphia, PA, 2001; Vol. 32, 17–34, ASTM STP 1406. Wallin, K. Effect of strain rate on the fracture toughness reference temperature T0 for ferritic steels. Recent Advances in Fracture, Orlando, FL, Feb 10–13, 1997; pp 171–182.
4.15
Ceramic Breeder Materials
J. G. van der Laan, A. V. Fedorov, and S. van Til Nuclear Research and Consultancy Group, Petten, The Netherlands
J. Reimann Karlsruhe Institute of Technology, Karlsruhe, Germany
ß 2012 Elsevier Ltd. All rights reserved.
4.15.1 4.15.1.1 4.15.1.2 4.15.2 4.15.2.1 4.15.2.2 4.15.2.3 4.15.2.4 4.15.2.5 4.15.3 4.15.3.1 4.15.3.2 4.15.3.2.1 4.15.3.2.2 4.15.4 4.15.4.1 4.15.4.2 4.15.4.3 4.15.4.3.1 4.15.4.3.2 4.15.4.3.3 4.15.4.3.4 4.15.4.4 4.15.4.4.1 4.15.4.4.2 4.15.4.5 4.15.4.5.1 4.15.4.5.2 4.15.4.6 4.15.5 4.15.5.1 4.15.5.1.1 4.15.5.1.2 4.15.6 4.15.6.1 4.15.7 4.15.8 4.15.8.1 4.15.8.2 4.15.8.3 4.15.8.4 4.15.8.5 4.15.8.6 4.15.8.7
Introduction Tritium Breeding Breeding Blankets Ceramic Breeder Blankets Pellets/Pins/Blocks Pebble-Bed Concepts Blanket Design Parameters Testing of Blanket Modules in ITER Ceramic Breeder Requirements Ceramic Breeder Fabrication Base Properties Fabrication of Shapes Pellets or blocks Pebbles Pebble and Pebble-Bed Thermomechanics Introduction Single Pebble Testing Properties of Pebble Beds Pebble-bed density and packing factor Mechanical behavior of pebble beds Thermal creep Cyclic loading Heat Transfer Properties Thermal conductivity Heat transfer Pebble-Bed Modeling Continuum models Discrete-element modeling In-Pile Behavior Tritium Production and Release Tritium Release Out-of-pile In-pile testing Irradiation Parameters Irradiation Damage Activation and Waste Issues Summary and Outlook Microstructure High Burnup High Fluence High Temperature Effects of Transients and Off-Normal Conditions Accident Behavior (Safety and Investment Integrity) Development of Tools
465 465 465 466 466 467 469 469 470 470 472 473 473 473 475 475 476 476 477 479 479 481 481 481 483 484 484 485 486 491 491 492 495 503 503 504 504 505 505 506 506 506 506 506 463
464
Ceramic Breeder Materials
4.15.8.8 4.15.8.9 References
Compatibility with Structure Waste Management and Reuse/Recycling
Abbreviations AECL BIT BOT BU CB CEA DEM DEMO DPC model D–T fusion EBR EOL EXOTIC FEM FFTF FPR FPY FW FZK HCCB HCPB HFR INTOR ITER
JAEA JAERI JRR KIT Li4SiO4 Li2TiO3 Li2ZrO3 MAPI NET NRG PIE
Atomic Energy of Canada Limited Breeder-in-tube Breeder-out-of-tube Lithium Burn-Up Ceramic breeder Commissariat a` l’Energie Atomique Discrete-element modeling DEMOnstration reactor for power generation Drucker Prager Cap model (from soil mechanics) Deuterium–tritium fusion Experimental breeder reactor (US) End of life Extraction of tritium in ceramics Finite-element modeling Fast Flux Test Facility Fusion power reactor Full power year First wall Forschungszentrum Karlsruhe, Germany, changed into KIT 2009 Helium-cooled ceramic breeder Helium-cooled pebble bed High Flux Reactor at Petten, the Netherlands Precursor of ITER project is Latin for ‘‘The Way’’ (before that, ITER was an acronym for International Thermonuclear Experimental Reactor) Japan Atomic Energy Agency (former JAERI) Japan Atomic Energy Research Institute (now JAEA) Japan Research Reactor Karlsruhe Institute of Technology, Germany, formerly FZK Lithium orthosilicate Lithium metatitanate Lithium metazirconate Mitsubishi Atomic Power Industry Next European Torus Nuclear Research & consultancy Group Postirradiation examinations
506 506 507
R&D SBZ SCW TBM TBR TPD UCT WCCB WCSB
Research & development Schlu¨nder Bauer Zehner model Super-critical water Test blanket module Tritium breeding ratio Temperature programmed desorption Uniaxial compression test Water-cooled ceramic breeder Water-cooled solid breeder
Symbols d d 1, d2 D E G h H I kpb ks Ø Pp R t T Tew TM Tw e ecr g rk rp rpb
rs s t
Diameter Pebble diameters in binary bed Pebble-bed diameter Elastic modulus Tritium production rate Heat transfer coefficient Pebble-bed height Tritium inventory Pebble-bed thermal conductivity Thermal conductivity of pebble material Diameter Total pebble porosity ¼ average fraction of porosity within a single pebble Tritium release rate Time Temperature Temperature extrapolated from bulk pebblebed conductivity Melting temperature Undisturbed wall temperature Strain Creep strain Packing factor ¼ ratio of pebble volume to pebble-bed volume Contact surface ratio Pebble density Pebble-bed density ¼ ratio of pebble-bed mass mpb to pebble-bed volume Vpb, also called ‘tap density’ or ‘apparent density’ Density of the pebble base material Stress Tritium residence time
Ceramic Breeder Materials 6
4.15.1 Introduction 4.15.1.1
465
Li density can be raised from the natural 7.42% to about any desired value.
Tritium Breeding
The fusion reaction of tritium and deuterium is considered one of the most suitable options for nearterm large-scale fusion power generation, through D þ T ! Heð3:56MeVÞ þ nð14:03MeVÞ 4
Deuterium is a hydrogen isotope with an abundance of 1 out of 6500 atoms in seawater, implying virtually boundless resources. Tritium is the next hydrogen isotope, and it is radioactive with a half-life of 12.3 years under emission of a b-particle; it cannot be obtained from natural resources. Therefore, the D–T fuel cycle requires the breeding of tritium from lithium using one of the following reactions: n þ 6 Li ! T þ 4 He þ 4:78MeV n þ 7 Li ! T þ 4 He 2:47MeV þ n The neutron supplied by the D–T fusion reaction shown above is also the one that provides useful energy. The reaction with 6Li is exothermic, providing a small energy gain; on the other hand, the reaction with 7Li is endothermic but does not consume the neutron, though a more thermalized neutron is released. Natural lithium contains 7.42% 6Li and 92.58% 7Li. In fact, lithium has been identified as the only viable element to breed tritium. 6Li has a very high cross-section to capture a neutron (see Figure 1), and through the use of isotope enrichment, the effective
4.15.1.2
Breeding Blankets
In the development of magnetic fusion power plants, the tritium breeding function is effectively integrated in the blanket, which also serves as the main thermal power conversion system and an effective shield for the adjacent reactor components from neutrons and g-rays. An integrated tritium breeding blanket acts as a shield, as a heat exchanger, and as a breeding zone of tritium fuel, as pictured in Figure 2. The blanket has a primary or first wall that faces the plasma, and this component is in direct contact with the edge of the fusion plasma. The latter is typically designed to remove the plasma radiative power and part of the nuclear heating, either with or without a cooling circuit separate from the blanket system. In order to obtain a closed D–T fuel cycle for a fusion power plant, it is mandatory that the tritium production rate is, at least effectively, equal to its consumption rate and accounts for decay and losses at scheduled or unscheduled plant outages; this principle is usually called ‘tritium self-sufficiency.’ These conditions will not be achieved in near-term fusion devices, where tritium resources available from fission plants can be used and where production from a so-called ‘driver’ blanket is an additional or alternative source for fuel supply. Effective tritium production requires that the lithium compounds are located in such a way that the Blanket
Shield
Vacuum vessel
103 6Li(n,α)t 7Li(n,α)t
X-section (barn)
102
Radiation 101
DT plasma 100
Neutrons
10–1
10–2 100
First wall 102
104 Energy (eV)
106
Figure 1 Tritium production nuclear cross-sections for Lithium-6 and Lithium-7 isotopes. Calculated with JANDL4.0.(http://www.oecd-nea.org/janis/).
Coolant
Magnets
Breeding zone Figure 2 Schematic of tritium breeding blanket configuration: the breeding zone is between the plasma facing wall and the neutron shield protecting the vessel and magnets.
466
Ceramic Breeder Materials
maximum capture of D–T neutrons is obtained in the so-called tritium breeding blanket. As most fusion devices require partial use of the plasmafacing area for plasma heating, plasma diagnostics, plasma control, and fuel exhaust, the effective capture of D–T neutrons for breeding tritium requires the use of neutron multipliers in the blanket. Net tritium breeding ratios (TBRs) foreseen for power plants should be about 1.05–1.1. The breeder material used in blanket designs that have attractive thermal efficiency for magnetic confinement power plants should conform to certain requirements. It should 1. breed tritium in a relatively small volume with a high production rate 2. release tritium in a manner that allows fast processing into plasma fueling 3. possess physical and chemical stability at high temperature 4. display compatibility with adjacent structures and other blanket components 5. exhibit adequate irradiation behavior 6. not pose specific safety risks under off-normal and accidental conditions 7. have activation characteristics allowing recycling or treatment as low-active waste. Lithium-based ceramics are recognized as attractive tritium breeding materials for the first generation of fusion power plants, due to their inherent thermal stability and chemical inertness. This chapter describes the development of ceramic breeder (CB) blankets and material production routes applied or investigated and summarizes the properties and R&D results for a number of lithium-based ceramic materials.1 In this chapter the chemical formulas are used, though the actual composition are very often non-stoichiometric, which is more evident when a larger fraction of lithium has been burned. Most of the work presented in this chapter is subject to rapid evolutions in local, national or international programs. The authors like to stress that any of the activities, e.g. those concerning ITER can be quite different in their evolution.
4.15.2 Ceramic Breeder Blankets When ceramic breeder research was initiated in the 1970s, relevant data for lithium-based ceramics were scarce or nonexistent. Initial screening of candidates was mainly based on examination of the physical and
chemical characteristics and neutronic behavior. An extensive R&D effort focused on determining the properties of unirradiated materials and on designing irradiation experiments to understand and quantify the effect of neutron irradiation on material properties and on recovery of generated tritium. With the publication of these data, the relative merits of the candidates became known and interest changed accordingly. With the evolution of the INternational TOkamak Reactor (INTOR) and International Thermonuclear Experimental Reactor (ITER) projects, much of the international R&D effort was focused on the opportunities for implementing breeding units or driver blankets in such a device and its role in the technological development of power reactor blanket systems. Several breeder blanket design options had been developed such as high-temperature water-cooled and helium-cooled concepts for DEMO and power reactors, and low-temperature water-cooled concepts for ITER.2–21 The ceramics under consideration exhibit different characteristics, which can make one ceramic more adaptable to a specific blanket concept. In this chapter, the issues being addressed in R&D in support of current blanket design studies are highlighted. In this chapter no reference is made to R&D for other fusion systems like inertial confinement. Two major types of ceramic breeder material configurations have been developed based on pressed and sintered pins or pellets or as a collection of packed spheres or pebble beds. The actual arrangement of pebbles though may be in tube, which may lead to some confusion when ‘breeder-in-tube’ (BIT) is mentioned. The paper by Ihli et al.21 provides an overview of blanket design developments and references ongoing work by various parties. Until 2010, pebble-bed concepts were the preferred options for all parties involved in ceramic breeder test blanket module (TBM) programs for ITER as reported by Giancarli and coworkers.23,24 Typically, inert gas with hydrogen addition, whose characteristics are discussed later in this chapter, is used for the extraction of the tritium produced from lithium ceramics.25–28 4.15.2.1
Pellets/Pins/Blocks
Early developments of ceramic breeder blanket concepts were based on fission reactor technologies, and the R&D activities concentrated on pellet or pin type specimens. An example is the European BIT project, using machined annular pellets or pins.2–5 Figure 3 shows a schematic view of the BIT concept, based on stacking annular pellets of
Ceramic Breeder Materials
Purge gas
Central dummy rod
Spacer
Purge gas channel
467
Be/Li4SiO4 pebble bed breeder zone
Inlet
Coolant systems Outlet Inlet
Breeder ceramic annular pellet
Beryllium
Flow separator (baffle) Beryllium cladding
Pressure vessel
Figure 3 Breeder inside tube (BIT) concept based on lithium ceramic pellets. Reproduced from Dalle Donne, M.; Anzidei, L. Fusion Eng. Des. 1995, 27, 319–336.
LiAlO2, enriched up to 90% 6Li, operating at temperatures between 400 and 600 C.5 Li2ZrO3 was considered as an alternative material, which has been further developed at the Commissariat a` l’Energie Atomique (CEA).29 Other breeding concepts using machined shapes have been developed, for example, in the Russian Federation,30 and recently Sharafat et al.31 proposed shaping the breeder ceramic as foams. In this chapter, only a few R&D results with pellets or pins are mentioned. 4.15.2.2
Diffusion-welded first wall
Stiffening plate
Figure 4 Breeder-out-of-tube (BOT) with a mixed bed of Li4SiO4 and beryllium pebbles. Reproduced from Dalle Donne, M.; Anzidei, L. Fusion Eng. Des. 1995, 27, 319–336.
Pebble-Bed Concepts
In the 1970s, alternative fission fuel technologies had been developed based on packing of spheres to reduce the problems associated with excessive swelling and fragmentation of pellets.32 One of the early pebble-bed blanket designs was developed by Dalle Donne and coworkers6,7,9 at Forschungszentrum Karlsruhe (FZK), now called Karlsruhe Institute of Technology (KIT), Germany. In this concept, breeder ceramic and neutron multiplier were both shaped as small spheres or pebbles and arranged in a so-called mixed bed (see Figure 4). The concept was based on small (0.1–0.2 mm diameter) pebbles of Li4SiO4 and a binary mixture of beryllium pebbles (0.1–0.2 mm and 1.5–2.3 mm diameter) (see Figure 5), taken from the extraction of tritium in ceramics 7 (EXOTIC-7) irradiation project.33 It was found that the compatibility of Li4SiO4 and beryllium was drastically reduced under neutronirradiation conditions.34 This initiated the separation of breeder and neutron multiplier in different pebble beds in further blanket design evolution.
0
1
2
Figure 5 Mixed-bed test configuration of Li4SiO4 pebbles with small and large beryllium pebbles packed to high density used in EXOTIC-7 irradiation experiment. Reproduced from van der Laan, J. G.; Kwast, H.; Stijkel, M.; et al. J. Nucl. Mater. 1996, 233–237, 1446–1451.
Extensive R&D on breeder pebbles has also been performed by the Canadian Atomic Energy of Canada Limited (AECL), where in particular Li2ZrO3 and Li2TiO3 were developed.9,35 With growing insight into thermodynamics and the experimental results obtained from neutronirradiation testing, the European breeder out of tube (BOT) concept evolved into the helium-cooled pebble-bed (HCPB) concept,7 in view of preparing a test module program for ITER. This concept evolved further in Europe within the scope of the Power Plant Conceptual Study.15,16 The key features of this early HCPB concept are given in Figure 6. It consists of
468
Ceramic Breeder Materials
Fir
st
wa
g zone
Breedin
10 mm
ll
ed
bble b
C
ed bble b
40 mm
der pe
bree eramic
TS S L
m pe
Berylliu
ld
nifo
BZ
Ma
HT
Figure 6 The Helium Cooled Pebble Bed (HCPB) blanket concept from the European Power Plant Conceptual Study (PPCS), model B. Reproduced from EFDA, A. Conceptual Study of Commercial Fusion Power Plants; Final Report of the European Fusion Power Plant Conceptual Study (PPCS); Report EFDA-RP-RE-5.0; 2005.
To FW
From caps and grid
Ceramic breeder
Be
To caps and grid From FW
MF 2 MF 3 MF 4 Bypass
Manifold (MF) 1
Purge gas He inlet
He coolant Figure 7 Evolved helium-cooled pebble-bed concept. Reproduced from Poitevin. Y; et al. Fusion Eng. Des. 2010, 85, 2340–2347.
alternating beds of ceramic breeder and beryllium pebbles, perpendicular to the plasma-facing wall, between flat coolant plates of Eurofer-97, a so-called reduced activation steel based on conventional 9Cr steels.36,37 In this study, the blanket box is considered a consumable component, with (1) the maximum irradiation damage of primary wall structures set at 150 dpa (about 5 FPY (full power year)) and probably the limiting factor of the box lifetime; and (2) the burnup of the ceramic breeder and swelling of the beryllium neutron multiplier depending on the design.38 Further evolution of the HCPB line in Europe concentrated on the strategies for the ITER TBM, as explained by Poitevin and coworkers.17,39–41
The internal structure of the blanket is given in Figure 7. All structures contain dense patterns of cooling channels, with beds of Be and ceramic breeder in the form of near-spherical particles (Ø 0.25–0.63 mm for Li4SiO4, Ø 1 mm for alternative breeder Li2TiO3, and Ø 1 mm for beryllium) separated by cooled steel plates and bed heights sufficiently low (about 10 mm) to conduct heat to the cooling plates without exceeding material temperature limits. Tritium is removed from the pebble beds by a slow purge flow of helium at nearatmospheric pressure, with hydrogen (typically 0.1 vol%) and defined levels of other constituents such as H2O, and so on for optimized integral performance. Some alternative concepts were explored such as those using a 9-Cr steel variant with higher
Ceramic Breeder Materials
469
~5
00
m
He in
m pol. rad.
He turn 2
tor.
He out Beryllide
First wall
beryllide
Cooling panel
Module box (F82H)
Neutron multiplier pebble (Be, beryllide) Tritium breeder pebble (Li2TiO3) Figure 8 Schematic of water-cooled ceramic breeder concept developed in Japan. Reproduced from Akiba, M.; Enoeda, M.; Tsuru, D.; et al. Fusion Eng. Des. 2009, 84, 329–332.
temperature resistance, see Hermsmeyer et al.,18 or SiC-based composite structure.21 The Japanese designed water-cooled solid breeder blanket consists of two submodules (see Figure 8).20 One submodule consists of a module box, tritium breeding pebbles, neutron multiplier pebbles, and cooling panels. The tritium breeding pebbles and neutron multiplier pebbles are separated by the cooling panels, and the beds are oriented parallel to the plasma-facing wall. The module box and the cooling panels are made of reduced activation ferritic– martensitic steel, F82H.36,42 For the tritium breeding pebbles, Li2TiO3 is selected as the primary candidate material. These pebbles are about 0.2–2 mm in diameter, with either a monodisperse or binary size distribution. Among the pebble-bed concepts, there are also designs for a low-pressure water-cooled ITER driver blanket (see, e.g., Lorenzetto et al.10) and the ITER 1998 design document by Ioki and coworkers.11,13 Nardi et al.13 developed ideas for a driver blanket for the reduced size ITER-FEAT. Recent work reported by Ihli et al.21 also included mixed bed options for DEMO blankets, as shown in Figure 9.
Ceramics Tube with thermal insulation
Figure 9 Breeder inside tube concept with ceramic pebbles. Reproduced from Ihli, T.; Basu, T. K.; Giancarli, L. M.; et al. Fusion Eng. Des. 2008, 83, 912–919.
4.15.2.3
Blanket Design Parameters
Table 1 provides a non-exhaustive summary of ceramic breeder blanket designs, and key parameters, compiled from the references in this chapter. Most of the concepts have a ferritic–martensitic steel as the structural material. Typical values for the expected blanket neutron wall load are about 3 MW m2, and its lifetime is mostly considered to be limited by about 150 dpa for the structural material. For a DEMO reactor, which is not intended to be a power plant with high availability, typical lithium burnups are 11% (Li4SiO4) or 17% (Li2TiO3), and fast neutron damage in the reduced activation ferritic/ martensitic (RAFM) steels is about 70 dpa.47,48 A typical lifetime for a power reactor blanket is estimated to be of the order of 5 years, implying about ten replacements during a 60-year reactor life. The blanket structure should, therefore, be designed with low-activation materials, enabling it to be recycled typically in a period up to 100 years. Such activation requirements have led to Li4SiO4 and Li2TiO3 as the preferred ceramic breeder systems for the European HCPB concept. A case study for Li4SiO4 has been elaborated by Fischer and Tsige-Tamirat.49 4.15.2.4
Testing of Blanket Modules in ITER
Breeding blanket and associated systems in fusion power plants have to ensure tritium breeding selfsufficiency, show a sufficient power conversion efficiency, and withstand high neutron fluences.50 TBMs for ITER should be representative for such blanket modules (see Giancarli et al.23 and Chuyanov et al.24).
470
Ceramic Breeder Materials
Among the technical objectives of ITER, it is specifically stated that ‘‘ITER should test tritium breeding module concepts that would lead in a future reactor to tritium self-sufficiency and to the extraction of high-grade heat and electricity production.’’51 The main testing objectives shall be Validation of structural integrity theoretical predictions under combined and relevant thermal, mechanical, and electromagnetic loads Validation of tritium breeding predictions Validation of tritium recovery process efficiency and T-inventories in blanket materials Validation of thermal predictions for strongly heterogeneous breeding blanket concepts with volumetric heat sources Demonstration of the integral performance of the blanket systems. Table 1 also provides values for the ITER TBM loading parameters and tentative requirements for DEMO as a next step. It is seen that the neutron wall load in ITER is relatively small, which requires specific measures to make meaningful nuclear tests with TBMs. Four versions of a TBM are considered with specific objectives as follows: 1. H-phase and H–He-phase: focus on electromagnetic behavior; 2. D-phase: focus on thermal and neutronic behavior; 3. First D–T-phase: DEMO-relevant data acquisition on neutronics, tritium production and management, and thermomechanics; 4. Second D–T-phase: DEMO-relevant data acquisitionwith high duty, long pulses with an integral TBM. All ITER parties consider helium-cooled ceramic breeder (HCCB) blankets. This type of blanket requires beryllium as the neutron multiplier and ferritic–martensitic steel as the structural material. The ceramic breeder is a Li-based compound, either Li2TiO3 or Li4SiO4, and is used in pebble beds. A water-cooled ceramic breeder (WCCB) blanket is proposed for the Japanese TBM. ITER prepares three horizontal ports for TBM testing, and two TBMs can be installed in one port. Though TBMs will be installed and tested as part of the ITER activities, each TBM will be developed under the responsibility of the distinct ITER party. Four of the ITER parties, China, European Union (EU), Japan, and Russian Federation (RF), have made TBM design proposals with solid breeder materials, while the United States and Korea propose to test submodules integrated into one of them. All ceramic
breeder-based TBMs use pebble beds and ferritic– martensitic steel structures and He coolant at 8 MPa with inlet temperature of 300 C and outlet up to 500 C depending on the operating conditions. Only the Japanese party proposes a water-cooled concept in addition. Figure 10 shows the typical arrangement of a port cell in ITER to accommodate TBMs. 4.15.2.5
Ceramic Breeder Requirements
The development of any ceramic breeder blanket concept toward demonstration and realization of fusion power must ensure that the ceramic breeder material meets the following specific requirements: Though the ceramic breeder has no structural function in the blanket, the pebbles or pellets must withstand the stresses induced under reactor operating conditions (pressures, temperatures, temperature gradients and thermal shocks, irradiation-induced swelling, creep) without an excessive fragmentation, which might result in degradation of the heat transfer parameters and purge gas flow, up to end of life (EOL) peak burnup and displacement damage. Stability of the ceramic at the maximum operating temperature with regard to lithium transport (e.g., by evaporation or redistribution). Compatibility between the ceramic and the structural material in the reference purge gas conditions under neutron irradiation. Compatibility is one of the criteria defining the maximum interface temperature between ceramic breeder and the structural material. Sufficiently low tritium residence time to minimize the tritium inventory in blanket and auxiliary systems that determine source term in off-normal and accidental conditions. Activation as low as possible under neutron irradiation, including activation from impurities, so as to reduce the D–T fuel cycle back-end issues (includes the materials’ recycling aspects).
4.15.3 Ceramic Breeder Fabrication Because significant quantities of ceramics will be needed in the near future for the fabrication of ITER TBMs and for a potential ITER driver blanket, various efforts have been initiated to evaluate fabrication process development. One of the fabrication issues is the hygroscopic nature of several candidate
Table 1 Some design and loading parameters for a number of fusion reactor blanket concepts utilizing ceramic breeder materials. For comparison values for the EU ITER Test Blanket Modules are given also. In addition to the references mentioned in the table, reader can turn to references in the text like e.g. 16, 21, 24 and 38. Design
BIT
BOT
HCPB
WCSB
SSTR
A-SSTR2
PPCS-B (HCPB)
HCPB-TBM in ITER
References Ceramic breeder (CB)
EU 6,7,8 LiAlO2
EU 2,3,4,8 Li4SiO4
Japan 218 Li2TiO3
EU 16 Li4SiO4
EU 17–19 Li4SiO4
25
Japan 20 Li2TiO3 or other 90
Japan 217 Li2O
90
EU 15 Li4SiO4 or Li2TiO3 30% (Li4SiO4), 60% (Li2TiO3)
30-40
90
Li2TiO3
Li2TiO3
Pebble Bed H2O or SCW Be-pebbles
Pebble Bed He Be-pebbles
Pebble Bed He Be-pebbles
Pebble Bed He Be-pebbles
450
700
400
250
Li-6 enrichment
%
alternative ceramic breeder Shape Coolant Neutron multiplier
(Li2ZrO3; Li2TiO3) Pellet water Be-pellets
Pebble Bed He Be-pebbles
Pebble Bed He Be-pebbles
Li2TiO3
C
410
300
400
C
590
660
890
900
750
800
920
900
316 SS
MANET
Eurofer
F82H
SiCf/SiC
Eurofer
Eurofer
Surface heat flux Neutron wall loading Lifetime fluence
MW/m2 MW/m2 MWa/m2
0.5 2.4
1 3.2
F82H (RAFS) 1 3–5 >10
<1 50 10
0.4 / 0.5 (peak) 2.2 / 3.5(peak) 7–8
0.27/ 0.5 (peak) 0.78 0.1
Ceramic Breeder Materials
Minimum ceramic breeder operation temperature Maximum ceramic breeder operation temperature Structural material
Pebble Bed water Be-pebbles, or Be12Ti 400
471
472
Ceramic Breeder Materials
Bio-shield plug
Cryostat plug
Breeder concentric pipe
Vacuum vessel port extension
Transporter
FW
TBM frame and shield plug
Cryostat extension
Vacuum vessel plug
Drain pipe
Figure 10 View of a typical test blanket module port cell arrangement in ITER. Reproduced from http://www.iter.org/mach/ tritiumbreeding.
lithium ceramics. Sensitivity to moisture increases as the lithium oxide content increases and as the specific surface area increases. The research activity initially involved g-LiAlO2, Li2O, Li2SiO3, and Li2ZrO3; see, for example, Johnson et al.25 and Roux et al.52 Later work concerned Li4SiO4, Li8ZrO6, and Li2TiO3.26,33,53–55,185,189 Currently, most blanket concepts are based on Li4SiO4 or Li2TiO3, though recently, work on other systems such as Li3TaO456 and Li8PbO657 as well as composites of Li2TiO3 with Li2O or Li4TiO4 additives58 has been reported.57 Breeder development has also started in Korea and India.59–61 4.15.3.1
Base Properties
Table 2 summarizes lithium compounds considered as breeding material and lists some key properties. Table 2
Lithium oxide was favored in early blanket concepts, in particular because of its very high lithium density and good thermal conductivity. Its biggest disadvantage is its strong sensitivity to moisture and the evidence it shows of significant swelling under irradiation. The silicates studied most widely are Li2SiO3 and Li4SiO4. In practice, traces of Li2SiO5 or Li6Si2O7 can be found as well. A variety of lithium zirconates have also been studied, such as metazirconate Li2ZrO3 and Li8ZrO6 octazirconate by Roux and coworkers.26,33,35,52,53,62,63,156 Because of activation issues for Zr in the fusion blanket spectra, these compounds became less attractive. Zirconates have been shaped in pellets as well as pebbles, and their irradiation performance at high burnups was considered promising in terms of pellet integrity and tritium release characteristics.33
Overview of most relevant basic ceramic lithium compounds and their major characteristics
Material Lithium-
Composition
Li density (g cc1)
Theoretical density (g cc1)
Melting temp ( C)
Specific heat at 400 C (J gK1)
Thermal conductivity at 400 C (W m1 K1)
Linear expansion at 400 C, relat. to 25 C (%)
Oxide Octo-zirconate Ortho-silicate Ortho-tantalate Meta-titanate Hexa-zirconate Meta-silicate Meta-zirconate g-Aluminate Octo-plumbate
Li2O Li8ZrO6 Li4SiO4 Li3TaO4 Li2TiO3 Li6Zr2O7 Li2SiO3 Li2ZrO3 LiAlO2 Li8PbO6
0.93 0.69 0.54 0.46 0.44 0.43 0.39 0.38 0.27 0.66
2.01 3.01 2.39 5.87 3.55 3.43 2.53 4.15 2.55 4.24
1432 1295 1255 1679 1270 1535 1201 1600 1610
2.5 1.5 1.9
6 1.5 2.5
1 0.7 0.9
1.4 1.3 1.6 1 1.3
1.7 1.7 2.4 1.4 2.6
0.5 0.6 0.6 0.4 0.4
Ceramic Breeder Materials
Lithium titanates were initially introduced at AECL by Kopasz and coworkers,64,65 and early tritium release experiments have shown promising results. Li2TiO3 has been studied by Japanese and EU parties. Hoshino et al.66–70,179–181,190 found that not only oxygen-deficient but also lithium oxide-deficient defects form on changing the atmosphere from hydrogen to oxygen. Thus, the doubly nonstoichiometric composition, Li2–xTiO3–y, has been confirmed. Further, it has been shown by thermal diffusivity measurement that 95Li2TiO3 has a higher thermal conductivity than 100Li2TiO3.66 Lithium aluminates have been studied mostly in the form of g-LiAlO2. Data accumulated have been summarized by Billone and Grayhack.71 Due to its low lithium density and modest tritium release performance, the material has gradually begun to receive less consideration.62 More recently, Zhu et al.56 in China started investigations of lithium tantalates such as Li3TaO4, which has a reasonably high lithium density. Sedano and coworkers57 report on the development of lithium plumbates such as Li8PbO6, revisiting earlier work of Hayashi et al.72 4.15.3.2
Fabrication of Shapes
The first ceramic breeder blanket concepts used blocks or plates; later, pellet designs emerged. The larger the breeder component, the greater the threat to its integrity because of cracking or fragmentation due to thermal stresses and irradiation-induced swelling and embrittlement. This is the major reason for favoring pebble-bed concepts.
473
4.15.3.2.2 Pebbles
For TBR considerations, the density of the pebbles should be high and enable a dense packing to achieve a high lithium density. Further comments on pebble shapes are given in a later section. The presently used or developed processes are as follows: 1. A melting–spraying process was used at KIT (formerly FZK), in collaboration with Schott Glaswerke, for the production of 0.25–0.63 mm Li4SiO4 and Li4SiO4–SiO2 pebbles76 (see Figure 11). After annealing, spherical pebbles of 95–96% of theoretical density (TD) exhibiting satisfactory mechanical strength were obtained. Long-term annealing experiments on various candidates ceramic breeder materials were performed by Piazza et al.77 An alternative route avoiding use of carbonate and using hydroxide was developed by Knitter et al.,78 with slightly lower density. The reference composition is Li4SiO4 þ 2.5 wt SiO2, resulting in a two-phase Li4SiO4 þ Li6Si2O7 structure in as-melted condition, and Li4SiO4 þ Li2SiO3 after heat treatment (see Figure 12). A melting-dropping process was investigated by Tsuchiya et al.79 in collaboration with Mitsubishi to produce 1-mm Li2O spheres. 2. Sol–gel type processes were developed at Japan Atomic Energy Agency (JAEA), with Nuclear Fuel Industries, to produce 1 mm Li2O and 1.6 mm Li2TiO3 pebbles80 (see Figure 13). Similarly, Muis and coworkers81 at Energy Research Centre of The Netherlands (ECN) produced 0.5–1.0 mm Li2TiO3
5
4.15.3.2.1 Pellets or blocks
Pellet or block fabrication makes use of proven technologies in the ceramic industry. Pressing and sintering of ceramic powders is a widely used and cost-effective industrial process. Pellets and rectangular blocks can be manufactured up to some centimeters in size with excellent material homogeneity and controlled density. Thus, LiAlO2, Li2ZrO3, and Li2TiO3 pellets meeting dimensional, microstructural, and purity characteristics were produced by Pechiney in collaboration with CEA.52 Similar results were obtained by ENEA, SCK/CEN, UKAEASpringfields, and US laboratories.65,73–75 Kapychev et al.30 fabricated pellets of Li4SiO4, metasilicate (Li2SiO3), and aluminate (LiAlO2), with a diameter of about 10 mm and heights of 5, 10, and 14 mm.
1 2
3
4
(a)
6 (b)
Figure 11 Illustration of the melt drop and jet spraying processes developed for production of Li4SiO4 pebbles at KIT. Reproduced from Kolb, M. H. H.; Knitter, R.; Kaufmann, U.; Mundt, D. Fusion Eng. Des. 2011, doi: 10.1016/j.fusengdes.2011.01.104.
474
Ceramic Breeder Materials
(a)
100 mm
(b)
50 mm
Figure 12 Cross-section micrograph of pebble fabricated by a melt spray process featuring large domains of dendritically grown crystals composed of Li4SiO4 (light) and Li6Si2O7 (dark): (A) overview and (B) detail. Reproduced from Kolb, M. H. H.; Knitter, R.; Kaufmann, U.; Mundt, D. Fusion Eng. Des. 2011, doi: 10.1016/j. fusengdes.2011.01.104.
Fabrication process Li2TiO3 solvent (H2O2 etc.) (binder)
In air Gelation solvent
pebbles. In these cases, the pebble densities were <80% TD. Further work led to pebbles with Li2TiO3 þ 5% TiO2 composition82 (see Figure 14). Wu et al.83 started the development of a sol–gel type process for Li4SiO4 and achieved 75% of TD for 1.2 mm diameter pebbles. 3. A process consisting of extrusion, spheronization, and sintering has, for several years, been used by AECL to produce 1.2 mm LiAlO2, Li2ZrO3, and Li2TiO3 pebbles in collaboration with Ceramics Kingston.33 Material densities are in the 80–90% TD range. Similar process trials were made by Lulewicz and Roux53 at CEA, with Pechiney, to produce 1 mm Li2ZrO3 pebbles, and later work concerned Li2TiO384,193 (Figure 15). 4. An agglomeration-sintering process has been used by JAEA, in collaboration with Kawasaki Industries,
Fabrication parameters Selection of solvent
Dissolving (mixing)
Condition of dissolving Selection of solution concentration
Dropping
Selection of gelation solvent Temperature and time for aging
Atmosphere for drying Drying
Temperature and time for drying Heating Calcination sintering
Temperature and time for calcination Temperature and time for sintering
Figure 13 Manufacturing process for Li2TiO3 by sol–gel method at Japan Atomic Energy Agency. Reproduced from Tsuchiya, K.; Kawamura, H.; Uchida, M.; Casadio, S.; Alvani, C.; Ito, Y. Fusion Eng. Des. 2003, 69, 449–453.
ECN SEI 5.0 kV
⫻50 100 m WD15 mm
ECN SEI 5.0 kV
⫻500 10 m WD15 mm
Figure 14 Scanning electron micrographs of Li2TiO3 pebbles produced by a sol-gel route (see text).
Ceramic Breeder Materials
ECN SEI 15.0 kV
⫻50 100 m WD36 mm
ECN SEI 15.0 kV
475
⫻500 10 m WD36 mm
Figure 15 Scanning electron micrographs of Li2TiO3 pebbles produced by an extrusion-spheronization method (see text).
for producing 1 mm Li2O, Li4SiO4, and Li2ZrO3 pebbles. Pebble densities in the 90% TD range were obtained.85 This process has also been investigated at CEA for producing 1 mm Li2TiO3 pebbles. Pebble density of 90% TD and good mechanical strength were obtained.84 5. Zhu et al. developed a wet process for fabrication of Li3TaO4 ceramic pebbles. Typical pebble diameters are about 0.7–1.0 mm, and the density achieved is over 90% TD, with crush loads more than 40 N.56 X-ray diffraction (XRD) patterns showed 99% of b-Li3TaO4 and traces of LiTaO3. The necessity to recycle ceramic breeders after service imposes specific requirements on pebble manufacturing technologies. This reprocessing aspect may become a significant driver in fusion power economics on a longer term. See Sections 4.15.7 and 4.15.8.8 for further discussion.
4.15.4 Pebble and Pebble-Bed Thermomechanics 4.15.4.1
Introduction
One of the issues that need to be addressed is the thermal and mechanical behavior of constrained pebble beds under (cyclic) nuclear loading conditions experienced in a tritium breeding blanket. The way a pebble bed responds to a thermal load depends primarily on the thermal transport properties of the bed, such as the packing factor of the bed, and the thermal conductivities of the pebble material and the surrounding purge gas. In addition, due to differences in the thermal expansion coefficients between the pebble-bed material and its surrounding structure, stresses will be induced in both the pebble bed and the structural material, and the contact
areas between pebbles and pebbles and walls become an important parameter. Also, during longer term operation in the neutron-irradiation environment, swelling of the breeder material will generate stresses. Furthermore, the pebble-bed thermal transfer properties may deteriorate with irradiation dose and lithium burnup. These induced stresses may directly or indirectly affect the functional operation if the mechanical integrity of the blanket element is endangered or if heat or tritium removal is significantly deteriorated due to pebble fracture, sintering, or melting. Various creep phenomena will affect the actual evolution of stresses, like thermal creep and irradiation creep. In addition, chemical reactions in between pebbles and between structures and pebbles may be enhanced under high contact pressures and the compositional changes arising from long-term operation by lithium burnup and other transmutation reactions. When the (constrained) pebbles experience stresses, by either compression due to thermal expansion or irradiation-induced swelling, these stresses will be (partially) relieved through thermal creep, that is, irreversible deformation of the pebbles. Under high stresses and at high temperatures, typically 0.6–0.8 times the melting temperature TM, these relaxation processes include sintering of the material. This irreversible strain and deformation in the material can lead to embrittlement and, ultimately, to fragmentation of the pebbles. The fracture of pebbles results in an inhomogeneous pebble-bed density and contact area and would lead to an inhomogeneous temperature distribution that is hard to model or predict. The lack of predictability of the temperature distribution can be a major safety issue. A mechanically stable pebble bed with a high packing factor is desired.
476
Ceramic Breeder Materials
4.15.4.2
Single Pebble Testing
pebbles experience much larger forces than the average value. Thermal annealing results in a significant grain size growth in Li2TiO3 pebbles; a slight decrease in crush load during aging was found for the Li4SiO4 and Li2TiO3 pebbles.77 Neutron irradiation will have the strongest effect to reduce the pebble strength, as mentioned here.
The simplest way to determine the mechanical strength of pebbles is by performing crush tests with individual pebbles. These tests are very useful for optimization of pebble properties and for quality control during pebble production. In crush tests, pebbles are arranged between two plane plates. In a pebble bed, the pebbles in contact with the cavity wall experience similar conditions because there is only one contact at the corresponding pebble hemisphere. Pebbles that are in contact only with other pebbles have on an average six contacts86,87 and the contact forces are in general smaller than on the pebble-wall contact. The crush loads of 0.5-mm diameter Li4SiO4 and 1-mm diameter Li2TiO3 pebbles used for the European HCPB project are in the range of 4–5 and 35–50 N, respectively, as shown by Knitter et al.191 and Roux et al.192 (see Figure 16).90 The crush loads of the Li2TiO3 pebbles were independent of 6Li enrichment. Assuming coverage of the wall by pebbles of about 0.7, the lower values of 4 and 21 N for the 0.5 and 1.0 mm diameter pebbles correspond to pressures on the wall pebbles of about 12 and 15 MPa. These values are significantly above the assessed maximum pressures of about 6 MPa in the HCPB ceramic breeder pebble bed.91 However, there are several effects that decrease this margin:
Fortunately, there is one mechanism, as outlined below, that is expected to alleviate the problem considerably: thermal creep. A general remark to be made here is that care is to be taken with respect to statistics: postirradiation tests from the EXOTIC-7 high burnup experiment showed fragmented pebbles as well as intact pebbles maintaining average strength levels.33,93 This statistical aspect is to be addressed more rigorously in optimization of pebble-bed technologies serving predictable and reliable operation of breeding blankets. 4.15.4.3
Properties of Pebble Beds
As the ceramic breeder material is used in the form of pebble beds, the macroscopic properties of the granular material are of particular interest. Most of these properties (mechanical, thermal, etc.) cannot be deduced directly or precisely from the properties of the base material or the single pebble; therefore, dedicated experiments are required. In these experiments, particular attention is paid to reproduce representative conditions in terms of, for example, pebble bed typical dimensions, packing factor, reducing atmosphere, and so on, as envisaged for the blanket component.
In pebble beds, more generally, in granular materials, an inhomogeneous branching of forces occurs (see Jaeger et al.92), with the effect that some
65
10
Ti 1100 CEA CTI 13 B0
Li4SiO4 01/3-3 (ex carbonate) Li4SiO4 01/3-4 (ex hydroxide)
55 Crush load (N)
Crush load (N)
8
6
4
45
35
2 Initial Cond. material material
3
24
48
Annealing time (days)
96
25 Initial Cond. material material
3
24
48
Annealing time (days)
Figure 16 Typical results of single pebble crush tests as a function of annealing time. Published by courtesy of Dr. R. Knitter, KIT.
96
Ceramic Breeder Materials
This section mostly follows the approach developed for the European blanket program and the work Reimann and coworkers94–106,110,115 have performed from the late 1990s onwards. 4.15.4.3.1 Pebble-bed density and packing factor
The following definitions are used: rs ¼ density of the pebble base material rp ¼ pebble density Pp ¼ total pebble porosity ¼ average fraction of porosity within (macroscopic, single) pebble volume rpb ¼ pebble-bed density ¼ ratio of pebble-bed mass mpb to pebble-bed volume Vpb g ¼ packing factor ¼ ratio of pebble volume to pebble-bed volume The quantities are connected by rpb ¼ mpb =Vpb ¼ ð1 Pp Þrs g
½1
In literature, the term ‘fraction of theoretical density’ is also used. This term is identical to (1Pp). The pebble-bed density rpb is often named ‘apparent’ or ‘tap’ density. The pebble-bed density rpb is an important quantity for nuclear calculations as these require the specific density of lithium and other constituents as inputs. In pebble-bed engineering, the packing factor g (generally given in percentage) is the characteristic quantity. The packing factor is influenced by the following parameters: filling procedure, pebble shape, surface roughness diameter d, diameter distribution, and container (cavity) dimensions: height H, diameter D. Filling procedure: The selection of an optimum procedure is not trivial, especially for larger mock-ups: filling should be assisted by vibration in such a way that a particle flow within the cavity is prevented or, in case of too small vibration energies, friction forces between pebbles and walls can be overcome. Pebble shape and surface roughness: ceramic breeder pebble shapes are almost spherical, partly with indentations (melting, spraying processes) or egg-like (extrusion, spheronization, sintering processes). The pebbles originating from melting have a smoother surface than the others. The packing factor is influenced by the surface roughness but negligibly by the pebble shape. Diameter distribution: The diameter d is of influence with respect to the cavity dimensions, discussed later. Ideally, the packing factor increases with increasing diameter distribution; see, for example, McGeary.87
477
However, sedimentation during the filling is a critical issue. For blanket pebble beds, two groups are of interest: (1) Pebble beds with a relatively small d variation: for example, Li2TiO3 pebbles with a nominal diameter of 1 mm ranging from 0.8 to 1.2 mm and Li4SiO4 pebbles ranging from 0.25 to 0.63 mm. Production costs generally favor a wider spread, which is not detrimental in respect to the packing factor, at least for the diameter ranges cited earlier. Packing factors in the range of 63–66% are reached; (2) Binary beds: here, large pebbles, dl, and small pebbles, ds, are used, with ds 0.1dl. First, the large pebbles are vibrated in the cavity. Then the pebbles are fixed in their position, for example, by a sieve in the filling pipe, to avoid sedimentation during the following pouring in of the small pebbles, which fill the spaces between the large pebbles. Packing factors of about 82% were obtained for beryllium beds with dl 2 mm, ds 0.2 mm.94 Critical issues are (1) the stability of such pebble packing structure during thermal cycles in the blanket and (2) the buildup of large stresses due to irradiation-induced swelling causing pebble degradation, leading to fracture. Diameter d, bed height H, diameter D: For finite cavities, there exist two characteristic ratios: H/d and D/d. The reason why g depends on H/d and D/d is the existence of two different characteristic pebble packing structures as demonstrated by tomography experiments101,105,106: in the pebble bulk, there is no preferential direction of contacts between pebbles, whereas at the wall, a zone with a thickness of about 4d exists, where regular packing structures are observed and contact zones are no longer homogeneously distributed. Figure 17 shows tomography results106 for a cylindrical container (D ¼ 49 mm, H ¼ 50 mm) filled with spherical pebbles (d ¼ 2.3 mm). Distinct pebble layers close to the walls can be clearly seen in Figure 17(a) where the positions of the pebble centers are plotted in a horizontal plane (left) and a vertical plane (right). Figure 17(a) shows axial void fraction distributions. The degree of regularity is largest for the pebbles in contact with the wall and decreases with increasing wall distance. The regularity is most expressed for pebbles in contact with plane walls (D ¼ 1); here, large isles with a dense hexagonal pebble arrangement are observed; see Reimann et al.105 The average packing factor of a plane wall zone is about equal to the bulk packing factor if the bed height is sufficiently large for the two wall zones
478
Ceramic Breeder Materials
Horizontal position of sphere centers
Vertical position of sphere centers
0.025
0.00
ROI height (m)
y(m) from the ROI center
0.01
0.000
0.02
0.03
0.04
–0.025 –0.025
0.000
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0.000
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0.010
0.015
0.020
ROI radius (m)
50
1.0
45
0.9
E0: p = 0 MPa
40
0.8
E2: p = 9.4 MPa
35 30
E0: p = 0 MPa E2: p = 9.4 MPa
25 20 15 10
Radial void fraction (1)
Vertical distance (mm)
(a)
E4: p = 9.2 MPa
0.7 0.6 0.5 0.4 0.3 0.2
5 0
0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1.0
(b)
Axial void fraction (1)
0
2.5
5.0
7.5
10.0 12.5 15.0
17.5 20.0 22.5
25.0
Radial wall distance (mm)
Figure 17 Visualization of pebble-wall layers by tomography: (a) Positions of sphere centers in region of interest (ROI) and b) Void fraction distributions. Reproduced from Pieritz, R. A.; Reimann, J.; Ferrero, C. Adv. Eng. Mater. 2011,13, Nr.3, 145–155.
Best location for filling
Y X Z
Figure 18 Pebble-bed filling experiments relevant for ITER Test Blanket Modules. Reproduced from Abou-Sena, A.; Neuberger, H.; Ihli, T. Fusion Eng. Des. 2009, 84, 355–358.
to interfere. This is an important result because it demonstrates that the often cited statement that wall zones are characterized by reduced packing factors is not generally valid. However, with decreasing D, both the regularity of the packing and the packing factor decrease, as seen clearly in Figure 17(b). Corners in pebble-bed cavities and inserts, for example, thermocouples, generally decrease the packing factor. This is also a reason why in small experimental setups the achieved packing factor is often distinctively below the maximum value. Pebble beds in HCPB blanket concepts are typically ‘thin’ pebble beds, that is, one dimension; the bed height H, is small compared with the other dimensions. For the ceramic breeder pebble beds, H 10 mm; for beryllium, H 30 mm. For a height of H 10 mm and a pebble diameter of 1 mm, the maximum packing factor is already difficult to achieve.
Ceramic Breeder Materials
A disadvantage of a large pebble diameter is also the smaller heat transfer coefficient, as outlined below. The blanket modules will probably be filled with pebbles through small pipes, as schematically shown in Figure 18. For simple cubic cavities, a homogeneous filling was achieved only by carefully vibrating the container in different tilted positions.97,100 The filling of a TBM mock-up using glass balls for the beryllium pebble beds was investigated by Abou-Sena et al.,107 by varying the number and positions of the filling pipes. The best results were obtained by using a single filling pipe close to the box corner and, keeping the filling pipe at the highest position, tilting the box around two axes. Here, a packing factor of 63.6% was achieved, which is considered to be very close to the maximum obtainable value. In the HELICHETTA experiment (see Dell’Orco et al.108,109) where elongated pebble-bed cavities (10 100 480 mm3) were filled through a pipe at the top (largest dimension orientated vertically), packing factors of 60% and 62% were achieved for Li2TiO3 pebbles with diameters between 0.8 and 1.2 mm, and 0.6 and 0.8 mm, respectively. For Li4SiO4 pebbles with d between 0.25 and 0.63, the packing factor was more than 64%. 4.15.4.3.2 Mechanical behavior of pebble beds
Reimann et al.94,96,110 obtained typical load displacement data for the ceramic pebble beds introduced above. The typical setup used is given in Figure 19.110 The pebble beds, contained in cylindrical cavities, are compressed by a piston and the pressure (equal to the uniaxial stress s) and the bed deformation (strain e) are measured, and the modulus E of the pebble bed is derived. These uniaxial compression tests (UCTs) should not be performed with bed height H to diameter D ratios larger than 1 so as to avoid a significant influence of friction effects at the cylindrical wall. Figure 20110 shows a typical result in terms of stress and strain. Key features of the pebble-bed deformation under monotonous isothermal mechanical loading are as follows: Nonlinear elasticity: the pebble-bed stiffens with higher degrees of deformation. Irreversible deformation is observed after initial unloading due to pebble relocation and plastic deformation at pebble–pebble and pebble–wall contacts (which disappears after multiple loadingunloading cycles).
479
Displacement transmitter (in total 4) Furnace
Piston
Al2O3-disc
TE
Pebble bed Container
Figure 19 Set-up for Uniaxial Compression Testing of pebble-beds. Reproduced from Reimann, J.; Harsch, H. In CBBI-12, 12th International Workshop on Ceramic Breeder Blanket Interactions, Karlsruhe, Germany, 2004; FZKA 7078.
Thermal creep at constant load with further irreversible deformation. Further stiffening observed at loading and unloading cycles. In addition, friction forces may provide some hysteresis at any loading-unloading cycle. 4.15.4.3.3 Thermal creep
Characteristic differences of thermal creep effects between homogeneous materials and pebble beds are: In pebble beds, the contact surfaces between the pebbles increase with time. Blanket relevant creep time periods of greatest interest are in the order of less than a day because stress relaxation effects are expected to be quite fast.104 Pebble beds are loaded first with relatively small stress gradients ds/dt; therefore, it is not differentiated between instantaneous plastic deformations and conventional thermal creep effects (the thermal creep correlations given below include both effects). Thermal creep strains are measured by UCTs by keeping the uniaxial stress s constant at a given temperature T. Figure 21110 shows creep strains for Li4SiO4 pebble beds for different values of s and T; the thermal creep strain occurring during the stress increase period has been taken into account in this representation.
480
Ceramic Breeder Materials
5.0
5 Li2TiO3-M 1st cycle
4 Uniaxial stress s (MPa)
Uniaxial stress s (MPa)
4.0 (b) (b)
3.0
(a)
2.0 (a)
Li4SiO4 ex hydroxide T(°C) 25 850
1.0
2nd cycle 3rd cycle
3
2
1
0
0.0 0.0 (a)
1.0
2.0 3.0 Uniaxial strain e (%)
4.0
0
5.0
0.2
0.4
0.6
0.8
1
Uniaxial strain e (%)
(b)
Figure 20 Example of uniaxial compression test results for Li4SiO4 pebble beds. (a) Li4SiO4 pebble beds; (b) Li2TiO3 pebble beds. Reproduced from Reimann, J.; Wo¨rner, G. In Proceedings of the 12th International Workshop on Ceramic Breeder Blanket Interactions, Karlsruhe, Germany, FZKA; 2005; p 7078.
5 FZK-silicate C1/s0.85 = 12.12exp(-10 220/T(K)) p = 8.6 MPa p = 4.3 MPa p = 2. 2 MPa p = 6.4 MPa
0.001 s = 2 MPa 750 ⬚C: Li4SiO4 ex hyd 900 ⬚C: Li4SiO4 ex hydr 750 ⬚C: Li2TiO3 900 ⬚C: Li2TiO3 750 ⬚C: Li4SiO4 ex silic 900 ⬚C: Li4SiO4 ex silic
3
2
C/s 0
Creep strain (%)
4
CEA-titanate C1/s0.85 = 0.67exp(-7576/T(K))
1
JAERI-titanate C1/s0.85 = 0.37exp(-6947/T(K))
0.0001 0.88
0 0
1000
2000
3000
4000
5000
Time (min) Figure 21 Thermal creep strain as a function of time for Li4SiO4 & Li2TiO3 pebble-beds using UCT. Reproduced from Reimann, J.; Wo¨rner, G. In Proceedings of the 12th International Workshop on Ceramic Breeder Blanket Interactions, Karlsruhe, Germany, FZKA; 2005; p 7078.
The data are well fitted by a constant exponent n for the time dependence. A constant exponent n was also found for most of the Li2TiO3 pebble beds.95 For some batches, during the first several hours, the same exponent n was observed; however, there was a subsequent increase in the creep rates. The Li2TiO3 batches that showed this behavior were in general characterized by low sintering temperatures, large
0.92
0.96
1
1.04
1.08
1.12
1.16
1000/T(K) Figure 22 Pebble-bed creep: temperature dependence for a constant value of the stress exponent. Reproduced from Reimann, J.; Wo¨rner, G. In Proceedings of the 12th International Workshop on Ceramic Breeder Blanket Interactions, Karlsruhe, Germany, FZKA; 2005; p 7078.
porosities, and small grain sizes, (for details, see Reimann et al.98). Eventually, impurities could play a role, too. The data are fairly well plotted in Arrhenius graphs by straight curves, as shown in Figure 22.95 For a selection of ceramic breeder material candidates, the influence of pebble-bed thermal conductivity as a function of, for example, packing factor and pebble-bed compression was studied by
Ceramic Breeder Materials
means of UCTs. These were performed for the ceramics Li4SiO4, Li2TiO3, Li2ZrO3, and Li2O with temperatures up to 480 C and pressures up to 8 MPa, with packing factors varying between 56% and 63.5%. Creep strains in the pebble beds are identified to be functions of temperature, stress, and time and are found to be of the form ecr ¼ AðT Þsm t n where m and n are independent of temperature and need to be deduced experimentally.97 This research was extended for Li4SiO4 pebble beds with temperatures up to 850 C and pressures up to 9 MPa and concludes that thermal creep effects are negligible at temperatures below 600 C.96,97 Creep behavior is also determined by the pebble properties: as mentioned earlier, lower creep strains were found for Li2TiO3 with small grain sizes (5 mm) and high sintering temperatures.98 4.15.4.3.4 Cyclic loading
An ITER TBM will experience not a constant heat load but a cyclic heat load behavior due to burn pulses of the plasma. Dalle Donne et al.111 performed exploratory thermal cycling experiments with densified Li4SiO4 pebble beds in a horizontal tube (D ¼ 20 mm, L ¼ 110 mm). After about 500 cycles (350–600 C), fractions of 1.8% and 3.5 wt% of broken pebbles in two measurement campaigns were found. The pressure drop in the helium purge flow was found to saturate after some hundred cycles, indicating that no further pebble fragmentation occurred. Thermal shock is not considered an issue for ceramic breeder pebbles; the experiments showed that pebbles only fail for dT/dt >60 C s1, that is, values much larger than expected in the blanket. Cycles of UCTs up to 4 MPa at ambient temperatures lead to negligible plastic deformations in the Li4SiO4 and Li2TiO3 pebble beds. The residual compaction is likely to be due to pebble relocation or pebble cracking. Multiple 6 MPa compression cycles at 750 C led to irreversible strains of 3.5% and 4% for Li2TiO3 and Li4SiO4 pebble beds, respectively. When 4 MPa compression cycles are performed on an Li4SiO4 pebble bed at 840 C, the creep strains reach values of 6%. In these last tests, the influence of thermal creep is clearly visible.110 Similar cyclic compression tests on Japanese Li2TiO3 pebble beds reveal a compression ‘equilibrium’ of 1.5% only after several cycles of 10 MPa at 400 C.112 Performing these tests at 600 C gives lower strain values in the pebble bed, likely due to the increased compaction.113 This is in agreement with the above-mentioned results.110
481
Several HCPB mock-up experiments have been performed in the HE-FUS3 facility at ENEA, Brasimone by Dell’Orco et al.108,109,114 In the HELICHETTA and HELICA experiments, the thermomechanical behavior of Li4SiO4 or Li2TiO3 pebble beds under thermal cycling loads was investigated. The following larger HEXCALIBUR type experiments were characterized by two beryllium pebble beds and one ceramic breeder pebble bed in between. Nuclear heating was simulated by using electrically heated plates within the pebble beds. In the HELICA mock-up tests, the thermal cycling of the pebble beds showed a saturation of the pebble-bed strains after 30 cycles. The bed heights were then reduced by 4% and 1.5% for the Li4SiO4 and Li2TiO3 pebble beds, respectively. For constant power levels, the system did not change thermally with increasing cycle number. During demounting, submicron pulverized material was released, especially from the Li4SiO4 pebble beds, and pebble fragmentation was observed. An important objective of these experiments was the validation of pebble-bed thermomechanical models as outlined below. 4.15.4.4
Heat Transfer Properties
4.15.4.4.1 Thermal conductivity
There are several advanced models to describe the thermal conductivity of pebble beds, which take into account the relevant parameters such as the thermal conductivity of the pebble material ks as a function of temperature T, the thermal conductivities of the surrounding gas kg as a function of temperature T and pressure p, the pebble diameter d, packing factor g, contact surface ratio rk2, and several other secondorder effect parameters. In the following section the Schlu¨nder Bauer Zehner model (SBZ model)114 is used to demonstrate the influence of some parameters. Figure 23 shows ks as a function of temperature for both Li4SiO4 and Li2TiO3 and kg for helium and other gases at 0.1 MPa.99 For most ceramic breeder materials, ks decreases first with increasing T, reaches at high temperatures a plateau, or increases again slightly; for details, see Abou-Sena et al.116 The helium conductivity increases strongly with increasing T. Figure 24(a) and 24(b) from Reimann et al.99 shows the influence of T and rk2 on the pebble-bed conductivity k of an Li4SiO4 pebble bed with mean diameter of d ¼ 0.4 mm and g ¼ 64%. For a noncompressed bed, rk2 ¼ 0, k increases moderately with T because kg increases with T. For a moderately compressed bed,
Ceramic Breeder Materials
0.5
3.5 0.4 3 2.5
0.3
2 1.5 1
Orthosilicate Metatitanate Helium
0.2
Air Argon
0.1
0.5 0 0
200
400 600 800 Temperature T(⬚C)
7
0 1000
Figure 23 Thermal conductivity of ceramic breeder pebble beds in helium and air. Reproduced from Reimann, J.; Hermsmeyer, S. Fusion Eng. Des. 2002, 61/62, 345–351.
0.3 0.3 Orthosilicate pebble bed T = 25 ⬚C
6
Uniaxial stress (MPa)
4
Gas conductivities (W m–1 K-1)
Breeder material conductivities (W m–1 K-1)
482
5 0.29 4 k = 0.29 W mK-1 3 0.28 0.28 0.27 0.27 0.26 0.26 0.24
2 1
0.4
0
0.8 1.2 Uniaxial strain (%)
(a)
rk2 ¼ 0,
0.64 Orthosilicate pebble bed T = 800 ⬚C
6
Uniaxial stress (MPa)
k is larger than that for but the temperature dependence is quite small. The thermal conductivity of noncompressed Li4SiO4 and Li2TiO3 pebble beds was measured by several authors. Figure 25 from Enoeda et al.86 shows that there is a good agreement among different authors. The Li4SiO4 pebble-bed data are best fitted with a correlation established by Dalle Donne204; the Li2TiO3 data were well predicted by the SBZ model using a value rk2 ¼ 0.0049.117 Li2TiO3 pebble-bed data for 2 mm pebbles from Abou-Sena116 are characterized by the tendency of a decrease in k with increasing T. Results for noncompressed beds including further ceramic breeder pebble materials were summarized by Abou-Sena et al.116: k should increase with increasing ks in the sequence Li2ZrO3, Li4SiO4, Li2TiO3, Li2O for equal values of the other parameters. Because the other parameters differed, this tendency was masked. For compressed pebble beds, the SBZ model contains the parameter rk2, which is a priori not known. If the pebble-bed strain is measured, it is easier to use this quantity as the relevant parameter. Figures 26(a) and 26(b) from Reimann and Hermsmeyer99 show results for Li4SiO4 and Li2TiO3 pebble beds, different gas conditions, and strain values e. For noncompressed beds, rk2 ¼ 0, the measured data agree fairly well with SBZ model predictions. With increasing strain, k increases; however, only very moderately compared with beryllium pebble bed. Even for a strain of about 4%, obtained in air at 800 C because
0.26
0
7
rk2 ¼ 0.02,
0.26
0.63
1.6
2
0.65 0.67
Stress = constant for 1500 min
5 k = 0.60 W mK-1 4 3 0.59 2 0.59
1
0.55 0 0 (b)
1
2
3
4
Uniaxial strain (%)
Figure 24 Stress–strain dependence of Li4SiO4 pebble beds at (a) 25 and (b) 800 C and evolution of thermal conductivity. Reproduced from Reimann, J.; Hermsmeyer, S. Fusion Eng. Des. 2002, 61/62, 345–351.
of significant thermal creep, k is only increased by about 20% for both types of pebble beds. For helium atmosphere, this difference becomes even smaller. Figure 26(b) also contains some results for binary Japanese Li2TiO3 pebble beds (0.2 and 2 mm pebbles, g ¼ 81.5%) in air atmosphere at ambient temperature. Compared with the monodisperse pebble bed with d ¼ 2 mm pebbles and g ¼ 64.3%, k is increased by a factor of 2. For blanket relevant conditions, this factor reduces to 1.3 for T ¼ 600 C and helium atmosphere.
483
2 FZK Li4SiO4+SiO2 0.25–0.63 mm 60.13%PF, hot wire method FZK Li4SiO4 0.5 mm 52%PF by M.D.D. & G.Sondon, fusion tech. volt 7 697(1990) Dr. Plazza FZK Li4SiO4+TeO2 66%PF SZB mod HM mod
Thermal conductivity (W mK-1)
Thermal conductivity (W mK-1)
Ceramic Breeder Materials
1
He pressure = 101-103 kPa Contact area fraction = 1. Ce-0 Accommodation coeff. = 0.2
0 0
200
400
600
800
1000
4 Li2O data Contact area fraction = 4.9 × 10−3 CEA Li2TiO3 60.8%PF Accommodation coeff. = 0.2 CEA Li2TiO3 50.4%PF CEA Li2ZrO3 54.42%PF SZB mod HM mod
3
2
1
0 0
200
Temperature (⬚C)
400
600
800
1000
Temperature (⬚C)
Figure 25 Comparison of pebble bed thermal conductivity data after Enoeda, M.; Ohara, Y.; Roux, N.; Ying, A.; Malang, S. In Proceedings of the CBBI-8, Colorado Springs, CO, Oct 6–8, 1999.
1.1
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1
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Gas T(⬚C) Helium; 475
0.4 0.3
Helium; 25 Air; 750 Air; 25 Air; 800
0.2 0.1 1
2
3
Strain e(%)
Type Gas Ti-J Air Ti-D Air Argon Ti-D Ti-E Air Air Ti-D Helium Ti-D Ti-J-bin Air Ti-D Air
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Figure 26 Evolution of thermal conductivity of compressed Li4SiO4 and Li2TiO3 pebble-beds as a function of imposed strain. Reproduced from Reimann, J.; Hermsmeyer, S. Fusion Eng. Des. 2002, 61/62, 345–351.
Experiments with compressed Li2TiO3 pebble beds (d ¼ 2 mm, g ¼ 65–67%) were also performed by Tanigawa et al.113 For a strain of about 1%, T ¼ 600 C, and helium at 0.1 MPa, k increased only by 3% compared with the noncompressed pebble bed. After annealing the pebble bed at 700 C without compression for 1 day, larger bed strains (factor 2) were obtained in the subsequent cycles and with this an additional increase in conductivity. 4.15.4.4.2 Heat transfer
Close to the wall, the pebble packing differs significantly from that in the bulk, as demonstrated in Figure 17. For noncompressed pebbles, the void fraction at the wall surface is close to 100%. Heat transfer characteristics in the wall zone are, therefore, different from those in the bulk. This fact is taken into account by using the heat transfer coefficient h,
which is defined with the temperature difference (TewTw), where Tew is obtained by extrapolating the bulk pebble-bed thermal conductivity k up to the wall, and Tw is the undisturbed wall temperature. In measurements, (TewTw) is very small and is sensitively dependent on extrapolated temperature profiles. Accurate measurements are, therefore, extremely difficult, and the discrepancies in experimental data are significant. The smaller the pebble diameter, the thinner the wall zone, and with this the difficulty to obtain accurate data increases. Again, different models exist to predict h exist, which have been validated in better suited experiments. Figure 27 shows h ¼ f (T ) for Li4SiO4 and Li2TiO3 pebble beds for noncompressed beds calculated with the model from Yagi and Kunii118: the strong increase of h with decreasing pebble diameter d is obvious. With progressing compression, the wall
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Figure 27 Heat transfer coefficients h from Yagi and Kunii model for different Li4SiO4 and Li2TiO3 pebble diameters.
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4.15.4.5
Pebble-Bed Modeling
There are different types of models to describe the behavior of pebble beds (see, e.g., Reimann et al.97): finite-element models based on continuum mechanics, also called finite-element modeling (FEM), and the so-called discrete-element models, DEM, for description of mechanics at micromechanics level (i.e., individual pebble–pebble interactions). The development of computational tools at these two different length scales (macro- and microscales) allows for a better description of the thermomechanics of the pebble beds.
p-hor (MPa) 0.4 0.12
0.08 0.06
Figure 28 Bi-axial pebble-bed deformation tests: experimental results and calculations with ABAQUS code system. Reproduced from Hermsmeyer, S.; Reimann, J. Fusion Eng. Des., 2002, 61–61, 367–373.
200 SCATOLA calculations
Plates axial displacement (mm)
contact surfaces increase and with this h. Again, this increase is expected to be much smaller than that in beryllium pebble beds because of the small thermal conductivity of ceramic breeder materials. At present, mechanistic models are developed taking into account the pebble arrangements close to the wall as determined by tomography101,105,106,199 or by discrete-element modeling (DEM), outlined later. A typical example is the work of Gan et al.133
for 3.0 mm plates
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Figure 29 SCATOLA bench-mark calculation with a continuum model.127 Reproduced from van der Laan, J.; et al. Proceedings of the 8th International Workshop on Ceramic Breeder Blanket Interactions (CBBI-8), Colorado Springs, CO, 1999; Ying, A., Ed.; UCLA Report.
4.15.4.5.1 Continuum models
The macroscopic behavior of pebble beds is described by constitutive equations commonly used in soil mechanics, considering the granular material as a continuum material that can undergo reversible elastic deformations, inelastic volume compaction (consolidation), and pressure-dependent shear failure. To account for these properties, different models have been developed, which are implemented in structural computational programs.91,115,117,120–125 Pebble-bed data (see sections above) had to be implemented, and user-defined subroutines had to be written, for example, for the thermal creep laws.115
The codes were first validated with fairly simple experiments.104,123,126,127 Figures 28 and 29 show examples of results from the biaxial experiment126 and the SCATOLA experiment for calculational benchmarking127 (Figure 29). These codes were later validated with more complex mock-up experiments such as HELICA and HEXCALIBUR91,108,114,128,129 and, as outlined below, aimed to be validated by the pebble bed assembly (PBA) experiment.130 Figures 30(a)117 and 30(b)91 show a comparison between measured and calculated HELICA temperatures.
Ceramic Breeder Materials
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Figure 30 (a) Measured temperatures in the HELICA experiment, reprinted from Dell’Orco, G.; di Maio, P. A.; Giamusso, R.; Tincani, A.; Vella, H. Fusion Eng. Des., 2007, 82, 2366–2374. (b) Calculated temperatures for the HELICA experiment, reprinted with courtesy from Gan, Y. Ph.D. Thesis, Universita¨t Karlsruhe, 2008.
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Figure 31 Stress–strain behavior of granular materials in (a) a rectangular box under uniaxial compaction and (b) packing density effect. Reproduced from An, Z.; Ying, A.; Abdou, M. Fusion Eng. Des. 2007, 82, 2233–2238.
The final goal of these codes is to determine the thermomechanical behavior of pebble beds in TBMs for ITER and DEMO blanket modules. At present, the codes are set up for this task. Up to now, only small portions of a DEMO blanket have been modeled. In these calculations, maximum stresses (very localized) of 6 and 2 MPa were obtained for the ceramic breeder and beryllium pebble beds, respectively. Because of thermal creep, these values were reduced after 2 h to 75% and 25%, respectively, of the initial value. 4.15.4.5.2 Discrete-element modeling
DEM is used to study the interparticle force distribution and translate the microscopic information into macroscopic information such as stress–strain response. This method is important for estimating the overall properties of pebble beds, such as yield strength and crush probability. The use of DEM for
fusion reactor blanket analyses was started at the University of California, Los Angeles (UCLA)124,131,205 and has been continued by KIT.91,119,132,133 A standard experiment, used by both groups, to validate DEM is UCTs; see corresponding section above. Figure 31 from An et al.205 shows three cycles of a pebble bed with an initial packing factor of 60.3%. With increasing cycle number, the pebble bed becomes stiffer. The strong dependence of the stress–strain curves on the packing factor was also stated by Gan and Kamlah.119 DEM analyses were also compared with the SCATOLA experiments.124 Some features were well described (see Figure 32), but thermal creep was not satisfactorily predicted. With increasing load on the pebble bed, the number of contacts between pebbles, Nc, increases, as shown in Figure 33.91 An important quantity to assess the fraction of crushed pebbles during blanket operation is the
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Experimental data (increasing T ) Experimental data (decreasing T ) Numerical estimations (increasing T ) Numerical estimations (decreasing T ) Numerical estimations (increasing T ) Numerical estimations (decreasing T )
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Figure 32 Calculational results for SCATOLA experiments. Reproduced from Ying, A.; Huang, H.; Abdou, M.; Lu, Z. In: Proceedings of the 9th International Workshop on Ceramic Breeder Blanket Interactions, Toki, Japan, Oct. 27–29, 2000.
Coordination number
6.6 6.5 nc = 5.80 ⫹ 0.490 p0.288
6.3 6.2 H-1 H-2 H-3 M-1 M-2 M-3
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Figure 34 The probability of maximum contact forces between pebbles in a pebble-bed increases with load. Reproduced from An, Z.; Ying A.; Abdou, M. Fusion Eng. Des. 2007, 82, 2233–2238.
Although the present DEM codes have proven to be very helpful for the understanding of the interaction between pebbles, there is still considerable development work required until quantitative results for small thermally loaded pebble-bed geometries can be expected. The method will certainly not be applicable for large components in the near future because of computational costs, but the improved understanding of the micromechanism will be beneficial for the improvement of the continuum codes.
6.7
6.4
1
4.15.4.6
5.8
In-Pile Behavior
8
Ceramic breeder materials have been tested under various conditions to determine their irradiation response in terms of tritium production and release and their microstructural, thermal, chemical, and mechanical stability. The nuclear loading of ceramic breeder affects their performance in various ways, of which the most important are the following:
evaluation of the probability distribution of the maximum contact forces for different loading conditions (see Figure 34).205 A comparison of results obtained by DEM and tomography is shown in Figure 35133: the packing factor distributions agree quite well, and the radial and vertical positions of particles show the same structure as shown in Figure 17(b). Other approaches are found like those of Aquaro and coworkers.123,199,202,203
Lithium burnup and other neutron transmutation reactions gradually change the material composition and affect the chemical and physicochemical interactions. Transmutation reactions of major and minor constituents gradually change the radioactivity levels relevant for in-tokamak operation, hot-cell operations, and management of waste, including recycling options.
0
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2 3 4 5 6 Hydrostatic pressure (MPa)
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Figure 33 The number of contacts between pebbles in a pebble-bed increases with load, reproduced with courtesy from Gan, Y. Ph.D. Thesis, Universita¨t Karlsruhe, 2008.
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Figure 35 Comparison of discrete-element modeling calculation with tomography results. Reproduced from Gan, Y.; Kamlah, M.; Reimann, J. Fusion Eng. Des. 2010, 85, 1782–1787.
Atomic displacements are induced by neutron impact, most effectively by fast neutrons, causing significant lattice defects, damage, swelling, and so on. Generated tritium affects the physicochemical and mechanical behavior, and the inert, 3He decay product stays within the material if it is not desorbed earlier into the purge gas flow. Generated helium resides in the material, and through formation of clusters of bubbles will generate stresses leading to macroscopic effects such as swelling, altered thermal transport, and/or fracture. Numerous irradiation experiments have been performed on ceramic breeder pebbles using material test reactors with thermal, mixed, or irradiation phenomena fast neutron spectrum. As the 6Li crosssection, in particular, is much higher for thermal neutrons, many of the irradiation phenomena can be effectively studied in thermal or mixed-spectrum reactors, that is, without using 14 MeV neutrons. This
section concentrates on the thermal–mechanical behavior, while tritium production and transport are dealt with in the following section. High fluence and high lithium burnup were achieved in fast reactor irradiations at experimental breeder reactor II (EBR-II) and fast flux test facility (FFTF) facilities, as reported by Hollenberg and coworkers.28,134 The bulk of these experiments concerned Li2O, LiAlO2, and LiZrO2, with only a few data on Li4SiO4 shaped as annular pellets and LiZrO3 pebbles. Good tritium release behavior of Li2O and LiZrO3 has been reported, even for temperatures higher than 1000 C; pellet thermal conductivity of Li2O and LiAlO2 was decreased at lower irradiation temperatures but appeared fairly unaffected when operated over 400 C.28,134,163,165,166,186–188 AECL tests at NRU and JAEA tests at JMTR addressed the impact of neutron irradiation on pebble-bed properties, such as conductivity. In these cases, the constraints were modest: either higher
488
Ceramic Breeder Materials
As a major step in the preparation of the European HCPB TBM program in ITER, an in-pile test of pebble-bed assemblies was defined. This experiment was designed to address the neutron-irradiation effects on the thermal–mechanical behavior of a breeder pebble bed at HCPB DEMO representative levels of temperature and defined thermal–mechanical loads.121,130,141 A schematic is given in Figure 37. The core of each test element is a horizontal cylindrical bed of ceramic breeder pebbles, either Li4SiO4 (OSi) or Li2TiO3 (MTi), with an outer diameter of about 45 mm and bed thickness of about 10 mm, sandwiched between two beryllium pebble beds. The breeder and beryllium pebble beds are separated by Eurofer-97 steel plates. The heat flow is managed so as to have a radial temperature distribution in the ceramic breeder pebble bed as flat as reasonably possible. The test element design and test matrix required extensive pretesting, improved pebble-bed modeling, design curves for the material characteristics, and performance analyses allowing in-reactor operation in High Flux Reactor (HFR) Petten121,130,141 (Figure 38). A specific pretest compaction procedure by pressing and heating has been developed, in particular to increase the thermal conductivity of the beryllium pebble beds and provide conditions that would result in limited changes during in-pile operation.130,141 The compaction procedure consisted of a subsequent loading of the pressure plate of the total assembly to 3 MPa. The X-ray pictures combined with the
6 30
: 50 MW Reactor power Sweep gas flow rate : 200 cm3 min–1 : 4 ⬚C min Temp. raising rate 4
Temperature difference (⬚C)
Effective thermal diffusivity (⫻10-3 cm2 s–1)
burnup with few pebbles in the heat flow direction or low volumetric heat loads and low lithium burnup.135–139 Verrall et al.138 seem to the first to report on bubble formation in a ceramic breeder. They observed this in Li2O irradiated in NRU up to 1% lithium burn-up up, see Figure 45. The effective thermal diffusivity of a Li2TiO3 pebble bed was studied in an in-pile irradiation experiment by Kawamura et al.140 at the JMTR test reactor. The cylindrical assembly of the Li2TiO3 pebble-bed was instrumented with a number thermocouples to determine the radial temperature profile. The derived thermal diffusivity as function of temperature is shown in Figure 36. The dimension of the pebble bed was 20 mm in diameter and 260 mm in length. The effective thermal diffusivity of the Li2TiO3 pebble bed was found to decrease with increasing irradiation temperature. This tendency remained up to a thermal neutron fluence of 1 1024 n m2, while the thermal conductivity at given temperatures also remained constant. None of these experiments addressed the pebblebed deformation behavior at the strong temperature gradients envisaged for breeding blankets. Outof-pile testing is less representative as it requires the use of heater plates, with their specific impact on reduced bed thickness, and running a ceramicheater interface at the highest temperature. Also, when material is irradiated in a stressed state, there is an additional phenomenon of irradiationinduced creep.
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Figure 36 The effective pebble-bed conductivity derived from in-pile Li2TiO3 pebble-bed irradiation in JMTR, as performed by Kawamura and coworkers.140 Reproduced from Kawamura, H.; Kikukawa, A.; Tsuchiya, K.; et al. Fusion Eng. Des. 2003, 69, 263–267.
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HCPB pebble-bed assembly
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Threaded ring, Sealing plate: all Eurofer-97
Figure 37 Schematic of HCPB pebble-bed assembly (PBA) test-element for in-pile testing in the High Flux Reactor at Petten, The Netherlands. Each test-element has a cylindrical ceramic breeder section, either Li4SiO4 or Li2TiO3, in between two cylindrical shaped beryllium pebble beds and separated by Eurofer plates.
Figure 38 Picture of PBA test element during assembly: Ceramic breeder section (Li2TiO3 pebbles from CEA) with penetrating thermocouple tubes; note pebble alignment at circumference.
dimensional measurements taken during assembly allowed the determination of actual bed size and compaction values. An example of a postassembly X-ray picture is shown in Figure 39.
The design and safety requirements for in-pile operation required the development of a full-coupled thermomechanical model in the MARC finite-element code. In this way, the pressure buildup and stress relaxation in the pebble beds could be simulated in detail to guide the required reactor startup profile.121,130,141 The two test elements with Li4SiO4 pebbles were irradiated at nominal temperatures of 600 and 800 C in the breeder bed, to see any effects of thermal creep. The other two test elements contained Li2TiO3 pebbles with different grain sizes and were irradiated at the same temperature, the nominal temperature of 800 C. The pebble beds were typically purged with a helium–hydrogen mixture of reference composition (0.1% H2). The gas purge entered at the lower beryllium bed and exited at the upper beryllium beds. The PBA has been irradiated for 294 FPD (full power days), achieving lithium burnups of 1.5–2.2% for Li4SiO4 and 2.8–2.9% for Li2TiO3, without enriching in 6Li. The dose in the Eurofer-97 parts ranged from 2 to 3 dpa.142 Extensive analyses of the in-pile data using FEM calculations showed that the maximum
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After inspection of assembled testelement: Stepped precompaction, 1, 2, 3 MPa plus 24 h at 350 ⬚C Purge line splitter
Pt clad on TCtubes
Last inspections prior and after sealing.
Large pebble: low density near wall
Pebbles well aligned with plate Figure 39 X-ray picture of PBA as assembled.
temperatures in the Li4SiO4 pebble beds of the top and bottom test-elements are about 600 and 800 C, respectively. The average temperatures in the Li4SiO4 beds are about 550 and 740 C, respectively.142 In postirradiation examinations of both Li4SiO4 samples, a little sintering and a significant amount of cracking or fragmentation have been observed. No significant difference between the lower and higher temperature case was found. Most of the evidence of cracking and fragmentation in the Li4SiO4 pebbles is observed toward the middle of the bed (highest temperature, highest deformation). This is visible in scanning electron microscopy images (Figure 42). There is some evidence of grain growth. Reactions of Li4SiO4 pebbles with Eurofer were found to be very small. The maximum temperatures in the Li2TiO3 pebble beds of both the two test-elements in the middle are about 780 C. The average temperatures in the Li2TiO3 beds are about 690 and 720 C, respectively.142 In both test-elements, Li2TiO3 pebbles showed a significant amount of sintering and necking, which was found most significant in the test-element #3. The average temperature of this test-element was higher by about 35 C. Almost no fracture or fragmentation was seen. There appeared to be a small reaction layer, distributed uniformly along the Eurofer (Figures 43–44).
Figure 40 Neutron radiograph of PBA taken after few irradiation cycles.
In the high burnup irradiation EXOTIC-7 (see details in Section 4.15.5.1.2), the pellet stacks and pebble beds were found to be essentially intact by neutron radiography analyses after irradiation, and except for one capsule containing Li2ZrO3 pellets, three out of five were found intact after unloading. Fragmentation of the 0.1–0.2 mm Li4SiO4
Ceramic Breeder Materials
pebbles was also observed but was very difficult to quantify33,143–145 (Figure 46). The pores observed in the images are related to the pebble fabrication method rather than to neutron irradiation. Chikhray et al. irradiated Li2TiO3 þ 5 mol% TiO2 in the Kazakhstan water water research reactor (WWRK) reactor to lithium burnups of about 20% at 760 and 920 K; see details in the next section and Chikhray et al.82 Pebble crush tests showed reduction of strength, whereas microhardness tests also revealed ingrowth of soft phases. X-ray diffraction measurements showed traces of LiTi2O4, LiTiO2, and Li4Ti5O12; see Chikhray et al.82 for more details. Temperature (⬚C)
827 727
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In an IEA-framed international collaboration, European Li4SiO4 and Li2TiO3 and Japanese Li2TiO3, reference pebble materials were tested in a high fluence irradiation project at the HFR in Petten, named high neutron fluence irradiation of pebble stacks for fusion (HICU).146–148,194 The neutronic analyses as reported by Fischer and coworkers149,150 demonstrated that relevant nuclear irradiation parameters such as the displacement damage accumulation, the lithium burnup, and the damage production function W(T) are met with the selected neutron shielding and 6Li enrichments chosen. This project is conceived to irradiate ceramic breeder pebble stacks at high temperatures under blanket prototypical ratios of fast neutron damage (dpa) and lithium burnup. Compared with the PBA experiment, the pebble stacks are smaller, but capsule dimensions are up to about 10 mm; X-ray tomography was used for detailed mapping of pebble location prior to irradiation.148,151
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4.15.5 Tritium Production and Release
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4.15.5.1
427 327 227 127 Test module #4 Figure 41 Temperature fields in the PBA as calculated with NRG’s fully coupled (cylindrical) pebble-bed thermo-mechanical model.
(a)
Tritium Release
A wide range of mechanisms play a role in the tritium transport and release processes of the lithiumcontaining ceramics, of which an impression is given in Figure 48.152 Tritium generated from neutron capture is first transported to the grain boundary by bulk diffusion. The bulk diffusion and trapping inside the grains are affected by the neutron radiation-induced defects. Via the intergranular diffusion, the tritium is then delivered to the grain surfaces, which are exposed to open and closed porosity. The closed porosity fraction provides another means to build up inventory in
(b)
Figure 42 Optical micrographs of cracked Li4SiO4 from PBA postirradiation examinations.
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(a)
(b)
Figure 43 Optical micrographs of sintered Li2TiO3 from PBA postirradiation examinations.
Figure 44 Optical micrograph on Li2TiO3 interaction with structure from PBA postirradiation examinations.
the material. At the surface isotope exchange with hydrogen (H2) and water (H2O) lead to desorption of tritium in molecular forms of HT and HTO, respectively. Further, the tritium in molecular form is transported through the interconnected pores and enters the flow of the purge gas. In order to assess the tritium retention in the candidate ceramic breeder material, one needs to know which of the steps are rate determining and which operation parameters are the most relevant for facilitation of the tritium release (Table 3). The tritium release characteristics of lithium ceramics are typically studied in two parameter ranges: 1. Out-of-pile: Tritium production through exposure to neutron irradiation, followed by out-of-pile tritium desorption through stepwise isothermal or ramp annealing tests in laboratory setups, also known as temperature programmed desorption (TPD). If irradiation doses are very low, such activity is typically called ‘tritium doping,’ and tritium transport parameters reflect beginning of life (BOL) conditions only because irradiation damage and lithium burnup remain negligible.
1 mm
Figure 45 Scanning electron micrograph showing small bubbles, randomly distributed, amid larger bubbles in Li2O after irradiation in NRU.138
2. In-pile: In case of in-pile experiments, typically steady-state tritium production and release conditions. In general, such parameters are closer to breeding blanket conditions, as they allow the application of a wide range of temperatures and purge gas conditions, and the study of long-term performance issues such as irradiation damage and lithium burnup. At present, such data are limited in terms of fast neutron damage doses (thermal and mixed spectra materials test reactor (MTR) only).
4.15.5.1.1 Out-of-pile
In the majority of cases, lithium ceramic samples are irradiated in research reactors with thermal or mixed neutron spectra. Extensive studies on tritium retention in neutron irradiated lithium ceramics using the TPD
Ceramic Breeder Materials
method have been reported in the literature.152–178,201 The chemical form of released tritium from Li4SiO4 (from FZK), LiAlO2 (from JAERI), Li2TiO3 (from CEA), and Li2ZrO3 (from MAPI) was studied in the out-of-pile tritium release experiment under various purge gas conditions. The pebbles were irradiated for a few minutes in the fluxes 4 1017 m2 s1 at Japan Research Reactor 4 ( JRR-4) or 1.65–2.75 1017 m2 s1 at the Kyoto University Reactor. The tritium was
released in the purge gas of dry nitrogen, nitrogen with 0.1% of helium, and nitrogen with 0.1% water vapor. Even if hydrogen was added to the purge gas, a considerable fraction of tritium was released in the molecular form of water, HTO. Addition of water vapor to the purge gas greatly enhanced the release rate of tritium. It was concluded that the isotope exchange of tritium with water at the exposed surfaces of the grains is much faster than the isotope exchange of tritium with hydrogen. The water is adsorbed at the grain surfaces from the water vapor present in the purge gas. Even small traces of water of 30 ppm in dry purge gas can be enough to promote tritium release in the HTO molecular form. Another source of water at the surfaces of the grains is the reduction reaction that takes place in the H2 reducing atmosphere. Among the studied materials, Li2TiO3 showed the largest water formation capacity.161 Figure 49 shows outof-pile annealing tests for Li4SiO4 with 0.1% H2/N2 sweep gas and 0.1% H2O/N2 sweep gas, and tritium doped for 2 min at a neutron flux of 2.75 1013 cm2 s1.158 Later work on palladium deposited as catalyst on orthosilicate indicated that almost all tritium was released as tritiated water vapor from lithium orthosilicate pebbles and tritium at higher temperatures remains slow.160 In contrast, it was also found that a considerably larger amount of tritium was released as the molecular form (HT) from the lithium
Figure 46 High burn-up Li4SiO4 pebble irradiated to 11% lithium burn-up, with fracture features, and large pores that originate from the manufacturing process.
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orthosilicate pebbles already deposited with palladium at lower temperatures (see Figure 50). Alvani and coworkers172,173 correlated TPD after short irradiations in the Casaccia reactor ( 2 1021 m2) with those from long-term irradiation in the HFR; Petten (thermal 0.5 1025 m2) revealed two peaks, at 770 and 941 K (b ¼ 5 K min1) (see Figure 51). It was proposed that the second peak is related to tritium trapping at the oxygen vacancies located along the grain boundary interface. The concentration of these trapping sites is significantly increased by the reduction effect of the R-gas, which results in a shift in the release peak to higher temperatures for the pretreated pebbles. The observed effect is more pronounced for the pebbles with finer
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H2O
HTO (LiOT + H2O LiOH + HTO)
LiOT
T
Li(n, a)T 7 Li(n, n*a)T
6
Neutron
Breeder pebble
LiOT
T T
LiOH Defects LiOT
Solid breeder grain
H2
HT (LiOT + H2 LiOH + HT)
T 2O (2LiOT
L i2O + T2O )
HTO (LiOT + LiOH Li2O+ HTO)
He Dry He purge
Figure 48 Schematic of tritium transfer phenomena through the ceramic breeder structure. Reproduced from Nishikawa, M.; Kinjyo, T.; Nishida, Y. J. Nucl. Mater. 2004, 325, 87–93.
grains. It was also found that a thermal pretreatment at 473 K for 2 h removes only the environmental H2O and CO2 contamination from the surface of the pebbles without affecting the H2O desorption at higher temperatures. The observed H2O release above 1173 K is ascribed to the reduction of Li2TiO3 to Li2TiO3x, with x reaching a steady-state value of xeq ¼ 0.01. Later work by Casadio et al.174 concerned a batch of Li2TiO3 pebbles (ENEA code FN5) prepared following the ‘citrate’ route. Analysis of TPD spectra gave the correct order of magnitude of the time constants characterizing the main desorption sites, in rough agreement with the residence times obtained by the in-pile step-perturbation methods performed during EXOTIC-8/9 experiment (see Figure 52).175 Pure helium purge increases the tritium inventory; during the last cycle of this irradiation experiment, variations of the H2 concentration in the He purge showed an increase in tritium release rate from Li2TiO3 pebbles that was found to be proportional at 473 C, Figure 53. to PH0:34 2 The effect of open and closed porosity on tritium release behavior in Li2O single crystal and sintered pellets was studied by Tanifuji et al.176–179 The pellets had densities in the range of 70–98% and grain sizes from 10 to 60 mm. Irradiations were performed by thermal neutrons in JRR-4, JAERI, up to 2 1023 m2. The porosity dependence of tritium release behavior from the Li2O sintered pellets has been investigated through isothermal heating tests, and the results are shown in Figure 54. For 88% TD specimens irradiated up to neutron fluences of 2 1022 and 2 1023 n m2, no
Table 3 Comparison of surface reactions on ceramic breeder grains associated with tritium release. Reproduced from Nishikawa, M.; Kinjyo, T.; Nishida, Y. J. Nucl. Mater. 2004, 325, 87–93. Dry purge gas
Purge gas with hydrogen
Purge gas with water vapor
573 K
Adsorption/desorption
Adsorption/desorption
573–473 K
Adsorption/desorption
773–973 K
Adsorption/desorption
973 K
Adsorption/desorption
Adsorption/desorption Water formation Isotope exchange 2 Isotope exchange I Adsorption/desorption Water formation Isotope exchange 2 Isotope exchange I Surface condition change Adsorption/desorption Water formation Isotope exchange I Isotope exchange 2 Surface condition change
Isotope exchange 2 Adsorption/desorption Isotope exchange 2 Adsorption/desorption Isotope exchange 2 Adsorption/desorption
Isotope exchange 2 Adsorption/desorption
495
Ceramic Breeder Materials
irradiation effects on the tritium residence time have been observed.
neutron damage doses (thermal and mixed spectra MTR only). Several irradiation programs have been executed around the world, involving thermal, mixedspectrum, and fast reactors [as in Dido (Germany), Siloe (France), EBR-2 (US), FFTF (US), HFR (The Netherlands), JMTR (Japan), and WWRK (Kazakhstan)]. The European irradiation projects under the acronym EXOTIC, for extraction of tritium in ceramics, commenced during the mid 1980s at the HFR in Petten26,33,63,144,175,180–183,195 (see Table 4). The initial series concerned both closed capsule and vented capsule operation. The materials investigated were mainly Li2O, Li2SiO3, LiAlO2, LiZrO3, Li8ZrO6, and Li4SiO4 in the series EXOTIC-1 to-6.26 The objective of the EXOTIC-7 experiment has been to irradiate candidate ceramic breeder materials in the HFR to a high lithium burnup (target 10%) and to determine the effects on the mechanical integrity of pellets and pebble-bed configurations, and those on tritium-inventory and-release characteristics.33 The experiment concerned 8 capsules, and during 11 HFR cycles (261 FPD), lithium burnups of 6–18% were achieved. The test matrix comprised pellets of Li2ZrO3, Li8ZrO6, and LiAIO2 and pebbles of Li2ZrO3 and Li4SiO4. Two capsules contained a mixture of Li4SiO4 and beryllium pebbles. To obtain a high lithium burnup within a reasonable irradiation time, the target materials were enriched with 6Li to about 50%. The tritium residence time of Li2ZrO3 pellets a higher temperatures was not changed during the irradiation. After a lithium burnup of about 5%, the residence time at lower temperatures decreases significantly with increasing burnup.
In-pile experiments have the strong advantage that the tritium release characteristics can be studied, as a function of neutron damage and lithium burnup in combination with thermal–mechanical behavior under neutron irradiation. Such in-pile experiments allow steady-state tritium production and release conditions to be achieved. In general, such parameters are closer to breeding blanket conditions, as they allow the application of a wide range of temperatures and purge gas conditions and the study of long-term performance issues such as irradiation damage and lithium burnup. At present, such data are limited in terms of fast 1000 Temperature 0.004
800
0.1% H2O/N2
Water vapor
0.003
0.1% H2/N2
600
0.002
400 Background
0.001
Temperature (⬚C)
200
0.000 1
0
2
3
4
0
Time (h)
Figure 49 Out-of-pile annealing tests for Li4SiO4 with 0.1% H2/N2 sweep gas and 0.1% H2O/N2 sweep gas (amount of breeder, 0.3 g); flow rate, 100 ml min1; irradiation time, 2 min; neutron flux, 2.75 1013 cm2 s1. Reproduced from Munakata, K.; Yokoyama, Y.; Baba, A.; Penzhorn, R. D.; Oyaidzu, M.; Okuno, K. Fusion Eng. Des. 2005, 75–79, 673–678.
1000 Breeder:Li4SiO4 Irradiation:12.0 min Sweep gas:0.1% H2/N2 Flow rate:100 ml min–1
0.03
800
600
(0.05% H2 +0.5 % H2O)/N2
0.02
Released HTO + HT
400
Released HT
0.01
200
0.00 1
2
3
4 Time (h)
5
6
Breeder:0.2%Pd/Li4SiO4 Irradiation:12.0 min
Sweep gas:0.1% H2/N2 Flow rate:100 ml min–1
0.03
800 (500 ppmH2 +5000 ppmH2O)/N2
Released HTO + HT
600
0.02 400 Released HT
0.01
200
0.00
0 0
(a)
1000
0.04
Temperature (⬚C)
Tritium concentration (mCi cm-3)
0.04
Tritium concentration (mCi cm-3)
Tritium concentration (mCi cm-3)
0.005
7
0
(b)
1
2
3
4 Time (h)
5
6
0 7
Figure 50 Out-of-pile annealing experiment of (a) Li4SiO4 and (b) Pd-coated Li4SiO4. Reproduced from Munakata, K.; Shinozaki, T.; Inoue, K.; et al. J. Nucl. Mater. 2009, 386–388, 1091–1094.
Temperature (⬚C)
4.15.5.1.2 In-pile testing
496
Ceramic Breeder Materials
ceo_clean
1.0 1
0.8
0.6
(a.u.)
Normalized desorption rate (Cn)
TRIGA casaccia lazy susan capsule Tirr about 330 K
0.4
0.2 HFR petten EXOTIC 8.1 capsule Tirr about 800 K
0
0.0 400
(a)
600
800 Temperature (K)
1000
1200
400
600 800 Temperature (K)
(b)
1000
Figure 51 (a) Normalized temperature programmed desorption spectra of tritium release from CEA pebbles after tritium doping in TRIGA (circles) and medium-term irradiation HFR-EXOTIC-8/1 (squares), 5 K min1 in both cases, and (b) image of the deconvolution-fit analysis performed on the spectrum in TRIGA reactor. Reproduced from Alvani, C.; Casadio, S. T.; Casadio, S. Fusion Eng. Des. 2003, 69, 275–280.
100
He
Tritium residence time (h)
Tritium residence time (h)
100
10
10
1 1.2
1.3
1.4 1000/T
1.5
1.6
(K−1)
1 10
100 1000 10 000 H2 concentration in purge gas (vpm)
Figure 52 Arrhenius plot of the tritium residence time for Li2TiO3 (FN5) pebbles irradiated in EXOTIC-8/9 experiment under reducing atmosphere; the black point was obtained in pure helium purge gas174. Reproduced from Casadio, S.; van der Laan, J. G.; Alvani, C.; Magielsen, A. J.; Stijkel, M. P. J. Nucl. Mater. 2004, 329–333, 1252–1255.
Figure 53 Tritium residence time in Li2TiO3 (FN5) pebbles at 473 C as a function of H2 concentration in the helium purge. The black point results from the best fitting line obtained for the reference composition176. Reproduced from Casadio, S.; van der Laan, J. G.; Alvani, C.; Magielsen, A. J.; Stijkel, M. P. J. Nucl. Mater. 2004, 329–333, 1252–1255.
After selection of the HCPB as the single solid breeder concept in the European Blanket Project, the EXOTIC-8 and-9 series concentrated on Li4SiO4 and Li2ZrO3 pebbles and a range of Li2TiO3 products.175 The irradiation test program concentrated on two types of experiments:
The typical designs for these experiments are given in Figure 55, with the general layout and an example of a cross-section from postirradiation testing. Figure 56 shows a sample temperature and tritium quantities in the purge gas for a typical irradiation cycle. Tritium residence time (t):
1. Tritium release to low or moderate lithium burnups 2. High lithium burnup and mechanical integrity
t¼
I G
Ceramic Breeder Materials
Two types of experiments were targeted:
Average residence time (min)
105 104
1000
-2
2 ⫻ 10
23
nm
2 ⫻ 10
22
n m -2
4 ⫻ 10
20
n m -2
4 ⫻ 10
20
n m -2
1. Major focus on tritium release characteristics by determining differential tritium inventories by thermal transients and achieving low to medium lithium burnups of 1–3%. 2. Major focus on high lithium burnup experiments using pebbles with 50% 6Li enrichment and achieving 11% lithium burnup for Li4SiO4 and 17% for Li2TiO3, at relatively constant temperatures, and only few data on release, mostly from postirradiation annealing tests.
-1
Ea = 160 kJ mol
-1
Ea = 184 kJ mol
Ea = 138 kJ mol
-1
100
10
1
(a)
(b)
(c) 81% TD (d) 81% TD As-irrad. Preadsorbed HTO
88% TD As-irrad. 0.1 1.2
497
1.3
1.4
1.5
1.6
1000/T
1.7
1.8
1.9
2
( K-1 )
Figure 54 Arrhenius plots of the average residence time t. Reproduced from Tanifuji, T.; Yamaki, D.; Jitsukawa, S. Fusion Eng. Des. 2006, 81, 595–600.
I, tritium inventory (Bq); G, tritium production rate (Bq min1). At steady state: Tritium release rate (R) ¼ Tritium generation rate (G) Temperature transients are performed: DT Difference in tritium inventory (area): DI ¼ I2I1 Difference in residence time: Dt ¼ DI/G Data set is processed to obtain: t(T) See Figure 57. In the EXOTIC-8 program, the tritium release characteristics and mechanical stability of the reference breeder materials for the European HCPB project have been studied, including the effect of long-term neutron irradiation and high lithium burnup.175 The EXOTIC-8 program started in 1997 and ended in 2002. It consisted of ten experiments, named EXOTIC-8/1 to EXOTIC-8/10. The irradiations were carried out in the HFR in Petten in peripheral core positions with the typical neutron fluence rate of about 9 1017 m2 s2 (fast, En > 0.1 MeV) and 5 1017 m2 s1 (thermal). The following materials were used: Li2TiO3 pebbles produced by agglomeration – sintering and extrusion – sintering, provided by CEA; Li2TiO3 pellets produced by cold pressing, provided by CEA; Li2TiO3 pebbles produced by wet processing and sintering, provided by ENEA; Li2ZrO3 pebbles produced by extrusion – sintering process, provided by CEA; and Li4SiO4 pebbles produced by melt spray process, provided by FZK Karlsruhe.
Tritium release characteristics are measured both in situ by applying the temperature transients and after irradiation in the TPD setup. The tritium release is characterized by the tritium residence time t with the Arrhenius temperature behavior, as depicted in the summary graph; Figure 58. A correlation has been established between the pebble density and the measured residence time. The obtained results confirm the understanding that open porosity and small grain size are favorable for faster tritium release. The corresponding activation energies derived from the temperature transients and from the TPD measurements are in fair agreement. For Li2TiO3 pebbles, Q ¼ 82–93 kJ mol1 and for Li4SiO4 pebbles, Q ¼ 112–123 kJ mol1. The activation energy for Li2ZrO3 pebbles is derived only from the temperature transients: Q ¼ 84 kJ mol1. In TPD experiments, it was shown that the tritium release can involve multiple release processes. Moreover, in some instances the release can be limited by recombination at the grain surface, which introduces uncertainties in the measured values of the activation energy. The results of in-pile tritium behavior during normal operation and during transients in temperature and gas chemistry measured in the latest irradiation experiment from the EXOTIC series, EXOTIC-9/1, are reported in Peeters et al.180 The Li2TiO3 pebbles produced by extrusion–spheronization sintering at CEA were irradiated in the HFR in Petten (thermal 0.5 1018 m2 s1) for 301 FPD and achieved a burnup of 3.8–4.1%. The temperature varied between 613 and 853 K. Based upon the in-pile tritium release measurements and the analysis of the tritium residence time, it was concluded that tritium release in the new batch of the high-density Li2TiO3 pebbles (93.0% TD) is rather slow compared with the ceramics irradiated in the EXOTIC-8 irradiation campaign.174 Thus, the tritium residence time measured at 773 K in the EXOTIC-9/1 experiment was 30 h, whereas
Materials
Shape
Li-6 (%)
Li-6 burnup (%)
FUBR-IA
Li2O, LiAlO2, Li2ZrO3, Li4SiO4
pellets
56, 95
FUBR-IB
Li2O, LiAlO2, Li2ZrO3, Li4SiO4 Li2O, LiAlO2, Li2ZrO3, Li4SiO4
pellets
0.07,7.5,56,95
pellets
0.07,7.5,56,95
MOTA-2A (BEATRIX-II)
Li2O, LiAlO2, Li4SiO4, Li2ZrO3
rings, pebbles
MOTA-2B (BEATRIX-II)
Li2O, LiAlO2, Li4SiO4, Li2ZrO3
rings, pebbles
CRITIC-3 COMPLIMENT (ELIMA-2, DELICE-3) In-Pile PebbleBed EXOTIC-1 (BEATRIX)
Li2TiO3 LiAlO2, Li2SiO3, Li4SiO4, Li2O, Li2ZrO3 Li2TiO3
pebbles pellets, pebbles
0.07, 0.2, 7.5, 56, 95 0.2 85 95 1.85 7.5
LiAlO2, Li2SiO3, Li2O
pellet
EXOTIC-2
LiAlO2, Li2SiO3, Li2O Li2ZrO3, Li2O, Li2SiO3 Li2ZrO3, Li6Zr2O7, Li8ZrO6, Li2O, Li2SiO3 LiAlO2, Li2ZrO3, Li4SiO4, Li6Zr2O7, Li8ZrO6
pellet
BEATRIX-I
EXOTIC-3 EXOTIC-4 EXOTIC-5
Total Li burn up (%)
Fluence thermal (1020 cm2)
Fluence E > 1 MeV (1022 cm2)
EFPD (Days)
1.1–1.2 2.0–2.2 3.1–3.5 3.9–4.3
1.4–1.45 3.9–4.1 2.5–2.7 4.7–5.1
96.7, 177.6, 274.3 341.5
6.7–7.4 10.3–11.3 10.3–11.3 1.7, 7.6–9.9
8.2–8.9 13–14 13–14 0.2, 7.4–8.8
599.3 940.8 936.9 299.7
1.27, 5.3–7.9
0.2, 7.7–7.9 0.9 0.25 1
pebbles
Temp ( C)
Reference
[186] 450–1225
[186] [186]
550–640 440–1000
[186]
203.3
530–640 440–1100
[186]
0.57
334 178 (HFR) 75 (OSIRIS)
200–900 400–450 650–700
[208] [209]
0.01
JMTR
400–610
[140]
0.06 0.6 7.5 0.55–0.6
0.004 0.035 0.28 0,1
0.01
0.013–0.025
25
350–700
[26],[210]
0.02–04
0.025–0.05
50
350–725
[26],[210]
pellet, solid pellet, solid
0.55–0.6
0.12–0.13
0.03–0.06
0.03–0.08
75
385–650
[26]
0.55–0.6, 7.7
0.13–0.15, 0.8
0.04–0.07
0.05–0.1
97
310–680
[26]
annular pellets and pebbles
7.5
1.8–2.1
0.07–0.13
0.08–0.17
166
325–630
[26]
Ceramic Breeder Materials
Experiment ID
498
Table 4 Overview of ceramic breeder irradiation experiments performed in fast or mixed spectrum fission reactors to relatively high levels of thermal and/or fast neutron fluence. Values for the experimental parameters quoted were taken from the mentioned references. Readers are also referred to the text throughout the whole chapter.
EXOTIC-6
LiAlO2, Li2ZrO3, Li6Zr2O7, Li8ZrO6, Li4SiO4
EXOTIC-7
Li2ZrO3, Li8ZrO6, Li4SiO4, LiAlO2
EXOTIC-8/1 EXOTIC-8/2
Li2TiO3 Li2TiO3
EXOTIC-8/3
Li4SiO4 + 2% TeO2, Be Li4SiO4 + 2% TeO2 Li2TiO3 Li2ZrO3 Li4SiO4, Be Li2TiO3, Be Li2TiO3 Li4SiO4 Li4SiO4, Li2TiO3
EXOTIC-8/4 EXOTIC-8/5 EXOTIC-8/6 EXOTIC-8/7 EXOTIC-8/8 EXOTIC-8/9 EXOTIC-8/10 PBA EXOTIC-9 K-578 (WWRK)
Li4SiO4 Li2TiO3
7.5
3.0–3.2
0.08–0.15
0.1–0.2
199
315–520
[26]
38–50
5.8–18.1
0.27
0.13
261
410–745
[33]
7.5 7.5
1.86 1.34
0.046 0.039
201 173
310–570 420–640
[175] [175]
50
2.6–4
0.003–0.013
201
200–700
[175]
pebbles
7.5
1.38
0.033
201
280–590
[175]
pebbles pebbles pebbles pebbles pebbles pebbles pebbles
7.5 7.5 51.5 50 7.5 7.5 7.5
2.1 2.5 8–17.5 4.1–10.9 3.51
0.06 0.071 0.025–0.064 0.11–0.38 0.1 0.09 0.04–0.07
0.11–0.17
200 299 648 450 448 448 294
500–660 400–640 400–690 400–700 420–560 520–695 550–800
[175] [175] [175] [175] [175] [175] [142]
pebbles
7.5, 30.5 7.5, 20 96
0.15
0.09
300
340–580
[143]
220
500–900
[212], [82]
403
600–900
[211]
pebbles, pellets pebbles
0.06, 7.5, 20 0.06, 11, 30
1.5–2.9 (for natural 6Li 3.1 (for natural 6Li) 18–23 0.7–13
1.5
0.7
Ceramic Breeder Materials
HICU
Li2TiO3 Li4SiO4 Li2TiO3
annular pellets and pebbles annular pellets and pebbles pebbles annular pellets pebbles
499
500
Ceramic Breeder Materials
Four channel irradiation rig External fluence detector Purge gas tube Second containment gas gap
Filler element Tritium breeder Central tube First containment Thermocouple Second containment Third cont. gas supply
24.1 mm 37.2 mm
(b)
(a)
Figure 55 Cross section of the EXOTIC series of ceramic breeder irradiation tests: a) schematic drawing, and b) macrograph from postirradiation microscopic investigations.
800
2 ¬ 7 ¬1
600
¬4
1
400
¬2
¬6
0.5
200 ¬3
0
Temperature (⬚C)
Tritium release (mCi min–1)
1.5
0
5
¬8
¬5
10
15
20
25
30
0
Time (days) Figure 56 Examples of tritium release from the EXOTIC-8 experiment series with a number of programmed temperature transients (including reactor power scram at day 19+20).
the same characteristic measured on the Li2TiO3 pebbles obtained from CEA and ENEA in the EXOTIC-8 campaign was 1.3 and 5.2 h, respectively (Figure 59). Changes in the tritium inventory resulting from the variation of the H2 concentration in the purge gas (from 0.1% to 1.0%) appeared to be much smaller
than those resulting from the temperature transients. From this observation, it was concluded that the tritium inventory was determined by the thermally activated processes taking place in the bulk of the material (dissociation from traps, diffusion) rather than by recombination and isotope exchange with hydrogen at the surface.
Ceramic Breeder Materials
501
Tritium release rate (mCi min-1)
0.40
0.35
T2
T1
0.30
0.25
0.20
0.15 0.0
Steadystate release
0.2
Steadystate release
Difference in inventory
0.4
0.6
0.8 Time (h)
1.0
1.2
1.4
1.6
Figure 57 Determination of differential tritium inventory from a small negative temperature transient in a steady state tritium production/release experiment.
104
103
t (h)
102
350
400
450
500
550
600
E-8/1: Li2TiO3 E-8/2: Li2TiO3 E-8/4: Li4SiO4 E-8/5: Li2TiO3 E-8/6: Li2ZrO3 E-8/9: Li2TiO3 E-8/10: Li4SiO4
104
103
102
101
101
100
100
10-1
10-1
10-2 1.1
1.2
1.3
1.4 1000/T (K)
1.5
1.6
10-2 1.7
Figure 58 Summary graph of tritium residence time obtained from the temperature transient experiments for different pebble materials used in the EXOTIC-8 program.
Tritium recovery characteristics of a binary bed containing 0.3 and 2 mm diameter Li2TiO3 pebbles were studied under continuous (20 h) and pulsed (200, 400, and 800 s) neutron irradiation in the JMTR.185 From the temperature transients from 573 to 623 K, the tritium residence time was estimated as 3 h (63.2% of the steady-state value). The complete recovery of the steady-state conditions was achieved after 20 h. The tritium recovery behavior under the
pulsed operation was almost the same as under continuous operation, except for the modulations introduced by the pulse operation, which did not exceed 20% of the total signal variation (Figures 60–62). Effects of irradiation temperature, purge gas flow rate, and hydrogen content in the purge gas on the tritium release characteristics of the Li2TiO3 pebbles were studied in the in-pile irradiation experiment in the JMTR.185 The Li2TiO3 pebbles were fabricated
502
Ceramic Breeder Materials
1000
Li2TiO3 CEA extrus; TD = 93; OP = 1.7 (E-9/1)
Li2TiO3 CEA extrus; TD = 94; OP = ? (E-8/9)
Li2TiO3 CEA extrus; TD = 87; OP = 7 (E-8/5)
Li2TiO3 CEA extrus; TD = 89; OP = 6 (E-8/7)
Tritium residence time (h)
E-9/1
100
Li2TiO3 ENEA wet pr; TD = 93; OP = 5.5 (E-8/9) E-8/9-wet E-8/1 E-8/5
10
E-8/2
E-8/7-ex
E-8/9-ex E-8/5
E-8/7-aggl
E-8/9-wet
Li2TiO3 CEA agglom; TD = 84; OP = 11 (E-8/7)
1
E-8/9-spinel E-8/1 E-8/2
E-8/7-aggl
0.1 1.1
1.2
Li2TiO3 CEA agglom; TD = 84; OP = 10 (E-8/1) Li4Ti5O12 (spinel) CEA agglom; TD = 90; OP = ? (E-8/9) Li2TiO3 CEA cold press, pellets; TP = 83; OP = 15 (E-8/2)
1.4
1.3
1.5
1.6
1.7
Temperature (1000/T (K)) Figure 59 Accumulated data for Li2TiO3 from EXOTIC series up to EXOTIC-9/1, including values from postirradiation inventory measurements180. Reproduced from Peeters, M. M. W.; Magielsen, A. J.; Stijkel, M. P.; van der Laan, J. G. Fusion Eng. Des. 2007, 82, 2318–2325.
109
: 200 cm3 min-1 Sweep gas flow rate : <1 ppm Moisture concentration Hydrogen content in sweep gas : 1000 ppm
Total tritium release (Bq min-1)
Total tritium release (Bq min-1)
109
108
Center temperature of Li2 TiO3 pebble bed
723 107
0
50
773 723 773 100
673
Center temperature of Li2TiO3 pebble bed Moisture concentration Hydrogen content in sweep gas
108
Sweep gas flow rate
773 673 [K]
150 200 250 Elapsed time (h)
300
350
Figure 60 Tritium release rate variation as function of the central temperature of the Li2TiO3 pebble bed irradiated in JMTR-experiment185. Reproduced from Tsuchiya, K.; Kikukawa, A.; Yamaki, D.; Nakamichi, M.; Enoeda, M.; Kawamura, H. Fusion Eng. Des. 2001, 58–59, 679–682.
by the rotating granulation method. The irradiation continued during one cycle of 25 days with the tritium generation rates of 6 1010 Bq d1 or 11 1010 Bq d1 depending on the position in the core. The following findings were reported: an increase in the purge gas flow rate accompanied by a temporal increase in the tritium release, which was followed by a swift recovery (<10 h at 773 K). A decrease in the flow rate produced
: 773 K : <1 ppm : 1000 ppm
950 107
0
200 50
100 200 950 200 (cm3 min-1) 100 150 Elapsed time (h)
200
250
Figure 61 Example of change of tritium release on variation of sweep gas flow rate185. Reproduced from Tsuchiya, K.; Kikukawa, A.; Yamaki, D.; Nakamichi, M.; Enoeda, M.; Kawamura, H. Fusion Eng. Des. 2001, 58–59, 679–682.
an opposite effect. The hydrogen concentration in the purge gas varied from 10 to 10 000 ppm. It was found that at hydrogen partial pressures <100 Pa, the tritium desorption is controlled by the surface reactions and at higher partial pressures by bulk diffusion. Chikhray et al.82 performed irradiation tests of Japanese Li2TiO3 ceramics with 96% enrichment of isotope 6Li in the WWRK reactor. Three types
Ceramic Breeder Materials
Overall rate constant of tritium desorption (Bq min-2)
of ceramic samples were examined simultaneously using a system for in-pile tritium monitoring: one (pebbles) – under constant temperature of 650 C, and two (pebbles and pellets) – within temperature change ranges from 500 to 900 C. Lithium burnup reached 23% for the active ampoule pebbles, 20% for the passive ampoule pebbles, and 18% for the pellets. The tritium measurement system permitted the
108 Moisture concentration : <1 ppm Center temperature of Li2TiO3 pebble bed : 773 K
107
106
105 0 10
101 102 103 Hydrogen partial pressure (Pa)
104
Figure 62 The relation found between hydrogen partial pressure and overall rate constant of tritium desorption185. Reproduced from Tsuchiya, K.; Kikukawa, A.; Yamaki, D.; Nakamichi, M.; Enoeda, M.; Kawamura, H. Fusion Eng. Des. 2001, 58–59, 679–682.
503
tritium yield rate to be determined under long-term irradiation of lithium ceramic Li2TiO3. Postirradiation testing also included mechanical testing.
4.15.6 Irradiation Parameters 4.15.6.1
Irradiation Damage
Greenwood calculated damage parameters for solid lithium tritium breeder materials irradiated in the fusion breeder reactor (FUBR) and breeder exchange matrix (BEATRIX) series of irradiations in EBR-II and FFTF reactors.186 He showed that most of the displacement damage arises from fast neutron reactions such as elastic and inelastic scattering rather than from the 6Li (n, a) T reaction. The displacements per atom is only an estimate of the total energy available for creating displacements in a compound and does not directly relate to observable damage effects due to the high probability of recombination effects. Leichtle and coworkers149,150 evaluated damage correlation parameters in fusion and fission reactor systems for the ceramic breeder materials Li2TiO3, Li4SiO4, and Li2O. A sophisticated damage calculation taking into account all relevant primary knock-on atoms (PKAs) was performed to obtain displacement damage parameters for these low-mass polyatomic materials. They arrived at a concise fusion–fission correlation of irradiation effects. Such comparisons have been performed for typical fission neutron flux spectra of a thermal and a fast reactor system. In
HCPB DEMO blanket : DPA versus lithium burnup
100 1 FPY 2 FPY 4 FPY
90 80 70
Li4SiO4 (30 at.%)
dpa
60 50
Li2TiO3 (45 at.%)
Li2O (15 at.%)
40 30 20 10 0 0
5
10 Lithium burnup (at.%)
15
20
Figure 63 Displacement per atom (dpa) accumulation versus lithium burn-up for the breeder materials Li2O, Li4SiO4, and Li2TiO3 in a helium-cooled pebble bed DEMO blanket, as analyzed by Fischer et al.150 The markers represent points in the breeder zone as a function of distance from the plasma-facing wall, with the lowest dpa values at the blanket backside.
504
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Former Li4SiO4
1E+1
Li4SiO4 pebbles
Contact dose rate (Sv h-1)
6Li
Li4SiO4 Pure Li4SiO4
1E-1 Recycling limit 1E-3
1E-5
Hands-on-limit
1E-7 1E-4
1E-2
1E+0 1E+2 Time after shutdown (years)
1E+4
Figure 64 Contact dose rate versus time after shutdown of pure Li4SiO4 without impurities, and the main contributors. Reproduced from Knitter, R.; Fischer, U.; Herber, S.; Adelhelm, C. J. Nucl. Mater. 2009, 386–388, 1071–1073.
preparation for an IEA-framed international program on high fluence breeder irradiation, Fischer et al.150 performed analyses of lithium burnup and displacement damage in Li2O, Li4SiO4, and Li2TO3 for a HCPB-type DEMO blanket design (Figure 63).
4.15.7 Activation and Waste Issues During neutron irradiation of ceramic breeders, not only does the transmutation of Li to T and He take place but other isotopes are formed as well, and the impurities contained will all contribute to the induction of radioactivity from exposure to the neutron irradiation environment. This enables assessment of the feasibility of recycling ceramic breeder material from blanket components having reached EOL conditions. Recycling options for ceramic breeder material not only avoid large waste volumes requiring long-term storage but also contribute to resource efficiency of valuable constituents such as lithium. From a ceramic breeder perspective, Li2O, Li2TiO3, and Li4SiO4 are more attractive than Li2ZrO3 due to their long-term activation characteristics. Knitter et al.187 assessed contact dose rates of different Li4SiO4 materials for high radiation levels expected from a fusion power reactor for 1 FPY. As an example, Figure 64 displays the contact dose rate versus time after shutdown for pure Li4SiO4, which specifically results from the production of 28Al, 24Na, 7 Be, and 26Al. Due to the activity of 7Be, the recycling limit for remote handling (10 mSv h1) and the
hands-on limit (10 muSv h1) are reached after <1 and <4 years, respectively. The activation of as manufactured Li4SiO4 appears to be strongly influenced by impurities, of which Co and Pt appear to be the most important.187 Knitter et al. also simulated remelting of lithium depleted Li4SiO4 experimentally.188 Pebble manufacturing routes by wet processes require different considerations for reprocessing.196 In the powder preparation stage the powders need to be produced from pre-existing pebbles by dissolution and precipitation steps; see, for example, Tsuchiya et al.184
4.15.8 Summary and Outlook Ceramic breeder materials offer a wide range of possibilities for the development of fusion energy based on the deuterium–tritium fuel cycle. Currently, most ceramic breeder R&D is focused on the lithium orthosilicate and metatitanate systems in the form of pebble beds. This chapter started with blanket designs, material requirements, manufacturing routes, pebble and pebble-bed thermomechanics, tritium production and release properties, neutron-irradiation behavior, chemistry, and modeling. One of the most important missions of ITER is to provide a test bed for breeding blanket modules, the so-called test blanket modules (TBMs). However, because ITER testing is a cost-intensive exercise and is most likely the only opportunity to test a fusion
Ceramic Breeder Materials
505
Economic, recycling, radioisotope removal Fabrication process/science Material characterization Out-of-pile pebble bed TM performance and chemical stability
Grain size
<5 mm
Pebble density
>80–85%
Surface area
i.e., >0.1 mg m–2
Pebble size Open porosity
0.25 - ~2 mm i.e., 4–8%
Sphericity/roundness
>0.7
% of pebble breakage, i.e., <0.5% CH composition change, i.e., <1%
In-pile T release at high burnup
T residence time, pebble integrity
In-pile PBA tests at 2–3% BU
Procurement specifications for ITER ceramic breeding materials
ITER testing
T release at high temperatures, Li burn up and dpa
DEMO application
PBA performance at DEMO TM conditions and burnup per dpa
HICU and others
Figure 65 Outline of a ceramic breeder material development roadmap, prior to Iter testing, as proposed by Ying et al.200 The road connects actual experiments like PBA (pebble-bed assembly), HICU (small pebble stacks), using parameters as fluence (dpa), lithium burn-up/BU, and temperature in order to cover the thermo-mechanical (TM) loadings of a pebble-bed anticipated in a DEMO power plant.
blanket component prior to a DEMO, the questions of the choice of materials for the ITER TBM and the definition of a set of requirements (and the related qualification program) to ensure safety, reliability, and test performances become particularly important. Accordingly, Ying et al.190 proposed a roadmap outlining the necessary development steps for qualifying and accepting the pebbles for ITER and fusion applications (see Figure 65). For each development step, a set of criteria is presented as a means for initial screening before proceeding to the next evaluation tests to reduce development costs. However, it is important to recognize that ITER conditions (neutron fluence about 1.52 orders of magnitude lower) are far from sufficient to qualify any specific breeder material to be used in DEMO. Thus, parallel with ITER and subsequent to ITER testing, tests such as HICU or in fusion relevant neutron sources such as International Fusion Materials Irradiation Facility (IFMIF) for any candidate ceramic breeders under typical reactor blanket conditions with relevant nuclear environment are necessary for this purpose.
Though a significant R&D effort on ceramic breeder development has already been made and a vast amount of data on material performance have been obtained, the knowledge to date on the limiting factors in blanket designs for long-term operation is still modest. These limitations are addressed here: 4.15.8.1
Microstructure
Only a few types of microstructures have been studied until today, preventing clear conclusions regarding their suitability for reactor application. Both for pebble-bed and pellet/block type designs, significant opportunities still exist to tailor stoichiometric composition, grain size, and porosity shape and size as well as surface treatments, coatings, and other pre-conditioning treatments. 4.15.8.2
High Burnup
Though some high lithium burnup samples have been achieved in experimental programs, the findings are not yet conclusive as to whether the
506
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ceramic still functions adequately under normal operation, and whether interaction with the blanket structure satisfies, for example, future power plant safety requirements. 4.15.8.3
High Fluence
Few data were obtained in fast breeder experiments; the influence of fast neutron damage on the integrity and functionality of the breeder ceramics is yet to be addressed, in particular for the newly developed pebble types and compositions. 4.15.8.4
High Temperature
The effect of long-term exposure in a high radiation environment, causing damage, induced swelling, with increasing burnup, may also affect the vaporization behavior and even induce lithium transport within the breeder area. 4.15.8.5 Effects of Transients and Off-Normal Conditions Though work has been done already on thermal cycling of pebble beds, there is uncertainty with regard to the effects of disruption and electromagnetic (EM) loads on, for example, pebble relocation, fragment relocation, and ratcheting in case of inclined or vertical beds. Specific attention is required on the change of heat transfer properties and free flow of purge gas. 4.15.8.6 Accident Behavior (Safety and Investment Integrity) In addition to the integrity and safety requirements according to which a blanket component has to be designed in order to be licensed, there is another matter of economics: in case of transient events or accidents not affecting safety, they may prohibit further operation of the blanket and require replacement. Those blanket concepts that are more tolerant to design base and other accidents will be preferred by utilities. In this context, the use of beryllium-based neutron multipliers, the type of coolant, and the tritium issues for the purge and coolant processing units (isotope separation, purification, steam generator operation, toxicity at accidental conditions) also appear to be very important aspects associated with the breeder material use.
4.15.8.7
Development of Tools
Thermomechanical codes to describe the interaction of pebble beds with the structural material have achieved significant progress in recent years. For complete TBMs and, even more, DEMO blanket modules, computational codes based on continuum mechanics will be the first choice in the near future. These codes should be quickly developed to enable fairly good predictions for the blanket behavior at the BOL of a DEMO plant or power reactor. The database for irradiated material is still very small and the amount of relevant engineering data in the near future will also be very limited, as material development is still ongoing. Pebble-bed experiments that require a considerable amount of material will be costly and hardly possible for some time. This is the significant challenge for discrete-element methods (DEM). These codes must be improved in order to describe more realistically the interaction between nonirradiated pebbles (taking into account pebble shape, surface condition, and material properties, including thermal creep). If this goal is achieved, it should be possible to do this for irradiated materials, because a small pebble mass is required to make corresponding experiments. The results obtained then with DEM codes should be fed into formal correlations in the continuum codes in order to assess the blanket behavior toward EOL conditions. 4.15.8.8
Compatibility with Structure
Though experimental evidence is now accumulating for the irradiation behavior of current candidates for ceramic breeder and structure, there is no clear insight into the extent of chemical/physicochemical interactions taking place in long-term operations under a reducing atmosphere in a DEMO or power reactor. While these may affect the breeder properties, they may also affect the blanket components structural integrity. 4.15.8.9 Waste Management and Reuse/Recycling The large volume of ceramic breeder and multiplier required for breeding blankets in future power reactors necessitate ecological and sound economic solutions for intermediate storage and back end. Cost effectiveness and sound nuclear industrial practices will promote the selection and qualification of ceramic breeder technologies with full processing and recycling capabilities.
Ceramic Breeder Materials
Acknowledgments The authors acknowledge the support of their home organizations NRG and KIT. This chapter is also based on work partially supported by the European Union. NRG activities have been financially supported by The Netherlands Ministry of Economy, Agriculture and Innovation. The content of the publication is the sole responsibility of the authors and it does not necessarily represent the views of the European Commission or its services and the views from other sponsors.
References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10.
11. 12. 13. 14. 15. 16.
17. 18. 19. 20.
Ying, A.; Akiba, M.; Boccaccini, L. V.; et al. J. Nucl. Mater. 2007, 367–370, 1281–1286. Proust, E.; Anzidei, L.; Roux, N.; et al. Fusion Technol. 1992, 21, 2089–2098. Giancarli, L.; Petrizzi, L.; Diop, C.; Rado, V. Fusion Eng. Des. 1991, 17, 145–151. Eid, M. Ed. Helium-Cooled Ceramic-Breeder-In-Tube Blanket Line; CEA Report DMT 94/576 SERMA 1682; Dec 1994. Gohar, Y.; Attaya, H.; Billone, M.; et al. Fusion Technol. 1991, 19(3), 1538–1545. Dalle Donne, M.; Bojarsky, E.; Fisher, U.; et al. In Proceedings of the 17th SOFT, Rome, 1992; p 1326. Dalle Donne, M. (Comp.), European DEMO BOT Solid Breeder Blanket, KfK-Report, Forschungszentrum Karlsruhe, KfK 5429; 1994. Dalle Donne, M.; Anzidei, L. Fusion Eng. Des. 1995, 27, 319–336. Gierszewski, P.; Dalle Donne, M.; Kawamura, H.; Tillack, M. Fusion Eng. Des. 1995, 27, 167–178. Lorenzetto, P.; Gierszewski, P.; Chiocchio, S.; Daenner, W.; Federici, G.; et al. Fusion technology 1992. In Proceedings of the 17th Symposium on Fusion Technology, Rome, 1992; p 1414. Ioki, K.; Ferrari, M. Final Design Report, Design Description Document WBS1.6B, Tritium Breeding Blanket System, ITER Doc G16 DDD 2 98-0-10 W0.4; 1998. Ferrari, M.; et al. Fusion Eng. Des. 1999, 46, 177. Nardi, C.; Petrizzi, L.; Piazza, G. Fusion Eng. Des. 2003, 69, 315–319. Gorokhov, V. A.; Gryaznov, N. S.; Davydov, D. A.; et al. Atom. Energy 2000, 89(2), 638–645. Hermsmeyer, S.; Malang, S.; Fischer, U.; Gordeev, S. Fusion Eng. Des. 2003, 69, 281–287. EFDA, A. Conceptual Study of Commercial Fusion Power Plants; Final Report of the European Fusion Power Plant Conceptual Study (PPCS); Report EFDA-RP-RE-5.0; 2005. Poitevin, Y.; Boccaccini, L. V.; Dell’Orco, G.; et al. Fusion Eng. Des. 2007, 82, 2164–2170. Hermsmeyer, S.; Fischer, U.; Fuetterer, M.; Schleisiek, K.; Schmuck, I.; Schnauder, H. Fusion Eng. Des. 2001, 58–59, 689–693. Boccaccini, L. V.; Fischer, U.; Gordeev, S.; Malang, S. Fusion Eng. Des. 2000, 49–50, 491–497. Akiba, M.; Enoeda, M.; Tsuru, D.; et al. Fusion Eng. Des. 2009, 84, 329–332.
21.
507
Ihli, T.; Basu, T. K.; Giancarli, L. M.; et al. Fusion Eng. Des. 2008, 83, 912–919. 22. Hermsmeyer, S.; et al. In Proceedings of CBBI-7, 7th International Workshop on Ceramic Breeder Blanket Interactions, Petten, The Netherlands, NRG Report 21099/99.23482/P, Mar 1999. 23. Giancarli, L.; Chuyanov, V.; Abdou, M.; et al. Fusion Eng. Des. 2006, 81, 393–405. 24. Chuyanov, V. A.; Campbella, D. J.; Giancarli, L. M. Fusion Eng. Des. 2010, 85, 2005–2011. 25. Johnson, C. E.; Kummerer, K. R.; Roth, E. J. Nucl. Mater. 1988, 155–157, 188–201. 26. Kwast, H.; Stijkel, M.; Muis, R.; Conrad, R. EXOTIC: Development of Ceramic Tritium Breeding Materials for Fusion Reactor Blankets; ECN Report ECN-C-95-123; Dec 1995. 27. Dienst, W.; Schild, D.; Werle, H. Tritium Release of Li4SiO4, Li2O and Beryllium and chemical compatibility of Beryllium with Li4SiO4, Li2O and Steel (SIBELIUS Irradiation), KfK Bericht 5109, Karlsruhe, Dec 1992. 28. Roux, N.; Hollenberg, G.; Johnson, C.; Noda, K.; Verrall, R. Fusion Eng. Des. 1995, 27, 154–166. 29. Roux, N.; Johnson, C.; Noda, K. J. Nucl. Mater. 1992, 191–194, 15–22. 30. Kapychev, V.; Davydov, D.; Gorokhov, V.; et al. J. Nucl. Mater. 2000, 283–287, 1429–1433. 31. Sharafat, S.; Ghoniem, N.; Sawan, M.; Ying, A.; Williams, B. Fusion Eng. Des. 2006, 81, 455–460. 32. Sens, P. Y.; Majoor, E. B. M. In Nuclear Energy Maturity, Vol. 3&4, Nuclear Fuel Performance and Management; Sher, R., Ed.; Pergamon: Oxford, 1976; pp 152–160. 33. van der Laan, J. G.; Kwast, H.; Stijkel, M.; et al. J. Nucl. Mater. 1996, 233–237, 1446–1451. 34. Kleykamp, H. In Proceedings of CBBI-7, 7th International Workshop on Ceramic Breeder Blanket Interactions, Petten, The Netherlands, NRG Report 21099/99.23482/P, Mar 1999. 35. Gierszewski, P.; Hamilton, H.; Miller, J.; et al. Fusion Eng. Des. 1995, 27, 297–306. 36. CNM, Chapter 66: FM Steels – Internal CNM. 37. van der Schaaf, B.; Tavassoli, F.; Fazio, C.; et al. Fusion Eng. Des. 2003, 69, 197. 38. Malang, S. Fusion Eng. Des. 1999, 46, 193–206. 39. Cismondi, F.; Kecske´s, S.; Ilic, M.; et al. Fusion Eng. Des. 2009, 84, 607–612. 40. Zmitko, M.; et al. ICFRM-9, Sapporo, 2009. 41. Poitevin. Y; et al. Fusion Eng. Des. 2010, 85, 2340–2347. 42. Kohyama, A.; Seki, M.; Abe, K.; et al. J. Nucl. Mater. 2000, 283–287, 20. 43. Seki, Y.; Kikuchi, M.; Ando, T.; Ohara, Y.; Nishio, S.; et al. In Proceedings of the 13th International Conference on Plasma Physics and Controlled Nuclear Fusion Research, Washington, DC, 1990; IAEA: Vienna, 1991; Vol. 3, p 473. 44. Hermsmeyer, S.; Gordeev, S.; Kleefeldt, K.; et al. In Improved Helium Cooled Pebble Bed Blanket; FZK report, FZKA 6399; Dec 1999. 45. Nishio, S; et al. Fusion Eng. Des. 1998, 41, 357–364. 46. Akiba, M; et al. Fusion Eng. Des. 2009, 84, 329–332. 47. Ehrlich, K.; Gasparotto, M.; Giancarli, L.; Le Marois, G.; Malang, S.; van der Schaaf, B. In European Material Assessment Meeting; Report EFDA-T-RE-2.0; 2001. 48. Boccaccini, L. V.; Giancarli, L.; Janeschitz, G.; et al. J. Nucl. Mater. 2004, 329–333, 148–155. 49. Fischer, U.; Tsige-Tamirat, H. J. Nucl. Mater. 2002, 307–311, 798–802. 50. Cook, I.; et al. Plasma Phys. Control. Fusion 2002, 44, B121.
508 51. 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. 77. 78.
79.
Ceramic Breeder Materials Summary of the ITER Final Design Report, July 2001, Report GAO FDR 4 01-07-26 R 0.4; p 9. Roux, N.; Avon, J.; Floreancig, A.; Mougin, J.; Rasneur, B.; Ravel, S. J. Nucl. Mater. 1996, 233–237, 1431. Lulewicz, J. D.; Roux, N. Fusion Eng. Des. 1998, 39–40, 745–750. Kwast, H.; Conrad, R.; May, R.; Casadio, S.; Roux, N.; Werle, H. J. Nucl. Mater. 1994, 212–215, 1010–1014. Casadio, S.; Alvani, C.; Johnson, C. E.; Pierdominici, F. Proceedings of the 20th SOFT, Marseille, 1998; p 1211. Zhu, D.; Peng, S.; Chen, X.; Gao, X.; Yang, T. J. Nucl. Mater. 2010, 396, 245–250. Palermo, I.; Gomez-Ros, J. M.; Sanz, J.; Sedano, L. In Presented at Tritium 2010, Nara, 2010. Hoshino, T.; Tsuchiya, K.; Hayashi, K.; Terai, T.; Tanaka, S.; Takahashi, Y. Fusion Eng. Des. 2005, 75–79, 939–943. Chaudhuri, P.; et al. In Proceedings of CBBI-15, 15th International Workshop on Ceramic Breeder Blanket Interactions; Tanigawa, H., Enoeda, M., Eds.; Sapporo, Japan, JAEA-Conf 2009-006. Mandal, D.; et al. In Proceedings of CBBI-15, 15th International Workshop on Ceramic Breeder Blanket Interactions; Tanigawa, H., Enoeda, M., Eds.; Sapporo, Japan, JAEA-Conf 2009-006. Cho, S.; Ahn, M.-Y.; Ku, D. Y.; et al. Fusion Sci. Technol. 2009, 56, 216–220. Roux, N.; Tanaka, S.; Johnson, C.; Verrall, R. Fusion Eng. Des. 1998, 41, 31–38. Magielsen, A. J.; et al. Private communication. Kopasz, J.; Miller, J.; Johnson, C. In Proceedings of the 6th International Conference on Fusion Reactor Materials, Stresa, Italy, 1993. Kopasz, J. P.; Johnson, C. E.; Baldwin, D. L. J. Nucl. Mater. 1995, 219, 259–264. Hoshino, T.; Dokiya, M.; Terai, T.; Takahashi, Y.; Yamawaki, M. Fusion Eng. Des. 2002, 61–62, 353–360. Hoshino, T.; et al. Fusion Eng. Des. 2007, 82, 2269. Hoshino, T.; et al. J. Nucl. Mater. 2009, 386–388, 1098. Tsuru, D.; Enoeda, M.; et al. Fusion Sci. Technol. 2009, 56, 875–882. Hoshino, T.; Kato, K.; Natori, Y.; et al. Fusion Eng. Des. 2009, 84, 956–959. Billone, M. C.; Grayhack, W. T. Summary of Mechanical Properties Data and Correlations for Li2O, Li4SiO4, LiAlO2 and Be; Argonne National Laboratory Report ANL/FPP/TM-218; Apr 1990. Hayashi, T.; Konishi, S.; Okuno, K. J. Nucl. Mater. 1990, 170, 60–65. Alvani, C.; Bruzzi, L.; Casadio, S.; Rondinella, V.; Tucci, A.; Toscano, E. H. J. Eur. Ceram. Soc. 1989, 5, 295–302. Kennedy, P.; Gilchrist, K. E.; Walker, D. E.; Broughton, S. Fusion Technol. 1986, 2, 1013. Flipot, J.; Brauns, E.; Diels, P. H. Adv. Ceram. 1989, 27, 95. Pannhorst, W.; Geiler, V.; Ra¨ke, G.; Speit, B.; Sprenger, D. In Proceedings of 20th SOFT, CEA, 1998; p 1441. Piazza, G.; Reimann, J.; Gu¨nther, E.; Knitter, R.; Roux, N.; Lulewicz, J. D. Fusion Eng. Des. 2001, 58–59, 653–659. Knitter, R.; et al. In Proceeding of the 11th International Workshop on Ceramic Breeder Blanket Interactions; Enoeda, M., Ed.; Japan Atomic Energy Research Institute, 2004, 108–119. Tsuchiya, K.; Watarumi, K.; Saito, S.; Fuchinoue, K.; Furuya, T.; Kawamura, H. In Presented at 5th International Workshop on Ceramic Breeder Blanket Interactions, Rome, 1996.
80. 81. 82. 83. 84. 85. 86. 87. 88. 89.
90. 91. 92. 93. 94. 95. 96. 97. 98.
99. 100. 101. 102. 103.
104. 105. 106. 107. 108. 109.
Tsuchiya, K.; Kawamura, H.; Uchida, M.; Casadio, S.; Alvani, C.; Ito, Y. Fusion Eng. Des. 2003, 69, 449–453. van der Laan, J.; Muis, R. J. Nucl. Mater. 1999, 271–272, 401–404. Chikhray, Y.; Shestakov, V.; Maksimkin, O.; et al. J. Nucl. Mater. 2009, 386–388, 286–289. Wu, X.; Wen, Z.; Xu, X.; Gu, Z.; Xu, X. J. Nucl. Mater. 2008, 373, 206–211. Lulewicz, J. D.; Roux, N. Fusion Eng. Des. 1998, 39–40, 745–750. Suzuki, T.; Murata, O.; Hirata, S. Ceram. Trans. 1992, 27, 17–36. Enoeda, M.; Ohara, Y.; Roux, N.; Ying, A.; Malang, S. In Proceedings of the CBBI-8, Colorado Springs, CO, Oct 6–8, 1999. McGeary, R. K. J. Am. Ceram. Soc. 1961, 44(10), 513–522. Dalle Donne, M.; Piazza, G.; Weisenburger, A.; et al. In Proceedings of the 19th SOFT, Lisboa, Portugal, 1996. Tsuchiya, K.; Watarumi, K.; Saito, S.; Fuchinoue, K.; Furuya, T.; Kawamura, H. In Presented at 5th International Workshop on Ceramic Breeder Blanket Interactions, Rome, 1996. Tsuchiya, K.; Kawamura, H.; Tanaka, S. Fusion Eng. Des. 2006, 81, 1065–1069. Gan, Y. FZKA 7455, Oct. 2008. Jaeger, H. M.; Nagel, S. N.; Behringer, R. B. Physics Today 1996, 32–38. Piazza, G.; Erbe, A.; Rolli, R.; Romer, O. J. Nucl. Mater. 2004, 329–333, 1260–1265. Reimann, J.; Arbogast, E.; Behnke, M.; Muller, S.; Thomauske, K. Fusion Eng. Des. 2000, 49–50, 643–649. Reimann, J.; Wo¨rner, G. CBBI-9, 9th International Workshop on Ceramic Breeder Blanket Interactions, Toki, Japan, Sept 27–29, 2000. Reimann, J.; Worner, G. Fusion Eng. Des. 2001, 58–59, 647–651. Reimann, J.; Boccaccini, L.; Enoeda, M.; Ying, A. Y. Fusion Eng. Des. 2002, 61/62, 319–331. Reimann, J.; Lulewicz, J. D.; Roux, N.; Wo¨rner, G. In CBBI-10, 10th International Workshop on Ceramic Breeder Blanket Interactions, Karlsruhe, Germany, Oct 22–14, 2001; FZKA 6720. Reimann, J.; Hermsmeyer, S. Fusion Eng. Des. 2002, 61/62, 345–351. Reimann, J.; Ericher, D.; Worner, G. Fusion Eng. Des. 2003, 69, 241–244. Reimann, J.; Pieritz, R. A.; di Michiel, M.; Ferrero, C. Fusion Eng. Des. 2005, 75–79, 1049–1053. Reimann, J.; Harsch, H. Fusion Eng. Des. 2005, 75–79, 1043–1047. Reimann, J. In Proceedings of CBBI-7, 7th International Workshop on Ceramic Breeder Blanket Interactions, Petten, The Netherlands, NRG Report 21099/99.23482/P, Mar 1999. Buehler, L.; Reimann, J. J. Nucl. Mater. 2002, 307–311, 807–810. Reimann, J.; Pieritz, R. A.; Ferrero, C.; di Michiel, M.; Rolli, R. Fusion Eng. Des. 2008, 83, 1326–1330. Pieritz, R. A.; Reimann, J.; Ferrero, C. Adv. Eng. Mater. 2011, 13(3), 145–155. Abou-Sena, A.; Neuberger, H.; Ihli, T. Fusion Eng. Des. 2009, 84, 355–358. Dell’Orco, G.; Ancona, A.; Di Maio, P. A.; Sansone, L.; Simoncini, M.; Vella, G. Fusion Eng. Des. 2003, 69, 233–240. Dell’Orco, G.; Di Maio, P. A.; Giammusso, R.; et al. Fusion Eng. Des. 2006, 81, 169–174.
Ceramic Breeder Materials 110. Reimann, J.; Harsch, H. In Proceedings of the 12th International Workshop on Ceramic Breeder Blanket Interactions, Karlsruhe, Germany, FZKA 7078, 2005. 111. Dalle Donne, M.; et al. In Proceedings of the 3rd IEA International Workshop on Beryllium Technology for Fusion, Mito, Oct 22–24, 1997. 112. Tanigawa, H.; Enoeda, M.; Akiba, M. Fusion Eng. Des. 2007, 82, 2259–2263. 113. Tanigawa, H.; Hatano, T.; Enoeda, M.; Akiba, M. Fusion Eng. Des. 2005, 75–79, 801–805. 114. Dell’Orco, G.; Ancona, A.; DiMaio, A.; Simoncini, M.; Vella, G. J. Nucl. Mater. 2004, 329–333, 1305–1308. 115. Bu¨hler, L. Forschungszentrum Karlsruhe Continuum Models for Pebble Beds in Fusion Blankets-FZKA 6561; 2002. 116. Abou-Sena, A.; Ying, A.; Abdou, M. Fusion Sci. Technol. 2005, 47, 1094–1100. 117. Dell’Orco, G.; di Maio, P. A.; Giamusso, R.; Tincani, A.; Vella, H. Fusion Eng. Des. 2007, 82, 2366–2374. 118. Yagi, S.; Kunii, D. AIChE J. 1960, 6(1), 97–104. 119. Gan, Y.; Kamlah, M. J. Mech. Phys. Solids 2010, 58, 129–144. 120. Vella, G.; Di Maio, P. A.; Oliveri, E.; Dalle Donne, M.; Piazza, G.; Scaffidi-Argentina, F. Fusion Eng. Des. 2001, 58–59, 635–640. 121. Fokkens, J. Thermo-Mechanical Finite Element Analyses for the HCPB In-Pile Test Element; NRG Report 21477/ 02.50560/P; Jun 2003. 122. Hofer, D.; Kamlah, M. Fusion Eng. Des. 2005, 73, 105–117. 123. Zaccari, N.; Aquaro, D. Fusion Eng. Des. 2008, 83, 1282–1286. 124. Ying, A.; Huang, H.; Abdou, M.; Lu, Z. In Proceedings of the 9th International Workshop on Ceramic Breeder Blanket Interactions, Toki, Japan, Oct. 27–29, 2000. 125. Poitevin, Y.; Duchesne, A.; Eid, M.; Giancarli, L., In Proceedings of the 20th Symposium on fusion technology, SOFT-20, Marseille, 2008; pp 1195–1198. 126. Hermsmeyer, S.; Reimann, J. Fusion Eng. Des. 2002, 61–61, 367–373. 127. van der Laan, J.; et al. In Proceedings of the 8th International Workshop on Ceramic Breeder Blanket Interactions, Colorado Springs, CO, 1999; Ying, A., Ed.; UCLA Report. 128. Di Maio, P. A.; Dell’Orco, G.; Giammusso, R.; et al. Fusion Eng. Des. 2008, 83, 1287–1293. 129. Fokkens, J. H. Helica Calculations; NRG report; 2010. 130. van der Laan, J. G.; et al. Fusion Eng. Des. 2000, 51–52, 909–918. 131. An, Z.; Ying, A.; Abdou, M. In 7th International Symposium on Fusion Nuclear Technology, Tokyo, Japan, May 22–27, 2005. 132. Gan, Y.; Kamlah, M. Fusion Eng. Des. 2010, 85, 24–32. 133. Gan, Y.; Kamlah, M.; Reimann, J. Fusion Eng. Des. 2010, 85, 1782–1787. 134. Ethridge, L.; Baker, D. E. Adv. Ceram. 1989, 25, 165. 135. Nagaoa, Y.; Tsuchiya, K.; Ishida, T.; Kawamura, H.; Niimi, M. Fusion Eng. Des. 2006, 81, 619–623. 136. Tsuchiya, K.; Kawamura, H.; Nakamichi, M.; et al. J. Nucl. Mater. 1995, 219, 240. 137. Verrall, R. A.; Miller, J. M.; Macdonald, D. S.; Bokwa, S. R. J. Nucl. Mater. 1988, 155–157, 553–557. 138. Verrall, R. A.; Rose, D. H.; Miller, J. M.; Hastings, I. J.; MacDonald, D. S. J. Nucl. Mater. 1991, 179–181, 855–858. 139. Verrall, R. A.; Slagle, O. D.; Hollenberg, G. W.; Kurasawa, T.; Sullivan, J. D. J. Nucl. Mater. 1994, 212–215, 902–907.
140.
509
Kawamura, H.; Kikukawa, A.; Tsuchiya, K.; et al. Fusion Eng. Des. 2003, 69, 263–267. 141. van der Laan, J. G.; Boccaccini, L. V.; Conrad, R.; et al. Fusion Eng. Des. 2002, 61/62, 383–390. 142. van Til, S. In Analysis of the HCPB Pebble Bed Assemblies, EFDA task TW1-TTBB-004 Deliverable D2, NRG-21844/11.107003. 143. van Til, S. Final report on EXOTIC-9/1 including PIE, EFDA Task TW2-TTBB-004b, NRG-21372/11.106196. 144. Bakker, K.; et al. In Proceedings of CBBI-7, 7th International Workshop on Ceramic Breeder Blanket Interactions, Petten, The Netherlands, NRG Report 21099/99.23482/P, Mar 1999. 145. Bakker, K; Magielsen, A. J. Private communication. 146. Hegeman, J. B. J.; van der Laan, J. G. High Fluence Irradiation of Ceramic Pebble Beds (HICU) Pre-tests and Pre-irradiation Characterization; NRG Report 20295/ 03.55382/P; July 2004. 147. van der Laan, J. G. et al. In Proceedings of CBBI-11, 11th International Workshop on Ceramic Breeder Blanket Interactions, Tokyo, Japan; Enoeda, M. Ed.; JAERI-Conf 2004-012. 148. Magielsen, A. J.; et al. In Proceedings of CBBI-15, 15th International Workshop on Ceramic Breeder Blanket Interactions, Sapporo, Japan; Tanigawa, H., Enoeda, M., Eds.; JAEA-Conf 2009-006. 149. Leichtle, D.; Fischer, U. Fusion Eng. Des. 2000, 51–52, 1–10. 150. Fischer, U.; Herring, S.; Hogenbirk, A.; et al. J. Nucl. Mater. 2000, 280, 151–161. 151. Hegeman, J. B. J.; van der Laan, J. G.; Beemsterboer, C.; Kamer, S.; Peeters, M. In Proceedings of CBBI-14, the 14th International Workshop on Ceramic Breeder Blanket Interaction, Petten, The Netherlands, Sept 2006 152. Nishikawa, M.; Kinjyo, T.; Nishida, Y. J. Nucl. Mater. 2004, 325, 87–93. 153. Kinjyo, T.; Nishikawa, M.; Yamashita, N.; Koyamaa, T.; Tanifuji, T.; Enoeda, M. Fusion Eng. Des. 2007, 82, 2147–2151. 154. Kinjyo, T.; Nishikawa, M.; Enoeda, M.; Fukada, S. Fusion Eng. Des. 2008, 83, 580–587. 155. Munakata, K.; Baba, A.; Kawagoe, T.; et al. Fusion Eng. Des. 2000, 49–50, 621–628. 156. Munakata, K.; Yokoyama, Y.; Baba, A.; et al. Fusion Eng. Des. 2001, 58–59, 683–687. 157. Munakata, K.; Koga, A.; Yokoyama, Y.; et al. Fusion Eng. Des. 2003, 69, 27–31. 158. Munakata, K.; Yokoyama, Y.; Baba, A.; Penzhorn, R. D.; Oyaidzu, M.; Okuno, K. Fusion Eng. Des. 2005, 75–79, 673–678. 159. Munakata, K.; Shinozaki, T.; Inoue, K.; et al. Fusion Eng. Des. 2008, 83, 1317–1320. 160. Munakata, K.; Shinozaki, T.; Inoue, K.; et al. J. Nucl. Mater. 2009, 386–388, 1091–1094. 161. Nishikawa, M.; Kinjyo, T.; Ishizaka, T.; et al. J. Nucl. Mater. 2004, 335, 70–76. 162. Kinjyo, T.; Nishikawa, M.; et al. Fusion Sci. Technol. 2005, 48, 646. 163. Kinjyo, T.; Nishikawa, M.; et al. Fusion Eng. Des. 2006, 81, 573–577. 164. Nishikawa, M.; Kinjyo, T.; Yamashita, N.; Koyama, T. In Proceedings of the 14th International Workshop on Ceramic Breeder Blanket Interactions, Petten, The Netherlands, Sept 6–8, 2006. 165. Kinjyo, T.; Nishikawa, M.; et al. In Proceedings of the 14th International Workshop on Ceramic Breeder Blanket Interactions, Petten, The Netherlands, Sept 6–8, 2006.
510
Ceramic Breeder Materials
166. Kinjyo, T.; Nishikawa, M.; Enoeda, M. J. Nucl. Mater. 2007, 367–370, 1361. 167. Nishikawa, M.; et al. J. Nucl. Mater. 2007, 367–370, 1371. 168. Kinjyo, T.; Nishikawa, M.; et al. Fusion Eng. Des. 2007, 82, 2147. 169. Kinjyo, T.; Nishikawa, M.; et al. Fusion Sci. Technol. 2008, 54, 557. 170. Suematsu, K.; Nishikawa, M.; et al. Fusion Sci. Technol. 2008, 54, 561. 171. Kinjyo, T.; Nishikawa, M.; Enoeda, M. Fusion Eng. Des. 2008, 83, 580. 172. Alvani, C.; Carconi, P. L.; Casadio, S.; Roux, N. Fusion Eng. Des. 2001, 58–59, 701–705. 173. Alvani, C.; Casadio, S. t.; Casadio, S. Fusion Eng. Des. 2003, 69, 275–280. 174. Casadio, S.; van der Laan, J. G.; Alvani, C.; Magielsen, A. J.; Stijkel, M. P. J. Nucl. Mater. 2004, 329–333, 1252–1255. 175. Fedorov, A. V.; Peeters, M. M. W.; van der Laan, J. G. EXOTIC-8: Tritium Release Measurements in Ceramic Tritium Breeder Materials, In-pile Operation and Tritium Release Analysis; NRG Report 21100/08.92317; 2010. 176. Tanifuji, T.; Yamaki, D.; Nasu, S.; Noda, K. J. Nucl. Mater. 1998, 258–263, 543–548. 177. Tanifuji, T.; Yamaki, D.; Takahashi, T.; Iwamoto, A. J. Nucl. Mater. 2000, 283–287, 1419–1423. 178. Tanifuji, T.; Yamaki, D.; Jitsukawa, S. J. Nucl. Mater. 2002, 307–311, 1456–1460. 179. Tanifuji, T.; Yamaki, D.; Jitsukawa, S. Fusion Eng. Des. 2006, 81, 595–600. 180. Peeters, M. M. W.; Magielsen, A. J.; Stijkel, M. P.; van der Laan, J. G. Fusion Eng. Des. 2007, 82, 2318–2325. 181. van Til, S.; Magielsen, A. J.; Stijkel, M. P.; Cobussen, H. L. Fusion Eng. Des. 2010, 85, 1143–1146. 182. Conrad, R.; DeBarberis, L. J. Nucl. Mater. 1991, 179–181, 1158–1161. 183. Conrad, R.; Bakker, K.; Chabrol, C.; et al. In Proceedings of the 8th International Workshop on Ceramic Breeder Blanket Interactions, Colorado Springs, CO, 1999; Ying, A., Ed.; UCLA Report. 184. Tsuchiya, K.; Kikukawa, A.; Hoshino, T.; et al. J. Nucl. Mater. 2004, 329–333, 1248–1251. 185. Tsuchiya, K.; Kikukawa, A.; Yamaki, D.; Nakamichi, M.; Enoeda, M.; Kawamura, H. Fusion Eng. Des. 2001, 58–59, 679–682. 186. Greenwood, L. R. Radiation Damage Calculations for the FUBR and BEATRIX Irradiations of Lithium Compounds in EBR-II and FFTF; Report PNNL-12200; May 1999. 187. Knitter, R.; Fischer, U.; Herber, S.; Adelhelm, C. J. Nucl. Mater. 2009, 386–388, 1071–1073. 188. Knitter, R.; Lo¨bbecke, B. J. Nucl. Mater. 2007, 361, 104–111. 189. Herber, S.-C.; Fischer, U.; Knitter, R.; Foßhag, E. ICFRM14. Fus. Eng. Des. 2009 (submitted). 190. Ying, A.; Akiba, M.; Boccaccini, L. V.; et al. J. Nucl. Mater. 2007, 367–370, 1281–1286. 191. Knitter, R.; Alm, B.; Roth, G. J. Nucl. Mater. 2007, 367–370, 1387–1392. 192. Roux, N.; Tanaka, S.; Johnson, C.; Verrall, R. Fusion Eng. Des. 1998, 41, 31–38. 193. Lulewicz, J. D.; Roux, N. In Proceedings of CBBI-7, 7th International Workshop on Ceramic Breeder Blanket Interactions, Petten, The Netherlands, NRG Report 21099/99.23482/P, Mar 1999.
194. 195.
196. 197. 198. 199. 200. 201. 202. 203. 204.
205. 206. 207.
208.
209. 210. 211. 212. 213. 214. 215.
216.
217.
Hegeman, J. B. J.; van Essen, E. D. L.; Jong, M.; van der Laan, J. G.; Reimann, J. Fusion Eng. Des. 2003, 69, 425–429. van der Laan, J. G.; Bakker, K.; Conrad, R.; Stijkel, M. P.; Werle, H. In Proceedings of CBBI-7, 7th International Workshop on Ceramic Breeder Blanket Interactions, Petten, The Netherlands, NRG Report 21099/99.23482/P, Mar 1999. Tsuchiya, K.; Kawamura, H.; Saito, M.; Tatenuma, K.; Kainose, M. J. Nucl. Mater. 1995, 219, 246–249. Billone, M. C.; Dienst, W.; Flament, T.; Lorenzetto, P.; Noda, K.; Roux, N. ITER Solid Breeder Blanket Materials Data Base; Report ANL/FPP/TM-263; 1993. Abou-Sena, A.; Ying, A.; Abdou, M. J. Mater. Process. Technol. 2007, 181, 206–212. Aquaro, D.; Zaccari, N. J. Nucl. Mater. 2007, 367–370, 1293–1297. Ying, A.; Abdou, M.; Calderoni, P.; et al. Fusion Eng. Des. 2006, 81, 659–664. Okazaki, M.; Yamaski, T.; Gotoh, S.; Toei, R. J. Chem. Eng. Jpn. 1981, 14(3), 183–189. Zaccari, N.; Aquaro, D. Fusion Eng. Des. 2007, 82, 2375–2382. Zaccari, N.; Aquaro, D. J. Nucl. Mater. 2009, 386–388, 1078–1082. Dalle Donne, M.; Goraieb, A.; Piazza, G.; ScaffidiArgentina, F. In Proceedings of the 5th International Symposium on Fusion Nuclear Technology, Roma, 15–19 Sept., 1999. An, Z.; Ying, A.; Abdou, M. Fusion Eng. Des. 2007, 82, 2233–2238. Takatsu, H.; et al. Fusion Eng. Des. 1998, 39–40, 645–650. Nishio, S. Prototype fusion reactor based on SiC/SiC composite material focusing on easy maintenance. In Proceedings of the IAEA-TCM on Fusion Power Plant Design, Culham, March 24–27, 1998. Verrall, R. A.; Miller, J. M.; Gierszewski, P. In Proceedings of the 8th International Workshop on Ceramic Breeder Blanket Interactions (CBBI-8), Colorado Springs; 1999; p 66. Kummerer, K. R.; Ritzhaupt-Kleissl, H.-J. J. Nucl. Mater. 1991, 179–181, 831–834. Kwast, H.; et al. J. Nucl. Mater. 1988, 155–157, 558–562. van Til, S. HICU: Second irradiation progress report, subtask TW1-TTBB-006-D6, Report NRG-20295/ 10.101223; 2011. Chikhray, Y.; et al. J. Nucl. Mater. 2007, 367–370, 1028–1032. http://www.iter.org/mach/tritiumbreeding. Kolb, M. H. H.; Knitter, R.; Kaufmann, U.; Mundt, D. Fusion Eng. Des. 2011, doi: 10.1016/j. fusengdes.2011.01.104. Reimann, J.; Wo¨rner, G. In Proceedings of the 12th International Workshop on Ceramic Breeder Blanket Interactions, Karlsruhe, Germany, FZKA; p 7078, 2005. van der Laan, J.; et al. In Proceedings of the 8th International Workshop on Ceramic Breeder Blanket Interactions (CBBI-8), Colorado Springs, CO 1999, Ying, A., Ed.; UCLA Report. Schlu¨nder, E. V. In Proceedings of the 7th International Heat Transfer Conference, Mu¨nchen, Germany; 1982; Vol. 1, RK10, pp 192–212.
4.16 Tritium Barriers and Tritium Diffusion in Fusion Reactors R. A. Causey, R. A. Karnesky, and C. San Marchi Sandia National Laboratories, Livermore, CA, USA
ß 2012 Elsevier Ltd. All rights reserved.
4.16.1
Introduction
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4.16.2 4.16.2.1 4.16.2.2 4.16.2.3 4.16.2.4 4.16.2.5 4.16.2.6 4.16.2.7 4.16.3 4.16.3.1 4.16.3.1.1 4.16.3.1.2 4.16.3.1.3 4.16.3.2 4.16.3.2.1 4.16.3.2.2 4.16.3.2.3 4.16.3.2.4 4.16.3.2.5 4.16.3.3 4.16.3.3.1 4.16.3.3.2 4.16.3.3.3 4.16.3.3.4 4.16.3.3.5 4.16.4 4.16.4.1 4.16.4.2 4.16.4.3 4.16.5 References
Background Equation of State of Gases Diffusivity Solubility Trapping Permeability Recombination Irradiation and Implantation Fusion Reactor Materials Plasma-Facing Materials Carbon Tungsten Beryllium Structural Materials Austenitic stainless steels Ferritic/martensitic steels V–Cr–Ti alloys Zirconium alloys Other structural metals Barrier Materials Oxides Aluminides Nitrides Carbides Low permeation metals Application of Barriers Expected In-Reactor Performance How Barriers Work and Why Radiation Affects Them Why Barriers Are Needed for Fusion Reactors Summary
513 513 513 514 514 516 517 518 518 518 518 521 524 527 527 528 532 534 536 536 536 537 538 539 541 542 542 543 545 545 546
Abbreviations bcc CLAM CVD fcc HFR HIP ITER
Body-centered cubic China low activation martensitic steel Chemical vapor deposition Face-centered cubic High flux reactor Hot isostatically pressed International Thermonuclear Experimental Reactor
PCA Prime candidate alloy PRF Permeation reduction factor RAFM Reduced activation ferritic/martensitic steel
4.16.1 Introduction As fusion energy research progresses over the next several decades, and ignition and energy production 511
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are attempted, the fuel for fusion reactors will be a combination of deuterium and tritium. From a safety point of view, these are not the ideal materials. The reaction of deuterium with tritium produces a-particles and 14.1 MeV neutrons. These neutrons are used not only to breed the tritium fuel, but also interact with other materials, making some of them radioactive. Although the decay of tritium produces only a lowenergy b-radiation, it is difficult to contain tritium. Additionally, being an isotope of hydrogen, tritium can become part of the hydrocarbons that compose our bodies. From the tritium point of view, the fusion facility can be divided into three components: the inner vessel area where the plasma is formed, the blanket where tritium production occurs, and the tritium exhaust and reprocessing system. There is the potential for tritium release in all the three sections of the facility. The tritium cycle for a fusion reactor begins in the blanket region. It is here that the tritium is produced by the interaction of neutrons with lithium. Specifically, the reaction is given symbolically as 6Li(n,a)3H. A neutron that has been thermalized, or lowered in energy by interaction with surrounding materials, is absorbed by 6 Li to produce both an a-particle (helium nucleus) and a triton. Elemental lithium contains 7.5% 6Li. As a breeder material in a fusion plant, lithium is enriched in the 6Li isotope to various degrees, depending on the particular blanket design. The 7Li isotope also produces a small amount of tritium via the 7Li (n,a)3H þ n reaction. The cross-section for this endothermic reaction is much smaller than that for the 6Li reaction. Upon release from the lithium breeder, the tritium is separated from other elements and other hydrogen isotopes. It is then injected as a gas or frozen pellet into the torus, where it becomes part of the plasma. A fraction of the tritium fuses with deuterium as part of the fusion process, or it is swept out of the chamber by the pumping system. If tritium is removed from the torus by the pumping system and sent to the reprocessing system, it is again filtered to separate other elements and other isotopes of hydrogen. All through the different steps, there is the potential for permeation of the tritium through the materials containing it and for its release to the environment. The probability of this occurring depends on the location in the tritium cycle. This chapter describes hydrogen permeability through two categories of materials that will be used in fusion reactors: candidate plasma-facing and structural materials. The plasma-facing materials in future fusion devices will be heated by high-energy neutrons, by
direct interaction of the plasma particles, and by electromagnetic energy released from the plasma. These plasma-facing materials must be cooled. It is primarily through the cooling tubes passing through the plasma-facing materials that tritium losses can occur in the primary vacuum vessel. The three materials typically used for plasma-facing applications are carbon, tungsten, and beryllium. In this report, we describe the behavior of these materials as plasmafacing materials and how tritium can be lost to the cooling system. The term ‘structural material’ is used here to describe materials that serve as the vacuum boundary in the main chamber, as the containment boundary for the blanket region, and as the piping for cooling and vacuum lines. These materials can be ferritic and austenitic steels, vanadium alloys, and zirconium alloys, as well as aluminum alloys in some locations, or potentially ceramics. We give a complete list of the different types of structural materials and review their tritium permeation characteristics. Materials with a low permeability for tritium are being considered as barriers to prevent the loss of tritium from fusion plants. There are a few metals with relatively small values of permeability, but as a whole, metals themselves are not good barriers to the transport of tritium. Ceramics, on the other hand, are typically very good barriers if they are not porous. In most cases, the low permeation is due to extremely low solubility of hydrogen isotopes in ceramic materials. Bulk ceramics, such as silicon carbide, may one day be used as tritium permeation barriers, but most of the current barrier development is for coatings of oxides, nitrides, or carbides of the metals themselves. We show in this review that many such oxides and nitrides may exhibit extremely good permeation behavior in the laboratory, but their performance as a barrier is significantly compromised when used in a radiation environment. We review the permeation parameters of materials being considered for barriers. This report begins with a review of the processes that control the uptake and transport of hydrogen isotopes through materials. The parameters used to define these processes include diffusivity, solubility, permeability, trapping characteristics, and recombination-rate coefficients. We examine the transport of hydrogen isotopes in plasma-facing materials, discuss the conditions that exist in the main torus, and look at the ways in which tritium can be lost there. Next, we consider the tritium transport properties of structural materials, followed by the transport properties in barrier materials, including
Tritium Barriers and Tritium Diffusion in Fusion Reactors
oxides, nitrides, and carbides of structural metals, as well as low-permeation metals. The application of tritium barriers is discussed in some detail: both the theoretical performance of barriers and their observed performance in radiation environments, as well as an example of tritium permeation in the blanket of a fusion reactor. We conclude by summarizing the tritium permeation properties of all the materials, providing the necessary parameters to help designers of fusion reactors to predict tritium losses during operation.
4.16.2 Background Hydrogen and its isotopes behave similarly in many regards. Gaseous protium, deuterium, and tritium are all diatomic gases that dissociate, especially on metal surfaces, and dissolve into the metal lattice in their atomic form (in some materials, such as polymers and some ceramics, the molecules may retain their diatomic character during penetration of the material). The isotope atoms readily recombine on the free surfaces, resulting in permeation of the gaseous hydrogen isotopes through metals that support a gradient in hydrogen concentration from one side to the other. In order to understand this process, it is necessary to characterize the source of the hydrogen isotope as well as its transport within the materials. For the purposes of the presentation in this section, we focus on tritium and its transport through materials. Much of the discussion is equally valid for the deuterium and protium as well (and subsequent sections normalize data to protium). In this section, we provide background on the diffusivity and solubility of tritium in metals and relate these thermodynamic parameters to the permeability. In addition, we discuss the role of trapping of hydrogen isotopes on transport of these isotopes, as well as kinetically limited transport phenomena such as recombination. 4.16.2.1
Equation of State of Gases
In the case of gaseous exposure, the ideal gas equation of state characterizes the thermodynamic state of the gas: Vm0 ¼ RT =p Vm0
½1
where is the molar volume of the ideal gas, T is the temperature of the system in Kelvin, p is the partial pressure of the gaseous species of interest, and R is the universal gas constant equal to 8.31447 J mol1 K1. The ideal gas equation of state provides
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a good estimate for most gases, particularly at low pressures (near ambient) and elevated temperatures (greater than room temperature). In the context of materials exposed to hydrogen isotopes in fusion technologies, the assumption of ideal gas behavior is a reasonable estimate for gaseous hydrogen and its isotopes. More details about the equation of state for real gaseous hydrogen and its isotopes can be found in San Marchi et al.1 4.16.2.2
Diffusivity
Tritium diffusion in metals is simply the process of atomic tritium moving or hopping through a crystal lattice. Tritium tends to diffuse relatively rapidly through most materials and its diffusion can be measured at relatively low temperatures. Diffusivity, D, is a thermodynamic parameter, and therefore, follows the conventional Arrhenius-type dependence on temperature: D ¼ D0 expðED =RT Þ
½2
where D0 is a constant and ED is the activation energy of diffusion. Measuring tritium diffusion is nontrivial because of the availability of tritium. Therefore, hydrogen and deuterium are often used as surrogates. From the classic rate theory, it is commonly inferred that the ratio of diffusivities of hydrogen isotopes is equivalent to the inverse ratio of the square root of the masses of the isotopes: rffiffiffiffiffiffi DT mH ¼ ½3 DH mT where m is the mass of the respective isotope, and the subscripts tritium and hydrogen refer to tritium and hydrogen, respectively. When this approximation is invoked, the activation energy for diffusion is generally assumed to be independent of the mass of the isotope. Diffusion data at subambient temperatures do not support eqn [3] for a number of metals;2 however, at elevated temperatures, the inverse square root dependence on mass generally provides a reasonable approximation (especially for face-centered cubic (fcc) structural metals).3–9 Although eqn [3] provides a good engineering estimate of the relative diffusivity of hydrogen and its isotopes, more advanced theories have been applied to explain experimental data; for example, quantum corrections and anharmonic effects can account for experimentally observed differences of diffusivity of isotopes compared to the predictions of eqn [3].3,10 For the purposes of this report, we assume that eqn [3] is a good approximation
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Tritium Barriers and Tritium Diffusion in Fusion Reactors
for the diffusion of hydrogen isotopes (as well as for permeation) unless otherwise noted, and we normalize reported values and relationships of diffusivity (and permeability) to protium. 4.16.2.3
Solubility
The solubility (K) represents equilibrium between the diatomic tritium molecule and tritium atoms in a metal according to the following reaction: 1=2T2 $ T
½4
The solubility, like diffusivity, generally follows the classic exponential dependence of thermodynamic parameters: K ¼ K0 expðDHs =RT Þ
½5
where K0 is a constant and DHs is the standard enthalpy of dissolution of tritium (also called the heat of solution), which is the enthalpy associated with the reaction expressed in eqn [4]. A word of caution: the enthalpy of dissolution is sometimes reported per mole of gas (i.e., with regard to the reaction T2 $ 2T as in Caskey11), which is twice the value of DHs as defined here. Assuming a dilute solution of dissolved tritium and ideal gas behavior, the chemical equilibrium between the diatomic gas and atomic tritium dissolved in a metal (eqn [4]) is expressed as pTT 0 ¼ m0t þ RT ln c0 ½6 1=2 mTT þ RT ln 0 pTT where c0 is the equilibrium concentration of tritium dissolved in the metal lattice in the absence of stress, m0TT is the chemical potential of the diatomic gas at a 0 , and m0T is the reference partial pressure of pTT chemical potential of tritium in the metal at infinite dilution. This relationship is the theoretical origin of Sievert’s law: c0 ¼ K ðpTT Þ1=2
½7
where to a first approximation, the solubility is equivalent for all isotopes of hydrogen. It is important to distinguish between solubility and concentration: solubility is a thermodynamic property of the material, while the concentration is a dependent variable that depends on system conditions (including whether equilibrium has been attained). For example, once dissolved in a metal lattice, atomic tritium can interact with elastic stress fields: hydrostatic tension dilates the lattice and increases the concentration of tritium that can dissolve in the metal, while
hydrostatic compression decreases the concentration. The relationship that describes this effect in the absence of a tritium flux12–14 is written as VT cL ¼ c0 exp ½8 RT where cL is the concentration of tritium in the lattice subjected to a hydrostatic stress ( ¼ ii =3), and VT is the partial molar volume of tritium in the lattice. For steels, the partial molar volume of hydrogen is 2 cm3 mol1,15 which can be assumed to first order to be the same for tritium. For most systems, the increase of tritium concentration will be relatively small for ordinary applied stresses, particularly at elevated temperatures; for example, hydrostatic tension near 400 MPa at 673 K results in a 15% increase in concentration. On the other hand, internal stresses near defects or other stress concentrators can substantially increase the local concentration near the defect. It is unlikely that local concentrations will significantly contribute to elevated tritium inventory in the material, but locally elevated concentrations of hydrogen isotopes become sites for initiating and propagating hydrogenassisted fracture in structural metals. 4.16.2.4
Trapping
Tritium can bond to microstructural features within metals, including vacancies, interfaces, grain boundaries, and dislocations. This phenomenon is generally referred to as ‘trapping.’15–18 The trapping of hydrogen and its isotopes is a thermally governed process with a characteristic energy generally referred to as the trap binding energy Et. This characteristic energy represents the reduction in the energy of the hydrogen relative to dissolution in the lattice16,19 and can be thought of as the strength of the bond between the hydrogen isotope and the trap site to which it is bound. Oriani16 assumed dynamic equilibrium between the lattice hydrogen and trapped hydrogen yT yL Et ½9 ¼ exp 1 yT 1 y L RT where yT is the fraction of trapping sites filled with tritium and yL is the fraction of the available lattice sites filled with tritium. According to eqn [9], the fraction of trap sites that are filled depends sensitively on the binding energy of the trap (Et) and the lattice concentration of tritium (yL). For example, traps in ferritic steels, which are typically characterized by low lattice concentrations and trap energy <100 kJ mol1, tend to be depopulated at high temperatures (>1000 K).
Tritium Barriers and Tritium Diffusion in Fusion Reactors
For materials with strong traps and high lattice concentration of tritium, trapping can remain active to very high temperatures, particularly if the trap energy is large (>50 kJ mol1). The coverage of trapping sites for low and high energy traps is shown in Figure 1 for two values of K: one material with relatively low solubility of hydrogen and the other with high solubility. The absolute amount of trapped tritium, cT, depends on yT and the concentration of trap sites, nT15: cT ¼ anT yT
½10
where a is the number of hydrogen atoms that can occupy the trap site, which we assume is one. If multiple trapping sites exist in the metal, cT is the sum of trapped tritium from each type of trap. A similar expression can be written for the tritium in lattice sites, cL: cL ¼ bnL yL
½11
where nL is the concentration of metal atoms and b is the number of lattice sites that hydrogen can occupy per metal atom (which we again assume is one). Substituting eqns [10] and [11] into eqn [9] and recognizing that yL 1, the ratio of trapped tritium to lattice tritium can be expressed as cT nT ¼ cL ½cL þ nL expðEt =RT Þ
½12
515
Therefore, the ratio of trapped tritium to dissolved (lattice) tritium will be large if cL is small and Et is large. Conversely, the amount of trapped tritium will be relatively low in materials that dissolve large amounts of tritium. The transport and distribution of tritium in metals can be significantly affected by trapping of tritium. Oriani16 postulated that diffusion follows the same phenomenological form when hydrogen is trapped; however, the lattice diffusivity (D) is reduced and can be replaced by an effective diffusivity, Deff, in Fick’s first law. Oriani went on to show that the effective diffusivity is proportional to D and is a function of the relative amounts of trapped and lattice hydrogen: Deff ¼
D cT 1 þ ð1 yT Þ cL
½13
If the amount of trapped tritium (cT) is large relative to the amount of lattice tritium (cL), the effective diffusivity can be several orders of magnitude less than the lattice diffusivity.20 Moreover, the effective diffusivity is a function of the composition of the hydrogen isotopes, depending on the conditions of the test as well as sensitive to the geometry and microstructure of the test specimen. Thus, the intrinsic diffusivity of the material (D) cannot be measured directly when tritium is being trapped. Equation [13] is the general form of a simplified expression that is commonly used in the literature:
Fractional coverage of traps q T
1
0.8 Et = 50 kJ mol–1 0.6
0.4
0.2
0 200
‘Low’ solubility ‘High’ solubility Et = 10 kJ mol–1
300
400
500
600
700
800
Temperature (K) Figure 1 Fraction of filled traps as a function of temperature for ‘low-solubility’ and ‘high-solubility’ materials (modeled as reduced activation ferritic/martensitic steel and austenitic stainless steel, respectively, using relationships from Table 1). The pressure is 0.1 MPa, the molar volume of the steels is approximated as 7 cm3 mol1 and there is assumed to be one lattice site for hydrogen per metal atom.
516
Tritium Barriers and Tritium Diffusion in Fusion Reactors
Deff ¼
D nT Et exp 1þ nL RT
½14
Equation [14] does not account for the effect of lattice concentration, and is therefore inadequate when the concentration of tritium is relatively large. For materials with high solubilities of tritium (such as austenitic stainless steels), trapping may not affect transport significantly and Deff D as shown in Figure 2. For materials with a low solubility and relatively large Et, the effective diffusivity can be substantially reduced compared to the lattice diffusivity (Figure 2). The wide variation of reported diffusivity of hydrogen in a-iron at low temperatures is a classic example of the effect of trapping on hydrogen transport2,20: while the diffusivity of hydrogen at high temperatures is consistent between studies, the effective diffusivity measured at low temperatures is significantly lower (in some cases by orders of magnitude) compared to the Arrhenius relationship established from measurements at elevated temperatures. Moreover, the range of reported values of effective diffusivity demonstrates the sensitivity of the measurements to experimental technique and test conditions. For these reasons, it is important to be critical of diffusion data that may be affected by trapping and be cautious of extrapolating diffusion data to experimental conditions and temperatures
different from those measured, especially if trapping is not well characterized or the role of trapping is not known. 4.16.2.5
Permeability
Permeability of hydrogen and its isotopes is generally defined as the steady-state diffusional transport of atoms through a material that supports a differential pressure of the hydrogen isotope. Assuming steady state, semi-infinite plate, and Fick’s first law for diffusion J ¼ Dðdc=dxÞ, we can express the steadystate diffusional flux of tritium as ðcx cx1 Þ ½15 J1 ¼ D 2 x2 x1 where cx is the concentration at position x within the thickness of the plate. Using chemical equilibrium (eqn [7]) and assuming that the tritium partial pressure is negligible on one side of the plate of thickness t, the steady-state diffusional flux can be expressed as DK 1=2 p t TT and the permeability, F, is defined as: J1 ¼
F DK
½16
½17
Substituting eqns [2] and [5] into eqn [17], the permeability can be expressed as a function of temperature in the usual manner:
100 nT = 10–7
nT = 10–5
10–1 Deff / D
nT = 10–3
10–2
10–3 200
300
400
500 Temperature (K)
600
700
800
Figure 2 Ratio of effective diffusivity to lattice diffusivity (Deff/D) as a function of temperature for ‘low-solubility’ material (squares with varying nT, modeled as reduced activation ferritic/martensitic steel with Et ¼ 50 kJ mol1) and ‘high-solubility’ material (triangles, modeled as austenitic stainless steel with Et ¼ 10 kJ mol1 and nT ¼ 103 traps per metal atom). The pressure is 0.1 MPa, the molar volume of the steels is approximated as 7 cm3 mol1 and there is assumed to be one lattice site for hydrogen per metal atom.
Tritium Barriers and Tritium Diffusion in Fusion Reactors
F ¼ K0 D0 exp½ðDHs þ ED Þ=RT
½18
Permeability is a material property that characterizes diffusional transport through a bulk material, that is, it is a relative measure of the transport of tritium when diffusion-limited transport dominates; see LeClaire21 for an extensive discussion of permeation. By definition, the permeability (as well as diffusivity and solubility) of hydrogen isotopes through metals is independent of surface condition, since it is related to diffusion of hydrogen through the material lattice (diffusivity) and the thermodynamic equilibrium between the gas and the metal (solubility). In practice, experimental measurements are strongly influenced by surface condition, such that the measured transport properties may not reflect diffusion-limited transport. Under some conditions (such as low pressure or due to the presence of residual oxygen/moisture in the measurement system), the theoretical proportionality between the square root of pressure and hydrogen isotope flux does not describe the transport;21,22 thus, studies that do not verify diffusion-limited transport should be viewed critically. In particular, determination of the diffusivity of hydrogen and its isotopes is particularly influenced by the surface condition of the specimen, since diffusivity is determined from transient measurements. While permeation measurements (being steady-state measurements) are relatively less sensitive to experimental details, the quality of reported solubility relationships depends directly on the quality of diffusion, since solubility is typically determined from the measured permeability and diffusivity.1 In addition, trapping affects diffusivity and must, therefore, be mitigated in order to produce solubility relationships that reflect the lattice dissolution of hydrogen and its isotopes in the metal. These characteristics of the actual measurements explain the fidelity of permeation measurements between studies in comparison with the much larger variation in the reported diffusivity and solubility. 4.16.2.6
Recombination
As shown earlier, steady-state permeation of hydrogen through materials is normally governed only by solubility and diffusivity. It has been shown23 that at low pressures, permeation can also be limited by dissociation at the surface. Due to limited data in the literature on this effect (and questions about whether this condition ever really exists), we do not
517
consider this effect in this chapter. It is also possible for permeation to be limited by the rate at which atoms can recombine back into molecules. With the exception of extremely high temperatures, this recombination is necessary for hydrogen to be released from a material. Wherever the release rate from a surface is limited by recombination, the boundary condition at that boundary is given by: Jr ¼ kr c 2
½19
where kr is the recombination-rate constant and c is the concentration of hydrogen near the surface (for this discussion, we assume that there is no surface roughness). The units for k and c are m4 s1 per mol of H2 and mol H2 m3, respectively. There are two specific types of conditions that can lead to the hydrogen release being rate limited by recombination. One of them occurs for plasmafacing materials in which the recombination-rate coefficient is relatively low, and the implantation rate is high. With this condition, the concentration of hydrogen in the very-near plasma-exposed surface will increase to the point at which Jr is effectively equal to the implantation rate. It is not exactly equal to the implantation rate because there is permeation away from that surface to the downstream surface. The other condition that can lead the hydrogen release being controlled by recombination is when the rate of ingress at the upstream boundary is very low. This condition can occur either when the upstream pressure is extremely low or a barrier is placed on the upstream surface, and the downstream surface has a relatively low recombination-rate constant. In the extreme case, the release rate at the downstream side is so slow that the hydrogen concentration becomes uniform throughout the material. The release rate from the downstream surface will be krc2, where c now represents the uniform concentra1=2 tion. From c ¼ KpTT and eqn [19], it can easily be shown that the recombination-limited permeation is linearly dependent on pressure, rather than having the square root of pressure dependence of diffusionlimited permeation. There are various derivations and definitions of the recombination-rate constant. In the case of intense plasma exposure in which extreme near-surface concentrations are generated, Baskes24 derived the recombination-rate constant with the assumption that the rate was controlled by the process of bulk atoms jumping to the surface, combining with surface atoms, and then desorbing. His expression for the recombination-rate constant is
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Tritium Barriers and Tritium Diffusion in Fusion Reactors
kr ¼ C
8 mT
1=2
s 2DHs EX exp RT K02
½20
where C is a constant, s is the sticking constant, which depends on the cleanliness of the material surface, and EX ¼ DHs þ ED > 0, otherwise EX ¼ 0. The sticking constant can be anywhere from 1 for clean surfaces to 104 or smaller for oxidized surfaces. Pick and Sonnenberg25 solved the recombinationrate constant for the case where the near-surface concentration of hydrogen is small. In the limit of low surface concentration, the rate of atom jump to the surface does not play an important role in the recombination rate, thus eliminating EX from the exponential. The sticking constant in the Pick and Sonnenberg model is thermally activated: s ¼ s0 expð2EC =RT Þ, where s0 is the sticking coefficient and EC is the activation energy for hydrogen adsorption. Wampler26 also studied the case of low nearsurface concentration to arrive at an expression for the recombination-rate constant. He assumed equilibrium between hydrogen atoms in surface chemisorption sites and atoms in solution, deriving the recombination-rate constant as ns n 2DHs ½21 exp kr ¼ RT ðbnL Þ2 where ns is the area density of surface chemisorption sites, and n is the jump frequency. These expressions differ, but also display many similarities. Unfortunately, the surface cleanliness dominates the rate of recombination and these theoretical relationships are relevant only for sputtercleaned surfaces and very low pressures. For example, Causey and Baskes27 showed that the Baskes24 model predicts fairly accurate results for plasma-driven permeation of deuterium in nickel. Comparison with values in the literature for nickel showed other results to differ by as much as four orders of magnitude and to have significantly different activation energies. 4.16.2.7
Irradiation and Implantation
Irradiation and implantation can affect the transport of hydrogen isotopes in materials. Since these effects can be complex and depend on the conditions of the materials and the environment, it can be difficult to draw broad conclusions from the literature. Nevertheless, changes in apparent transport properties are generally attributed to damage and the creation of hydrogen traps28–31 (see also Chapter 1.03, Radiation-Induced Effects on Microstructure). Therefore, the effects of
irradiation and implantation will depend sensitively on the characteristics of the traps that are created by these processes. The density of damage is an important consideration: for example, it has been shown that helium bubbles are not effective trapping sites for steels,32 likely because in these experiments, the density of helium bubbles was relatively low. The energy of the trap will determine the coverage as a function of temperature (eqn [9]): generally, the effect of trapping will be stronger at low temperatures, especially in materials with a low solubility (Figure 2), which can result in substantial increases in hydrogen isotope inventory compared to hydrogen content predicted from lattice solubility. Additionally, irradiation may increase ionization of hydrogen isotopes, thus enhancing apparent permeation.29 Reactor environments can defeat permeation barriers, for example, by damaging the integrity of oxide layers; this is discussed at the end of this chapter.
4.16.3 Fusion Reactor Materials 4.16.3.1
Plasma-Facing Materials
Tritium generated in the fusion-reactor blanket will be fed directly into the plasma in the main vacuum chamber. There, the tritium will be partially consumed, but it will also interact with the materials composing the first wall. Materials used to line the first wall will be exposed to energetic tritium and deuterium escaping from the plasma. Particle fluxes in the range of 1021 (D þ T) m2 s1 will continuously bombard the plasma-facing materials. Materials used for the divertor at the top and/or bottom of the torus will be exposed to lower energy particles with a flux of 1023 (D þ T) m2 s1 or higher. While the neutral gas pressure of tritium will be relatively low at the outer vacuum wall boundary, some minor permeation losses will occur. In reality, the primary concerns in the plasma-facing areas are tritium inventory and permeation into the coolant through coolant tubes. While future power reactors are likely to have primarily refractory metals such as tungsten, present-day devices are still using carbon and beryllium. In this section on plasma-facing materials, we examine the interaction of tritium with carbon, tungsten, and beryllium. 4.16.3.1.1 Carbon
In many ways, carbon is ideal for fusion applications. It is a low-Z material with a low vapor pressure and excellent thermal properties. The carbon used in fusion applications comes in two forms, graphite and carbon
Tritium Barriers and Tritium Diffusion in Fusion Reactors
composites. Graphite is described in Chapter 2.10, Graphite: Properties and Characteristics; Chapter 4.10, Radiation Effects in Graphite, and Chapter 4.18, Carbon as a Fusion Plasma-Facing Material and is typically made using the Acheson process.33 Calcined coke is crushed, milled, and then sized. The properties of the graphite are determined by the size and shape of these particles. Coal tar is added to the particles and the batch is heated to 1200 K. This process is repeated several times to increase the density of the compact. The final bake is at temperatures between 2900 and 3300 K and takes 15 days. The final product is quite porous, with a density of around 1.8–1.9 g cm3 (compared to a theoretical density of 2.3 g cm3). Graphite is composed of grains (from the original coke particles) with a size of 5–50 mm, which are in turn composed of graphite subgrains with a typical size of 5 nm. Carbon composites are made by pyrolyzing a composite of carbon fibers in an organic matrix. These fibers have a high strength-to-weight ratio and are composed of almost pure carbon. As with graphite, carbon composites are quite porous with a density of <2 g cm3. Carbon (as either graphite or carbon composite) will not serve as a vacuum or coolant boundary. Therefore, permeation through the carbon will not directly affect tritium release into the environment. Coolant tubes inside the graphite will control the tritium release to the coolant system. The ways in which hydrogen isotopes can interact with carbon are explored in the following sections. It is difficult to consider hydrogen isotope permeation in graphite in the same manner as one would for the metals. As was shown by Kiyoshi et al.,34 hydrogen passes rather readily through graphite in the molecular form. It is when one considers the uptake or retention of atomic hydrogen in graphite or a carbon composite that it becomes more interesting. Most graphites have a Brunauer–Emmett–Teller (BET) or specific surface area of 0.25–1.0 m2 g1.35 This presents a lot of surface area for the absorption of hydrogen isotopes. Barrer,36 Thomas,37 and Bansal38 are only a few of the many who have reported hydrogen retention on carbon surfaces. Activation energies varying from 24 to 210 kJ mol1 have been reported. Once absorbed on a carbon surface, the hydrogen can migrate by jumping from one chemically active site to another. In a fusion reactor, this process occurs in carbon at lower temperatures at which the plasma provides ionic and atomic hydrogen for direct absorption. Molecular hydrogen can dissociate on carbon surfaces, but only at elevated temperatures (>1000 K). As a plasma-facing material, graphite will be exposed
519
to atomic tritium and deuterium, and these hydrogen isotopes will migrate inward along the open porosity. Several research groups have measured the diffusion coefficient for hydrogen on carbon surfaces. Robell et al.39 inferred an activation energy of 164 kJ mol1 for the diffusion during measurements on the uptake of hydrogen on platinized carbon between 573 and 665 K. Olander and Balooch40 used similar experiments to determine the diffusion coefficient for hydrogen on both the basal and prism plane: 6 109 exp (7790/T) m2 s1 for the basal plane and 6 1011 exp(4420/T) m2 s1 for the prism plane. Causey et al.41 used tritium profiles in POCO AXF-5Q graphite exposed to a tritium plasma to extract a diffusivity for tritium on carbon pores of 1.2 104 exp(11 670/T) m2 s1. An example of the deep penetration of hydrogen isotopes into the porosity of graphite was reported by Penzhorn et al.42 Graphite and carbon composite tiles removed from the Joint European Torus (JET) fusion reactor were mechanically sectioned. The sections were then oxidized, and tritiated water was collected for liquid scintillation counting. Relatively high concentrations of tritium were detected tens of millimeters deep into the tiles. The diffusion or migration of hydrogen isotopes on carbon surfaces occurs at lower temperatures. The solubility, diffusion, and trapping of hydrogen in the carbon grains are higher temperature processes than adsorption and surface diffusion. At higher temperatures (>1000 K), hydrogen molecules can dissociate and be absorbed at chemically active sites on carbon surfaces. Some of these sites are located on the outside of the grains, but many exist along the edges of the subgrains that make up the larger grains. Hydrogen isotopes dissociating on the outer grain boundary are able to migrate along the subgrain boundaries, entering into the interior of the grain. It is the jumping from one moderate energy site (240 kJ mol1) to another that determines the effective diffusion coefficient. Traps on the grain boundaries pose a binding energy barrier (175 kJ mol1) that must be overcome in addition to this normal lattice activation. Atsumi et al.43 used the pressure change in a constant volume to determine the solubility of deuterium in ISO 88 graphite. They determined the solubility to be given by K ¼ 18.9 exp (þ2320/T) mol H2 m3 MPa1/2 over the temperature range of 1123–1323 K. This solubility is shown in Figure 3 along with two data points by Causey44 at 1273 and 1473 K. A negative heat of solution is seen in both sets of data, suggesting the formation of a bond between hydrogen and carbon.
520
Tritium Barriers and Tritium Diffusion in Fusion Reactors
Solubility (mol m–3 MPa–1/2)
1000
Atsumi et al.
100
Causey et al. 10 0.65
0.7
0.8 0.75 Temperature, 1000/T (K–1)
0.85
0.9
Figure 3 Solubility of hydrogen in carbon. Adapted from Atsumi, H.; Tokura, S.; Miyake, M. J. Nucl. Mater. 1988, 155, 241–245; Causey, R. A. J. Nucl. Mater. 1989, 162, 151–161.
10–11 Causey (best estimate)
Diffusivity (m2 s–1)
10–13
Rohrig et al.
10–15
10–17
Atsumi et al.
10–19 Malka et al.
Causey et al. 10–21
0.4
0.6
0.8
1
1.2
1.4
Temperature, 1000/T (K–1) Figure 4 Diffusivity of hydrogen in graphite. The bold line is the best estimate given in Causey44 based on uptake data for tritium in POCO AFX-5Q graphite. Adapted from Atsumi, H.; Tokura, S.; Miyake, M. J. Nucl. Mater. 1988, 155, 241–245; Causey, R. A. J. Nucl. Mater. 1989, 162, 151–161; Ro¨hrig, H. D.; Fischer, P. G.; Hecker, R. J. Am. Ceram. Soc. 1976, 59, 316–320; Causey, R. A.; Elleman, T. S.; Verghese, K. Carbon 1979, 17, 323–328; Malka, V.; Ro¨hrig, H. D.; Hecker, R. Int. J. Appl. Radiat. Isot. 1980, 31, 469.
The variation in the diffusivities of hydrogen in graphite determined by various researchers is extreme. This variation results primarily from differences in interpretation of the mechanism of diffusion (e.g., bulk diffusion or grain boundary diffusion). Representative values for the diffusion is shown in Figure 4. Ro¨hrig et al.45 determined their diffusion coefficient using the release rate of tritium from nuclear grade graphite during isothermal anneals. They correctly used the grain size as the real diffusion
distance. Causey et al.46 measured the release rate of tritium recoil injected into pyrolytic carbon to determine the diffusivity. Malka et al.47 used the release rate of lithium-bred tritium in a nuclear graphite to determine the diffusion coefficient. Atsumi et al.43 used the desorption rate of deuterium gas from graphite samples that had been exposed to gas at elevated temperatures to determine a diffusivity. Building on the work of others, Causey44 proposed an alternative expression for the diffusivity that he labeled ‘best estimate.’
Tritium Barriers and Tritium Diffusion in Fusion Reactors
The result was based on uptake experiments for tritium into POCO AFX-5Q graphite, and the assumption that the total uptake is determined by the product of the diffusivity and the solubility. The uptake data were analyzed assuming that the expression given by Atsumi et al.43 for the solubility was correct. The expression ‘best estimate’ was used because it properly took into consideration that the grain size was the effective diffusion distance, and that diffusivity was more properly determined by uptake than release. Release rates are strongly affected by trapping. The trapping of hydrogen isotopes at natural and radiation-induced traps has been examined by several research groups.41,48–50 Causey et al.41 exposed POCO-AFX-5Q graphite to a deuterium/tritium mixture at elevated temperatures in an examination of the kinetics of hydrogen uptake. For temperatures above 1500 K, it was discovered that increasing the pressure did not increase the retention. It appeared that the solubility hit an upper limit at 17 appm. Analysis of the data revealed that the 17 appm did not represent solubility, but a trap density. The trap was determined to have a binding energy of 175 kJ mol1. Atsumi et al.48 found that radiation damage increased the apparent solubility of deuterium in graphite by a factor of 20–50 with saturation at a radiation damage level of 0.3 dpa. A significant decrease in the apparent diffusivity was also noted. In a later study, Atsumi et al.49 reported graphites and composites to vary significantly in their natural retention values. They saw that the saturation retention was inversely proportional to the lattice constant (which relates to the degree of graphitization, and so grain size). Radiation damage was seen to decrease the apparent lattice constant and increase the saturation retention. 6 MeV Cþ ions were used by Wampler et al.50 to simulate neutron damage to different graphites. The trap density increased with damage levels up to 0.04 dpa at which the saturation retention was 650 appm. In a defining set of experiments, Causey et al.28 examined tritium uptake in unirradiated and radiated pitch-based carbon composites. Pitch-based carbon composites have a large lattice parameter due to the sheet-like configuration of the fibers. These composites retained significantly less tritium before and after irradiation than other carbon materials. From these results, it is apparent that highenergy trapping occurs at the edges of the hexagonal crystals on the prism plane. Hydrogen isotope permeation in the normal sense does not apply to graphite and carbon composites. There is inward migration of atomic hydrogen isotopes along porosity at lower temperatures. To calculate
521
the potential tritium inventory for this process, one can obtain an upper bound by assuming monolayer coverage of the pore surfaces. Typical nuclear grade graphite has a specific surface area of 1 m2 g1. Complete loading of that amount of surface area yields 2 1025 T m3, or about 2 g for a 20 m2 carbon wall that is 10-mm thick. At higher temperatures, molecular hydrogen isotopes, which are moving through the graphite pore system, are able to dissociate and enter the multimicron-sized graphite grains. This migration into the grains occurs along the edges of the nanometer scale subgrains. As the hydrogen migrates inward, it decorates high-energy trap sites. The density of these trap sites is higher than the effective solubility derived from the migration rate. Permeation into graphite will not lead to tritium release from a fusion device, but will affect tritium inventory. If one assumes that radiation damage from neutrons has increased the concentration of traps with a binding energy of 175 kJ mol1 up to 1000 appm, and that tritium is occupying 100% of those traps, the same 20 m2 wall 10-mm thick listed above would now contain an additional 100 g of tritium. Occupation of all of the traps is difficult to achieve: at low temperatures, kinetics makes it impossible to achieve saturation, while at substantially higher temperatures, the traps do not remain filled. There is another process called carbon codeposition that can strongly affect tritium inventory in a fusion device. In the codeposition process, carbon eroded from the walls of the tokamak is redeposited in cooler areas along with deuterium and tritium from the plasma. Because carbon codeposition is not a diffusion or permeation process, it will not be covered in this review. The interested reader is referred to a review of this process by Jacob.51 4.16.3.1.2 Tungsten
Tungsten is another of the plasma-facing materials, described in Chapter 4.17, Tungsten as a PlasmaFacing Material. Like carbon, it will not be a vacuum barrier. Thus, permeation through the tungsten will not lead to tritium release directly into the environment. It can lead to tritium permeation into the coolant through the coolant tubes inside the tungsten facing materials. Permeation will also affect the tritium inventory of the fusion device. Tungsten has excellent thermal properties with a very high melting point of 3683 K. The problem that tungsten presents to the tokamak designer is the deleterious radiation losses if tungsten is present in the plasma. Fortunately, the energy threshold for sputtering by
522
Tritium Barriers and Tritium Diffusion in Fusion Reactors
hydrogen ions is quite high, 700 eV for tritium.52 For that reason, tungsten will be used primarily in the divertor region where the energy of the impacting particles can be limited. There are a limited number of reports on the diffusivity of hydrogen isotopes in tungsten. Frauenfelder53 measured the rate of hydrogen outgassing from saturated rolled sheet samples at temperatures over the wide range 1200–2400 K. His material was 99.95% pure tungsten. Zakahrov and Sharapov54 used 99.99% pure tungsten samples in their permeation techniques to determine the hydrogen diffusivity over a limited temperature range of 900–1060 K. In the Benamati et al.55 experiments using tungsten containing 5% rhenium, a gaseous permeation technique was also used. These experiments were performed over a very limited temperature range of 850–885 K. Reported diffusivities are shown in Figure 5. There are a couple of reasons why the diffusion coefficient reported by Frauenfelder53 is widely accepted as most correct. The first of these reasons is the wide temperature range over which experiments were performed. The second reason is that the experiments were performed at a temperature above that where trapping typically occurs. It can be seen in Figure 5 that Zakahrov’s54 diffusivity agrees quite well with Frauenfelder’s at the highest temperatures, but falls below his values at lower temperatures, where trapping would occur. The database on hydrogen solubility in tungsten is also limited. The results of the two experimental
studies are shown in Figure 6. In the same experiments used to determine the diffusivity, Frauenfelder53 also measured solubility. Over the temperature range 1100–2400 K, samples were saturated at fixed pressures and then heated to drive out all of the hydrogen. Over a more limited temperature range of 1900–2400 K, Mazayev et al.56 also examined hydrogen solubility in tungsten. The agreement with the Frauenfelder’s53 data is quite good in magnitude, but not good in apparent activation energy. As with his diffusivity, the solubility reported by Frauenfelder is typically the value used in predicting the migration of hydrogen in tungsten. Hydrogen trapping in tungsten has been studied by several research groups. van Veen et al.57 used bombardment by 2 keV protons in their study of the bonding of hydrogen to voids in single-crystal tungsten. Thermal desorption from the samples with appms of voids revealed a broad release peak at 600–700 K. It was stated that the release could be modeled as gas going back into solution from the voids with a trap binding energy of 96.5–135 kJ mol1 controlling the process. Eleveld and van Veen,58 in a similar study, used a lower fluence of 30 keV Dþ ions in desorption experiments. In these samples containing vacancies but no voids, the release occurred at 500–550 K. The authors reported a value of 100 kJ mol1 for the trap binding energy of vacancies. Pisarev et al.59 used lower fluences of 7.5 keV deuterons into 99.94% pure tungsten samples. During thermal desorption ramps, peaks in the release rates were seen at 350, 480, 600, and 750 K. The release at the highest temperature was seen only in the highest
10–7
Diffusivity (m2 s–1)
10–8
Frauenfelder Zakharov et al.
10–9
10–10 Benamati et al.
10–11 0.4
0.6
0.8
1
1.2
1.4
Temperature, 1000/T (K–1) Figure 5 Diffusivity of hydrogen in W. Adapted from Frauenfelder, R. J. Vac. Sci. Technol. 1969, 6, 388–397; Zakahrov, A. P.; Sharapov, V. M. Fiziko-Khimicheskaya Mekhanika Materialov 1973, 9, 29–33; Benamati, G.; Serra, E.; Wu, C. H. J. Nucl. Mater. 2000, 283–287, 1033–1037.
Tritium Barriers and Tritium Diffusion in Fusion Reactors
523
Solubility (mol m−3 MPa−1/2)
10
1
Mazayev et al.
Frauenfelder
0.1
0.01 0.4
0.5
0.6
0.7
0.8
0.9
1
Temperature, 1000/T (K–1) Figure 6 Solubility of hydrogen in W. Adapted from Frauenfelder, R. J. Vac. Sci. Technol. 1969, 6, 388–397; Mazayev, A. A.; Avarbe, R. G.; Vilk, Y. N. Russian Metallurgy-Metally-USSR 1968, 6, 153–158.
fluences. Garcı´a-Rosales et al.60 used 100 eV deuterium implantation to study the trapping and release rate of hydrogen isotopes from wrought and plasma-sprayed tungsten. Two broad desorption peaks at 475–612 K and 670–850 K were seen in the thermal desorption spectra. Modeling of the release data suggested the lower temperature peak to be controlled by both diffusion and trapping at a binding energy of 44 kJ mol1. The second release peak was reported to correspond to trapping at defects with a binding energy of 97 kJ mol1. In experiments with 99.99% pure tungsten and tungsten with 1% lanthanum oxide, Causey et al.61 examined tritium retention in plasma-exposed samples. Modeling of the results suggested two traps, one with a binding energy of 97 kJ mol1 and another with 204 kJ mol1. The density of the trapped tritium averaged 400–500 appm. Anderl et al.62 used deuterium implantation into polycrystalline tungsten to determine the correlation between dislocation density on cell walls and deuterium trapping. Annealing tungsten at 1673 K reduced the dislocation density by a factor of 7, subsequently reducing the deuterium trapping by a similar factor. The binding energy of these traps was estimated to be 88–107 kJ mol1. As-received 99.95% pure tungsten was used by Sze et al.63 in experiments with intense deuterium plasma exposure. Exposure at 400 K resulted in blisters with diameters of tens of microns. Elevating the temperature to 1250 K eliminated the blisters. Venhaus et al.64 used high-purity foils in experiments to examine the effect of annealing temperature on blistering by deuterium plasma exposure. An unannealed sample and one annealed at 1473
K both exhibited blisters after the plasma exposure. The sample annealed at 1273 K did not blister. There have been a multitude of other reports on blister formation on tungsten samples exposed to various forms of hydrogen implantation.65–68 Anderl et al.62 used 99.95% tungsten in 3 keV Dþ 3 ion implantation to determine the recombinationrate coefficient. Over a temperature range of 690– 825 K, the recombination rate coefficient was given as kr ¼ 3.85 109 exp(13 500/T )m4 s1 per mol of H2. This expression is shown in Figure 7, where it is plotted along with the expression given by the Baskes24 model. It can be seen that there is very little correlation between the measured Anderl value and the calculated Baskes value. This is not entirely unusual. Impurities on the surface, especially oxide layers, can have a very strong effect on this coefficient. While tungsten has excellent low permeability for gaseous tritium, it will be used only in fusion devices as a plasma-facing material. As a plasmafacing material, tungsten will be exposed to intense fluxes of energetic tritium and deuterium. With traps for hydrogen at binding energies of 97 and 203 kJ mol1(57–62) at natural and radiation-induced defects, it would appear that a substantial tritium inventory could be generated in divertor tungsten. There are several reasons why this high inventory is not likely to occur. The first reason is the high recombinationrate coefficient given earlier. For a recombinationrate constant of 101 m4 s1 per mol of H2 or higher (see Figure 7), the recombination rate on the surface is so rapid as to generate the equivalent of c ¼ 0 at
Tritium Barriers and Tritium Diffusion in Fusion Reactors
Recombination-rate coefficient (m4 s–1 per mol H2)
524
107 106
Baskes model
105 104 1000 Anderl et al. 100 10 1 0.8
1
1.2
1.4
1.6
Temperature, 1000/T
1.8
2
2.2
(K–1)
Figure 7 Recombination-rate coefficient of hydrogen in W. Adapted from Baskes, M. I. J. Nucl. Mater. 1980, 92, 318–324; Anderl, R.; Holland, D. F.; Longhurst, G. R.; et al. Fusion Technol. 1991, 21, 745–752.
the boundary. With the very limited penetration distance of energetic hydrogen in the dense tungsten, most of the implanted material is immediately released back out of the surface. There are also recent reports suggesting that ruptured blisters and very fine cracks near the surface69–71 will even further reduce the inward migration of deuterium and tritium into the tungsten. 4.16.3.1.3 Beryllium
Beryllium is a low-Z material with good thermal characteristics, described in Chapter 2.11, Neutron Reflector Materials (Be, Hydrides) and Chapter 4.19, Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices. Additionally, it is a good getter for oxygen impurities in the plasma. The low-Z minimizes the radiation losses from the plasma, and the oxygen removal keeps the plasma clean. For these reasons, beryllium has been used in the JET fusion reactor and will be the first wall material for the International Thermonuclear Experimental Reactor (ITER). Beryllium has interesting hydrogen retention behavior. Beryllium may also be used as a neutron multiplier in the blanket area of future fusion devices to increase the tritium breeding ratio. Abramov et al.72 used two grades of beryllium in their permeation–diffusion experiments. These were high-purity (99%) and extra grade (99.8%). Adding to the validity of their experimental result was the fact that the authors used multilayer permeation theory analysis to take permeation through the outer oxide layer into consideration. For a lower
purity material (98%), Tazhibaeva et al.73 also used the multilayer permeation analysis to determine diffusivity. Jones and Gibson74 studied tritium diffusivity and solubility for arc-cast beryllium in the temperature range of 673–1173 K. Beryllium was exposed to tritium gas for various temperatures, durations, and pressures during isothermal anneals. After removing the samples to another experimental system, the samples were heated to various temperatures. For the initial heating, the tritium release would rise, but soon fall to zero. Elevating the temperature would reestablish the tritium release, but again the release would fall. While this behavior is not typical of diffusion controlled release, the data were analyzed to extract an effective diffusivity. The different reported diffusivities are shown in Figure 8. It can be seen that the diffusivity reported by Abramov et al.72 is considerably larger than those of Tazhibaeva et al.73 and Jones and Gibson.74 It is apparent that the purity of the beryllium played a strong role in determining the effective diffusivity. Oxygen, the primary impurity in beryllium provides a strong trap for hydrogen. Thompson and Macaulay-Newcombe75,76 examined the diffusion of deuterium in single-crystal and polycrystalline beryllium. The effective diffusivity in the single-crystal material was lower than that for the polycrystalline material. The polycrystalline results agreed quite well with the results reported by Abramov et al.72 They suggested that the lower diffusivity seen for the single-crystal samples was the true diffusivity for beryllium, and that the polycrystalline results represented diffusion along the grain boundaries.
Tritium Barriers and Tritium Diffusion in Fusion Reactors
525
10–10
Diffusivity (m2 s–1)
Abramov et al. Extra grade High grade 10–11
10–12
Jones and Gibson Tazhibaeva et al.
10–13 0.8
1
1.2
1.4
1.6
1.8
Temperature, 1000/T (K–1) Figure 8 Diffusivity of hydrogen in beryllium. ‘Extra’ grade is 99.8% pure and ‘high’ is 99.0% pure. Other authors did not specify purities. Adapted from Abramov, E. I. L.; Riehm, M. V. P.; Thompson, D. A.; et al. J. Nucl. Mater. 1990, 175, 95427– 95430; Tazhibaeva, I. L.; Shestakov, V. P.; Chikhray, Y. V. In Proceedings of the 18th Symposium of Fusion Technology; Elsevier: Karlsruhe, Germany, 1990; pp 427–430; Jones, P. M. S.; Gibson, R. J. Nucl. Mater. 1967, 21, 353–354.
10
Solubility (mol m–3 MPa–1/2)
Shapovalov and Dukel’skii Jones and Gibson
1 Swansiger
0.1 0.6
0.8
1
1.2
1.4
1.6
1.8
2
Temperature, 1000/ T (K–1) Figure 9 Solubility of hydrogen in beryllium. Adapted from Jones, P. M. S.; Gibson, R. J. Nucl. Mater. 1967, 21, 353–354; Shapovalov, V. I.; Dukel’skii, Y. M. Izvestiva Akademii Nauk SSR Metally 5, 201–202; Swansiger, W. A. J. Vac. Sci. Technol. A 1986, 4, 1216–1217.
If hydrogen isotopes migrate along the grain boundaries, it is logical that the rate of migration would be affected by oxygen segregated to those boundaries. The very limited results for hydrogen isotope solubility in beryllium are shown in Figure 9. In the earlier described experiments by Jones and Gibson,74 the solubility was seen to be effectively independent of temperature in the temperature range 550–1250 K. For sintered, distilled a-beryllium, Shapovalov and Dukel’skii77 reported similar values
of solubility for the temperature range 673–1473 K. In experiments using 98.5% and 99.8% pure beryllium samples, Swansiger78 used gaseous uptake of tritium to determine the solubility. The amount of tritium uptake did not increase with increasing sample size. The solubility for the two purity materials was also seen to be the same. For temperatures below 650 K, the apparent solubility increased; this strange effect was attributed to trapping. It is interesting to note the fact that the apparent
526
Tritium Barriers and Tritium Diffusion in Fusion Reactors
solubility over temperatures at which the three research groups74,77,78 performed their experiments varied by less than one order of magnitude even though the activation energies varied by 96 kJ mol1. It should be questioned whether the reported values really represent the solubility of hydrogen isotopes in bulk beryllium. For all plasma-facing materials, there is concern that implantation of energetic deuterium and tritium could lead to excessive retention and permeation. Implantation of hydrogen isotopes into a material with a low recombination-rate constant can lead to a majority of the hydrogen being pushed into the bulk of the material. In the limiting case of slow recombination, 50% of the hydrogen exits the front face and 50% exits the rear face. Langley79 implanted 25 keV deuterium into 99.1% pure hot isostatic pressed beryllium. The retention was seen to be 100% until the particle fluence reached 2 1022 D m2. The retention flattened to a limit of 2.8 1022 D m2. Wampler80 recorded similar results for his implantation of 0.5 and 1.5 keV deuterium into 99.6% pure beryllium samples. Saturation occurred at 0.31 D/Be in the implant zone. Yoshida et al.81 used 99% pure beryllium in his implantation experiments with 8 keV deuterons. Transmission electron microscopy revealed bubble formation at all temperatures between room temperature and 873 K. The bubbles were not removed even by annealing at temperatures up to 973 K. Plasma exposure was used by Causey et al.82 and by Doerner et al.83 in low-energy, high-fluence deuterium exposures to beryllium. In both studies, the fractional retention was extremely low and decreased with increasing temperature. Open porosity in the implant zone was listed as the likely cause of the low retention. Chernikov et al.84 and Alimov et al.85 showed bubbles and microchannels to be responsible for the behavior of implanted hydrogen in beryllium. At 300 K, very small bubbles with a high volume density are formed even at low fluences. As the fluence is increased, the bubbles agglomerate into larger bubbles and then form microchannels that eventually intersect with the surface. For irradiation at 500–700 K, small facetted bubbles and large oblate, gas-filled cavities are formed. This microstructure was seen to extend well beyond the implant zone. Alimov et al.85 postulated that the hydrogen retention in the porous region was due to binding to the beryllium oxide that forms on the pore surfaces. Beryllium is known as a neutron multiplier because of the reaction 9Be þ n ! 8Be þ 2n. Another neutron reaction for beryllium is 9Be þ n ! 4He þ 6He,
followed by 6He decaying to 6Li. 6Li has a very large cross-section to absorb a thermal neutron and produce a helium atom and a tritium atom. Baldwin and Billone86 calculated the amount of tritium that could be produced in a large fusion device of the future. In an experiment, they exposed beryllium to a neutron fluence of 5 1026 n m2 with 6% of the neutrons having energy >1 MeV. The resulting tritium level was determined to be 2530 appm. Scaling up to a fusion reactor with 50 Mg of beryllium exposed to 3 MWy m2 results in the production of 5.5 kg of tritium. This is a sizeable quantity of tritium. The relevant question is whether this tritium would be released during normal operation of the fusion plant. Baldwin and Billone86 examined exactly that question in their experiment. The samples containing the 2530 appm of tritium were heated in stepped anneals to determine the release rate of tritium from beryllium materials with different densities. The annealing began with a very long anneal at 773 K, and the temperature was increased in increments of 100 K. For each temperature, there was a nondiffusional burst of release followed by a rapid decrease in the release rate. The release behavior for the different materials was similar, but the fractional release was greater for the less dense materials. Andreev et al.87 irradiated hot-pressed beryllium at 373 K. After neutron irradiation, thermal desorption spectroscopy was used with a heating rate of 10 K s1. Release began to occur at 773 K. The temperature at which maximum release occurred depended on the neutron fluence. The sample irradiated to a fluence of 3 1025 n m2 had a peak release at 1080 K, while the sample irradiated to the higher fluence of 1 1026 n m2 exhibited a peak release at a lower temperature of 1030 K. The authors examined the microstructure of the samples after the release anneals. If the anneal was stopped at 973 K, pores with a diameter of 2–16 mm were formed. If the anneal was taken to 1373 K, the pore diameters increased to 25–30 mm. Due to the toxicity of beryllium, there have been relatively small numbers of experiments performed on the behavior of hydrogen isotopes in beryllium. The apparent diffusion coefficient of hydrogen in beryllium is strongly affected by purity levels. The values determined for the solubility of hydrogen in beryllium all fall within one order of magnitude even though the apparent activation energy differs by 96 kJ mol1. Implantation of hydrogen into beryllium results in the formation of bubbles and eventually open channels or porosity. Connection of
Tritium Barriers and Tritium Diffusion in Fusion Reactors
the porosity to the surface facilitates the release of hydrogen from the beryllium as the particle fluence is increased. The tendency to form bubbles would suggest that the solubility of hydrogen in beryllium is extremely small. It is possible that the values determined for the solubility of hydrogen in beryllium actually represent the amount of hydrogen absorbed on the external surface and on the grain boundaries. The measured diffusivity may represent migration along the grain boundaries. More experiments, and experiments with single crystals, are needed to answer these questions. For beryllium used for long times in future fusion devices, tritium produced by neutron reactions on the beryllium is likely to dominate tritium retention in beryllium. Tritium inventory from eroded beryllium codeposited with tritium may play a strong role in tritium inventory, but that effect is not covered in this review. 4.16.3.2
Structural Materials
Structural materials for a fusion reactor are simply those that comprise a majority of the plant. They are not directly exposed to the plasma, but most are exposed to high doses of neutrons and electromagnetic radiation. Many of these materials are used in the reactor blanket where the tritium is bred by the nuclear reaction 6Li(n,a)3H. It is in the blanket and in the fuel reprocessing area that the structural materials are most likely to be exposed to tritium. The following sections review the structural materials that have been considered for fusion reactors. 4.16.3.2.1 Austenitic stainless steels
Austenitic stainless steels, particularly type 316, have been used extensively as a construction material for nuclear reactors (see Chapter 2.09, Properties of Austenitic Steels for Nuclear Reactor Applications). The type 300-series austenitic stainless steels (Fe–Cr–Ni) have relatively high nickel content (8–12 wt% for the 304 family of austenitic stainless steels and 10–14 wt% for 316 alloys), which is a detriment for fusion applications for several reasons including the susceptibility of nickel to activation (induced radioactivity).88–90 The solubility and the diffusivity of gaseous hydrogen and its isotopes through type 300-series austenitic stainless steels have been extensively studied and reviewed in San Marchi et al.1 Higher strength austenitic stainless steels (such as the Fe–Cr–Ni–Mn alloys, which have not been widely considered for fusion
527
applications) feature solubility and diffusivity that differ by a factor of about 2 compared to the type 300-series alloys.1 The so-called prime candidate alloy (PCA) is a variant of type 316 austenitic stainless steel modified for fusion applications (although interestingly enough with higher nickel content); from a permeation perspective PCA is anticipated to behave in a manner essentially similar to conventional type 316 alloys.91 The Fe–Cr–Mn austenitic stainless steels have been considered as a substitute for the more common grades of austenitic stainless steels since they have only a nominal nickel content,88,89,90 although low-activation ferritic/martensitic steels have received more attention (see subsequent section). Alloys that have been considered typically contain both chromium and Mn in the range 10–20 wt%, often with small amounts of other alloying elements (Sahin and Uebeyli90 provides a list of a number of alloys that have been explored for fusion applications). Unlike the Fe–Cr–Ni austenitic stainless steels, there are few reports of transport properties for the Fe–Cr–Mn austenitic alloys; data for oxidized Fe–16Cr–16Mn are reported in Gromov and Kovneristyi.92 Austenitic stainless steels can contain ferritic phases in the form of residual ferrite from alloy production, ferrite in welds formed during solidification, and in some cases, strain-induced martensite from deformation processing. The ferritic phases can result in a fast pathway for the transport of hydrogen and its isotopes at a relatively low temperature because the ferritic phases have a much higher diffusivity for hydrogen and its isotopes than austenite.93,94 In the absence of ferritic second phases, however, hydrogen transport in austenitic stainless steels is independent of whether the material is annealed or heavily cold-worked95–97 and relatively insensitive to composition for the type 300-series alloys.1 Reported values of hydrogen diffusivity in austenitic stainless steels are less consistent than permeability as a consequence of surface effects and trapping, as mentioned earlier and elsewhere.1 Figure 10 shows the reported diffusivity of hydrogen from a number of studies in which special precautions were taken to control surface conditions. The activation energy for diffusion is relatively large for austenitic stainless steels (ED ¼ 49.3 kJ mol1), and thus the diffusivity is sensitive to temperature, approaching the values of the ferritic steels at very high temperatures (>1000 K), while being many orders of magnitude lower at room temperature.
528
Tritium Barriers and Tritium Diffusion in Fusion Reactors
Diffusivity (m2 s–1)
10–9
10–10
10–11
10–12
1
1.2
1.4
1.6
1.8
2
Temperature, 1000/T (K–1) Figure 10 Diffusivity of hydrogen in austenitic stainless steels from gas permeation studies that confirmed diffusion-limited transport. The bold line represents the average relationship determined in Perng and Altstetter93 for several austenitic stainless steels. Adapted from Quick, N. R.; Johnson, H. H. Metall. Trans. 1979, 10A, 67–70; Gromov, A. I.; Kovneristyi, Y. K. Met. Sci. Heat Treat. 1980, 22, 321–324; Perng, T. P.; Altstetter, C. J. Acta Metall. 1986, 34, 1771–1781; Louthan, M. R.; Derrick, R. G. Corrosion Sci. 1975, 15, 565–577; Sun, X. K.; Xu, J.; Li, Y. Y. Mater. Sci. Eng. A 1989, 114, 179–187; Grant, D. M.; Cummings, D. L.; Blackburn, D. A. J. Nucl. Mater. 1987, 149, 180–191; Grant, D. M.; Cummings, D. L.; Blackburn, D. A. J. Nucl. Mater. 1988, 152, 139–145; Mitchell, D. J.; Edge, E. M. J. Appl. Phys. 1985, 57, 5226–5235; Kishimoto, N.; Tanabe, T.; Suzuki, T.; et al. J. Nucl. Mater. 1985, 127, 1–9.
The exceptionally low diffusivity of hydrogen near room temperature results in austenitic stainless steels having significantly lower permeability of hydrogen than other structural steels. The solubility of hydrogen and its isotopes in the type 300-series austenitic stainless steels is high relative to most structural materials. Compilation of data from gas permeation studies shows that most studies are consistent with one another,93,95,96 while studies that considered a variety of alloys within this class show that the solubility of hydrogen is essentially the same for a wide range of type 300-series austenitic stainless.93,95,96 The heat of solution of hydrogen in austenitic stainless steels is relatively low (DHs ¼ 6.9 kJ mol1), and thus the equilibrium content of hydrogen in the metal remains high even at room temperature. The solubility of hydrogen and its isotopes is plotted in Figure 11, while Table 1 lists the recommended transport properties for austenitic stainless steels (and a number of other metals and alloys). The primary traps in type 300-series austenitic stainless steels are dislocations with relatively low binding energy 10 kJ mol1.112 Therefore, the amount of trapped hydrogen (in the absence of irradiation and implantation damage) is relatively low at elevated temperatures. Moreover, due to the high solubility of hydrogen and its isotopes in austenitic
stainless steels, the density of trapping sites would need to be impractically high to measurably increase the inventory of hydrogen and its isotopes in the metal.20 For these reasons, trapping from a microstructural origin is anticipated to have little, if any, impact on the transport and inventory of hydrogen and its isotopes in austenitic stainless steels at temperatures greater than ambient. The recombination-rate constant (kr) for austenitic stainless steels near ambient temperature is typically less than about 109 m4 s1 per mol of H2.113 At higher temperatures (700 K), the value varies between 105 and 107 m4 s1 per mol of H2, depending on the surface condition.80,113–116 4.16.3.2.2 Ferritic/martensitic steels
There is significant interest in reduced activation ferritic/martensitic (RAFM) steels to replace nickelbearing austenitic stainless steels in reactor applications117. There are many RAFM steels that have been proposed and investigated in the literature specifically for fusion applications; these typically contain between 7 and 12 wt% chromium, relatively low carbon (<0.15 wt% C), and controlled alloying additions to bolster structural properties, while minimizing activation (e.g., additions of W, Ta, and vanadium and reductions of nickel, molybdenum,
Solubility (mol m–3 MPa–1/2)
Tritium Barriers and Tritium Diffusion in Fusion Reactors
529
100
10
1
1.2
1.4 1.6 Temperature, 1000/T (K–1)
1.8
2
Figure 11 Solubility of hydrogen in austenitic stainless steels from gas permeation studies that confirmed diffusion-limited transport. The bold line represents the average relationship determined in Perng and Altstetter93 for several austenitic stainless steels. Adapted from Quick, N. R.; Johnson, H. H. Metall. Trans. 1979, 10A, 67–70; Gromov, A. I.; Kovneristyi, Y. K. Met. Sci. Heat Treat. 1980, 22, 321–324; Perng, T. P.; Altstetter, C. J. Acta Metall. 1986, 34, 1771–1781; Louthan, M. R.; Derrick, R. G. Corrosion Sci. 1975, 15, 565–577; Sun, X. K.; Xu, J.; Li, Y. Y. Mater. Sci. Eng. A 1989, 114, 179–187; Grant, D. M.; Cummings, D. L.; Blackburn, D. A. J. Nucl. Mater. 1987, 149, 180–191; Grant, D. M.; Cummings, D. L.; Blackburn, D. A. J. Nucl. Mater. 1988, 152, 139–145; Mitchell, D. J.; Edge, E. M. J. Appl. Phys. 1985, 57, 5226–5235; Kishimoto, N.; Tanabe, T.; Suzuki, T.; et al. J. Nucl. Mater. 1985, 127, 1–9.
and niobium content). The transport of hydrogen and its isotopes has been extensively studied in MANET (MArtensitic for NET, including the so-called MANET II) and modified F82H (generally referred to as F82H-mod). Some of the other designations of RAFM steels that can be found in literature include EUROFER 97, Batman, OPTIFER-IVb, HT9, JLF-1, and CLAM steel. In general, studies of RAFM steels report relatively consistent transport properties of hydrogen and its isotopes; some of these studies are reviewed in Serra et al.118 Despite the consistency of the data available in literature from several research groups, few studies verify the expected pressure dependence of the transport properties that is expected for diffusioncontrolled transport. Pisarev and coworkers119,120 have suggested that the literature data may underestimate diffusivity and solubility due to surface limited transport. Similar suggestions have been presented to explain some of the data for the austenitic stainless steels1; however, the work on austenitic stainless steels has been cognizant of the issues with surface effects; generally surface effects are mitigated by coating specimens with palladium or other surface catalyst. Such precautions have not been systematically employed for permeation studies of the RAFM steels, although the need to control the surface
condition (and confirm the square root dependence on pressure) has been widely acknowledged.29,30,118,121 While the apparent transport properties in the absence of trapping are relatively consistent for all the RAFM steels, the issue of surface effects and the suggestions of Pisarev et al. need further validation in the literature because the transport of tritium is less likely to be affected by surface conditions compared to deuterium and protium. The diffusivity of hydrogen is shown in Figure 12 along with an average relationship (Table 1). The literature data are generally within a factor of 2 of the average relationship. The MANET alloys tend to have lower diffusivity of hydrogen and its isotopes than F82H-mod. Differences in permeability between these two alloys has been attributed to Chromium content;29,30 however, a clear correlation of transport properties with Chromium content cannot be established on the basis of existing data.122 At temperatures less than about 573 K, the apparent diffusivity is significantly less than the exponential relationship extrapolated from higher temperatures. This is attributed to the effect of trapping on the transport of hydrogen and its isotopes. The reported values of apparent solubility of hydrogen and its isotopes in RAFM varies very little in the temperature range from 573 to 873 K. Pisarev
530
Tritium Barriers and Tritium Diffusion in Fusion Reactors
Table 1 Recommended diffusivity and solubility relationships for protium in various metals and classes of alloys in the absence of trapping Alloy
Solubility, F/D
Diffusivity
K ¼ K0 exp (DHs/RT)
D ¼ D0 exp (ED/RT) 2
1
D0 (m s )
1
ED (kJ mol )
Beryllium
3 1011
18.3
Graphite Aluminum Vanadium RAFM steelsc Austenitic stainless steel Nickel Copper Zirconium Molybdenum Silver Tungsten Platinum Gold
9 105 2 108 3 108b 1 107 2 107 7 107 1 106 8 107 4 108 9 107 6 104 6 107 5.6 108
270 16 4.3b 13.2 49.3 39.5 38.5 45.3 22.3 30.1 103.1 24.7 23.6
References
K0 (mol H2 m3 MPa1/2)
DHs (kJ mol1)
18.9a 5.9 106a 19 46 138 436 266 564 792 3.4 107 3300 258 1490 207 77 900d
16.8a 96.6a 19.2 39.7 29 28.6 6.9 15.8 38.9 35.8 37.4 56.7 100.8 46.0 99.4d
74, 43 78 43 98, 99 100, 101 93 102 103 104, 105 106 107, 108 53 109 111
a
Per text, the solubility of hydrogen in beryllium is very low and there is not good agreement between the few studies of the material. Data for isotopes other than protium does not scale as the square root of mass. c Values are averaged over the data presented in Figures 12 and 13. d Estimated using the permeability from Caskey and Derrick110 and the quoted diffusivity. b
and coworkers report values that are three to four times higher on the basis of their assessment of surface effects. Here we recommend a relationship for the apparent solubility (Table 1) that is consistent with the majority of the literature data with DHs ¼ 28.6 kJ mol1, which is based on a simple curve fitting of the data shown in Figure 13. The values of the solubility are about an order of magnitude less than the austenitic stainless steels in the temperature range between 500 and 1000 K, although the solubility of hydrogen is more sensitive to temperature for the RAFM steels since DHs is four times the value for the austenitic stainless steels. The trapping characteristics of the RAFM steels have been estimated for several alloys.19,118,121,123–126 Although binding energies and densities of hydrogen traps vary substantially, the majority of reported values for RAFM steels are in the range 40–60 kJ mol1 and 103–105 traps per metal atom, respectively. The traps are attributed primarily to boundaries118 and result in a significant reduction in the apparent diffusivity at temperatures less than about 573 K. At higher temperatures, the traps are essentially unoccupied and do not affect diffusion.20 The measured recombination coefficient is many orders of magnitude lower than theoretical predictions; moreover, the measured values can also vary
substantially from one study to another.118,127,128 Measured values for the recombination coefficient for deuterium on MANET alloys are approximately in the range 102–104 m4 s1 per mol of H2 for the temperature range 573–773 K.127,128 Oxidation of MANET was shown to induce surface-limited transport of deuterium and reduce the recombination coefficient kr 106 m4 s1 per mol of H2.128 Furthermore, it is suggested that structure and composition of the oxide may also affect the recombination coefficient and that oxidation can increase the energy barrier associated with dissociation of the gaseous diatomic hydrogen isotopes.128 In summary, the diffusivity and the solubility of hydrogen and its isotopes are consistently similar for all the RAFM steels that have been tested for fusion applications. RAFM steels show a relatively rapid diffusion and low solubility of hydrogen and its isotopes at ambient temperature. The diffusivity is six orders of magnitude greater than that of the austenitic stainless steels at 300 K, while the solubility is more than three orders of magnitude lower than that of the austenitic stainless steels. The diffusivity of hydrogen and its isotopes is not strongly sensitive to temperature compared to most other metals. On the other hand, the heat of solution (DHs) for the RAFM steels is quite large,
Tritium Barriers and Tritium Diffusion in Fusion Reactors
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Figure 12 Diffusivity of hydrogen in reduced activation ferritic/martensitic steels from gas permeation studies that confirmed diffusion-limited transport. The bold line represents an approximate relationship estimated from the plotted data. Adapted from Forcey, K. S.; Ross, D. K.; Simpson, J. C. B.; et al. J. Nucl. Mater. 1988, 160, 117–124; Serra, E.; Perujo, A.; Benamati, G. J. Nucl. Mater. 1997, 245, 108–114; Serra, E.; Benamati, G.; Ogorodnikova, O. V. J. Nucl. Mater. 1998, 255, 105–115; Pisarev, A.; Shestakov, V.; Kulsartov, S.; et al. Phys. Scripta 2001, T94, 121–127; Esteban, G. A.; Perujo, A.; Douglas, K.; et al. J. Nucl. Mater. 2000, 281, 34–41; Dolinski, Y.; Lyasota, I.; Shestakov, A.; et al. J. Nucl. Mater. 2000, 283–287, 854–857; Kulsartov, T. V.; Hayashi, K.; Nakamichi, M.; et al. Fusion Eng. Des. 2006, 81, 701–705.
10
1
1
1.2
1.4 1.6 Temperature, 1000/T (K–1)
1.8
2
Figure 13 Solubility of hydrogen in reduced activation ferritic/martensitic steels from gas permeation studies that confirmed diffusion-limited transport. The bold line represents an approximate relationship estimated from the plotted data. Adapted from Forcey, K. S.; Ross, D. K.; Simpson, J. C. B.; et al. J. Nucl. Mater. 1988, 160, 117–124; Serra, E.; Perujo, A.; Benamati, G. J. Nucl. Mater. 1997, 245, 108–114; Serra, E.; Benamati, G.; Ogorodnikova, O. V. J. Nucl. Mater. 1998, 255, 105–115; Pisarev, A.; Shestakov, V.; Kulsartov, S.; et al. Phys. Scripta 2001, T94, 121–127; Esteban, G. A.; Perujo, A.; Douglas, K.; et al. J. Nucl. Mater. 2000, 281, 34–41; Dolinski, Y.; Lyasota, I.; Shestakov, A.; et al. J. Nucl. Mater. 2000, 283–287, 854–857; Kulsartov, T. V.; Hayashi, K.; Nakamichi, M.; et al. Fusion Eng. Des. 2006, 81, 701–705.
and thus the solubility of hydrogen approaches that of austenitic stainless steels at temperatures >1000 K. Consequently, at elevated temperatures (e.g., >700 K), the permeability is less than an
order of magnitude greater than that of the austenitic stainless steels and within a factor of 5 at temperature >1000 K. Trapping is significant in the RAFM steels at temperatures less than about
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Tritium Barriers and Tritium Diffusion in Fusion Reactors
573 K, and thus the apparent diffusivity is much lower than expected from tests that are performed at higher temperatures. 4.16.3.2.3 V–Cr–Ti alloys
Body-centered cubic (bcc)-structured V–Cr–Ti alloys (particularly composition ranges of around V–4Cr–4Ti and V–15Cr–5Ti) have low neutron cross-sections and the isotopes that do form with neutron capture have short half lives (51V has a half-life of <4 min). As noted in Chapter 4.12, Vanadium for Nuclear Systems, these characteristics, along with reasonable operating temperatures (limited by radiation hardening and helium bubble formation to 575–775 K129), make V–Cr–Ti an attractive material for first walls and blankets, but the tritium retention characteristics of vanadium alloys leave much to be desired. Vanadium has a large solubility for hydrogen and a very large diffusion coefficient for hydrogen. These two traits make the permeability of hydrogen in vanadium comparable to that of titanium and palladium.130 Vanadium absorbs hydrogen exothermically. Additions of chromium tend to increase this energy, while titanium additions tend to decrease it and, to first order, alloys with roughly equal and small amounts of chromium and titanium (such as V–4Cr–4Ti) are assumed to react similarly to hydrogen isotopes as pure vanadium.129 Vanadium alloys form hydrides below 450 K,129 which is below the typical operating temperatures. The diffusivity of hydrogen in vanadium is 108 2 1 m s in the range of operating temperatures, a larger value than in most metals.129 There has been extensive experimental measurement of several V–Cr–Ti alloys using different hydrogen isotopes. These are summarized in Figure 14. Schaumann et al.131 and Cantelli et al.132 independently measured the diffusivity of both protium and deuterium in pure vanadium after charging them with gas at 775 K using the Gorsky anelastic relaxation effect.131 Both groups found that the prefactor did not depend strongly on the isotope, whereas the activation energy did. This is contrary to the common naı¨ve expectation, where the activation energies would be identical and the prefactors would differ by a factor of the square root of the mass. The two groups also each reported a deviation from exponential behavior at lower temperatures, which could be attributed to some combination of surface effects, trapping, and a V–H phase transition at 200 K.132 However, deviation from Arrhenius behavior is common in bcc metals133 and has been reported in a
number of studies of vanadium. The transition temperature from exponential behavior varies widely with the technique used to measure diffusivity and the group that measures it,131,132,134–139 and has been as high as 813 K when measuring uncharged specimens using the absorption technique.134 This supports the notion that much of the deviation reported in the literature may be due to surface recombination limitation at lower temperatures. Compounding the recombination limitation of the vanadium base metal is the fact that surface oxides (particularly TiO) form and limit recombination more.140,141 At lower temperatures, there waspan ffiffiffi increased deviation in the measured DH/DD from 2, which has also been observed by others100,138 and in the diffusivity of titanium.142 The heavier hydrogen isotopes do not diffuse much slower than predicted until temperatures below 373 K, thus the physics associated with deviations from the predictions of classical rate theory cannot be exploited for the use in fusion applications. The electrochemical pulse experiments of Boes and Zuchner143 derived an activation energy that is twice as large using the electrochemical pulse method, which was also supported by absorption experiments by Eguchi and Morozumi.134 Both techniques are influenced by the surface, while Gorsky effect measurements and electrical resistivity measurements are only influenced by the bulk.133 The absorption experiments also indicated that the hydrogen diffusion coefficient in vanadium alloys is decreased by additions of chromium (as well as iron and niobium), but that it could be increased by large titanium additions, which is thought to be due to electronic contributions.134 Most other experiments have found that moderate amounts of titanium decrease the hydrogen diffusion coefficient much more than chromium is able to.133,139 Ti’s strong ability to trap hydrogen isotopes may explain the discrepancy in these measurements. Increasing titanium content decreases the DH/DD ratio, while increasing chromium content increases the ratio.139 While the solubility of hydrogen in vanadium is lower than that in either zirconium or titanium, it is still very large, being greater than the value in palladium and much greater than the value in the other structural metals considered here (Figure 15). The reported hydrogen lattice solubilities in vanadium alloys are in reasonable agreement, regardless of composition.130 Alloying additions do, however, change trapping in the alloy. Titanium has a higher heat of solution for hydrogen than vanadium and titanium138 additions
Tritium Barriers and Tritium Diffusion in Fusion Reactors
533
10–7
Diffusivity (m2 s–1)
10–8
10–9
10–10
10–11
10–12
0
2
4
6
8
10
Temperature, 1000/T (K–1) Figure 14 Diffusivity of hydrogen in vanadium and its alloys. The bold line represents the relationship for pure vanadium, reported in Freudenberg et al.100 Adapted from Schaumann, G.; Vo¨lki, J.; Alefeld, G. Phys. Status Solidi B 1970, 42, 401–413; Cantelli, R.; Mazzolai, F. M.; Nuovo, M. J. Phys. Chem. Solid. 1970, 31, 1811–1817; Tanaka, S.; Kimura, H. Trans. Jpn. Inst. Met. 1979, 20, 647–658; Eguchi, T.; Morozumi, S. J. Jpn. Inst. Met. 1977, 41, 795–802; Hashizume, K.; Masuda, J.; Otsuka, K. T.; et al. Fusion Sci. Technol. 2008, 54, 553–556; Klepikov, A. K.; Romanenko, O. G.; Chikhray, Y. V.; et al. Fusion Eng. Des. 2000, 51–52, 127–133; Lottner, V.; Heim, A.; Springer, T. Zeitschrift fu¨r Physik B 1979, 32, 157–165; Masuda, J.; Hashizume, K.; Otsuka, T.; et al. J. Nucl. Mater. 2007, 363–365, 1256–1260; Pine, D. J.; Cotts, R. M. Phys. Rev. B 1983, 28, 641; Freudenberg, U.; Vo¨lkl, J.; Bressers, J.; et al. Scripta Metall. 1978, 12, 165–167; Qi, Z.; Volkl, J.; Lasser, R.; et al. J. Phys. F 1983, 13, 2053–2062; Boes, N.; Zu¨chner, H. Phys. Status Solidi A 1973, 17, K111–K114; Anderl, R. A.; Longhurst, G. R.; Struttmann, D. A. J. Nucl. Mater. 1987, 145–147, 344–347; Romanenko, O. G.; Tazhibaeva, I. L.; Shestakov, V. P.; et al. J. Nucl. Mater. 1996, 233–237, 376–380; Fujii, K.; Hashizume, K.; Hatano, Y.; et al. J. Alloys Compd. 1998, 270, 42–46; Hashizume, K.; Masuda, J.; Otsuka, T.; et al. J. Nucl. Mater. 2007, 367–370, 876–881; Heller, R.; Wipf, H. Phys. Status Solidi (a) 1976, 33, 525–529.
increase the lattice parameter of vanadium.133 However, titanium is a much stronger trap than other elements that increase the lattice parameter as much or more (including niobium, molybdenum, and zirconium).133 Pine and Cotts139 assert that titanium solute atoms trap not only hydrogen isotopes at nearest-neighbor interstitial sites, but also hydrogen substitutionally. They demonstrated that the binding energy varied from 3 kJ mol1 in V–3Ti to 9.84 kJ mol1 in V–8Ti. The trapping energy for D is larger than that for hydrogen for both alloys. However, it should also be noted that there is considerable shortrange ordering in V–Ti alloys with more than 4 at. % Ti.133 This ordering means that trapping will not obey an Oriani-type behavior, in which trapping would be linearly dependent on the number of solute atoms, because the solid solution is not random. The elements chromium, iron, and copper all reduce the lattice parameter of vanadium and the diffusivity change in alloys containing these elements is also much lower than that in alloys with Ti.133 In fits to the apparent diffusivity in tritium diffusion experiments in ternary V–Cr–Ti alloys, Hashizume et al.135
show that, in addition to single titanium atoms, the most likely secondary trap is not chromium. Instead, the secondary trap has much higher energy and a lower concentration when compared to the monomer titanium trap. They also speculated that this was due to solute dimers and larger clusters. Interstitial oxygen, carbon, and nitrogen are also common in vanadium alloys. One or more hydrogen atoms bind with single carbon or nitrogen atoms readily, and oxygen atoms tend to trap at least two hydrogen atoms each.144 Other defects, such as dislocations, may still be effective traps at 773 K.145 As with other materials, vanadium can be damaged by radiation, and this will likely be the dominant trap in fusion reactors.146,147 The recombination coefficient for hydrogen is over five orders of magnitude slower in vanadium than in nitrogen in the range of operating temperatures, and is relatively insensitive to the surface concentration of sulfur.129 Because of this and the high diffusivity of tritium, release is recombination limited in vanadium alloys. Deuterium ion-driven permeation experiments148 of V–15Cr–5Ti have
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Tritium Barriers and Tritium Diffusion in Fusion Reactors
Solubility (mol m–3 MPa–1/2)
106
105
104
1000 0.8
1
1.2
1.4
1.6
1.8
2
–1
Temperature, 1000/T (K ) Figure 15 Solubility of hydrogen in vanadium and its alloys. The bold line represents the relationship for pure vanadium, reported in Steward.101 Adapted from Klepikov, A. K.; Romanenko, O. G.; Chikhray, Y. V.; et al. Fusion Eng. Des. 2000, 51–52, 127–133; Heller, R.; Wipf, H. Phys. Status Solidi (a) 1976, 33, 525–529; Steward, S. A. Review of Hydrogen Isotope Permeability Through Materials; Lawrence Livermore National Laboratory: Livermore, CA, 1983; Buxbaum, R. E.; Subramanian, R.; Park, J. H.; et al. J. Nucl. Mater. 1996, 233–237, 510–512; Maroni, V. A.; Van Deventer, E. H. J. Nucl. Mater. 1979, 85–86, 257–269; Zaluzhnyi, A. G.; Tebus, V. N.; Riazantseva, N. N.; et al. Fusion Eng. Des. 1998, 41, 181–185.
estimated the recombination-rate coefficient to be 2.4 1029 m4 s1 (although this measurement is three orders of magnitude lower than measurements on more dilute alloys149 and two orders of magnitude higher than measured in pure vanadium150). It should be noted that most measured recombination rates are lower bounds due to surface oxides. In environments in which this native oxide layer may be damaged (such as by radiation in a fusion reactor), the actual recombination rate may be higher.147 V–Cr–Ti alloys have hydrogen permeabilities that are at least two orders of magnitude more than nearly any other blanket material and form detrimental hydrides.129,130,151–156 The ongoing studies of permeation barriers may allow mitigation of this significant disadvantage so that V’s positive traits in a high-energy neutron environment can still be utilized. 4.16.3.2.4 Zirconium alloys
Zirconium alloys are described more fully in Chapter 2.07, Zirconium Alloys: Properties and Characteristics. They are used in fusion reactors partly because of their corrosion resistance in aqueous environments and low neutron cross-sections.157 However, zirconium readily forms embrittling hydride precipitates. Zirconium alloys oxidize and the surface ZrO2 may be an effective permeation barrier, preventing both hydrogen release and formation of detrimental
hydrides. Andrieu et al.158 demonstrated that the rate of tritium release of zircaloy-4 (Zry4) decreased substantially upon oxide formation in tritiated water. Zirconium has multiple phases at temperatures of interest: for example, a-, b-, and g-Zr coexist in equilibrium at 833 K. Most solubility and diffusivity studies have been conducted on the single-phase a-Zr generally at 773 K and below (Figure 16). Above this temperature, zirconium alloys dissolve up to 50 at.% hydrogen and this solubility decreases rapidly with decreasing temperature, causing hydride precipitates within the alloys. The solubility has been found to vary slightly with the alloying content. Yamanaka et al.159 note that the solubility in the b-phase decreases with alloying additions, while the solubility in the a-phase increases with alloying additions. The solubility of hydrogen in ZrO2, regardless of the crystal structure (104 to 105 mol hydrogen per mol oxide), is much lower than in the base metal and is even lower than that in Al2O3. a-ZrO2 exhibits a solubility almost an order of magnitude lower than b-ZrO2.160 Greger et al.161 have reviewed hydrogen diffusion in zirconium. The diffusivities reported in studies they cite and in others is plotted in Figure 17. At 623 K, the diffusivity of hydrogen in zirconium is 1010 m2 s1,104,158,161–164 while the diffusivity in ZrO2 is only 1019 to 1020 m2 s1.158,163,165 Austin
Tritium Barriers and Tritium Diffusion in Fusion Reactors
535
Solubility (mol m–3 MPa–1/2)
106
105
104
1000 1.2
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1.5
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Temperature, 1000/T (K–1) Figure 16 Solubility of hydrogen in zirconium and its alloys. The bold line represents the average for 13 studies on pure zirconium and Zr-based alloys, reported in Kearns.105 Adapted from Mallett, M. W.; Albrecht, W. M. J. Electrochem. Soc. 1957, 104, 142–146; Kearns, J. J. J. Nucl. Mater. 1967, 22, 292–303; Giroldi, J. P.; Vizcaı´no, P.; Flores, A. V.; et al. J. Alloys Compd. 2009, 474, 140–146; Khatamian, D. J. Alloys Compd. 1999, 293–295, 893–899; Khatamian, D. J. Alloys Compd. 2003, 356–357, 22–26; Khatamian, D.; Pan, Z. L.; Puls, M. P.; et al. J. Alloys Compd. 1995, 231, 488–493; Sawatzky, A.; Wilkins, B. J. S. J. Nucl. Mater. 1967, 22, 304–310; Une, K.; Ishimoto, S.; Etoh, Y.; et al. J. Nucl. Mater. 2009, 389, 127–136; Vizcaı´no, P.; Rı´os, R. O.; Banchik, A. D. Thermochim. Acta 2005, 429, 7–11.
10–9
Diffusivity (m2 s–1)
10–10 10–11 10–12 10–13 10–14 10–15 10–16 10–17
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Temperature, 1000/T (K–1) Figure 17 Diffusivity of hydrogen in zirconium and its alloys. The bold line represents the relationship for pure zirconium, reported in Kearns.104 Adapted from Mallett, M. W.; Albrecht, W. M. J. Electrochem. Soc. 1957, 104, 142–146; Greger, G. U.; Mu¨nzel, H.; Kunz, W.; et al. J. Nucl. Mater. 1980, 88, 15–22; Austin, J. H.; Elleman, T. S.; Verghese, K. J. Nucl. Mater. 1974, 51, 321–329; Cupp, C. R.; Flubacher, P. J. Nucl. Mater. 1962, 6, 213–228; Kearns, J. J. J. Nucl. Mater. 1972, 43, 330–338; Gulbransen, E. A.; Andrew, K. F. J. Electrochem. Soc. 1954, 101, 560–566; Kunz, W.; Mu¨nzel, H.; Helfrich, U. J. Nucl. Mater. 1982, 105, 178–183; Khatamian, D.; Manchester, F. D. J. Nucl. Mater. 1989, 166, 300–306; Sawatzky, A. J. Nucl. Mater. 1960, 2, 62–68.
et al.163 were able to measure the diffusivity in both a- and b-phases by measuring the activity, due to tritium, in tomographic slices of samples. The diffusivity values do not have a very strong dependence on crystallographic orientation or on alloy composition.
On the basis of observations of tritium segregation to some precipitates,158,164 many authors158,166,167 argue that intermetallic precipitates in zircaloy could be paths for short-circuit diffusion due to large reported values of solubility and diffusivity in
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some of these phases. However, these quantities appear to be relatively large for the zirconium-matrix material. Further, autoradiography shows depletion in some iron-rich precipitates and at 623 K, the diffusivity in ZrFe2 is 2.5 1011 m2 s1, slower than in bare zirconium.168 The permeability values through hydrides might be larger because of the high solubility of hydrogen isotopes in the hydride phase. However, the volume fraction of hydrides tends to be small and the activation energy has been shown to be independent of the presence of the hydride.169 Zirconium alloys that lack an oxide layer are not useful in hydrogen environments that exceed the solubility of hydrogen in zirconium, because of hydride formation. At relatively low use temperatures (<623 K), and in aqueous or otherwise oxidizing environments, zirconium oxide is able to grow and is an effective barrier against the permeation of hydrogen. Above this temperature, the integrity of the oxide layer cannot be maintained and the effective permeation of hydrogen isotopes is increased substantially. 4.16.3.2.5 Other structural metals 4.16.3.2.5.1 Aluminum
The permeability of hydrogen isotopes through aluminum is smaller than that through most metals. This is due to a very low intrinsic solubility and a moderate hydrogen diffusivity. Young and Scully99 reviewed the diffusivity of hydrogen in aluminum and discussed the reasons for the wide range of reported values. A majority of hydrogen present in aluminum alloys at lower temperatures is trapped at defects and alloying content tends to increase the solubility for hydrogen in aluminum slightly.99,170,171 Application of aluminum alloys in fusion reactors is limited by their neutron activation, low melting temperature (933 K for pure aluminum or significantly lower for some alloys), the formation of embrittling intermetallics when layered on other metals, and the difficulty of joining aluminum alloys (because of the native oxide layer). The thermal limitation is significant because precipitates in conventional aluminum alloys coarsen or dissolve above 450 K and most conventional alloys are not very resistant to creep deformation. Few studies of aluminum alloys have been made to date for use at temperatures greater than about 450 K. Because many aluminum-containing intermetallics and aluminum oxide also have low hydrogen permeability and do not have the same thermal limitations as metallic aluminum, it has been proposed that they be used as hydrogen isotope barriers.123
4.16.3.2.5.2
Nickel
Nickel alloys are described in Chapter 2.08, Nickel Alloys: Properties and Characteristics. The permeation of hydrogen and its isotopes in nickel has been extensively studied. The results from Louthan et al.102 provide a good estimate of the transport properties. They found that cold-worked nickel displays higher permeability to hydrogen, speculating that diffusion is enhanced along dislocation networks. Louthan4 also measured permeation in high-pressure gaseous hydrogen isotopes, showing that the transport is not dependent on concentration (and that Sievert’s law is appropriate at elevated pressure). Normalization of nickel’s permeation4 and diffusion3 by the mass ratio (eqn [3]) has been shown to provide good agreement for all the three isotopes. 4.16.3.2.5.3
Copper
There are numerous reports on the permeability of hydrogen in copper; the gas permeation results of Begeal103 are suggested here. Caskey and coworkers172 found that the diffusivity of hydrogen in copper depends on the oxygen content, and consequently, that the effective diffusivity can be much lower than the lattice diffusivity. Nevertheless, the reported diffusivity is consistent among a number of reports.3,103,172–174 Copper is also described in Chapter 4.20, Physical and Mechanical Properties of Copper and Copper Alloys. 4.16.3.3
Barrier Materials
The effectiveness of permeation barriers is often assessed by the permeation reduction factor (PRF), which is the ratio of the effective permeability of specimen without a barrier to an equivalent specimen with a barrier; thus, the greater the PRF value, the more effective the permeation barrier. 4.16.3.3.1 Oxides
Most metals form a native oxide layer in the presence of oxygen. Generally, oxides have very low permeabilities for hydrogen isotopes, and native oxide layers may reduce permeability by about an order of magnitude.175 The coefficient of thermal expansion is often very different from that of the underlying metal. This can cause cracks and spalling, which can reduce the effectiveness of the barrier. In environments that are aqueous or are subjected to elevated temperatures in the presence of oxygen, the oxide layer may be replenished and may even grow and coarsen. In addition to native oxides, an oxide layer might be deposited onto the base metal or the base metal might be dipped or
Tritium Barriers and Tritium Diffusion in Fusion Reactors
otherwise coated with a second metal, which forms a low-permeability oxide. Such coatings may reduce hydrogen permeability by five orders of magnitude or more.176 Chromia, alumina, and rare-earth oxides have been studied extensively. The low dissociation pressure of Cr2O3 makes it a common native oxide on steels when allowed to form at elevated temperatures and relatively low oxygen partial pressures.177 Chromia is a better barrier (offering a permeation reduction of about an order of magnitude)175,178 than various Cr2MO4 spinels that may also form (M ¼ Ni, Fe, Co).177 Chromia is also present in mixed oxides in chemical densified coatings, which help to give reduction factors of four orders of magnitude.179 Aluminum forms a self-passivating native oxide that has been shown to be resistant to hydrogen isotope permeability, because of a very low solubility for hydrogen. Because this layer is very thin (4 nm) and the hydrogen permeability of the base metal is very low, it remains debatable whether this amorphous native oxide, which can be grown by anodization, has a lower permeability than aluminum or not.180,181 Cleaned stainless steel may be hot-dipped in aluminum (forming both a relatively pure surface layer and mixed aluminides between the surface and substrate) and then oxidized. This hot-dip aluminizing processing is simple and generally forms coatings that have excellent adhesion properties (although substantially different thermal expansion coefficients),182 which reduce permeation rates by at least one order of
magnitude, and sometimes more than five orders of magnitude.175,183 The basic properties of hydrogen transport in alumina have been characterized and are presented in Figures 18 and 19. Roy and Coble185 hot-isostatically pressed high-purity (>99.99%) alumina powders and charged the dense alumina with hydrogen at elevated temperatures to determine solubility. Fowler et al.184 obtained diffusion coefficients for single-crystal, polycrystalline, and powdered alumina, and for alumina that was doped with MgO. They observed faster diffusion in powdered specimens, suggesting that the grain boundaries may provide short-circuit diffusion paths. They also noted that the diffusivity of MgO-doped alumina was four to five orders of magnitude greater than that of pure alumina. This suggests that the purity of barrier coatings matters a great deal and transmutation of barriers in a fusion environment may increase the permeability from the ideal case measured in the laboratory. Yttria and erbia have been deposited on specimens through a number of physical deposition techniques, including plasma spray, arc deposition, and sol–gel deposition.186–188 The advantage of these oxides is not the magnitude of permeation reduction (one to three orders of magnitude), but their high thermal and mechanical stability in a reducing atmosphere. 4.16.3.3.2 Aluminides
In addition to forming Al2O3, which is known to decrease hydrogen permeation, aluminization of steels
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Temperature, 1000/T (K–1) Figure 18 Diffusivity of hydrogen in alumina. The bold line represents the average for many sintered, powdered, and single crystal aluminas, reported in Fowler et al.184 Adapted from Fowler, J. D.; Chandra, D.; Elleman, T. S.; et al. J. Am. Ceram. Soc. 1977, 60, 155–161; Serra, E.; Bini, A. C.; Cosoli, G.; et al. J. Am. Ceram. Soc. 2005, 88, 15–18; Roberts, R. M.; Elleman, T. S.; Iii, H. P.; et al. J. Am. Ceram. Soc. 1979, 62, 495–499.
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Solubility (mol m–3 MPa–1/2)
10
Serra et al. 1
Roy and coble
0.1 0.4
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0.6
0.7
Temperature, 1000/T
0.8
0.9
1
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Figure 19 Solubility of hydrogen in alumina. Adapted from Serra, E.; Bini, A. C.; Cosoli, G.; et al. J. Am. Ceram. Soc. 2005, 88, 15–18; Roy, S. K.; Coble, R. L. J. Am. Ceram. Soc. 1967, 50, 435–436.
forms aluminide intermetallics that are believed to also lower permeability. Most studies of aluminized samples either have intentionally grown an oxidized layer in order to achieve greater PRFs or at least have not attempted to suppress the formation of surface Al2O3 prior to permeation testing. To our knowledge, no permeation measurements on oxide-free aluminides have been performed. However, different processes lead to oxide scales of differing composition, thickness, and defect density, and the PRF may not be attributable to oxides alone. Steels have been aluminized by the hot dip process (described earlier), as well as by various chemical vapor deposition (CVD), spray, packed cementation bed, and hot isostatic pressing (HIP) techniques. For those techniques that lay down a substantial amount of materials that does not react with the matrix (such as HIP), an aluminumcontaining iron alloy can be used in preference to aluminum to offer a higher temperature barrier performance.189 Oftentimes, the aluminized layer will be made up of mixed FeAl, FeAl2, Fe2Al3, Fe2Al5, FeAl3, Fe3Al, and even Fe4Al13 intermetallics.190 Nickel, chromium, and mixed-aluminides are also formed.191,192 Due to the aluminum-rich intermetallics on the surface, aluminized material will often have a mixed oxide scale that is rich in Al2O3.193,194 PRFs are generally larger than for a pure aluminum layer, varying between 10 and 10 000,175,195 while barriers containing a clean Al2O3 surface often have the greatest PRF.196 It should be noted that aluminum additions can also stabilize ferrite in some austenitic stainless steels, producing a duplex microstructure and an increased permeation.197
4.16.3.3.3 Nitrides
As with oxides, either the native nitrides of the base metal or depositions of other nitrides can be made to serve as barriers. One of the most common native nitrides is Fe2N, which forms when the upstream face of steel samples are nitrided and reduces permeability by one to three orders of magnitude.198–200 After oxides and aluminide, TiN coatings are one of the most researched barriers because of their good adhesion and the ease of deposition.175 Reported PRFs for nitride barriers vary widely, from less than an order of magnitude to six orders of magnitude. TiN barriers reduce permeability the most when they are placed on the high-pressure side of samples, and much less change in permeation is observed when they are placed downstream.201,202 Boron nitride has been shown to reduce the permeability of hydrogen in 304SS by one to two orders of magnitude.175,203 It forms both cubic and hexagonal structures. Checchetto et al.204 noted that the hexagonal structures absorb a greater amount of hydrogen isotopes (because of a larger number of trapping sites; particularly dangling B and nitrogen bonds), but also display a greater diffusivity of hydrogen. The latter effect may be due to preferential diffusion along the a direction of the hexagonal lattice. Bazzanella et al.205 found that a 1.7 mm (Al,Ti)N coating reduces the deuterium permeability of 0.1-mm thick 316L by two to three orders of magnitude. From permeation transients, they speculated that this reduction was primarily due to the very low diffusivity of D in (Al,Ti)N.
Tritium Barriers and Tritium Diffusion in Fusion Reactors
4.16.3.3.4 Carbides
The only report on the solubility of hydrogen in boron carbide is that by Shirasu et al.206 They exposed crystals of boron carbide to hydrogen gas at various temperatures and pressures for 20 h, and subsequently outgassed them during anneals in which the temperature was linearly increased at the rate of 20 K min1 up to 1273 K. The uptake was seen to increase with the square root of pressure, and to decrease with increasing temperature (exothermic). Schnarr and Munzel207,208 measured the diffusivity of tritium in both irradiated and unirradiated boron carbide. While the actual expression for the diffusivity for each case was not given, it can be extracted from the figures. It was noted that the apparent diffusivity decreased with increasing radiation damage until the percentage of 10B exceeded 10%. Elleman et al.209 used the 6Li-neutron reaction to generate tritium profiles in samples of boron carbide. Diffusivity was determined by examining the rate of release of the tritium during isothermal anneals at elevated temperatures. The diffusivity and the solubility of hydrogen in silicon carbide (a material described in Chapter 2.12, Properties and Characteristics of SiC and SiC/ SiC Composites and Chapter 4.07, Radiation Effects in SiC and SiC-SiC) have been measured twice by Causey et al.210,211 In the first set of experiments,210 various grades of silicon carbide were implanted with tritium, using the neutron reaction with 6Li on the sample surfaces. The diffusivity for each material was then determined by fitting the release curves determined during isothermal anneal to those predicted by the analytical solution to the diffusion equation. The results were seen to differ strongly depending on the type and purity of the silicon carbide. As an example, the measured diffusivity in hot pressed and aluminum-doped a-silicon carbide was approximately five orders of magnitude greater than that in vapor-deposited b-silicon carbide at 1273 K. The lowest diffusivities were reported for vapor-deposited b-silicon carbide and single-crystal a-silicon carbide. In all cases, the activation energy of the diffusivity was >200 kJ mol1 (suggesting that chemical bonding plays a strong role in the diffusion). For the diffusion of tritium in vapor-deposited silicon carbide, the diffusivity was given as D ¼ 1.58 104 exp(37 000/T ) m2 s1. Deuterium solubility was also determined for the vapor-deposited silicon carbide. The values were determined by exposing samples at elevated temperatures to deuterium gas followed by outgassing to determine the amount of uptake. Because equilibrium retention was not obtained in the
539
experiments, inherent in the calculations was the assumption that the diffusivity values determined in the implantation experiments were valid in the gaseous uptake experiments. The amount of uptake was assumed to be the product of diffusivity, the solubility, the sample area, and the square root of pressure. The solubility was given as K ¼ 1.1 103 exp (þ18 500/T ) mol H2 m2 MPa1/2. Again, the negative value of the activation energy would suggest chemical bonding of the hydrogen to the host material. In the later work by Causey et al.,211 vapordeposited silicon carbide was again tested. In these experiments, the implantation of energetic particles into the silicon carbide was avoided. Samples were exposed to gas containing 99% deuterium and 1% tritium at a temperature of 1573 K for 1 h. The samples were subsequently outgassed at temperatures from 1373 to 1773 K. The outgassing rates were then fitted to release curves predicted by the solution to the diffusion equation to determine the diffusivity. In this case, the diffusivity was given by the expression D ¼ 9.8 108 exp(21 870/T ) m2 s1, one to two orders of magnitude faster than the values determined earlier with energetic particles.210 The solubility was also determined in this study. Samples were exposed to the deuterium/tritium gas at temperatures from 1273 to 1873 K for sufficient duration to achieve equilibrium loading. The samples were then outgassed to determine this equilibrium amount. The expression for the solubility in this case was K ¼ 2.2 102 exp(þ7060/T ) mol H2 m3 MPa1/2. This solubility is one to two orders of magnitude lower than the one determined in the earlier experiments.210 If one assumes the migration of hydrogen in silicon carbide to occur along active sites on the edges of the grains, it is not unexpected that radiation damage produced by the implantation of energetic particles would increase the apparent solubility and proportionately decrease the apparent diffusivity. If hydrogen can exist only on the grain boundaries by being attached to trap sites, higher trapping means higher apparent solubility. Conversely, higher trapping means slower diffusion. It was the apparent higher solubility on small-grained samples that led Causey et al.211 to propose the trap-controlled grain boundary diffusion model. The permeation of hydrogen isotopes through silicon carbide has been measured by several groups.212–214 Verghese et al.213 measured the permeation of a hydrogen/tritium mixture through a KT silicon carbide tube that was manufactured by wet extrusion and sintering. The permeability reported for the experiments
540
Tritium Barriers and Tritium Diffusion in Fusion Reactors
Permeability (mol H2 m–1 s–1 MPa–1/2)
is given by F ¼ 3.8 108 exp(66 000/T ) mol H2 m1 s1 MPa1/2. Sinharoy and Lange212 measured the permeation of hydrogen through a tungsten tube with a CVD coating of silicon carbide. The retarding effect of the tungsten was taken into consideration in the calculation. The recorded permeation for these experiments was F ¼ 2 104 exp(6830/T ) mol H2 m1 s1 MPa1/2. Yao et al.214 performed permeation experiments on a steel sample that had been RF sputter-coated with silicon carbide. The thickness of the coating was estimated to be 1.3 mm and contained several percent oxygen and traces of iron. The coating was seen to decrease the permeation rate of steel by about two orders of magnitude, but did not change the activation energy. In this case, the coating was clearly porous, and the reduction in permeation was simply due to a reduction in the effective permeation surface area. The plot of the permeation values for the vapordeposited silicon carbide by Causey et al.211 (calculated as the product of diffusivity times solubility), KT silicon carbide by Verghese et al.,213 and CVD silicon carbide by Sinharoy and Lange212 is shown in Figure 20. The differences in the absolute values of the permeability as well as the differences in the activation energy of the process are extreme. It is difficult to even imagine that the values are for the same material. In fact, the materials are not the same. As mentioned for the original study by Causey et al.,210 differences in impurities play a significant role in determining the behavior of hydrogen in silicon carbide. If hydrogen does migrate along the grain boundaries, impurity metals along those grain boundaries
reduce the fraction of migrating hydrogen chemically bound to the silicon carbide. Likewise, the apparent diffusivity would be much more rapid if hydrogen trapping at the grain boundaries is reduced. In the case of the permeability measured by Sinharoy and Lange,212 it is difficult to believe that the measured permeation is not really controlled by permeation through the underlying tungsten with the specific surface area limited by the porous silicon carbide coating. The activation energy for the permeation in the report by Verghese et al.213 is difficult to understand. The value of 555 kJ mol1 is even greater than the chemical bond of hydrogen to carbon.215 The permeation was seen to vary by as much as an order of magnitude at the same temperature. There is no apparent explanation for the rapid change in permeation with temperature. Titanium carbide has also been tested as a permeation barrier. Due to adhesion problems with direct deposition on steel, titanium nitride was used as an intermediate layer between the steel and titanium carbide. Forcey et al.202 measured deuterium permeation through 3-mm thick layers of TiC and TiN on steel, observing a PRF of ten. For the experiments performed over the temperature range of 550–740 K, extended defects were listed as the reason for the relatively small improvement over bare steel. Checchetto et al.201 used ion-beam assisted deposition of TiN–TiC films on steel in their permeation experiments. When the film was deposited on the downstream side, little reduction in permeation was seen. Using the deposited film on the upstream side
10–6 Sinharoy and Lange 10–8 Verghese et al. 10–10
10–12
10–14 0.55
Causey et al.
0.6
0.65
0.7
0.75
0.8
0.85
0.9
0.95
Temperature, 1000/T (K–1) Figure 20 Permeability of hydrogen in SiC. Adapted from Causey, R. A.; Wampler, W. R.; Retelle, J. R.; et al. J. Nucl. Mater. 1993, 203, 196–205; Sinharoy, S.; Lange, W. J. J. Vac. Sci. Technol. A Vac. Surf. Films 1984, 2, 636–637; Verghese, K.; Zumwalt, L. R.; Feng, C. P.; et al. J. Nucl. Mater. 1979, 85–86, 1161–1164.
541
Tritium Barriers and Tritium Diffusion in Fusion Reactors
did yield a PRF of 50. Shan et al.216 used a CVD process to deposit their 2.5-mm thick film on steel and noted a permeation reduction of five to six orders of magnitude. It is obvious from these three studies that the deposition of theoretically dense thin films is very difficult. There is also the question of cracking of such thin films during thermal cycling. This is discussed later in this chapter.
parameters.101,106,118 The reported permeability values are relatively consistent between the majority of studies, while the diffusivity and solubility values range over several orders of magnitude. The results of Tanabe et al.106 are proposed here as they appear to represent nearly upper bounds of both diffusivity and solubility, without overestimating permeability. The study of Tanabe and coworkers also has the advantage that permeability and diffusivity were measured over a wide range of temperature and pressure, confirming the appropriate pressure dependencies of permeability and diffusivity for diffusionlimited transport.
4.16.3.3.5 Low permeation metals
The permeation of hydrogen and its isotopes through many of the transition metals is lower than that displayed by iron and the ferritic steels; the notable exceptions include groups 4 and 5 as well as palladium. Figure 21 shows the permeability of several metals; the diffusivity and the solubility are listed in Table 1 for these metals. In general, the activation energy associated with permeability ðDHs þ ED Þ is larger for the materials with lower permeability and the permeability tends to converge at elevated temperatures. We do not attempt to comprehensively review the data for nonferrous metals. However, gas permeation studies are considered the standard for transport properties, particularly studies that report permeability, diffusivity, and solubility. Permeation of tritium through metals and alloys was reviewed by Steward.101
4.16.3.3.5.2
4.16.3.3.5.3
10–3
V
10–6
10–6 Zr
RAFM
10–9
Mo Cu
10
10–9
Ni Aus. SS
Ag
–12
1
1.2
1.4
1.6
Temperature, 1000/T (K–1)
1.8
Pt W
Al
2
1
Au
1.2
1.4
–12
1.6
1.8
2
10
Permeability (mol H2 m–1 s–1 MPa–1/2)
Several reviews of the literature on hydrogen transport in molybdenum have noted variability of the transport
Permeability (mol H2 m–1 s–1 MPa–1/2)
Platinum
There are relatively few gas permeation studies of platinum. Ebisuzaki et al.109 report the permeability, diffusivity, and solubility of both hydrogen and
4.16.3.3.5.1 Molybdenum
10–3
Silver
The available data for hydrogen permeation through silver are limited. The diffusivity of hydrogen is reported by Katsuta and McLellan.107 McLellan also reports the solubility of Group IB metals from saturation experiments.108 Although these saturation experiments do not appear to provide reasonable values for other Group IB metals and are not consistent with other reported solubility measurements,217 Steward, nevertheless, suggests estimating the permeability of hydrogen using these reported relationships.108
–1
Temperature, 1000/T (K )
Figure 21 Permeability of hydrogen in various metals using data from Table 1. Data is distributed across two separate plots for clarity. Adapted from Frauenfelder, R. J. Vac. Sci. Technol. 1969, 6, 388–397; Perng, T. P.; Altstetter, C. J. Acta Metall. 1986, 34, 1771–1781; Freudenberg, U.; Vo¨lkl, J.; Bressers, J.; et al. Scripta Metall. 1978, 12, 165–167; Steward, S. A. Review of Hydrogen Isotope Permeability Through Materials; Lawrence Livermore National Laboratory: Livermore, CA, 1983; Kearns, J. J. J. Nucl. Mater. 1967, 22, 292–303; Kearns, J. J. J. Nucl. Mater. 1972, 43, 330–338; Young, G. A.; Scully, J. R. Acta Mater. 1998, 46, 6337–6349; Louthan, M. R.; Donovan, J. A.; Caskey, G. R. Acta Metall. 1975, 23, 745–749; Begeal, D. R. J. Vac. Sci. Technol. 1978, 15, 1146–1154; Tanabe, T.; Yamanishi, Y.; Imoto, S. J. Nucl. Mater. 1992, 191–194, 439–443; Katsuta, H.; McLellan, R. B. Scripta Metall. 1979, 13, 65–66; McLellan, R. B. J. Phys. Chem. Solids 1973, 34, 1137–1141; Ebisuzaki, Y.; Kass, W. J.; O’Keeffe, M. J. Chem. Phys. 1968, 49, 3329–3332; Eichenauer, W.; Liebscher, D. Zeitschrift fur Naturforschung 1962, 17A, 355; Ransley, C. E.; Neufeld, H. J. Inst. Met. 1948, 74, 599–620.
542
Tritium Barriers and Tritium Diffusion in Fusion Reactors
deuterium through single crystals of high-purity platinum. The permeability of hydrogen in platinum is similar to that in copper.
The diffusivity shown in Table 1 is from Eichenauer and Liebscher,111 while the solubility is estimated from this diffusivity and the permeability reported by Caskey and Derrick.110
4.16.3.3.5.4 Gold
Caskey and Derrick110 report the permeability of deuterium through gold; Begeal103 reports a similar relationship. Diffusivity measurements, however, differ depending on the conditions of the measurement and the microstructural state of the gold.110,218 Cold-worked gold tends to give a higher activation for diffusion, suggesting that trapping is active to relatively high temperatures. Caskey and Derrick110 speculate that trapping is related to vacancies.
4.16.4 Application of Barriers 4.16.4.1
Expected In-Reactor Performance
As implied by Figures 21 and 22, Tables 1 and 2 and in the earlier sections, permeation barriers can be used to reduce the effective permeation in laboratory testing.175–179,183,195,196,198–203,205 PRFs from laboratory experiments have been reported to be many
Permeability (mol H2 m–1 s–1 MPa–1/2)
10–12
10–15
B4C UO2
10–18 α–ZrO2 10–21
SiC Al2O3
10–24
1
1.2
1.4
1.6
1.8
2
Temperature, 1000/T (K–1) Figure 22 Permeability of hydrogen in various ceramics using data from Table 2. Adapted from DiStefano, J. R.; De Van, J. H.; Ro¨hrig, D. H.; et al. J. Nucl. Mater. 1999, 273, 102–110; Spitzig, W. A.; Owen, C. V.; Reed, L. K. J. Mater. Sci. 1992, 27, 2848–2856; Tanabe, T.; Tamanishi, Y.; Sawada, K.; et al. J. Nucl. Mater. 1984, 122&123, 1568–1572; Forcey, K. S.; Ross, D. K.; Simpson, J. C. B.; et al. J. Nucl. Mater. 1989, 161, 108–116; Wolarek, Z.; Zakroczymski, T. Acta Mater. 2006, 54, 1525–1532; Perujo, A.; Kolbe, H. J. Nucl. Mater. 1998, 263, 582–586; Song, W.; Du, J.; Xu, Y.; et al. J. Nucl. Mater. 1997, 246, 139–143.
Table 2 Recommended diffusivity and solubility relationships for protium in various nonmetallic materials in the absence of trapping Material
Al2O3 a-ZrO2 UO2 B4C SiC
Diffusivity
Solubility, F/D
D = D0 exp (ED/RT)
K = K0 exp (DHs/RT)
D0 (m2 s1)
ED (kJ mol1)
K0 (mol H2 m3 MPa1/2)
DHs (kJ mol1)
1.1 108 4 1018 3.7 106 1.2 1011 9.8 108
132 30.1 59.8 80.8 182
5.5 2.5 102 9.6 104 3.8 2.2 102
22.5 28.2 100 29.8 58.7
References
184, 233 160, 163 234 206 211
Tritium Barriers and Tritium Diffusion in Fusion Reactors
thousands in certain barrier systems.175,183,195,196,201,202 However, while the data available in the open literature are quite limited, there is significant evidence that the effectiveness of the permeation barriers decreases in radiation environments. There were three sets of experiments219–222 performed in the high flux reactor (HFR) Petten reactor in the Netherlands. In the first of these experiments219,222 reported in 1991 and 1992, tritium was produced by the liquid breeder material Pb–17Li. Permeation of tritium through a bare 316 stainless steel layer was compared with that through an identical layer covered with a 146-mm thick aluminide coating. Over the temperature range 540–760 K, the barrier was reported to decrease the permeation by a factor of 80 compared to the bare metal, that is, PRF ¼ 80. In the LIBRETTO-3 experiments,220 three different permeation barrier concepts were tested with the tritium again produced by the liquid breeder material. One irradiation capsule for tritium breeding was coated on the outside with a 6–8-mm thick CVD layer of TiC. A second capsule was coated on the inside with a 0.5–1.5-mm thick layer of TiC followed by a 2–3-mm thick layer of Al2O3. The third barrier was an aluminide coating produced by the cementation process. The aluminum-rich layer was 120-mm thick with about 5 mm of Al2O3 on the outside. The single TiC layer reduced the tritium permeation by a factor of only 3.2, the TiC and Al2O3 layer reduced the permeation by 3.4, and the pack cementation aluminide coating reduced the permeation by a factor of 14.7. These are surprisingly small reductions PRFs compared to laboratory experiments. In a third set of experiments,221 the tritium production was achieved with the solid ceramic breeder materials. Both double-wall tubes and single-wall tubes with a permeation barrier were tested. The double wall configuration had an inner layer of copper. The permeation barrier on the other system was an aluminide coating with a thickness of 7 mm. The aluminide coating was reported to be 70 times more effective than the double-wall configuration in suppressing permeation. Unfortunately, different breeder materials were used for the two different experiments, and the results could have been strongly affected by the amount of tritium released from the ceramic as well as the form of release (T2 vs. T2O). The bottom line on the irradiation testing of barriers is that barriers do not perform as well in a reactor environment as expected from laboratory experiments: a PRF > 1000 has not been achieved in reactor environments.
543
4.16.4.2 How Barriers Work and Why Radiation Affects Them To understand the effects of radiation on the performance of permeation barriers, we need to first examine how barriers work. For tritium to permeate through a material with or without a coating, the tritium must absorb on the surface, dissociate into atoms, dissolve into the material, diffuse through the material, and then recombine into molecules on the downstream side. In the simple case in which diffusion through the structural material is rate limiting, the permeation rate is controlled by the ratio of the permeability and the thickness of the pressure boundary (eqn [16]); as described earlier, the permeability is the product of the diffusivity and the solubility, which can be thought of as the velocity times capacity. These parameters are dependent on temperature, and should not be affected by radiation effects or nominal surface cracking. In the simple case of diffusion-limited permeation through the structural boundary, the experimental result determined in the laboratory cannot be extrapolated to the radiation environment. Permeation barriers, by their basic nature, consist of a thin layer adhered to the structural material. The performance of barriers depends on the integrity of the barrier as well as the physical interaction of the barrier material with tritium. What is it about many barriers and how they operate that causes laboratory and reactor data to disagree? In their review, Hollenberg et al.175 considered these aspects of barriers and their performance in radiation, proposing three models that describe distinct physics of the interactions between tritium and the barrier material. The most basic model is the Composite Diffusion Model, in which hydrogen transport is diffusioncontrolled in both the barrier and the base metal. The steady-state permeation rate (Q1 ) through a pressure boundary in this case is pffiffiffiffiffiffiffiffi A p TT ½22 Q1 ¼ t tM B þ FB FM where A is the surface area of the boundary, and the subscripts B and M refer to the barrier and structural metal, respectively. Considering the intent of the barrier, the ratio tB =FB should be much larger than tM =FM , thus the permeation is controlled simply by the permeation through the barrier. The second model proposed by Hollenberg et al. considers the barrier to be effectively impermeable to
544
Tritium Barriers and Tritium Diffusion in Fusion Reactors
tritium and is called the Area Defect Model. In this case, hydrogen is transported through the metal, reaching the metal surface through a limited number of cracks or other defects in the barrier layer. The permeation rate for this case is Q1 ¼ Ad
FM pffiffiffiffiffiffiffiffi p TT teff
½23
where Ad is the area of the defects and teff is the effective distance the hydrogen isotope must traverse to reach the other side of the metal. The third model proposed by Hollenberg et al. is the Surface Desorption Model, in which case, permeation is controlled by the recombination rate of hydrogen isotope atoms into molecules on the back surface and the recombination-limited flux of tritium is described by eqn [19]. Surface desorption does not make sense by itself; as show, it is actually part of the Area Defect Model. As reported by Hollenberg et al.175 and as revealed by a review of the literature on barriers and oxides,196,223–225 the activation energy of permeation is generally not altered by the addition of the barrier layer onto the substrate. This means that, in practice, the permeation process itself is being controlled by the substrate, not the barrier, strongly supporting the Area Defect Model described earlier. In short, the barrier works simply by limiting the area of the metal exposed to the driving pressure. Pisarev et al.226 provide particularly intriguing insights into the effects of cracks on permeation barriers. Their report showed that permeation reduction for the Area Defect Model is difficult to achieve when the distance between defects is not larger than the combined thickness of the barrier and substrate. Inherent in this conclusion is the assumption that the dissociation rate at the defect is sufficiently fast to maintain the equilibrium concentration dictated by Sievert’s law. If this condition is not met, then the activation energy for the process would be that associated with the dissociation, and not that of permeation through the substrate. Thus, barriers that can provide a significant permeation reduction in the laboratory must be essentially defect free. The physics of hydrogen transport in metals with permeation barriers can be further understood by examining the pressure dependence of permeation. As discussed earlier, diffusion-controlled permeation through metals is proportional to the square root of the hydrogen partial pressure. Perujo et al.227 reported that the pressure dependence of permeation through MANET plasma sprayed with aluminum
changed from the classic square root dependence to linear as the pressure was decreased below 20 000 Pa. Mcguire228 also noted the transition to near-linear pressure dependence in the pressure range from 200 to 1000 Pa. Linear pressure dependence is symptomatic of permeation limited by absorption or recombination. For example, if recombination limits permeation, the concentration of hydrogen in the metal will be almost constant and uniform, and it will be established by equilibrium at the upstream side of the pressure boundary. Thus, Sievert’s law (eqn [7]) can be substituted into eqn [24], leading to linear pressure dependence: Jr ¼ kr K 2 p TT
½24
While the same relationship will be found if the permeation is limited by absorption on the upstream surface, known values for the recombination-rate constant for MANET can explain the linear pressure dependence seen in permeation measurements.128 The conclusion is that a combination of the Area Defect Model and the Surface Desorption Model is needed to properly model permeation though barrier materials. If barriers work by limiting the area available for the gas to contact the underlying metal surface, and possibly by creating low enough permeation to have recombination even further reduce the permeation, how does radiation affect this process? One possible answer is by increasing the porosity or cracking of the barrier. According to Arshak and Korostynska229 properties of metal oxide materials are directly or indirectly connected to the presence of defects, oxygen vacancies in particular. Oxygen vacancies are also known as color centers, and these color centers are stabilized by hydrogen trapped at the defects. The hydrogen can come from preexisting OH groups or from hydrogen isotopes migrating through the oxide, possibly increased by the enhanced electrical conductivity generated by the radiation damage and the oxygen vacancies. While cracking was not considered by Arshak and Korostynska, one can speculate that the radiation damage with increased oxygen vacancies and trapped hydrogen would lead to a more brittle oxide layer. In metals, lateral stress from hydrogen or helium trapping can lead to blisters.230 Without the required ductility to allow blistering, the oxide layer could experience significantly increased cracking. The cracking would then increase the area available for hydrogen to reach the metal surfaces.
Tritium Barriers and Tritium Diffusion in Fusion Reactors
4.16.4.3 Why Barriers Are Needed for Fusion Reactors In this chapter, the materials for the blanket region have been reviewed, and their permeation parameters described. In this section, the need for barriers is evaluated. For example, consider the tritium migration processes that might be associated with the liquid Pb–17Li systems. In a set of experiments, Maeda et al.231 found the solubility of hydrogen in Pb–17Li to be on the order of 107 Pa1/2 atom fraction (106 mol H2 m3 MPa1/2). As tritium is produced in the blanket, some of the tritium will be in solution and some will be in the vapor phase. It is the tritium in the vapor phase that will drive the permeation through the metal used to contain the liquid. For an 800 MW fusion reactor, Maeda et al. state that 1.5 MCi (150 g) of tritium will have to be bred each day. That means that 1.5 MCi of tritium will be flowing around in stainless steel or similar metal tubes at a temperature >600 K. To estimate tritium permeation in a generic 800 MW plant, we will scale the design parameters proposed by Farabolini et al.232 for a much larger plant. Approximately 10 000 m2 of surface area will be needed for the tubes passing through the liquid Pb–17Li to extract the heat. We will assume that a sufficient number of detritiation cycles per day are performed to keep the amount of tritium in the liquid breeder at 10% of the 150 g listed above. Scaling to 800 MW, the amount of Pb–17Li will be 750 000 kg. This leads to a molar fraction of tritium equal to 6.7 107. Using the solubility of Maeda et al.231 for Pb–17Li yields a tritium pressure of 45 Pa. Assuming the containment metal to be 1 mm of MANET with an aluminized coating, a temperature of 700 K and an effective PRF of 1000, a permeation rate of 2.7 1010 T2 mols m2 s1 will occur.29,118,122,127,128 With the 10 000 m2 surface area, the daily permeation rate is 0.23 mol or 1.4 g of tritium per day. To prevent subsequent permeation through the steam generator tube walls, a tritium clean up unit will have to be applied to this helium loop. Because the steam generator tube wall must be thin to permit effective heat transfer, the tritium cleanup loop will have to be extremely effective to limit release of tritium to the environment. This calculation was performed simply to show the extreme need for barriers in the blanket region of fusion reactors. Even with an active detritiation unit and a barrier providing a PRF of 1000, 1.4 g or 14 000 Ci of tritium end up in the cooling system each day. The situation is not
545
much better for the solid breeders. The same amount of tritium will obviously be required for that system. To minimize the tritium inventory in the ceramic breeder materials, temperatures equal to or greater than that of the liquid breeder will be maintained. The tritium will be released into the helium coolant as elemental tritium (T2) and tritiated water (T2O); the relative concentrations of these forms depend on the type of ceramic breeder. The steel or similar containment metal will be exposed to nontrivial pressures of tritium gas. We can conclude that effective barriers are needed for the blanket. It is difficult to imagine that, even with double-walled designs, fusion reactor facilities can meet radioactive release requirements for tritium without an effective barrier.
4.16.5 Summary In this chapter, we have presented tritium permeation characteristics and parameters for materials used in fusion reactors. These materials have included those used to face the plasma in the main chamber as well as materials used as structural materials for the main chamber and blanket. A description of the conditions that exist in those locations has also been provided. Reasons were given why direct contact of the plasma with the plasma facing materials would not lead to sizeable quantities of tritium being lost to the environment or to the cooling system. The same was not concluded for the blanket region. The need for permeation barriers there was stressed. A number of materials were listed as possible tritium barriers. These materials included a few metals with somewhat reduced permeation and a larger number of ceramics with very low tritium permeability. Due to the difficulty of lining large chambers with bulk ceramics, much of the tritium permeation barrier development around the world has been dedicated to thin ceramic layers on metal surfaces. Unfortunately, radiation testing219–222 of these materials has shown that these thin layers lose their ability to limit tritium permeation during exposure to radiation damage. It was suggested, but not proved, that this increase in permeation was due to cracking of the ceramics or the increase in defects. To make this chapter more useful to the reader with a need for permeation data, tables and plots of the permeation coefficients are provided. The coefficients for metals are presented in Table 1 and Figure 21 and for the ceramics in Table 2 and Figure 22.
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In summary, effective permeation barriers are needed for fusion reactors to prevent the release of sizeable quantities of tritium. Fusion is touted as a clean form of energy, and releasing tritium into the environment will eliminate any political advantage that fusion has over fission. Research is needed to find ways to place radiation-resistant ceramic permeation barriers on top of structural metals. The fusion community must find a way to make this happen.
28. 29. 30. 31. 32. 33. 34. 35.
References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19. 20.
21. 22. 23. 24. 25. 26. 27.
San Marchi, C.; Somerday, B. P.; Robinson, S. L. Int. J. Hydrogen Energy 2007, 32, 100–116. Voelkl, J.; Alefeld, G. In Diffusion in Solids: Recent Developments; Nowick, A. S., Burton, J. J., Eds.; Academic Press: New York, 1975. Katz, L.; Guinan, M.; Borg, R. J. Phys. Rev. B 1971, 4, 330–341. Louthan, M. R.; Derrick, R. G. Scripta Metall. 1976, 10, 53–55. Quick, N. R.; Johnson, H. H. Metall. Trans. 1979, 10A, 67–70. Swansiger, W. A.; Bastasz, R. J. Nucl. Mater. 1979, 85–86, 335–339. Forcey, K. S.; Ross, D. K.; Simpson, J. C. B.; Evans, D. S. J. Nucl. Mater. 1988, 160, 117–124. Shiraishi, T.; Nishikawa, M.; Yamaguchi, T.; Kenmotsu, K. J. Nucl. Mater. 1999, 273, 60–65. Esteban, G. A.; Legarda, F.; Perujo, A. Fusion Sci. Technol. 2005, 48, 617–620. Ebisuzaki, Y.; Kass, W. J.; O’Keeffe, M. J. Chem. Phys. 1967, 46, 1373–1378. Caskey, G. R. In Hydrogen Degradation of Ferrous Alloys; Oriani, R. A., Hirth, J. P., Smialowski, M., Eds.; Noyes: Park Ridge, NJ, 1985; pp 822–862. Wriedt, H. A.; Oriani, R. A. Acta Metall. 1970, 18, 753–760. van Leeuwen, H. P. Eng. Fract. Mech. 1974, 6, 141–161. Sofronis, P. J. Mech. Phys. Solids 1995, 43, 1385–1407. Hirth, J. P. Metall. Trans. 1980, 11A, 861–890. Oriani, R. A. Acta Metall. 1970, 18, 147–157. Perkins, W. G. J. Vac. Sci. Technol. 1973, 10, 543–556. Oriani, R. A. Fusion Technol. 1994, 26, 235–266. Serra, E.; Perujo, A.; Benamati, G. J. Nucl. Mater. 1997, 245, 108–114. San Marchi, C.; Somerday, B. P. In Materials Resarch Society Proceedings, Mar 24–28, 2008; Choudhury, B., Dillon, A., Keller, J., Moen, C., Eds.; Materials Research Society: San Francisco, CA, 2008; Vol. 1098E, pp HH1008–1001. LeClaire, A. D. Diffusion Defect Data 1983, 34, 1–35. Zarchy, A. S.; Axtmann, R. C. J. Nucl. Mater. 1979, 79, 110–117. Perkins, H. K.; Noda, T. J. Nucl. Mater. 1978, 71, 349–364. Baskes, M. I. J. Nucl. Mater. 1980, 92, 318–324. Pick, M. A.; Sonnenberg, K. J. Nucl. Mater. 1985, 131, 208–220. Wampler, W. R. Appl. Phys. Lett. 1986, 48, 405–407. Causey, R. A.; Baskes, M. I. J. Nucl. Mater. 1987, 147, 284–287.
36. 37. 38. 39. 40. 41. 42. 43. 44. 45. 46. 47. 48. 49. 50. 51. 52. 53. 54. 55. 56. 57. 58. 59. 60. 61. 62.
Causey, R. A.; Harbin, W.; Taylor, D.; Snead, L. Phys. Scripta 1996, T64, 32–35. Wedig, F.; Jung, P. J. Nucl. Mater. 1997, 245, 138–146. Schliefer, F.; Liu, C.; Jung, P. J. Nucl. Mater. 2000, 283–287, 540–544. Yao, Z.; Liu, C.; Jung, P. Fusion Sci. Technol. 2005, 48, 1285–1291. Binyukova, S. Y.; Chernov, I. I.; Kalin, B. A.; Swe, T. J. Nucl. Mater. 2007, 367–370, 500–504. Kelly, B. T. Physics of Graphite; Springer: London, 1981. Kiyoshi, T.; Namba, T.; Yamawaki, M. J. Nucl. Mater. 1988, 155–157, 230–233. Reynolds, W. N. Physical Properties of Graphite; Elsevier: Amsterdam, 1968. Barrer, R. M. Proc. R. Soc. Lond. 1935, 149, 253–269. Thomas, W. J. Journal De Chimie Physique Et De Physico-Chimie Biologique 1961, 58, 61–69. Bansal, R. C.; Vastola, F. J.; Walker, P. L. Carbon 1971, 9, 185–192. Robell, A. J.; Ballou, E. V.; Boudart, M. J. Phys. Chem. 1964, 68, 2748–2753. Olander, D. R.; Balooch, M. J. Catal. 1979, 60, 41–56. Causey, R. A.; Baskes, M. I.; Wilson, K. L. J. Vac. Sci. Technol. A 1986, 4, 1189–1192. Penzhorn, R.-D.; Bekris, N.; Berndt, U.; Coad, J. P.; Ziegler, H.; Na¨gele, W. J. Nucl. Mater. 2001, 288, 170–178. Atsumi, H.; Tokura, S.; Miyake, M. J. Nucl. Mater. 1988, 155, 241–245. Causey, R. A. J. Nucl. Mater. 1989, 162, 151–161. Ro¨hrig, H. D.; Fischer, P. G.; Hecker, R. J. Am. Ceram. Soc. 1976, 59, 316–320. Causey, R. A.; Elleman, T. S.; Verghese, K. Carbon 1979, 17, 323–328. Malka, V.; Ro¨hrig, H. D.; Hecker, R. Int. J. Appl. Radiat. Isot. 1980, 31, 469. Atsumi, H.; Iseki, M.; Shikama, T. J. Nucl. Mater. 1992, 191, 368–372. Atsumi, H.; Iseki, M.; Shikama, T. J. Nucl. Mater. 1994, 215, 1478–1482. Wampler, W. R.; Doyle, B. L.; Causey, R. A.; Wilson, K. J. Nucl. Mater. 1990, 176–177, 983–986. Jacob, W. Thin Solid Films 1998, 326, 1–42. Sakamoto, R.; Muroga, T.; Yoshida, N. J. Nucl. Mater. 1995, 220–222, 819–822. Frauenfelder, R. J. Vac. Sci. Technol. 1969, 6, 388–397. Zakahrov, A. P.; Sharapov, V. M. Fiziko-Khimicheskaya Mekhanika Materialov 1973, 9, 29–33. Benamati, G.; Serra, E.; Wu, C. H. J. Nucl. Mater. 2000, 283–287, 1033–1037. Mazayev, A. A.; Avarbe, R. G.; Vilk, Y. N. Russian Metallurgy-Metally-USSR 1968, 6, 153–158. van Veen, A.; Filius, H. A.; de Vries, J.; Bijkerk, K. R.; Rozing, G. J.; Segers, D. J. Nucl. Mater. 1988, 155–157, 1113–1117. Eleveld, H.; van Veen, A. J. Nucl. Mater. 1992, 191–194, 433–438. Pisarev, A. A.; Varava, A. V.; Zhdanov, S. K. J. Nucl. Mater. 1995, 220–222, 926–929. Garcı´a-Rosales, C.; Franzen, P.; Plank, H.; Roth, J.; Gauthier, E. J. Nucl. Mater. 1996, 233–237, 803–808. Causey, R.; Wilson, K.; Venhaus, T.; Wampler, W. R. J. Nucl. Mater. 1999, 266–269, 467–471. Anderl, R.; Holland, D. F.; Longhurst, G. R.; Pawelko, R. J.; Trybus, C. L.; Sellers, C. H. Fusion Technol. 1991, 21, 745–752.
Tritium Barriers and Tritium Diffusion in Fusion Reactors 63. 64. 65. 66. 67. 68. 69. 70. 71. 72. 73. 74. 75. 76. 77. 78. 79. 80. 81. 82. 83. 84. 85. 86. 87. 88. 89. 90. 91.
92. 93. 94.
Sze, F. C.; Doerner, R. P.; Luckhardt, S. J. Nucl. Mater. 1999, 264, 89–98. Venhaus, T.; Causey, R.; Doerner, R.; Abeln, T. J. Nucl. Mater. 2001, 290–293, 505–508. Haasz, A. A.; Poon, M.; Davis, J. W. J. Nucl. Mater. 1999, 266–269, 520–525. Wang, W.; Roth, J.; Lindig, S.; Wu, C. H. J. Nucl. Mater. 2001, 299, 124–131. Shimada, T.; Ueda, Y.; Nishikawa, M. Fusion Eng. Des. 2003, 66–68, 247–251. Shu, W. M.; Luo, G. N.; Yamanishi, T. J. Nucl. Mater. 2007, 367–370, 1463–1467. Causey, R. A.; Doerner, R.; Fraser, H.; et al. J. Nucl. Mater. 2009, 390–391, 717–720. Miyamoto, M.; Nishijima, D.; Ueda, Y.; et al. Nucl. Fusion 2009, 49, 065035. Wampler, W. R.; Doerner, R. P. Nucl. Fusion 2009, 49, 115023. Abramov, E.; Riehm, M. P.; Thompson, D. A.; Smeltzer, W. W. J. Nucl. Mater. 1990, 175, 95427–95430. Tazhibaeva, I. L.; Shestakov, V. P.; Chikhray, Y. V. In Proceedings of the 18th Symposium of Fusion Technology; Elsevier: Karlsruhe, Germany, 1990; pp 427–430. Jones, P. M. S.; Gibson, R. J. Nucl. Mater. 1967, 21, 353–354. Altstetter, C. J.; Abraham, D. In Hydrogen Effects in Materials; Thompson, A. W., Moody, N. R., Eds.; TMS: Warrendale, PA, 1996; pp 599–609. Macaulay-Newcombe, R. G.; Thompson, D. A.; Smeltzer, W. W. J. Nucl. Mater. 1992, 191–194, 263–267. Shapovalov, V. I.; Dukel’skii, Y. M. Izvestiya Akademii Nauk SSR Metally 1988, 5, 201–203. Swansiger, W. A. J. Vac. Sci. Technol. A 1986, 4, 1216–1217. Langley, R. A. J. Nucl. Mater. 1979, 85–86, 1123–1126. Myers, S. M.; Wampler, W. R.; Besenbacher, F. J. Appl. Phys. 1984, 56, 1561–1571. Yoshida, N.; Mizusawa, S.; Sakamoto, R.; Muroga, T. J. Nucl. Mater. 1996, 233–237, 874–879. Causey, R. A.; Longhurst, G. R.; Harbin, W. J. Nucl. Mater. 1997, 241–243, 1041–1046. Doerner, R. P.; Grossman, A.; Luckhardt, S.; et al. J. Nucl. Mater. 1998, 257, 51–58. Chernikov, V. N.; Alimov, V. K.; Markin, A. V.; et al. J. Nucl. Mater. 1996, 233–237, 860–864. Alimov, V. K.; Chernikov, V. N.; Zakharov, A. P. J. Nucl. Mater. 1997, 241–243, 1047–1051. Baldwin, D. L.; Billone, M. C. J. Nucl. Mater. 1994, 212–215, 948–953. Andreev, D. V.; Bespalov, V. N.; Birjukov, A. J.; Gurovich, B. A.; Platonov, P. A. J. Nucl. Mater. 1996, 233–237, 880–885. Fenici, P.; Boerman, D.; Coen, V.; Lang, E.; Ponti, C.; Schu¨le, W. Nucl. Eng. Des./Fusion 1984, 1, 167–183. Kohyama, A.; Grossbeck, M. L.; Piatti, G. J. Nucl. Mater. 1992, 191–194, 37–44. Sahin, S.; Uebeyli, M. J. Fusion Eng. 2008, 27, 271–277. Anderl, R. A.; Holland, D. F.; Struttmann, D. A.; Lonhurst, G. R.; Merrill, B. J. In Proceedings of the 11th Symposium on Fusion Engineering, Austin, TX, Nov 18–22, 1985; Evans, J., Jones, D., Wilkins, D. C., Eds.; IEEE: New York, 1985; Vol. 1, pp 644–649. Gromov, A. I.; Kovneristyi, Y. K. Met. Sci. Heat Treat. 1980, 22, 321–324. Perng, T. P.; Altstetter, C. J. Acta Metall. 1986, 34, 1771–1781. Perng, T. P.; Johnson, M.; Altstetter, C. J. Acta Metall. 1989, 37, 3393–3397.
95. 96. 97. 98. 99. 100. 101. 102. 103. 104. 105. 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.
547
Louthan, M. R.; Derrick, R. G. Corrosion Sci. 1975, 15, 565–577. Sun, X. K.; Xu, J.; Li, Y. Y. Mater. Sci. Eng. A 1989, 114, 179–187. Sun, X. K.; Xu, J.; Li, Y. Y. Acta Metall. 1989, 37, 2171–2176. Ransley, C. E.; Neufeld, H. J. Inst. Met. 1948, 74, 599–620. Young, G. A.; Scully, J. R. Acta Mater. 1998, 46, 6337–6349. Freudenberg, U.; Vo¨lkl, J.; Bressers, J.; Alefeld, G. Scripta Metall. 1978, 12, 165–167. Steward, S. A. Review of Hydrogen Isotope Permeability Through Materials; Lawrence Livermore National Laboratory: Livermore, CA, 1983. Louthan, M. R.; Donovan, J. A.; Caskey, G. R. Acta Metall. 1975, 23, 745–749. Begeal, D. R. J. Vac. Sci. Technol. 1978, 15, 1146–1154. Kearns, J. J. J. Nucl. Mater. 1972, 43, 330–338. Kearns, J. J. J. Nucl. Mater. 1967, 22, 292–303. Tanabe, T.; Yamanishi, Y.; Imoto, S. J. Nucl. Mater. 1992, 191–194, 439–443. Katsuta, H.; McLellan, R. B. Scripta Metall. 1979, 13, 65–66. McLellan, R. B. J. Phys. Chem. Solids 1973, 34, 1137–1141. Ebisuzaki, Y.; Kass, W. J.; O’Keeffe, M. J. Chem. Phys. 1968, 49, 3329–3332. Caskey, G. R.; Derrick, R. G. Scripta Metall. 1976, 10, 377–380. Eichenauer, W.; Liebscher, D. Zeitschrift fur Naturforschung 1962, 17A, 355. Thomas, G. J. In Hydrogen Effects in Metals; Bernstein, I. M., Thompson, A. W., Eds.; The Metallurgical Society of AIME: New York, 1981; pp 77–85. Langley, R. A. J. Nucl. Mater. 1984, 128–129, 622–628. Grant, D. M.; Cummings, D. L.; Blackburn, D. A. J. Nucl. Mater. 1987, 149, 180–191. Grant, D. M.; Cummings, D. L.; Blackburn, D. A. J. Nucl. Mater. 1988, 152, 139–145. Braun, M.; Emmoth, B.; Waelbroeck, F.; Wienhold, P. J. Nucl. Mater. 1980, 93–94, 861–865. Baluc, N.; Gelles, D. S.; Jitsukawa, S.; et al. J. Nucl. Mater. 2007, 367–370, 33–41. Serra, E.; Benamati, G.; Ogorodnikova, O. V. J. Nucl. Mater. 1998, 255, 105–115. Pisarev, A.; Shestakov, V.; Kulsartov, S.; Vaitonene, A. Phys. Scripta 2001, T94, 121–127. Shestakov, V.; Pisarev, A.; Sobolev, V.; Kulsartov, S.; Tazhibaeva, I. J. Nucl. Mater. 2002, 307–311, 1494–1497. Esteban, G. A.; Perujo, A.; Sedano, L. A.; Douglas, K. J. Nucl. Mater. 2000, 281, 34–41. Jung, P. J. Nucl. Mater. 1996, 238, 189–197. Benamati, G.; Donato, A.; Solina, A.; Valentini, R.; Lanza, S. J. Nucl. Mater. 1994, 212–215, 1401–1405. Valentini, R.; Solina, A.; Tonelli, L.; Lanza, S.; Benamati, G.; Donato, A. J. Nucl. Mater. 1996, 233–237, 1123–1127. Forcey, K. S.; Iordanova, I.; Yaneva, M. J. Nucl. Mater. 1997, 240, 118–123. Esteban, G. A.; Pen˜a, A.; Urra, I.; Legarda, F.; Riccardi, B. J. Nucl. Mater. 2007, 367–370, 473–477. Serra, E.; Perujo, A. J. Nucl. Mater. 1995, 223, 157–162. Serra, E.; Perujo, A. J. Nucl. Mater. 1997, 240, 215–220. Andrew, P. L.; Pick, M. A. J. Nucl. Mater. 1994, 212–215, 111–117. Dutton, R. Int. J. Hydrogen Energy 1984, 9, 147–155.
548
Tritium Barriers and Tritium Diffusion in Fusion Reactors
131. Schaumann, G.; Vo¨lki, J.; Alefeld, G. Phys. Status Solidi B 1970, 42, 401–413. 132. Cantelli, R.; Mazzolai, F. M.; Nuovo, M. J. Phys. Chem. Solid. 1970, 31, 1811–1817. 133. Tanaka, S.; Kimura, H. Trans. Jpn. Inst. Met. 1979, 20, 647–658. 134. Eguchi, T.; Morozumi, S. J. Jpn. Inst. Met. 1977, 41, 795–802. 135. Hashizume, K.; Masuda, J.; Otsuka, K. T.; et al. Fusion Sci. Technol. 2008, 54, 553–556. 136. Klepikov, A. K.; Romanenko, O. G.; Chikhray, Y. V.; Tazhibaeva, I. L.; Shestakov, V. P.; Longhurst, G. R. Fusion Eng. Des. 2000, 51–52, 127–133. 137. Lottner, V.; Heim, A.; Springer, T. Zeitschrift fu¨r Physik B 1979, 32, 157–165. 138. Masuda, J.; Hashizume, K.; Otsuka, T.; et al. J. Nucl. Mater. 2007, 363–365, 1256–1260. 139. Pine, D. J.; Cotts, R. M. Phys. Rev. B 1983, 28, 641. 140. Hayakawa, R.; Hatano, Y.; Fukumoto, K.-I.; Matsui, H.; Watanabe, K. J. Nucl. Mater. 2004, 329–333, 411–415. 141. Hirohata, Y.; Oda, T.; Hino, T.; Sengoku, S. J. Nucl. Mater. 2001, 290–293, 196–200. 142. Qi, Z.; Vo¨lkl, J.; La¨sser, R.; Wenzl, H. J. Phys. F 1983, 13, 2053–2062. 143. Boes, N.; Zu¨chner, H. Phys. Status Solidi A 1973, 17, K111–K114. 144. Chang, H. Y.; Wert, C. A. Acta Metall. 1973, 21, 1233–1242. 145. Ohnuki, S.; Yasuda, T.; Suda, T.; Watanabe, S.; Oliver, B. M. J. Nucl. Mater. 2004, 329–333, 481–485. 146. Hirohata, Y.; Yamada, T.; Yamauchi, Y.; Hino, T.; Nagasaka, T.; Muroga, T. J. Nucl. Mater. 2006, 348, 33–39. 147. Yamauchi, Y.; Yamada, T.; Hirohata, Y.; Hino, T.; Muroga, T. J. Nucl. Mater. 2004, 329–333, 397–400. 148. Anderl, R. A.; Longhurst, G. R.; Struttmann, D. A. J. Nucl. Mater. 1987, 145–147, 344–347. 149. Romanenko, O. G.; Tazhibaeva, I. L.; Shestakov, V. P.; et al. J. Nucl. Mater. 1996, 233–237, 376–380. 150. Yamawaki, M.; Yamaguchi, K.; Tanaka, S.; Namba, T.; Kiyoshi, T.; Takahashi, Y. J. Nucl. Mater. 1989, 162–164, 1071–1076. 151. Aoyagi, K.; Torres, E. P.; Suda, T.; Ohnuki, S. J. Nucl. Mater. 2000, 283–287, 876–879. 152. DiStefano, J. R.; De Van, J. H.; Ro¨hrig, D. H.; Chitwood, L. D. J. Nucl. Mater. 1999, 273, 102–110. 153. Ro¨hrig, H. D.; DiStefano, J. R.; Chitwood, L. D. J. Nucl. Mater. 1998, 258–263, 1356–1360. 154. Spitzig, W. A.; Owen, C. V.; Reed, L. K. J. Mater. Sci. 1992, 27, 2848–2856. 155. Wang, Y.; Kanedome, M.; Yasuda, T.; et al. J. Nucl. Mater. 2004, 329–333, 477–480. 156. DiStefano, J. R.; Pint, B. A.; DeVan, J. H.; Ro¨hrig, H. D.; Chitwood, L. D. J. Nucl. Mater. 2000, 283–287, 841–845. 157. Forty, C. B. A.; Karditsas, P. J. J. Nucl. Mater. 2000, 283–287, 607–610. 158. Andrieu, C.; Ravel, S.; Ducros, G.; Lemaignan, C. J. Nucl. Mater. 2005, 347, 12–19. 159. Yamanaka, S.; Higuchi, K.; Miyake, M. J. Alloys Compd. 1995, 231, 503–507. 160. Yamanaka, S.; Nishizaki, T.; Uno, M.; Katsura, M. J. Alloys Compd. 1999, 293–295, 38–41. 161. Greger, G. U.; Mu¨nzel, H.; Kunz, W.; Schwierczinski, A. J. Nucl. Mater. 1980, 88, 15–22. 162. Mallett, M. W.; Albrecht, W. M. J. Electrochem. Soc. 1957, 104, 142–146. 163. Austin, J. H.; Elleman, T. S.; Verghese, K. J. Nucl. Mater. 1974, 51, 321–329.
164. 165. 166. 167. 168. 169. 170. 171. 172. 173. 174. 175. 176. 177. 178. 179. 180. 181. 182. 183. 184. 185. 186. 187. 188. 189. 190. 191. 192. 193. 194. 195. 196.
Cupp, C. R.; Flubacher, P. J. Nucl. Mater. 1962, 6, 213–228. Gulbransen, E. A.; Andrew, K. F. J. Electrochem. Soc. 1954, 101, 560–566. Hatano, Y.; Hitaka, R.; Sugisaki, M.; Hayashi, M. J. Nucl. Mater. 1997, 248, 311–314. Lim, B. H.; Hong, H. S.; Lee, K. S. J. Nucl. Mater. 2003, 312, 134–140. Maier, C. U.; Kronmuller, H. J. Phys. Condens. Matter 1992, 4, 4409–4420. Kunz, W.; Mu¨nzel, H.; Helfrich, U. J. Nucl. Mater. 1982, 105, 178–183. Scully, J. R.; Young, G. A.; Smith, S. W. Mater. Sci. Forum 2000, 331–337, 1583–1600. Smith, S. W.; Scully, J. R. Metall. Mater. Trans. A Phys. Metall. Mater. Sci. 2000, 31, 179–193. Caskey, G. R.; Dexter, A. H.; Holzworth, M. L.; Louthan, M. R.; Derrick, R. G. Corrosion 1976, 32, 370–374. Eichenauer, W.; Loser, W.; Witte, H. Zeitschrift fur Metallkunde 1965, 56, 287–293. Tanabe, T.; Yamanishi, Y.; Sawada, K.; Imoto, S. J. Nucl. Mater. 1984, 122&123, 1568–1572. Hollenberg, G. W.; Simonen, E. P.; Kalinin, G.; Terlain, A. Fusion Eng. Des. 1995, 28, 190–208. Takagi, I.; Kobayashi, T.; Ueyama, Y.; et al. J. Nucl. Mater. 2009, 386–388, 682–684. Hecker, R.; Sto¨ver, D.; Jonas, H.; Buchkremer, H. P. J. Nucl. Mater. 1990, 171, 84–93. Terai, T.; Yoneoka, T.; Tanaka, H.; Kawamura, H.; Nakamichi, M.; Miyajima, K. J. Nucl. Mater. 1994, 212–215, 976–980. Nakamichi, M.; Nakamura, H.; Hayashi, K.; Takagi, I. J. Nucl. Mater. 2009, 386–388, 692–695. Bell, J.; Redman, J. J. Mater. Energ. Syst. 1983, 4, 217–221. Smithells, C. J.; Ransley, C. E. Proc. R. Soc. A 1935, 152, 706–713. Zhang, Y.; Pint, B. A.; Garner, G. W.; et al. Surf. Coating. Technol. 2004, 188–189, 35–40. Levchuk, D.; Koch, F.; Maier, H.; Bolt, H. J. Nucl. Mater. 2004, 328, 103–106. Fowler, J. D.; Chandra, D.; Elleman, T. S.; Payne, A. W.; Verghese, K. J. Am. Ceram. Soc. 1977, 60, 155–161. Roy, S. K.; Coble, R. L. J. Am. Ceram. Soc. 1967, 50, 435–436. Chikada, T.; Suzuki, A.; Yao, Z.; et al. Fusion Eng. Des. 2009, 84, 590–592. Levchuk, D.; Levchuk, S.; Maier, H.; Bolt, H.; Suzuki, A. J. Nucl. Mater. 2007, 367–370, 1033–1037. Yao, Z.; Suzuki, A.; Levchuk, D.; et al. J. Nucl. Mater. 2009, 386–388, 700–702. Benamati, G.; Chabrol, C.; Perujo, A.; Rigal, E.; Glasbrenner, H. J. Nucl. Mater. 1999, 272, 391–395. Bhuvaneswaran, N.; Mudali, U. K.; Shankar, P. Scripta Mater. 2003, 49, 1133–1138. Kalin, B. A.; Yakushin, V. L.; Fomina, E. P. Fusion Eng. Des. 1998, 41, 119–127. Barbier, F.; Manuelli, D.; Bouche, K. Scripta Mater. 1997, 36, 425–431. Fazio, C.; Stein-Fechner, K.; Serra, E.; Glasbrenner, H.; Benamati, G. J. Nucl. Mater. 1999, 273, 233–238. Stein-Fechner, K.; Konys, J.; Wedemeyer, O. J. Nucl. Mater. 1997, 249, 33–38. Forcey, K. S.; Ross, D. K.; Simpson, J. C. B.; Evans, D. S.; Whitaker, A. G. J. Nucl. Mater. 1989, 161, 108–116. Forcey, K. S.; Ross, D. K.; Wu, C. H. J. Nucl. Mater. 1991, 182, 36–51.
Tritium Barriers and Tritium Diffusion in Fusion Reactors 197. Adams, T. M.; Korinko, P.; Duncan, A. Mater. Sci. Eng. A Struct. Mater. Prop. Microstruct. Process. 2006, 424, 33–39. 198. Brass, A. M.; Chene, J.; Pivin, J. C. J. Mater. Sci. 1989, 24, 1693–1699. 199. Wolarek, Z.; Zakroczymski, T. Acta Mater. 2004, 52, 2637–2643. 200. Wolarek, Z.; Zakroczymski, T. Acta Mater. 2006, 54, 1525–1532. 201. Checchetto, R.; Bonelli, M.; Gratton, L. M.; et al. Surf. Coating. Technol. 1996, 83, 40–44. 202. Forcey, K. S.; Perujo, A.; Reiter, F.; Lolli-Ceroni, P. L. J. Nucl. Mater. 1993, 200, 417–420. 203. Itakura, A.; Tosa, M.; Ikeda, S.; Yoshihara, K. Vacuum 1996, 47, 697–700. 204. Checchetto, R.; Chayahara, A.; Horino, H.; Miotello, A.; Fujii, K. Thin Solid Films 1997, 299, 5–9. 205. Bazzanella, N.; Checchetto, R.; Miotello, A.; Patton, B.; Kale, A. N.; Kothari, D. C. Appl. Phys. Lett. 2002, 81, 3762–3764. 206. Shirasu, Y.; Yamanaka, S.; Miyake, M. J. Alloys Compd. 1992, 190, 87–90. 207. Schnarr, K.; Munzel, H. J. Nucl. Mater. 1990, 170, 253–260. 208. Schnarr, K.; Munzel, H. J. Chem. Soc. Faraday Trans. 1990, 86, 651–656. 209. Elleman, T. S.; Zumwalt, L. R.; Verghese, K. Hydrogen diffusion and permeation in ceramic materials. In Proceedings of the 3rd Topical Meeting on Techniques of Controlled Nuclear Fusion, Santa Fe, NM, 1978; Vol. 2, pp 763–766. 210. Causey, R. A.; Fowler, J. D.; Ravanbakht, C.; Elleman, T. S.; Verghese, K. J. Am. Ceram. Soc. 1978, 61, 221–225. 211. Causey, R. A.; Wampler, W. R.; Retelle, J. R.; Kaae, J. L. J. Nucl. Mater. 1993, 203, 196–205. 212. Sinharoy, S.; Lange, W. J. J. Vac. Sci. Technol. A Vac. Surf. Films 1984, 2, 636–637. 213. Verghese, K.; Zumwalt, L. R.; Feng, C. P.; Elleman, T. S. J. Nucl. Mater. 1979, 85–86, 1161–1164. 214. Yao, Z.; Suzuki, A.; Levchuk, D.; Terai, T. Fusion Sci. Technol. 2007, 52, 865–869. 215. Pauling, L. The Nature of the Chemical Bond, and the Structure of Molecules and Crystals: An Introduction to Modern Structural Chemistry; Cornell University Press: Ithaca, NY, 1940. 216. Shan, C.; Wu, A.; Li, Y.; et al. J. Nucl. Mater. 1992, 191–194, 221–225. 217. Verbetsky, V. N.; Mitrokhin, S. V. Solid State Phenom. 2000, 73–75, 503–517. 218. Ishikawa, T.; McLellan, R. B. J. Phys. Chem. Solids 1985, 46, 1393–1396. 219. Conrad, R.; Debarberis, L.; Coen, V.; Flament, T. J. Nucl. Mater. 1991, 179, 875–878. 220. Conrad, R.; Futterer, M. A.; Giancarli, L.; May, R.; Perujo, A.; Sample, T. J. Nucl. Mater. 1994, 215, 998–1002. 221. Magielsen, A. J.; Bakker, K.; Chabrol, C.; et al. J. Nucl. Mater. 2002, 307, 832–836. 222. Proust, E.; Leroy, P.; Franenberg, H. W. J. Nucl. Mater. 1992, 191, 186–189. 223. Perujo, A.; Kolbe, H. J. Nucl. Mater. 1998, 263, 582–586.
224. 225. 226. 227. 228.
229. 230. 231. 232. 233. 234. 235. 236. 237. 238. 239. 240. 241. 242. 243. 244. 245. 246. 247. 248. 249. 250. 251. 252. 253. 254.
549
Song, W.; Du, J.; Xu, Y.; Long, B. J. Nucl. Mater. 1997, 246, 139–143. Ishikawa, Y.; Yoshimura, T.; Arai, M. Vacuum 1996, 47, 701–704. Pisarev, A.; Tsvetkov, I.; Yarko, S. Fusion Eng. Des. 2007, 82, 2120–2125. Perujo, A.; Forcey, K. S.; Sample, T. J. Nucl. Mater. 1993, 207, 86–91. Mcguire, J. G. Hydrogen permeation resistant layers for liquid metal reactors. In Proceedings of the Conference on Tritium Technology in Fission, Fusion, and Isotopic Applications, 64–68, 1980. Arshak, K.; Korostynska, O. Mater. Sci. Eng. B 2006, 133, 1–7. Eernisse, E. P.; Picraux, S. T. J. Appl. Phys. 1977, 48, 9–17. Maeda, Y.; Edao, Y.; Yamaguchi, S.; Fukada, S. Fusion Sci. Technol. 2008, 54, 131–134. Farabolini, W.; Ciampichetti, A.; Dabbene, F.; et al. Fusion Eng. Des. 2006, 81, 753–762. Serra, E.; Bini, A. C.; Cosoli, G.; Pilloni, L. J. Am. Ceram. Soc. 2005, 88, 15–18. Sherman, D. F.; Olander, D. R. J. Nucl. Mater. 1989, 166, 307–320. Mitchell, D. J.; Edge, E. M. J. Appl. Phys. 1985, 57, 5226–5235. Kishimoto, N.; Tanabe, T.; Suzuki, T.; Yoshida, H. J. Nucl. Mater. 1985, 127, 1–9. Dolinski, Y.; Lyasota, I.; Shestakov, A.; Repritsev, Y.; Zouev, Y. J. Nucl. Mater. 2000, 283–287, 854–857. Kulsartov, T. V.; Hayashi, K.; Nakamichi, M.; et al. Fusion Eng. Des. 2006, 81, 701–705. Fujii, K.; Hashizume, K.; Hatano, Y.; Sugisaki, M. J. Alloys Compd. 1998, 270, 42–46. Hashizume, K.; Masuda, J.; Otsuka, T.; et al. J. Nucl. Mater. 2007, 367–370, 876–881. Heller, R.; Wipf, H. Phys. Status Solidi (a) 1976, 33, 525–529. Buxbaum, R. E.; Subramanian, R.; Park, J. H.; Smith, D. L. J. Nucl. Mater. 1996, 233–237, 510–512. Maroni, V. A.; Van Deventer, E. H. J. Nucl. Mater. 1979, 85–86, 257–269. Zaluzhnyi, A. G.; Tebus, V. N.; Riazantseva, N. N.; et al. Fusion Eng. Des. 1998, 41, 181–185. Giroldi, J. P.; Vizcaı´no, P.; Flores, A. V.; Banchik, A. D. J. Alloys Compd. 2009, 474, 140–146. Khatamian, D. J. Alloys Compd. 1999, 293–295, 893–899. Khatamian, D. J. Alloys Compd. 2003, 356–357, 22–26. Khatamian, D.; Pan, Z. L.; Puls, M. P.; et al. J. Alloys Compd. 1995, 231, 488–493. Sawatzky, A.; Wilkins, B. J. S. J. Nucl. Mater. 1967, 22, 304–310: doi: 10.1016/0022-3115(67)90048-7. Une, K.; Ishimoto, S.; Etoh, Y.; et al. J. Nucl. Mater. 2009, 389, 127–136. Vizcaı´no, P.; Rı´os, R. O.; Banchik, A. D. Thermochim. Acta 2005, 429, 7–11. Khatamian, D.; Manchester, F. D. J. Nucl. Mater. 1989, 166, 300–306. Sawatzky, A. J. Nucl. Mater. 1960, 2, 62–68. Roberts, R. M.; Elleman, T. S.; Iii, H. P.; Verghese, K. J. Am. Ceram. Soc. 1979, 62, 495–499.
4.17
Tungsten as a Plasma-Facing Material
G. Pintsuk Forschungszentrum Ju¨lich, Ju¨lich, Germany
ß 2012 Elsevier Ltd. All rights reserved.
4.17.1 4.17.2 4.17.3 4.17.3.1 4.17.3.2 4.17.3.2.1 4.17.3.2.2 4.17.3.2.3 4.17.3.2.4 4.17.3.2.5 4.17.3.2.6 4.17.3.2.7 4.17.3.3 4.17.4 4.17.4.1 4.17.4.1.1 4.17.4.1.2 4.17.4.1.3 4.17.4.1.4 4.17.4.1.5 4.17.4.1.6 4.17.4.2 4.17.4.2.1 4.17.4.2.2 4.17.4.3 4.17.4.3.1 4.17.4.3.2 4.17.4.3.3 4.17.4.3.4 4.17.4.4 4.17.4.4.1 4.17.4.4.2 4.17.4.4.3 4.17.5 References
Introduction Functional Requirements Material Selection Fabrication and Microstructure Advantages and Limitations for Fusion Application High atomic number: material erosion/melting Recrystallization Machinability, mechanical properties, and DBTT Component fabrication: CTE mismatch with heat sink Neutron embrittlement Neutron activation and radiological hazards Material availability Tungsten Grades Influence of In-Service Conditions Thermal Shock Resistance Microstructure, composition, and mechanical properties Power density and pulse duration Base temperature Repetition rate Thermal shock during off-normal events: disruptions Thermal shock during normal operation: ELMs Thermal Fatigue Resistance ITER Prototype and commercial reactors Neutron Irradiation Thermophysical properties and swelling Mechanical properties Thermal shock on irradiated W Thermal fatigue on irradiated W components Ion Irradiation and Retention He-irradiation Hydrogen-irradiation and retention Combined loading conditions Conclusion
Abbreviations APS AUG CFC CTE CVD DBTT DEMO
Atmospheric plasma spraying ASDEX-upgrade Carbon fiber composite Coefficient of thermal expansion Chemical vapor deposition Ductile to brittle transition temperature Demonstration fusion reactor
ECAE ECAP ELMs fpy FTU ICRH IFE
552 553 554 554 555 556 556 556 557 557 558 558 559 561 561 561 562 562 563 563 564 566 566 567 568 568 569 569 570 570 570 571 575 576 576
Equal-channel angular extrusion Equal-channel angular pressure Edge localized modes Full power years Frascati tokamak upgrade (Frascati, Italy) Ion cyclotron resonance heating Inertial Fusion Experiment
551
552
Tungsten as a Plasma-Facing Material
IFMIF
International Fusion Materials Irradiation Facility ITER Tokamak, Latin for ‘the way’ JET Joint European Torus (Culham, UK) LPPS Low-pressure plasma spraying MIM Metal injection molding NIF National Ignition Facility (Livermore, CA, USA) PFC Plasma-facing component PFM Plasma-facing material PS Plasma spraying PVD Physical vapor deposition SC Single crystal SPS Spark plasma sintering TEXTOR Tokamak EXperiment for Technology Oriented Research (Ju¨lich, Germany) TZM Ti–Zr–Mo VPS Vacuum plasma spraying
Symbols cp Tm l r
The specific heat Melting temperature Thermal conductivity Density
4.17.1 Introduction Until the mid-1990s, only few fusion devices used high-Z elements in plasma-facing materials (PFMs).1 These devices either operated at high plasma currents and high plasma densities such as Alcator C-Mod2 and Frascati tokamak upgrade (FTU)3,4 or used high-Z materials only as test limiters such as Tokamak EXperiment for Technology Oriented Research (TEXTOR).5–9 Since then, high Z refractory metals have been attracting growing interest as candidates for PFMs because of their resistance against erosion and the need for low erosion and stability against neutron irradiation.10 Considerable effort has been made to study the behavior of high Z impurities in the core and edge plasmas, erosion/redeposition processes at the limiter/divertor surfaces, hydrogen isotope retention, and on material development and testing. In particular, the modification of ASDEX-upgrade (AUG) into a fully tungsten machine,11–17 which was achieved in 2007, provided positive answers to critical
questions on the reliability of tokamak operation with high-Z plasma-facing components (PFCs) and the compatibility with standard and advanced H-mode scenarios and with the available heating methods.10 Among the challenges, for tokamak devices, that still remain are the strong increase of the W source and W concentration resulting from ion cyclotron resonance heating (ICRH) and the need for rigorous modeling to support the extrapolation of current results to ITER conditions. Clearly, not all questions posed by ITER can be answered by AUG only. For example, the effects of material mixing with Be, the melt behavior under transients, or the change of the hydrogen retention due to damage by high-energy neutron irradiation18 cannot be addressed in AUG. Answers to some of these issues may be provided by the ITERlike wall project in Joint European Torus (JET), which is installing a bulk tungsten component for the strike point and physical vapor deposition (PVD)-W-coated carbon fiber composite (CFC) tiles for the remaining parts of the divertor.19–21 The remaining questions have to be answered by dedicated experiments in other plasma devices or can only be assessed by modeling. However, the results obtained so far do not exclude the use of W in ITER as a standard PFM.10 Further investigations related to future fusion power plants such as demonstration fusion reactor (DEMO) have to focus on the minimization of plasma heat loads to the PFCs to increase their lifetime. In particular, transient heat loads caused by instabilities significantly decrease the operation domain of PFCs, due to thermal stresses and consequent enhanced erosion.22 Therefore, it is also important to mitigate all instabilities, such as edge localized modes (ELMs), that cause significant plasma transient heat losses.23 Plasma scenarios need to be developed, such that the conditions for achieving the required fusion yield are maintained in steady state, while at the same time sustaining tolerable heat loads on the PFCs. The above-mentioned upgrades to the JET24 and AUG15 will allow further optimization of the plasma scenarios under these conditions, in particular with DEMO relevant tungsten PFCs.25 These investigations will show how the identified deficiencies of W can be overcome or how they have to be dealt with. In addition to the application of tungsten in ITER and in potential future tokamak devices such as DEMO,26–29 tungsten also became an interesting alternative for the divertor of stellarators, for example, Advanced Reactor Innovation Evaluation Studies – Compact Stellerator (ARIES-CS),30 and as a first
Tungsten as a Plasma-Facing Material
wall material for inertial fusion devices.31 Due to similar demands on the PFMs during the operation of all these devices, similar problems have to be solved for each application.
4.17.2 Functional Requirements In the current design of the ITER divertor32–34 for the start-up phase, tungsten has been selected as armor for the divertor dome and the upper part of the divertor vertical targets. In addition, due to excessive co-deposition of tritium in CFC raising regulatory concerns related to tritium inventory limits, a full tungsten divertor will be installed before the D–T phase of operation.32 The PFC design for ITER consists of bulk W bonded to an actively pressurized water-cooled Cu alloy heat sink. Here W has no primary structural function. However, due to the operating conditions listed in Table 1, the PFMs face large mechanical loads particularly at the interface to the heat sink material during cyclic steady state heat loads (see Section 4.17.4.2) and at the plasma-loaded surface during transient thermal events (see Section 4.17.4.1). Furthermore, the material response to these loads is influenced by the material damage or degradation due to neutron irradiation (see Table 1, Sections 4.17.4.3.3 and 4.17.4.3.4). Table 1
553
Along with thermally induced loads, the interaction of the PFM with the plasma, that is, the hydrogen isotopes D and T as well as the fusion product He, is of importance (see Section 4.17.4.4) because they have an influence on material erosion and nearsurface material degradation. The further development of the ITER design led to four conceptual designs for the DEMO divertor.25,35 These designs include either water (inlet 140 C/outlet 170 C) or, due to the higher achievable efficiency, more probably He-cooling (inlet 540 C/outlet 700 C). In all cases bulk W is foreseen as the armor material that will have to face peak steady state heat loads of 15 MW m2 in case of the water-cooled design and 10 MW m2 for the He-cooled designs. In contrast to ITER, off-normal events such as disruptions have to be avoided completely and transient thermal events during normal operation, for example, ELMs, have to be mitigated below the damage threshold of the material (see Section 4.17.4.1). This may be particularly important considering the expected neutron damage that will amount up to 40–60 dpa during the planned operation of the fusion reactor35 leading to a significant amount of transmutation products.36 However, the main limiting factor is expected to be the material’s erosion leading to a maximum lifetime of 2 years for the divertor armor.35
Operating conditions assumed for the design of the ITER PFCs during D–T operation
Material Number of replacements Baking temperature ( C) Normal operation Lifetime (number of cycles) Peak surface heat flux (MW m2) Peak particle flux (1023 m2 s1) ELM energy density (MJ m2) controlled/uncontrolled ELM duration (ms) ELM frequency (Hz) controlled/uncontrolled Maximum radiation damage (dpa) Operation temperature design window during normal operation ( C) Off normal operation: disruptions Peak surface heat load (MJ m2) Duration (ms) rise time/decay time Frequency (%)
Divertor target
Divertor baffle/dome
CFC/W 3 240
W 3 240
3000–10 000 10a 10 0.3–0.5/6–10 0.25–0.5 20–40/1–2 0.7b 200–1000
3000–10 000 3 <0.1 – –
4–40 1.5–3/1.5–6 <10
– –
0.6b 200–600
Source: Federici, G.; Wuerz, H.; Janeschitz, G.; Tivey, R. Fusion Eng. Des. 2002, 61–62, 81–94; Loarte, A.; Saibene, G.; Sartori, R.; et al. In Proceedings of the 22nd IAEA Fusion Energy Conference, Geneva, Switzerland, Oct 13–18, 2008; IT/P6-13; Raffray, A. R.; Nygren, R.; Whyte, D. G.; et al. Fusion Eng. Des. 2010, 85, 93–108. a Slow transients lasting 10 s up to 20 MW m2 (10%). b Without replacement.
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Tungsten as a Plasma-Facing Material
In comparison to tokamaks, calculations for a device such as ARIES-CS predict steady state heat loads between 5 and 18 MW m2.30,37 Similar to DEMO, a He-cooled W divertor is anticipated with a maximum heat removal capability of 10 MW m2. The design limits for neutron irradiation at the shield of ARIES-CS are up to 200 dpa at 40 fpy (full power years).38 The component lifetime limits are similarly dictated by the material’s expected erosion. Finally, tungsten or more specifically tungsten coatings find their application also in the dry wall concept for inertial fusion devices, for example, the National Ignition Facility (NIF). In future, inertial confinement devices, thermal loads will occur only in the form of transient thermal loads (P ¼ 0.1 MJ m2, t ¼ 1–3 ms, f ¼ 5–15 Hz, Tbase 500 C).31 These are similar to those expected during ELMs and almost identical to those occurring in an X-ray anode39 and, therefore, affect a thin surface layer only.
4.17.3 Material Selection 4.17.3.1
Fabrication and Microstructure
Tungsten and tungsten alloys are commercially available in many forms, for example, as bulk rods, plates and discs, or thin coatings on various kinds of substrates. For each of these tungsten products, optimized production routes exist involving mainly powder metallurgical techniques for bulk materials and PVD and chemical vapor deposition (CVD) as well as plasma spraying (PS) for coatings. Each of these processes has its own advantages and disadvantages as well as an individual influence on the material’s microstructure and subsequently the material properties. In addition to the fabrication method, the raw materials, the alloying elements and dopants/impurities, pre- and postthermomechanical treatments, and the final shape/geometry have a strong impact on the achieved microstructure. Focusing on the powder metallurgy fabrication route, tungsten powder is obtained from ammonium paratungstate ((NH4)2WO4), tungsten oxide (WO3), and tungsten blue oxide (WO3x) by hydrogen reduction at temperatures in the range of 700–1100 C. Various grain sizes can be produced depending on the reduction temperature and the hydrogen dew-point. The purity of the metal powder obtained is >99.97%. In the manufacture of doped or alloyed tungsten products, the dopants or alloying elements are either introduced into the raw materials before reduction or they can be added to the metal powder after reduction.
Following the reduction stage, the powder is sieved and homogenized. The initial densification of the powder in various plate and rod geometries takes place predominantly through die pressing and cold isostatic pressing. The pressed compacts are subsequently sintered at temperatures between 2000 and 2500 C (2273–2773 K), mostly using furnaces with hydrogen flow. This increases the density and the strength of the pressed blanks.40 After sintering, the products have a rather low density of about 80% of the theoretical value and poor mechanical properties. To increase density and improve mechanical properties, the sintered products are subject to a mechanical treatment such as rolling, forging, or swaging at temperatures up to 1600 C. Intermediate annealing, leading to recovery and recrystallization, is necessary to maintain sufficient workability. The working temperature can be reduced as the degree of deformation increases. In this way, forged parts such as rods and discs as well as sheets and foils are produced.40 The final step, that is, the mechanical treatment, changes the microstructure from isotropic with grain sizes determined by the initially used powder size into anisotropic. Depending on the deformation method, the grains may show either: an elongated, needle-like structure along the deformation direction for radially forged rods and uniaxially rolled plates (see Figure 1(a)), or (a)
200 mm (b)
200 mm Figure 1 Light microscopy images of etched cross-sections of (a) a deformed rod and (b) a rolled plate.
Tungsten as a Plasma-Facing Material
a flat disc-shaped structure for axially forged discs or blanks and cross-rolled plates (see Figure 1(b)). In addition to bulk materials, research and development is also directed on tungsten coatings. One possibility would be the plasma-spraying process, in which powders are injected into a plasma flame, melted, and accelerated toward the (heated) substrate. The deposited layers are splat-cooled, leading to a flat discshaped microstructure. Depending on the atmospheric conditions, the result may be layers with high porosity and oxygen content (water stabilized and atmospheric plasma spraying, APS, see Figure 2(a))41,42 or low porosity and good thermal contact (low-pressure or vacuum plasma spraying, LPPS/VPS).26,43–47 In contrast, PVD and CVD coatings show a columnar structure perpendicular to the coated substrate with grain sizes in the range of the coating thickness (see Figure 2(b)). PVD coatings, which are also used as thin intermediate layers below a plasma-sprayed tungsten top layer,48 are deposits of tungsten vapor on the substrate surface, which is in the source’s line of sight.43,49 CVD coatings are reactions of a W-containing gaseous phase and have the ability to coat complex geometries.6,50–52 In both cases, a high density (100%) of the coatings is achieved. The coated substrate can be graphite as used for AUG (PS),12,13 CFC as used for the ITER-like wall
(a)
project in JET (PVD),21,53–55 or copper and steel as it might be used for first wall applications in future fusion devices (PS, PVD, CVD).44,49,50,56–60 4.17.3.2 Advantages and Limitations for Fusion Application For fusion plasma-facing applications, the most essential properties are thermal conductivity, strength and ductility, thermal shock and thermal fatigue resistance, structural stability at elevated temperature, and stability of the properties under neutron irradiation. The advantages and disadvantages of tungsten for these conditions are manifold and opposed to each other as shown in Table 2. While the advantages of the material are mainly related to its high temperature-handling capability, the limitations are associated with manufacturing and handling at low temperatures (below ductile to brittle transition temperature, DBTT61–63), plasma compatibility including neutron irradiation, and radiological issues. However, with regard to other potential PFMs, for example, Be (see Chapter 4.19, Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices), CFC (see Chapter 4.18, Carbon as a Fusion Plasma-Facing Material), and Mo, tungsten is still the most promising, offering an advantageous combination of physical properties and, therefore, has become the material of choice for ITER and DEMO. Since this decision was made, R&D efforts for investigating newly developed tungsten grades
Table 2
50 mm (b)
Disadvantages
High melting point Low erosion (high
High Z (low allowed
Figure 2 Light microscopy images of etched crosssections of (a) atmospheric plasma spraying W and (b) chemical vapor deposition W on a graphite substrate.
Features of W armor materials
Advantages
300 mm
555
energy threshold for sputtering) High thermal stress resistance High thermal conductivity Low swelling Low tritium retention
concentration in plasma)
Potential loss of melt layer
during transient events Recrystallization Poor machinability High DBTT High CTE mismatch with Cu or stainless steel heat sink Neutron embrittlement Irradiation-induced transmutation High radioactivity (short-term waste, decay afterheat) Explosion dust potential Limited resistance to grain growth
556
Tungsten as a Plasma-Facing Material
and alloys that are able to overcome or at least mitigate some of the above-mentioned disadvantages have significantly increased. 4.17.3.2.1 High atomic number: material erosion/melting
As the high atomic number is an intrinsic material property that cannot be changed, the only possibility to avoid plasma contamination by tungsten is to adapt to the loading realities, that is, thermal loads and plasma wall interaction conditions, and the energy of the incident plasma particles. In particular, surface crack formation, loosening of particles, and particle ejection or melting are addressed (see Section 4.17.4). Concerning the latter, the addition of suitable alloying elements or dispersoids (see Section 4.17.3.3) reduces the material’s thermal conductivity causing a reduction of allowed applied heat fluxes. From this point of view only low-alloyed grades should be considered and the best grade is tungsten of high purity. 4.17.3.2.2 Recrystallization
Recrystallization is a thermally activated process. Therefore, it is expected that the activation energy of nucleation is dominated by small angle grain boundaries. The activation energy of grain growth is dominated by large angle grain boundaries.64 The temperature of recrystallization depends mainly on the deformation history, that is, the higher the degree of deformation, the lower the recrystallization temperature,65,66 and the chemical purity. When heated above the recrystallization temperature, the structure of tungsten is altered due to grain growth causing an increase in DBTT and reducing other mechanical properties, that is, strength and hardness.67 There are several possibilities for increasing the recrystallization temperature. Particle reinforcement and controlled formation of porosity are the best and most investigated options.68 For example, the higher recrystallization temperature of dispersion strengthened alloys results from the interaction between dispersoids and dislocations during hot-working; the higher the amount of hot-work, the finer are the dispersoid particles and the higher is the recrystallization temperature. During recrystallization, these particles prevent secondary grain growth and consequently, the recrystallization temperature of dispersion strengthened alloys may increase compared to pure W.67 Another example is highly creep-resistant doped/nonsag materials with aligned porosity acting as obstacles for dislocation movement as they are used in the lighting industry.69
Experience shows that incomplete recrystallization often helps to achieve the desired balance in material properties. If the operating temperature is well known, controlled recrystallization during application might be feasible as well.67 However, for operational conditions in nuclear fusion devices, it is expected that the very high thermal strain rates experienced in the thin layer heated by plasma disruption or any other transient thermal event will significantly affect the material’s microstructure and properties. 4.17.3.2.3 Machinability, mechanical properties, and DBTT
Mechanical properties of W strongly depend on variables such as production history, alloying elements, impurity level, thermomechanical treatment, and form of material. Depending on the production history and heat treatment, W and W-alloys could have anisotropic mechanical properties. This is expressed by showing significantly better properties in the direction of elongated grains (by rolling, forging, or due to deposition processes for coatings) but poorer properties in other directions.70 While reported data on single crystals (SCs) (e.g., Gumbsch62) and for isotropic materials (e.g., Kurishita et al.71) give a clear indication of the material’s performance, typically the reported data refer to the best orientation of the material as shown for fusion relevant tungsten grades in numerous publications.44,51,57,72–81 The properties in other directions, particularly the DBTT, could significantly differ.76 This will affect the operational performance, which is reflected by the orientationdependent thermal shock response.82 Tungsten is a body-centered cubic (bcc) refractory metal, with a comparatively low fracture toughness,61,83 high DBTT, and poor machinability, which is directly correlated to the material’s low ductility and low grain boundary strength.67 However, DBTT is an ill-defined property and depends strongly on purity, alloying elements, thermomechanical treatment, and, most essentially, the testing/loading conditions due to its deformation rate dependence.62,63 The obtained values vary over a broad temperature range from room temperature (RT) to several hundreds of degrees Celsius. The exact value depends on the stress state, for example, a three-dimensional state of stress in the sample leads to a lower DBTT. Although many other parameters influence the fracture of bcc metals, the DBTT is usually associated with the thermal activation of dislocation
Tungsten as a Plasma-Facing Material
kink pairs. Below this characteristic temperature the separation of a screw dislocation into three partial dislocations (which cannot easily recombine and are therefore more or less immobile) is responsible for the brittle behavior. Increasing temperature leads to thermal activation of the kink mechanism and increased ductility due to shielding of the crack tip.84 There is an empirical correlation between temperature and activation energy for brittle-to-ductile transitions in single-phase materials suggesting that the ratio between the activation energy and the DBTT gives approximately a value of 25.63 Another factor is the occurrence of interstitial solute elements, such as oxygen, carbon, and nitrogen, which even in very small amounts tend to segregate at grain boundaries, promoting intergranular brittleness and increasing the DBTT. Two ways can be used to get rid of or mitigate the negative effects of interstitial impurities: either a reduction of the grain size,84 to dilute their effect on a larger grain boundary surface, or the complete elimination of grain boundaries, as in SCs. The development of W-alloys essentially follows the first route, as the SC technique, although effective, is too costly. The conventional method to decrease the grain size of tungsten or tungsten alloys is to deform the material at an intermediate temperature, above the DBTT and below the recrystallization temperature.81,84–86 The formation of oxides and carbides of the alloy constituents helps to stabilize the grain boundaries and to dispersion strengthen the matrix at high temperature. Recently, mechanical alloying followed by powder densification has been applied to refractory alloys. Materials with a stabilized fine-grained structure and with the grain boundary strengthened by even finer dispersoids of TiC improve the low-temperature impact toughness of refractory alloys, leading to increased ductility even down to RT and create superplasticity at high temperatures.71,87–89 Another reliable method to increase the ductility at low temperatures and therefore reduce the DBTT is to alloy tungsten with the rather expensive element rhenium, which is a substitutional solute in the W lattice.67,83 As mentioned before, both material deformation and heat treatment influence the DBTT. A heat treatment slightly below the recrystallization temperature is able to significantly reduce the DBTT. In contrast, annealing above the recrystallization temperature reduces strength and hardness and increases the DBTT.67
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4.17.3.2.4 Component fabrication: CTE mismatch with heat sink
A mismatch between the coefficients of thermal expansion (CTEs) can lead to thermal stresses at the interface, which are detrimental to the component lifetime. This can occur with either Cu-based alloys or steels (steel is more likely to be used in case of coatings) such as that used for water-cooled designs, or to W and W-alloys in the He-cooled design. In particular W and W alloys, in the coldworked and stress-relieved condition, tend to delaminate in the direction parallel to the direction of deformation. Such delamination can occur during machining or during operation. To avoid failure due to delamination, the orientation of the texture has to be perpendicular to the surface of the joints,90 raising the question of the suitability of plasma-sprayed W coatings. Two possible options are recommended to mitigate the thermal stresses, that is, reducing the joint interface by introducing castellations or using smaller tiles,91–93 or introducing soft and chemically stable interlayers94,95 or graded layers.96–101 Despite the fact that surface finish has no direct effect on the performance of ITER-related components,94 it is recommended to avoid possible crack initiators in the armor design, such as castellations ending in the tile and to ensure accurate surface finishing.102–104 Designs that have been proven to reduce the tile and interface thermal stresses and to extend the component lifetime beyond the design limits are the macrobrush or the monoblock. The latter is the reference design for ITER105 because it provides the most reliable attachment and therefore a reduced risk of catastrophic cascade failure.106 Finally, the thermal treatment of W during joining manufacturing cycles might have an influence on the material’s properties. While the process temperatures during joining of W and Cu do not lead to any significant change of the W properties, in the case of high-temperature brazing of W to W alloys for the He-cooled divertor design,102 the recrystallization temperature of W has to be taken into account. 4.17.3.2.5 Neutron embrittlement
There are little data available for irradiated tungsten (see Chapter 1.04, Effect of Radiation on Strength and Ductility of Metals and Alloys). Based on results for other refractory alloys and limited data on tungsten, one would expect neutron irradiation to increase the strength and decrease the ductility of the tungsten armor largely through increases in the DBTT. To minimize the neutron-induced material
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degradation, it is reasonable to limit the operational conditions for components in a neutron environment to temperatures above 900 C where recovery of tungsten takes place,107 as the ductility loss is more pronounced below about 0.3 Tm. This is possible in the region close to the plasma-facing surface, but it is impossible in the heat sink region as tungsten will be in contact with materials that cannot operate at this temperature and stress concentrations in these ‘cold’ areas have to be avoided.108 In the case of ITER, Cu will be employed in the heat sink while steel is more likely to be used in DEMO, which has a higher operating temperature.29 Hence, a greater understanding of the irradiation response of tungsten at temperatures between 700 and 1000 C is needed.109,110 The effect of embrittlement is alleviated when operating above 250 C, although in the presence of He (produced by transmutation reactions) somewhat higher temperatures may be required.83 Although at intermediate temperatures (0.3–0.6 Tm), void swelling and irradiation creep are the dominant effects of irradiation, the amount of volumetric swelling associated with void formation in refractory alloys is generally within engineering design limits (<5%) even for high neutron fluences (10 dpa). Very little experimental data exist on irradiation creep of refractory alloys, but data for other bcc alloys suggest that the irradiation creep will produce negligible deformation for near-term reactor applications.110 4.17.3.2.6 Neutron activation and radiological hazards
The activation and transmutation of tungsten as a PFM is a critical issue, particularly concerning long-term storage and recycling times. Different studies on activation issues have been performed. These comprise the analysis of cross-sections for high-energy neutrons,111,112 studies on the heliumcooled lithium lead divertor for DEMO,113 the inertial fusion devices,109 other benchmark experiments,114,115 and modeling issues, for example, on the self-shielding ability of tungsten.116,117 Furthermore, it was shown that the long-term activation behavior is dominated by activation products of the assumed material impurities while the shortterm behavior is due to the activation of the stable W isotopes.113 For a short period of a few weeks, the latter causes a huge amount of decay-induced afterheat that has to be removed by continued active cooling.67 On the other hand, the accumulation of the highly radioactive transmutation product 186mRe was determined to be most critical, limiting the component
lifetime to a maximum of 5 fpy when using pure W or to 2 fpy when using Re-doped W before the limits for storage by shallow land burial could be exceeded.109 The dose rate limits for recycling after different applications are expected to be reached within 5 years115 to 50 years113 of storage or up to 75 years after end of plant life.118 Fischer et al.113 take a limit of 100 mSv h1 for remote handling into account, which might be a problem at the times when maintenance operations would be in progress. Taylor and Pampin118 give a value of 20 mSv h1 as the limit for allowing tungsten to be categorized as a recyclable material. The hands-on limit for tungsten should be achieved after about 200 years.115 Besides waste management, tungsten has also been investigated and evaluated according to characteristic radiological hazards that might occur when using it as PFM in tokamak fusion reactors. It was found that the tritium permeation into tungsten does not, in contrast to CFC, appear to be a major problem. However, due to neutron activation, the mobilization of activation products, for example, by forming volatile oxide species in the presence of steam and air, has to be limited by establishing shutdown requirements to avoid melting of tungsten in case of an accident. The potential exposure from mobilized activation products from the tungsten divertor may be modified by varying the operating conditions of fusion power and change-out time as well as the thickness of the divertor armor. The dose can be reduced by selecting shorter change-out times. However, the total lifecycle waste volume will be increased accordingly. A thinner divertor will produce less mobilized activation products while suffering a more restrictive shutdown requirement.119 4.17.3.2.7 Material availability
The quantity of W needed for the PFCs in a fusion device such as ITER or DEMO represents only a small fraction of the yearly production and the world’s reserves120 and its production can be easily satisfied by existing industrial capabilities. The same point is valid for stellarators and even more for inertial fusion devices, which only work with thin coatings. However, the issue of component lifetime has to be taken into account. Depending on the component lifetime, the recycling rate, and the storage time until a hands on level is achieved (see Section 4.17.3.2.6), the operation of numerous power plants may require an amount of tungsten that exceeds what is currently available from the market.
Tungsten as a Plasma-Facing Material
4.17.3.3
Tungsten Grades
Within current R&D programs for the selection and characterization of candidate grades of W and W alloys for fusion applications, many materials produced according to the schemes outlined above were investigated. These are discussed in the following section, which introduces some of their characteristics. The manifold production processes described below for pure W are also applicable to W alloys. Pure tungsten (undoped) Sintered W is the most readily available and cheapest grade with a grain size that depends on the initially used W powder. However, it is characterized by high porosity, low recrystallization temperature (1000–1200 C), and low strength at elevated temperature.96 The option of improving the sinterability by adding small amounts of activators (Ni, Fe)121 increases the radiological hazard due to additional activation products that have to be taken into account.119 Forged or swaged W offers an increased density and a refined microstructure compared to sintered material, resulting in higher ductility and mechanical strength. Forging and swaging are therefore the industrial production processes that are typically applied not only for pure tungsten but also for most kinds of tungsten alloys (see below). This grade of W is manufactured in block shape or more commonly in the form of rods with different diameters (90 mm)40 showing an anisotropic microstructure122 with elongated grains along the axial direction and an increasing grain size and porosity with increasing rod diameter. Thus, increasing rod diameter leads to a decrease in mechanical strength and ductility. For the production of monoblock tiles, such as those planned for ITER, rods with a minimum diameter of 30–35 mm are necessary. Rolled W is applied in the form of plates or foils with thicknesses from 0.02 to 20 mm.40,123,124 It offers a densified but layered microstructure that is strongly anisotropic, with flat disc-shaped grains parallel to the rolled surface affecting the mechanical properties (see Section 4.17.3.2.3) and resulting in the risk of delamination. Double-forged W is in the form of blanks with a diameter of 140 mm and a height of 45 mm. The double-forging process, first in the radial and then in the axial direction, provides a more
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isotropic microstructure than it is generated by single forging. This material should act as a reference grade for establishing a reliable materials database for finite element calculations.82 SC W provides higher ductility than polycrystalline W, higher thermal conductivity, lower neutron embrittlement, higher thermal fatigue resistance, and a more stable structure at elevated temperatures. The disadvantages are high cost and low industrial availability.96,125,126 Metal injection molded (MIM)-W127–129 provides a dense and isotropic microstructure with grain sizes on the order of the powder particle sizes used. A final densification by hot isostatic pressing (HIP) at temperatures >2000 C leads to an improvement of the mechanical properties; recrystallization and grain growth do not play a role. Furthermore, the production process offers the possibility of net shaping. Spark plasma sintered (SPS)-W and resistance sintering under ultra-high pressure.130–132 The material is characterized by a short fabrication time of only a few minutes keeping the initial fine microstructure determined by the powders used. The finer the grain size, the higher the microhardness and the bending strength but also the lower the achievable density. The application of alternatively uni-, two-, or threedirectional orthogonally applied forces for the material’s densification during the process leads to internal stresses, which have an influence on the recrystallization behavior. Recrystallization and grain growth occur at 1500 C. Depending on the amount of porosity, the finer the initial grain size of tungsten, the smaller is the grain growth. Severe plastically deformed W (and W alloys, see below) with ultra-fine grains in the nm range are produced by either high-pressure torsion at 400 C84,133 or by the equal-channel angular extrusion or pressure (ECAE or ECAP) process at high temperatures (1000–1200 C).134 The material shows stable, that is, deformationindependent, recrystallization temperatures and exhibits considerably enhanced ductility and fracture toughness.61,85,86,135,136 Plasma-sprayed W involves, in general, application of VPS, more precisely also called low-pressure plasma spraying (LPPS), which provides a significantly reduced oxygen content and improved thermophysical properties compared to atmospheric (APS) or water-stabilized
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Tungsten as a Plasma-Facing Material
plasma spraying.42 However, LPPS-W is typically characterized by a lower thermal conductivity (up to 60% of bulk tungsten is reported67) and a lower strength than bulk W particularly when deposited on large surfaces. The recrystallization temperature is similar to pure W.48,137 Although the thickness of the plasma-sprayed coatings required for fusion applications are flexible, coatings with 200 mm or thicker are commonly produced.26,43,67,138,139 Furthermore, PS is the only production method that offers the possibility to produce and repair W components.57,60,96 CVD W provides a microstructure with a columnar grain structure parallel to the surface, high thermal conductivity similar to bulk W, and a very high density and purity.6,140,141 Thicknesses up to 10 mm were produced,67 but its high cost is a significant drawback for practical applications.52,96 PVD W provides a featureless structure that is extremely dense and pore free. In contrast to plasma sprayed and similar to CVD-coatings, the deposition rates are low. Economic and process-related restrictions generally limit the deposited W thickness to 10–50 mm.13,54,55,67,142 W foam for Inertial Fusion Experiment (IFE) applications provides structural flexibility during quasivolumetric loading. The material is microengineered with a relative density of 21% and can be simultaneously optimized for stiffness, strength, thermal conductivity, and active surface area.143 Tungsten alloys Oxide dispersion strengthened W alloys such as W–La2O3, W–Y2O3, and W–CeO2 with oxide additions 2% are processed by powder metallurgy methods similar to pure W.40 The insoluble dispersoids, which are influenced in shape and distribution by the thermomechanical treatments during the production process,73,144 improve the grain boundary strength and machinability and play an important role in controlling recrystallization and the morphology of the recrystallized grains.68 This results in a higher recrystallization temperature by 100–350 K by suppression of secondary grain growth (i.e., grain boundary migration), lower grain size, higher strength after recrystallization, and better machinability than sintered W even at RT. This permits fabrication at
lower costs.67 The size of the dispersoids in commercially available alloys is 10 mm; however, research on mechanically alloyed materials using submicron dispersoids is currently being performed.145 However, the presence of oxide particles with a melting temperature below those of tungsten has a negative effect on the erosion resistance.146,147 W–3–5% Re is, compared to sintered pure W, characterized by a higher recrystallization temperature and strength even after recrystallization,148 better machinability, and improved ductility at low temperatures.67 The addition of Re, which has a high solubility in W, however, reduces thermal conductivity, increases embrittlement after neutron irradiation, and significantly increases the cost and safety concerns because of the high Re activation under neutron irradiation.96 W–1–2% Mo (þY and Ti) cast alloy. The addition of Mo and the reactive elements Y and Ti, which reduce the amount of free oxygen and carbon and form obstacles to grain growth, improves the mechanical properties compared to large grained pure cast W.67,73 W–TiC produced by mechanical alloying and slow deformation techniques provides, similar to all other W alloys, higher strength and recrystallization temperature, better machinability, and improved ductility compared to pure W with superplastic behavior at temperatures of 1400–1700 C.89 The addition of Ti-carbide particles stabilizes the grains during the material’s production process. This generates an isotropic grain structure and has the additional effect of keeping a fine grain structure even in the recrystallized condition, but the alloy is more expensive. After recrystallization, the finer dispersoids of TiC particles improve the lowtemperature impact toughness of refractory alloys following low-dose neutron irradiation.71,87–89,149–154 Other carbides, for example, ZrC155,156 or HfC (in combination with Re and Mo),96 can be used instead of TiC. K-doped W is a nonsag material that contains a maximum of 40 ppm of potassium.40 Originally known from the lighting industry, it provides high creep strength due to its aligned pore structure, high recrystallization temperature >1600 C, and good machinability.68,77,78,157 W–Si–Cr as a ternary or even by the addition of another element as a quarternary alloy is a
Tungsten as a Plasma-Facing Material
newly developed and not yet optimized material that is being investigated as a wall protection material due to its favorable oxidation resistance, preventing excessive material erosion in case of accidental air ingress.158,159 Severe plastically deformed W alloys offer, similar to pure tungsten (see above), significantly improved fracture toughness and ductility.61,84 The addition of alloying elements to the starting material (any developmental or commercial produced W alloy), such as Re or dispersoids, leads to an increasing stability of the grains and therefore a higher recrystallization temperature and less grain growth.133 Any of the bulk materials mentioned above could be used and are being investigated in its cold-worked, stress-relieved, or recrystallized state. The latter is of particular interest due to in situ recrystallization of surface near regions during operation.108 In spite of the fact that a large variety of tungsten grades and alloys already exist, the attempts to further optimize these materials are ongoing. The fabrication and successful testing of He-cooled divertor mock-ups for DEMO and ARIES-CS102,160 under a heat flux of 10 MW m2 are important driving forces for the present development of W alloys with improved performance in the fusion environment.25 However, R&D has to address many different issues related to the performance of the material when exposed to thermal loads, neutron irradiation, and the plasma; these will be discussed in the following section.
4.17.4 Influence of In-Service Conditions 4.17.4.1
Thermal Shock Resistance
Tungsten-armored PFCs will be subjected to different types of heat fluxes dependent on their field of application (see Section 4.17.2). Among others, this includes thermal transient loads (e.g., ELMs and disruptions). The behavior of the material under these conditions, that is, the combination of cyclic steady state and transient heat loads, is a key factor that has to be considered for the selection of a suitable grade of W. The machines simulating these operational conditions are electron and ion beam facilities, quasistationary plasma accelerators, plasma guns, and high-energy lasers. A most critical issue is the comparability of such simulations. Therefore, a round robin
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test involving some representative facilities was made for investigating the influence of the different time regimes and different power density levels. The results showed that when compared on the basis of a heat flux parameter P (MW m2 s1/2), which is directly proportional to the temperature increase, the cracking and melting thresholds are almost identical. This permits a direct transfer of the qualitative results obtained in any of these facilities.161 In contrast, quantitative results representative of the operational conditions in large fusion devices can only be obtained when the loads are applied in the desired time range. The reason for this is the heat penetration depth and the related stress field that is produced, which influences crack and melting depth. There are several parameters influencing the thermal shock behavior of tungsten that will be discussed in the following sections for the different materials under disruption and ELM-like loads. 4.17.4.1.1 Microstructure, composition, and mechanical properties
During thermal shock loads, steep temperature gradients of hundreds to several thousand degrees Celsius on a length scale of millimeter or even micrometer (depending on the pulse length) are formed, influencing only a limited volume near the loaded surface. While the heat load is applied, due to thermal expansion and the decreasing strength of the material at the surface compared to the bulk material, compressive stresses are formed in the surface plane. These stresses can lead to permanent plastic deformation that might, during cool down, generate tensile stresses sufficiently high to initiate crack formation perpendicular to the surface and thereby cause stress relaxation at the surface. Depending on the mechanical properties in the surface plane, the amount and starting point of crack formation can be influenced. Based on this and on the fact that the mechanical properties are strongly dependent on the material’s microstructure (see Section 4.17.3.2.3), a grain orientation parallel to the surface and therefore high strength in the surface plane might be preferred.162 However, grains oriented parallel to the surface, such as in rolled materials or plasma-sprayed coatings, might result in delamination (see Figure 3(a)), which causes overheating and subsequently surface melting if they have a lower strength in the depth direction and exhibit preferential cracking along the weak grain boundaries. Therefore, a grain orientation perpendicular to the surface and parallel to the direction of the heat
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(a)
200 mm
(b)
100 mm
500 mm
Figure 3 Light microscopy images of the etched cross-sections of thermal shock–loaded specimens with grains oriented (a) parallel and (b) perpendicular to the loaded surface; cracks follow the grain orientation/ deformation direction.
Figure 4 Light microscopy images of the etched cross-sections of thermal shock–loaded metal injection molding tungsten with isotropic grain structure.
flow is recommended.90 This will cause cracks to form along the grain boundaries toward the cooling structure (see Figure 3(b)) causing no degradation or only a negligible degradation of the material’s thermal transfer capabilities. Due to the lower mechanical properties in the surface plane, more or larger cracks will form during thermal shocks, running perpendicular to the surface and following the grain orientation. In contrast to deformed materials, crack formation and crack orientation in materials with isotropic or almost isotropic grain structures, for example, MIM-W or recrystallized W, is rather unstable and is strongly enhanced for the weakened recrystallized material. Depending on the applied power densities, the formed temperature gradient, and the resultant stress fields within the material, cracks initially running perpendicular to the surface might deflect at zones with compressive stresses and keep running parallel to the surface (see Figure 4).
operation) of the material, the damage and cracking threshold are determined mainly by the material’s mechanical properties. Damage here means that the material’s surface has undergone a visible and measurable modification, for example, by surface roughening, recrystallization, or pore/void formation.
4.17.4.1.2 Power density and pulse duration
The material’s response is strongly related to the applied temperature fields and by this to the absorbed power density and the pulse duration. This results in a material-related surface temperature increase and heat penetration depth.163 A classification of the impact of the temperature field is made by establishing three parameters: the damage, the cracking, and the melting threshold. While the latter depends on the thermal conductivity and the melting temperature (for alloys or mixed materials formed during tokamak
4.17.4.1.3 Base temperature
The base temperature influences the thermal shock behavior in two ways. First, a higher base temperature influences the damage, cracking, and melting threshold. All of them are essential because they limit the operational conditions and when exceeded cause enhanced material degradation. Therefore, lifetime estimates based on RT data will yield unrealistic conclusions. Second, crack formation strongly depends on the plastic deformation at high temperatures and even more on the stress developed during cool down. To understand the influence of a higher base temperature, one has to be aware of the typical shape of the yield and tensile strength curve for W or a W alloy.105,157 While the decrease in strength is rather high at low temperatures, the curve flattens at high temperatures despite a drop in strength when exceeding the recrystallization temperature. As a result, the high temperature plastic deformation induced by the combination of a heated surface and ‘cool’ base material can be significantly reduced by a small increase in base temperature. Combining this effect with the increased ductility of W at the given base temperature, brittle crack formation can be avoided when heating
Tungsten as a Plasma-Facing Material
the material above a certain threshold.82,157,164,165 This temperature threshold is related to the DBTT but is not necessarily identical to it. 4.17.4.1.4 Repetition rate
In addition to the parameters mentioned above, the damage, cracking, and melting thresholds are determined by the number of load repetitions, because of continuing material degradation such as hardening and recrystallization. This is of particular interest for short transient events with a high repetition rate in magnetic (ELMs) and inertial fusion devices. Up to now the simulation of submillisecond events (ELMs, IFE) has been performed only up to a relatively low number of cycles; large numbers of pulses (e.g., >106 ELM pulses during the life-time of the ITER divertor) are not feasible in the majority of the above-mentioned test facilities. 4.17.4.1.5 Thermal shock during off-normal events: disruptions
Disruptions still occur frequently in operating tokamaks, and therefore they are also expected for ITER with an anticipated occurrence in 10% of the ITER pulses (3000 pulses per expected component lifetime). During a disruption in which the plasma undergoes a partial or full thermal quench, most of the plasma thermal energy will be dumped on the divertor plates.166 Taking into account the resultant loading conditions (see Section 4.17.2), significant material loss from the tungsten plasmafacing surface should occur by melting and evaporation particularly in the dome area.167,168 In simulating these events, the amount of melting, the melt motion and subsequent roughening of the surface, the material erosion by droplet emission, the resolidification behavior, and finally, the crack formation occurring in the loaded area or at the boundary between melted and unmelted zone are the most important parameters to be determined. The underlying mechanisms for the abovementioned material degradation are well described (see Figure 5).169 Thermal loading of tungsten and metals, in general, at ‘moderate’ energy densities (up to a few MJ m2) will result in a homogeneous, localized melting of the sample surface. When higher energy densities are applied, surface evaporation occurs; the momentum transfer due to evaporating atoms from the surface generates an effective pressure on the melt layer, which finally results in the formation of a melting ridge. Increasing the incident energy density even further, the material’s response is characterized by intense
Incident beam
Cracking roughening
Homogeneous melting
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Incident beam
Melt ejection
Boiling and droplet formation
Increasing energy density
Figure 5 Performance of tungsten and metals in general under transient thermal loads.
boiling and convection of the melt layer resulting in droplet formation and ejection.170–172 Open pores in the recrystallized material have a strong impact on the thermophysical properties. The melting threshold and subsequently the amount of melt formation depend on the material’s thermal conductivity, which is lower for porous materials such as plasma-sprayed tungsten, and for tungsten alloys. In particular, it has to be taken into account that dispersoids such as La2O3 (Tm ¼ 2578 K) have a lower melting temperature than tungsten. This may result in early melting and increased evaporation causing the formation of a porous and depleted surface layer, which becomes even more important when applying loads below the melting threshold (see below and Section 4.17.4.1.2). On the other hand, the melting threshold is correlated with the base temperature of the PFM. When the base temperature increases, the melting threshold energy decreases and the amount of melt formation, the obtained crater depth, and the evaporation losses for the same applied loading conditions increase significantly.169 As it cools, the material resolidifies in a recrystallized state providing a columnar grain structure typical of PVD or CVD coatings. With further cooling, depending on the base temperature of the material/component (see ‘Base Temperature’ in Section 4.17.4.1.3), brittle crack formation will not take place above a certain threshold temperature. However, with fast cooling after loading below this temperature, the material will undergo severe cracking with crack lengths that can reach the order of millimeters.169 When the qualification of different W grades and alloys108,141,147 is done in combination with thermal fatigue loading,90 materials with high thermal conductivity in combination with superior mechanical properties, that is, with high ductility, performed best with regard to melt material loss and crack formation. This comprises low-alloyed W materials with increased ductility such as W–Re or W–Ta, or fine-grained pure W or W alloys.
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Disruption simulation experiments on bulk tungsten and tungsten coatings have also been described in the literature. These were performed not only to investigate the melting behavior but also for the purpose of characterizing the cracking behavior.26,42,60,101,122,130,131,162,165,173–176 Although these experiments are more related to those on the characterization of ELM conditions (see Section 4.17.4.1.2) and were often performed only at RT, the results indicate that the use of highly ductile SC materials is preferred.90,177 Alternately, in case of cheaper polycrystalline materials, it is necessary for the material to have the proper microstructure orientation as described above, that is, the grain orientation perpendicular to the loaded surface. The reason for this is that crack formation occurs mainly along the grain boundaries and follows the orientation of the deformed microstructure. The crack depth is, in general, related to the applied loading conditions and therefore the pulse length, which determines the heat penetration depth and the temperature and stress gradient induced during loading. The temperature gradient also determines the recrystallization zone, which is generated below the loaded area as a function of temperature (
plcp r P t¼ 2 a decrease in thermophysical properties consequently reduces the heat flux parameter and the melting threshold. As all performed investigations indicated that melting will cause increased material degradation
and the continuous erosion of the PFM, it will significantly limit the lifetime of the PFCs. Therefore, the safe and economic operation of a future fusion reactor requires that scenarios causing melt formation have to be limited to a minimum. 4.17.4.1.6 Thermal shock during normal operation: ELMs
In contrast to disruptions, ELMs occur during normal operation in the H-mode and are characterized as instabilities caused by the steep temperature and density gradients at the plasma edge, which deposit a significant amount of energy at a high repetition rate.181,182 In particular, it is the expected high repetition rate for ELMs during the lifetime of the PFC (>1 million of events at a frequency of 1–25 Hz183) that, although yet unexplored, will impose high demands on the PFMs. While it is the desire of plasma physicists to operate in H-mode regimes with high-energy ELM deposition (1 MJ m2), the response of bulk tungsten, tungsten coatings, and tungsten alloys to such loading conditions, that is, surface melting, melt motion, material erosion, and vaporization,170,171,184–189 is detrimental. To obtain further insight into material behavior under these conditions, modeling of experimental conditions was carried out.9,167,168,190–195 It has been shown that with regard to melt motion/ erosion, the results of the different facilities cannot be directly compared196 and none of the testing facilities used provides identical conditions to those that will occur in a tokamak. However, mitigation techniques have been explored for reducing the applied ELM energy, which, in general, can only be done at the expense of a higher repetition rate.183 The extent to which the ELMs have to be mitigated depends on the melt formation at tile edges due to the shallow plasma impact, which was experimentally found to be between 0.4 and 0.6 MJ m2 for pure forged tungsten.189,197 On the other hand, the effect of crack formation during ELMs on the lifetime behavior of the PFCs has to be taken into account. As mentioned before, this behavior is yet unexplored at high repetition rates. Typical investigations on various grades of W,82,187,189,198 coatings21,54,186,199 and alloys146,157,187 were in the range of 10–100 repetitions. In a few cases up to 1000 repetitions and in single experiments even on the order of tens of thousands of repetitions have been obtained, depending on the testing facility used. As the repetition rate is still rather low compared to the expected millions of events, the main interest of
Tungsten as a Plasma-Facing Material
Cracks
Surface modification
No damage
40 35 30
threshold
Heat flux parameter (MW m–2 s–1/2)
45
25 20 Cracking
15 10
Heat flux
Damage threshold
5 0 0
200
400 Temperature (⬚C)
600
200
400 Temperature (⬚C)
600
45 40 Cracking threshold
Heat flux parameter (MW m–2 s–1/2)
(a)
35 30 25
Heat flux
20 15 10 5 0 0
(b)
Figure 6 Thermal shock testing results of double forged W as a function of temperature and the heat flux parameter; grain orientation (a) perpendicular and (b) ‘parallel’ (one direction still perpendicular, indicated by the orientation of the large cracks) to the heat flux.
these investigations was the qualification of different W grades and alloys (see Section 4.17.3.3) with regard to their damage and cracking thresholds. The characterization was done as a function of the main parameters described in Section 4.17.4.1, that is, microstructure, power density, and base temperature. The results obtained so far showed that crack formation200 vanishes above a certain base temperature (see Figure 6).82,157,198 This temperature decreases with increasing material ductility, indicating that the use of W alloys or fine-grained W is preferred. In the case of an anisotropic microstructure, this effect strongly depends on the material’s orientation. Better results are obtained for grain orientations parallel to the loaded surface (see Section 4.17.4.1), yielding differences in the threshold temperature compared to the orthogonal direction of up to several hundred K (cf. Figure 6(a) and 6(b)). Recrystallization leads to a slight homogenization of the material’s
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microstructure and therefore the mechanical properties; however, there is no full convergency of the orientation-dependent thresholds.82 Despite the fact that for the currently limited number of applied pulses no crack formation was observed above a material and orientation-dependent temperature, the material is still damaged by plastic deformation and surface roughening. The evolution of this plastic deformation and of the related material hardening as a function of the applied number of loads is still unclear and has to be investigated. However, there are also heat load levels (at least up to Tbase 800 C), at which no visual material degradation could be determined and the future goal will be to investigate if these damage thresholds are still valid for high repetition rates, at higher base temperatures, and particularly in combination with neutron irradiation (see Section 4.17.4.3) and plasma wall interaction (see Section 4.17.4.4). All the information given above on the effect of ELMs is also directly transferable to the short transient events expected for inertial fusion applications and has been verified by IFE-related tests on different W-based materials.201–204 There are coating parameters of high interest besides those mentioned above; these include the manufacturing-induced residual stresses at the surface, which are dependent on the used substrate, and the coating thickness. As mentioned in Section 4.17.4.1.1, the applied loading conditions and therefore the pulse length determine the heat penetration depth.163 As a result, the temperature and stress gradient induced under IFE applications should be similar to those in X-ray anodes (see Section 4.17.2). In case of thin coatings, residual and induced stresses might affect the coating to substrate interface and could lead to interfacial crack formation and delamination. This leads to minimum requirements for coating thicknesses that depend on the applied loading conditions.54 For example, in industrially produced X-ray anodes, W–Re coatings are typically used with a thickness of 200–700 mm205,206 to provide better mechanical and thermal-shock properties compared to pure W.204 However, the first experience on the influence of ELMs on coatings under real plasma operational conditions will be gained in the ITER-like wall project in JET, which involves testing relatively thin PVDtungsten coatings (14–20 mm) on a CFC substrate that provides a strong and anisotropic CTE difference.19,142 The behavior of this material under the above outlined transient heat loads is of course a key factor for the lifetime assessment of PFCs. However, the
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Tungsten as a Plasma-Facing Material
results obtained for pure thermal shock testing might underestimate the material damage and by this overestimate its lifetime. Only a combination of thermal shock, thermal fatigue (see Section 4.17.4.2), neutron irradiation (see Section 4.17.4.3), and plasma wall interaction (see Section 4.17.4.4) will be able to give appropriate answers for the selection of suitable grades of W. 4.17.4.2
Thermal Fatigue Resistance
The thermal fatigue resistance of tungsten is strongly related to its performance as part of current inertially, and future actively cooled components for application in magnetic fusion devices. The functional requirements these components have to fulfill are listed in Section 4.17.2. State of the art inertially cooled components include W coatings on graphite, CFC, and TZM.12,13,46,47,54,140,207,208 These concepts are used or are going to be used in the large operating tokamaks, for example, AUG and JET. During thermal loading, they mainly suffer from the problem of high interfacial stresses as a result of the CTE difference between the W coating and the substrate. Furthermore, interfacial reaction products and their potential reduced power handling capability have to be taken into account. Besides coatings, recent development of an inertially cooled bulk tungsten divertor for JET20,123 showed that under thermal fatigue loads the W quality is of minor importance for the integrity of the component. The major issue for the tungsten PFM was found to be the necessary shadowing of the plasma-loaded surface to avoid overheating and melting at tile edges as a result of the shallow angle of the incident plasma. This was realized by surface shaping.209 In the design of actively cooled components, tungsten is joined to a water-cooled Cu-based heat sink (ITER) or He-cooled steel or W-based heat sinks (e.g., DEMO, ARIES-CS). Direct cooling of the tungsten armor should be avoided, particularly without castellation, as the induced stresses might cause catastrophic material failure with subsequent water or He-leakage.124 Therefore, the only performance requirements are a sufficiently good surface quality to reduce possible crack initiation points and therefore suitable fabrication and surface finishing technologies,103,210,211 the chemical compatibility with the heat sink and, if present, the joining interface material, and the cyclic stability of the joint(s). The latter is influenced by the temperature gradient
applied during steady state heat loads, the difference of the CTEs, the quality of the joining process and, perhaps most important for reducing induced stresses, the tile size, or the dimensions of the castellated segments (see Section 4.17.3.2.4). Smaller tile sizes significantly improve the stress situation at the interface, and also at the top surface of the PFM. This has to be taken into account when comparing the thermal fatigue results of various kinds of components and the response of different grades and alloys of W, as shown by Makhankov et al.,90 where smaller tile sizes resulted in little or no crack formation. Furthermore, variations in the size of the component investigated can often explain the contradictory results presented in the literature that show good behavior of a material in one test while it fails in another. However, there are limitations to the minimum size of tiles and a compromise between operational and economical needs has to be made. Despite the fact that design and manufacturing technique seems to be more important than the mechanical properties and the microstructure of the particular W grade or alloy, the latter should still be considered. Similar to thermal shock results, the risk of delamination parallel to the loaded surface at the interface90 or anywhere in the bulk material has to be minimized. Therefore, the grain orientation of the PFM microstructure should be perpendicular to the loaded surface, although this still bears the risk of crack formation toward the cooling structure.108 To avoid subsequent water or He-leakage in case of crack propagation into the heat sink material, particularly in the He-cooled divertor design (see Section 4.17.4.2.2), suitable material and design solutions still have to be found. Furthermore, shadowing of adjacent tiles similar to the JET bulk W divertor has not yet been included in the design of the actively cooled components. 4.17.4.2.1 ITER
The performance of bulk W for ITER has been investigated using water-cooled divertor designs, that is, flat tile, macrobrush, and, most relevant, monoblock options. One important factor in these design solutions is the maximum allowable distance between the front surface and coolant to accommodate the heat without melting96 and, if possible, to avoid recrystallization during normal operating conditions. The ability to estimate this parameter requires not only the thermal conductivity of the materials but also the amount of allowable damage at the interface. This requires knowing not only the damage produced during
Tungsten as a Plasma-Facing Material
operation but also understanding the manufacturing accuracy and reproducibility because tens of thousands of armor/heat sink joints will be produced. Studies on this issue have shown that the current W monoblock design with a defect extension up to 50 appears to be suitable for the upper part of the vertical target (P ¼ 10 MW m2), but is not well adapted to a heat flux of 20 MW m2, which is necessary for application at the strike point of the vertical target, as systematic defect propagation was observed. A tungsten flat tile design with 6-mm long defects in the material interface was studied and proved to be compatible with fluxes of 5 MW m2 but was unable to sustain cyclic fluxes of 10 MW m2.212 These results confirm that the monoblock geometry generally proves to have superior behavior under high heat flux testing when compared with flat tile geometry. However, it is worthwhile to continue the investigation of the flat W tile design for low-flux regions despite the hazard of cascade failure of the flat tiles106 for two reasons: cost and weight. Besides this characterization, a number of high heat flux tests have been carried out on mock-ups and prototypes without artificial defects representing the different design options to assess the ‘fitness for purpose’ of the developed technologies.33,90,161,213–218 The results obtained for small test mock-ups of the flat-tile and monoblock design can be transferred to large-scale prototypes for the divertor vertical target. Independent of the type of pure Wor W–La2O3 armor material used in these prototypes, the W parts survived in the nonneutron-irradiated condition up to 1000 cycles at 20 MW m2 in the monoblock design213,219 and up to 1000 cycles at 18 MW m2 in the flat tile
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design (see Figure 7).213 This is far beyond the design requirements for use in the upper part of the vertical target (P ¼ 5 MW m2) and, in case of the monoblock design, even meets the design requirements for the strike point area of the vertical target. Alternative concepts such as explosive bonding of tungsten to a heat sink material,220 PS on a Cu-alloy216 or on EUROFER steel44 could probably be of use in the divertor but even more for first wall applications for fusion machines beyond ITER. However, these concepts often suffer from high interfacial stresses as a result of the CTE difference between the W coating and the substrate. 4.17.4.2.2 Prototype and commercial reactors
There are many design proposals for a He-cooled first wall and divertor concept for DEMO and ARIES-CS.37,160,221 Among these, the He-cooled modular design with jet cooling (HEMJ)102 is the most developed and qualified in terms of microstructural response,103 having survived at reduced coolant temperatures of 450–550 C at least 100 cycles at 11 MW m2 without failure. In contrast to the results obtained for water-cooled components for ITER, no influence of grain orientation on the components performance was observed.102 This might be a result of the higher temperature, which was always above the DBTT. Nevertheless, some difficulties in the design still have to be resolved. First, there are problems related to temperature with a desired coolant temperature of 600 C; these include material recrystallization at the top surface and the necessary high temperature joining to the W-based heat sink material. Second,
W macrobrush 0 dpa:
1000 cycles at 18 MW m−2
0.6 dpa: 1000 cycles at 10 MW m−2 (increasing of Tsurf)
W monoblock 0 dpa:
1000 cycles at 20 MW m−2
0.6 dpa: 1000 cycles at 18 MW m−2 (no degradation) Figure 7 Thermal fatigue testing results of W macrobrush and W monoblock mock-ups before and after neutron irradiation.
Tungsten as a Plasma-Facing Material
issues related to the material’s mechanical properties must be solved, in particular for the ductility of W-based structural material and its neutron irradiation resistance (see Section 4.17.4.3). Finally, the manufacturing reproducibility has to be at a high level because of the large number of small units (1 unit 3 104 m2) necessary for cladding the DEMO divertor. 4.17.4.3
Neutron Irradiation
The irradiation of tungsten and tungsten alloys with energetic neutrons (14 MeV) resulting from the D–T reaction causes radiological hazards that were already discussed in Section 4.17.3.2.6. In addition, the neutron irradiation affects the material composition by transmutation of tungsten to Re and subsequently osmium (transmutation of W isotopes to Ta and Hf are negligible222). The amount of transmutation strongly depends on the applied neutron wall load and neutron spectrum223 and for the W to Re transmutation reaction reaches values between 0.3 and 5 at.% per dpa.222 The subsequent transmutation of Re to Os is expected to occur faster than the production of Re from W resulting in a steadily proceeding burnup of Re. The neutron fluence on the first wall varies strongly with location. For the full lifetime of ITER a maximum of 0.3 MWa m2 is achieved224 ( 1.35 dpa in tungsten225). As the divertor PFCs will be exchanged 3 times and only the last three will operate in a D–T environment, a neutron fluence of 0.1 MW m2 is expected during the lifetime of each PFC. For DEMO, an average neutron wall load of 2 MW m2 is assumed for the main chamber, which would result in 45 dpa after 5 full power operation years. These conditions yield a transmutation of 100% W into 75% W, 12% Re, and 13% Os.36 For geometrical reasons, that is, larger surface to angular extension ratio, it will be roughly a factor 2 less in the divertor region. Furthermore, neutron irradiation damages the material properties by the formation of vacancies and interstitials (see Chapter 1.03, Radiation-Induced Effects on Microstructure). Their behavior including analysis of displacement cross-sections,226,227 diffusion, mutual recombination, and clustering are being assessed by atomistic modeling.228–231 Both transmutation and defect generation influences the material properties and subsequently the material response to steady state and transient thermal loads.
4.17.4.3.1 Thermophysical properties and swelling
The influence of neutron irradiation on the thermophysical properties is related to the irradiation temperature and the number of defects generated in the crystal structure. At temperatures <1000 C, the electrical232–234 and thermal conductivity217 of tungsten and tungsten alloys decrease with increasing irradiation dose. However, at elevated temperatures such as those occurring in a fusion environment, the effect of neutron irradiation is strongly mitigated by annealing.107 Complete recovery of defect-induced material degradation should occur at temperatures >1200 C (see Figure 8). In addition to defect generation, material degradation is also related to the formation of transmutation products such as Re and Os, which in general exhibit poorer thermophysical properties. Transmutationinduced degradation increases with increasing temperature and irradiation dose, which makes it the most relevant process for the degradation of material properties for future fusion reactors such as DEMO. Despite the potential for full recovery of the material defects mentioned above, void-induced swelling occurs. The results235,236 of tungsten and tungsten alloys show that the material’s volume increases with increasing irradiation temperature (1050 C).237 W–Re alloys exhibit significantly improved swelling behavior compared to pure W, with a local maximum at 750 C. However, the swelling only amounts to 1.7% at 9.5 dpa.237 Therefore, a negligible effect of swelling can be expected for the operation of ITER. Experimental values do not exist at temperatures >1050 C as expected for the operation of DEMO.
70 Thermal diffusivity (mm2 s–1)
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65
Nonirradiated 0.6 dpa at 200 ⬚C
60 55 50 45 40 35 30 0
200
400
600
800
1000 1200 1400 1600
Temperature (⬚C)
Figure 8 Thermal diffusivity of W–1% La2O3 in nonirradiated and irradiated condition.
Tungsten as a Plasma-Facing Material
4.17.4.3.2 Mechanical properties
Data in the literature on mechanical properties of neutron-irradiated tungsten are very limited.234,238,239 However, in combination with experimental results obtained for other refractory metals, it has been shown that in metals with a bcc lattice structure, irradiation hardening causes a steep increase in yield stress and a decrease in ductility.110 Consequently, the material fails by brittle cleavage fracture as soon as the yield stress exceeds the cleavage strength. Therefore, the increase of the DBTT depends on the neutron fluence, the neutron spectrum (will be addressed by the International Fusion Materials Irradiation Facility, IFMIF), and the irradiation temperature. The radiation hardening in bcc alloys at low temperatures (<0.3 Tm) occurs even for doses as low as 0.15–0.6 dpa (irradiation of plasma facing materials for ITER and DEMO, PARIDE campaigns217), which corresponds to the expected ITER conditions. Therefore, operation of tungsten at temperatures >1000 C would be preferred as full or at least partial recovery of defect-induced material degradation is achieved by annealing at 1200 C.234 This implies that the nearsurface part of a W component will retain its ductility, which has a beneficial effect on the crack resistance at the plasma loaded surface. However, such temperatures are not feasible at the interface to the heat sink where tungsten will be in contact with copper (ITER) or steel (DEMO), which are limited to significantly lower operational temperatures. Hence, better understanding of the irradiation effects on tungsten at temperatures between 700 and 1000 C is needed, particularly related to reactor application in DEMO.109,110,240 In addition to the influencing factors on the DBTT mentioned above, that is, neutron fluence, neutron spectrum, and irradiation temperature, the material’s composition also plays an important role. While the addition of Re has a beneficial effect on the material’s ductility in the nonirradiated state, under neutron irradiation it results in more rapid and severe embrittlement than it is observed for pure W.239 Similarly, less mechanical strength and an increased loss of ductility compared to pure W is found for particlestrengthened W alloys (e.g., W–1% La2O3) when tested up to 700 C. The only exception among all explored tungsten alloys might be mechanically alloyed W–TiC (see Section 4.17.3.3) that showed no irradiation hardening as measured by Vickers hardness at 600 C.87
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Finally, the mechanical properties are influenced by neutron-induced He-generation and the transmutation of tungsten. While He generation in W is, compared to CFC and Be, very small (0.7 appm He per dpa) and its influence on the mechanical properties of W negligible,73,83,224 the transmutation of W into Re and subsequently Os significantly alters the material structure and its properties. The generation of significant amounts of ternary a and subsequently s-phases results in extreme material embrittlement and will cause shrinkage. In combination with thermally induced strains, this might produce high tensile stresses causing the extremely brittle material to extensively crack and perhaps even crumble to powder.36 4.17.4.3.3 Thermal shock on irradiated W
The simulation of disruptions (<20 MJ m2, t ¼ 5 ms) on VPS-W and W alloys irradiated to a dose of 0.2 dpa at 350 and 750 C resulted in heavy melting of the material241 but yielded no measurable degradation by neutron irradiation. This is understandable because the decrease in thermal conductivity, which is the most important material parameter for melting experiments, is almost negligible for the given irradiation conditions.217 Because of the continuous modification of the material composition by transmutation described above, with increasing levels of Re and Os, the thermal conductivity and the related melting threshold power density is expected to decrease steadily. In further investigations applying ELM-like loads on pure W and W–La2O3 irradiated to 0.6 dpa at 200 C, the crack pattern generated after irradiation exhibited an increased crack density in combination with a smaller crack width. Furthermore, in W–0.2% Re SC that was exposed to the same neutron irradiation conditions and exhibited no crack formation in the nonirradiated state, cracks were formed along the crystallographic planes (see Figure 9).177 The effect was more obvious in the results for the SC material, but in both cases the observed degradation was a result of mechanical property changes and a rise of the DBTT in particular. This would indicate a rise in the threshold temperature for crack formation (see also Section 4.17.4.3.3), which has not been verified yet. For the evaluation of the material’s performance in DEMO, both higher transmutation rates and significantly higher temperatures that are expected to stimulate defect recovery have to be taken into account.
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4.17.4.4
(a)
500 mm (b)
500 mm Figure 9 Thermal shock response of W–0.2% Re (HF ¼ 41 MW m2 s1/2, P ¼ 1.31 GW m2, t ¼ 1 ms, n ¼ 10); (a) before and (b) after neutron irradiation (0.6 dpa at 200 C).
4.17.4.3.4 Thermal fatigue on irradiated W components
Information on the thermal fatigue resistance of W components is limited to the experience obtained in two irradiation campaigns for ITER which reached neutron doses of 0.15 and 0.6 dpa at 200 C. Reference and irradiated actively cooled mock-ups with W–1% La2O3 as the PFM were exposed to 1000 cyclic steady state heat loads at power densities up to 18 MW m2.213,216,217 The results obtained indicate that at these neutron fluences the material changes occurring in tungsten do not have any significant influence on the component’s performance. However, mock-ups based on the macrobrush design experience a degradation of the maximum achievable power density from 18 to 10 MW m2, which is related to neutron embrittlement and subsequent cracking failure of the Cu-heat sink material. In contrast, monoblock mock-ups show identical high-level performance before and after irradiation, which makes it the favored design for ITER. Despite these positive results, based on the irradiation-induced mechanical property changes outlined above, the use of tungsten in any highly stressed component at low temperatures <500 C has to be avoided.108
Ion Irradiation and Retention
In addition to the impact of thermal loads and neutrons, plasma wall interaction comprises the contact of the PFMs with hydrogen isotopes, the helium ash and impurities originating from eroded surfaces. The interaction of these different particles with the PFM leads to near-surface material modification and degradation in the nm and low-mm range. The extent of the damage depends on the energy of the particles, their fluence, the surface temperature of the PFM and the temperature gradients within the PFC during steady state and thermal shock loading. Furthermore, the material’s microstructure, composition, and predamage, for example, cracking, have an influence on the material’s performance. Knowledge of all these parameters, particularly with regard to operational conditions, is essential to determine the material’s lifetime due to erosion, the amount of dust formation, which influences the plasma performance, and the safety hazards due to H and He-retention. 4.17.4.4.1 He-irradiation
The incident energy of He ions impinging on the wall of fusion devices is expected to vary between eV and MeV. Energies in the eV range are representative for the majority of particles in a magnetic fusion device that have lost most of their energy due to the interaction with the plasma. In contrast, MeV He-ions, carrying more or less all of their initial energy, are typical for inertial confinement facilities. This broad range of energetic particles results in varying amounts of sputtered material (erosion rate 1.8 times higher than for deuterium242) and further interaction of the ions with the tungsten wall leads to additional material degradation. This starts with the formation of vacancies along the path of the ions and continues via vacancy clustering, bubble formation, blistering, and the formation of spongy structures. However, there exists a lower energy limit for He-penetration, which is related to the surface barrier potential that was theoretically calculated to be about 6 eV,243 and with slight variations this value was verified by experiments.149,244 For faster ions penetrating into the material, the generation of particular defects depends on the ion energy, ion fluence, and temperature.26,245,246 The correlation between these parameters is summarized as follows: 4.17.4.4.1.1
Influence of ion energy and fluence
The ion energy determines the initial penetration depth of the particle and the vacancy concentration
Tungsten as a Plasma-Facing Material
varies as a function of the implantation fluence.247 The higher the ion energy, the lower the fluence and the temperature required to create material damage beyond vacancies and vacancy clusters. For example, very small blisters were observed for 8 keV ions at RT and a fluence of 4 1021 Heþ m2 and above.247 In contrast, for low energies (<30 eV), temperatures >1300 K and fluences of about 1026 Heþ m2 and more are necessary to form bubbles and surface holes.149,248 The reason for the lack of blistering at low temperature and low ion energy is assumed to be the trapping of He-ions at defects/vacancies in the very near-surface range. With increasing temperature, the defects and He-atoms debond and the He-atoms diffuse toward the bulk, agglomerate, and result in blistering.249 Similarly, with increasing energy, the penetration depth of the He-atoms increases from nm for eV-ions up to 1.7 mm for 1.3 MeV He-ions and the probability of blister formation correspondingly rises. In both cases, whether the process is driven by He-diffusion or high penetration depth, blistering and exfoliation are expected to occur when the amount of He locally reaches 4 at.% and 20–40 at.%, respectively.250 4.17.4.4.1.2 Influence of temperature
Vacancy mobility is dependent on temperature and starts at 523–573 K.251,252 As the mobility of vacancies and the formation of thermal vacancies are driving forces for the formation of bubbles, holes, and blisters, an increase in temperature increases the size and decreases the density of material damage.149,253 However, it is not only the temperature during ion irradiation but also the annealing temperature during experiments such as thermal desorption measurements, which can influence the damage characteristics.247 The formation of holes and porous structures observed after thermal treatments,254 particularly for temperatures above the material-dependent recrystallization temperature, is related to the movement of vacancies, accelerating the expansion and coalescence of He bubbles, their migration to the surface,253 and subsequently the release of He. The latter is also a function of temperature, showing several release peaks between 400 and 1600 K related to different trapping sites247 and determines the amount of He retained as a function of the incident fluence.242,247,255 However, helium retention may be mitigated by cyclic He-implantation and high temperature heating, for example, flash heating to 2000 C, because He flows
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away before critical amounts accumulate and form complex He-vacancy clusters with higher binding energy.250 4.17.4.4.1.3 Influence of material’s microstructure
The impact of the material’s microstructure is related to the amount of intrinsic defects at which He can be trapped and therefore determines the amount of He-retention.250 SC tungsten contains fewer defects than powder metallurgically (PM) produced tungsten (grain boundaries) and plasma-sprayed tungsten (grain boundaries and porosity), which is directly related to the thickness of porous/spongy structures (porosity about 90%,256 see Figure 10) that form depending on energy and temperature.257,258 However, investigations at 1650 K have shown that at such high temperatures there exists no difference between SC-W and PM-W, even for ion energies as low as 25 eV. The main trap sites at such high temperatures are thermal vacancies while intrinsic defects play a minor role.149,251 With regard to the material’s lifetime under He-exposure, the migration of He-bubbles toward the surface and the formation of pores and porous/ spongy structures seem to prevent the rupture and exfoliation that can accompany blistering. This is important as the exfoliation of blisters creates dust,199 which limits the plasma performance. For an anticipated flux of 2 1018 Heþ m2 s1 at 850 C in inertial confinement devices, this may lead to a removal of 20 mm year1 from the wall.109 Therefore, one way to increase the material’s lifetime might be to operate it at higher temperature.253 Another approach would be to develop advanced microengineered materials that have typical feature sizes less than the classical helium migration distance (20 nm).109 However, bubbles, holes, and porous/spongy structures significantly influence the material’s performance by reducing its thermal conductivity in the near-surface layer. This might play an important role when determining the erosion and melt formation under combined He-irradiation and transient thermal loads, which will be shortly addressed in Section 4.17.4.4.3. 4.17.4.4.2 Hydrogen-irradiation and retention
Besides He, hydrogen isotopes, particularly the fuel elements deuterium and tritium, are the main incident ion species contacting PFMs and PFCs. The energy of these particles corresponds with the plasma temperature at the edge, which is in the range of some eV, but also includes highly energetic particles (10 keV) escaping from the inner core of the
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SR W
(a)
W-Re(5% wt)
(d)
W-Re(10% wt)
(g)
SC W (100)
(b)
WLa2O3(1% wt)
ASTM B760 (ITER)
(c)
W-TiC(1.5% wt)
(e)
(f)
VPS-W (EAST)
RC-W
(h)
(i)
Figure 10 Cross-sectional scanning electron microscopic images for nine different grades of W relevant to fusion engineering practice. All target specimens were exposed to consistent pure He plasmas at 1120 K for 1 h. The Heþ impact energy was 40 eV; (a) PLANSEE stress-relieved W, (b) single crystal h100i W, (c) ITER ASTM B760 compliant W, (d) PLANSEE W–5% Re, (e) PLANSEE W–1% La2O3, (f) ultrafine-grained W–1.5% TiC, (g) ULTRAMET CVD W–10% Re, (h) VPS W (EAST), and (i) recrystallized W. Reproduced from Baldwin, M. J.; Doerner, R. P. J. Nucl. Mater. 2010, 404, 165–173, with permission from Elsevier.
plasma. The impact of the energetic hydrogen ions is influenced by the incident ion energy, the ion fluence, the temperature, and the material’s composition and microstructure. The resulting damage, that is, vacancy formation, vacancy clustering, bubble formation, and blistering, determine not only the amount of material degradation and erosion but also the hydrogen (tritium) retention in the material. For active control of the hydrogen retention, short-term thermal treatments of the surface are being investigated. However, the short thermal load required to effectively remove the deuterium and tritium may also destroy the thin material layer (nm to low-mm range) that is responsible for the majority of the retention.259 4.17.4.4.2.1 Influence of ion energy, fluence, and temperature
Independent of the ion energy, blistering (see Figures 11 and 12) due to H occurs only during irradiation at temperatures below 900–950 K and as a function of the ion fluence199,260 at 500 K.261–263 This temperature dependence of blistering is attributed to the formation, movement, and agglomeration
of vacancies containing trapped hydrogen,264 which is dominant at temperatures <500 K, while the detrapping of hydrogen from the defects is prevalent at temperatures >500 K.263 The fluence threshold for blister formation increases with decreasing ion energy and significantly increases to values >1025 Dþ m2 at ion energies <20 eV. This is assumed to be the result of thin (a few monolayers) which act as oxide diffusion barriers at the material’s surface.265 Furthermore, with increasing fluence, the size and number of blisters can increase up to a few 100 mm260 until saturation is reached, which is assumed to be related to the rupture of blisters.265–267 The rupture and related dust formation is the result of hydrogen accumulation and pressure build up, which is effectively released by the failure of the blister cap. However, whether the blister has vented or not, the thermal contact with the substrate has been significantly reduced (see Figure 11). This eventually results in melting or vaporization of the thin blister cap, particularly during transient thermal loading conditions as described in Section 4.17.4.1, which contributes to further material’s erosion and probably plasma contamination.268
Tungsten as a Plasma-Facing Material
573
(a)
(a)
Large blisters
Small blisters
(b)
(b)
Large blisters
(c) Crack/void along grain boundary (c)
Lids of small blisters partially or fully removed by FIB fabrication
Figure 11 Scanning electron micrographs of tungsten exposed to a hydrogen fluence of 1026 D m2 at 480 K (45 tilt). (a) Overall image; (b) cross-sectional image of a large blister; (c) internal image of small blisters. Reproduced from Shu, W. M.; Kawasuso, A.; Yamanishi, T. J. Nucl. Mater. 2009, 386–388, 356–359, with permission from Elsevier.
In contrast, from the point of view of H retention, blistering is favorable because tritium accumulates in blisters, which act as a diffusion barrier for hydrogen, and which can otherwise penetrate deep into the
Figure 12 Scanning electron micrographs of small blisters appearing at tungsten exposed to a hydrogen fluence of 1026 D m2 at 480 K (45 tilt). (a) Initial stage; (b) growing; (c) bursting. Reproduced from Shu, W. M.; Kawasuso, A.; Yamanishi, T. J. Nucl. Mater. 2009, 386–388, 356–359, with permission from Elsevier.
material even at RT269–274 until it finally ends up in the heat sink structure and the coolant. Accordingly, as tritium retention is, in general, strongly correlated with the generation of blisters,275 it shows a maximum at an irradiation temperature of about 500 K.45,261–263,266,268,276 However, the retention of tritium and deuterium is also dependent on the trapping sites existing in the material. These are, in ascending order of their trapping potential, residual impurities, from which slow desorption occurs even at RT,277 grain boundaries and dislocations,
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Tungsten as a Plasma-Facing Material
radiation-induced vacancies and vacancy clusters, and pores. Depending on the occurrence and dominance of particular sites, the temperature at which the maximum hydrogen retention is observed varies between 450 and 850 K.138,262,264,268,277–281 With increasing temperature (>1000 K), such as that occurring at the strike point of the divertor, and with the lack of blisters, continuously decreasing hydrogen retention is observed.2,255,278,282,283 The remaining amount of retained hydrogen might be attributed to the presence of hydrogen as a solute, which depends only slightly on the incident ion energy, but scales with the implantation fluence and which is assumed to be of the same order of magnitude as the trapped concentration. It decreases only slightly with increasing temperature and at 1600 K still amounts to about 10% of the initial hydrogen content retained at 300 K.274 In addition to blisters, the high amount of hydrogen out-gassing at temperatures of 873 K and above results in the formation of bubbles and pores.253,277,284,285 This effect depends on the ion energy and fluence, which determines the amount and penetration depth of trapped hydrogen. Even though a beneficial smoothing effect on the surface quality is observed in comparison to pure annealing without hydrogen impact, at high temperatures up to 2500 K the surface smoothening might be accompanied by detrimental crack formation.284 4.17.4.4.2.2 Influence of material’s composition and microstructure
As mentioned above, hydrogen retention depends on the trapping sites available in the material and their relative energies. Their existence and concentrations are influenced not only by the impinging H-ions, but also by the manufacturing process, thermal pretreatments, the material’s composition and microstructure, and the surface quality. Accordingly, the retention increases with the amount of porosity in the material, as it allows a deep penetration of hydrogen and the voids and pores provide the highest trapping energies264,275 with thermal desorption occurring at temperatures >700 K.45,262 Another material parameter that increases hydrogen retention is the number of dislocations,266,286 particularly those introduced during deformation processes used for material densification. However, the recrystallization of the material removes not only dislocations but also vacancies and vacancy clusters, which have been introduced by the impinging H-ions286,287 and as grain boundaries. This effectively
reduces the trapping sites for hydrogen retention and, consequently, the lowest retention is observed for high-purity SC materials, particularly due to the low diffusion rate compared to polycrystalline tungsten.275,288,289 This low diffusion rate results in a near-surface accumulation of hydrogen, which acts as a diffusion barrier and leads to a saturation of hydrogen retention with increasing fluence.290 Such saturation is not observed for pure polycrystalline tungsten due to the possible hydrogen migration along grain boundaries.291 Finally, the hydrogen retention is influenced by impurities277 and dopants. The addition of La2O3 and TiC particles as well as the formation of pores, for example, by potassium doping, not only introduce traps and increase hydrogen retention,293 but also decrease the diffusion rate.291 In contrast, alloying with up to 10% Re has no measurable effect on the H retention properties of the material,279 as it only creates a slightly deformed crystal lattice structure but no additional hydrogen traps. In addition to hydrogen retention, material damage and particularly blistering is influenced by the material’s microstructure. Blistering occurs preferentially when the crystal is oriented with the h111i direction perpendicular to the surface292 and the blisters develop in different shapes from low, large, and spherical to high, small, and dome or coneshaped.45,262,267 The blisters in recrystallized materials are mainly plateau-shaped, often multilayer structures, which indicate a step-wise build-up, and in few cases also small blisters on top of large ones are formed.263,293 However, it is significant that the blister size is commonly limited by the grain size45,294 indicating that the grain boundaries play an important role in the formation of blisters. Accordingly, SCs and nanostructured materials such as W–TiC provide the strongest resistance against blistering, although the particular reason is different. For SCs, the hydrogen diffusion and accumulation is limited and there is a fast desorption from lowenergy traps at elevated temperatures. In contrast, for nanostructured materials the size of individual grains is extremely small and so is the volume for blister formation. Furthermore, the migration of hydrogen is significantly increased by the large number of grain boundaries.292 Further material parameters that reduce blister formation are open porosity and the surface finish, particularly the number of random or artificially introduced scratches that might act similar to grain boundaries.45 In contrast, the introduction of
Tungsten as a Plasma-Facing Material
impurities and dopants in commercially available grades of tungsten increases the number of blisters and exfoliation in both their stress relieved and recrystallized states.277,293 4.17.4.4.3 Combined loading conditions
As described above, the damage mechanisms of hydrogen and He-irradiation are rather similar, although they occur in different temperature ranges. Accordingly, their mutual interaction is also strongly influenced by the implantation temperature. Therefore, the testing sequence plays a role in the behavior, as for He preirradiation followed by hydrogen implantation, the implantation temperature of He determines the amount and kind of produced damage and the He-retention, which subsequently influences the hydrogen uptake occurring as described in Section 4.17.4.4.2. For example, He-implantation at RT either does not change the retention or may increase it due to the formation of additional trapping sites295–297 and the lower diffusion rate of He compared to H. With increasing He implantation temperature up to 800 K, hydrogen retention significantly decreases compared to pure hydrogen irradiation.261,292,296 This may be attributed to the occupation of trap sites by He as a result of its increasing mobility.298 Potential trap sites are the numerous He-induced nanosized bubbles acting as a diffusion barrier.292 A further increase in temperature to 1600 K does create significant material damage by He due to pore and bubble formation or even blistering. This tremendously increases the number of trap sites in the material and leads to He desorption during implantation and accordingly increases the hydrogen retention.299 For simultaneous loading of He and hydrogen, the fraction of He should reach at least 5 at.% to observe significant changes in the material’s response.261,292 Furthermore, for implantation temperatures below 900 K, results similar to those described above are observed for sequential ion beam loading.299 However, due to desorption of hydrogen at high temperatures >1000 K, no hydrogen retention takes place and the damage mechanisms are dominated by the He-irradiation during such temperature excursions. Correlated with hydrogen retention, blister formation at temperatures <900 K decreases with decreasing hydrogen retention. This is valid until the number of voids and pores, which enhance hydrogen retention, start to form open porosity and thereby
575
generate small grain structures. These allow a fast hydrogen diffusion through the material and limit the agglomeration of hydrogen necessary to form blisters similar to the case of nano-structured materials such as W–TiC described above (Section 4.17.4.4.2). Investigations of the influence of radiation damage (highly energetic hydrogen, neutrons) and impurity irradiation, for example, by carbon atoms, resulted in the depth resolved and particle energy– dependent formation of dislocations, dislocation loops, and even small voids acting as effective trapping sites for hydrogen and influence blister formation.274,276,300–309 Upon annealing, the dislocations and dislocation loops were moved and/or annihilated,310,311 which is positive news as it would limit the tritium inventory,312 as long as no He is present in the system. In contrast, with the addition of He the dislocations, dislocation loops, and helium bubbles do not vanish at identical annealing conditions, which has a direct impact on the mechanical and thermophysical performance of tungsten. However, He positively influences hydrogen retention in an intermediated temperature range as described above and inhibits the formation of a W carbide layer, which is typical for combined hydrogen and carbon loading.311 Finally, the results obtained from the investigation of the mutual influence of ion irradiation and thermal loads are strongly correlated with the choice of the heat source. In the literature, electron beam guns were favored,313,314 which are characterized by heat deposition in a depth range of several micrometers for W, depending on the acceleration voltage. However, as the thickness of the ion-irradiationaffected surface layer is comparably thin, the majority of the electrons pass through the modified surface layer, which leads to most doubtful conclusions. In contrast, lasers are more reliable, as they apply only surface heat loads. The combination of He-irradiation and laser-induced thermal loads (DT ¼ 1400 K, n ¼ 18 000) at high base temperatures (1700 K), resulted in an affected layer thickness (13 mm) about 10 times larger than that without laser irradiation (1–2 mm), which might be attributed to the steep temperature gradient supporting the diffusion of He. This surface modification combined with laser-induced surface roughening, as observed in Section 4.17.4.1.2 for typical thermal shock loads, leads to an enhanced degradation of the thermal diffusivity of W, which further increases the surface roughness and results in local or full melting of the W surface.199
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4.17.5 Conclusion Along with other favorable properties, tungsten is characterized by the highest melting temperature among all metals, a low energy threshold for sputtering, and a low tritium inventory compared to carbonbased materials. These characteristics make tungsten the most promising material for the plasma-facing inner wall of future nuclear fusion devices based on the magnetic confinement principle, and it is also under consideration for inertial fusion applications. Accordingly, it has been selected as the PFM for a large part of the ITER divertor during its start-up phase and will be used for the full divertor as soon as tritium operation starts; in addition, it is the reference material for DEMO. However, tungsten also offers less favorable properties. Related to these, there are some material issues that have to be resolved before operating tungsten in a fusion environment in an economically reasonable way, which means in DEMO and beyond. These are recrystallization, which influences the mechanical properties by reducing the ductility and increasing the DBTT embrittlement as a result of neutron-induced damages and transmutation resistance to crack formation, depending on the mechanical properties, which is particularly important during transient thermal loads He-induced sputtering and modification of a thin surface layer, which is influenced by existing material damage as well as by temperature and temperature gradients, for example, those occurring during transient thermal loads melting, which is related to crack formation and the degradation of thermophysical properties as a result of He-irradiation-induced surface modification; melt splashing and droplet ejection will influence the stable operation of the fusion plasma. As all grades of tungsten investigated so far have their own individual drawbacks, R&D programs worldwide are aiming for a deeper understanding of the parameters that influence the degradation of tungsten, and the development of new tungsten grades that are capable of dealing with the above-mentioned requirements. Therefore, the materials are characterized and qualified with regard to their microstructure before and after recrystallization by mechanical tests: evaluation of the material’s strength and DBTT
thermal shock loading : determination of temperature- and power density–dependent damage, cracking and melting thresholds, which are related to the mechanical and physical properties thermal fatigue loading : evaluation of the material’s performance as part of an actively cooled component neutron irradiation: characterization of the degradation of the material’s strength and the DBTT as well as the thermal shock and thermal fatigue response He- and H-irradiation: determination of the damage mechanisms such as blister, void, and bubble formation as a function of ion energy, fluence, and temperature as well as addressing hydrogen retention issues. However, despite all these efforts, a clear answer on the suitability of tungsten for application in a real fusion environment can only be given by ITER, as it is the mutual interaction of all the different types of loading that determine its lifetime relative to the various material degradation mechanisms.
References 1. Noda, N.; Philipps, P.; Neu, R. J. Nucl. Mater. 1997, 241–243, 227–243. 2. Lipschultz, B.; Whyte, D. G.; Irby, J.; LaBombard, B.; Wright, G. M. Nucl. Fusion 2009, 49, 045009. 3. Maddaluno, G.; Pierdominici, F.; Vittori, M. J. Nucl. Mater. 1997, 241–243, 908–913. 4. Maddaluno, G.; Giacomi, G.; Rufoloni, A.; Verdini, L. J. Nucl. Mater. 2007, 363–365, 1236–1240. 5. Hirai, T.; Philipps, V.; Huber, A.; et al. J. Nucl. Mater. 2003, 313–316, 67–71. 6. Hirai, T.; Kreter, A.; Linke, J.; et al. Fusion Eng. Des. 2006, 81(1–7), 175–180. 7. Pospieszczyk, A.; Tanabe, T.; Philipps, V.; et al. J. Nucl. Mater. 2001, 290–293, 947–952. 8. Sergienko, G.; Bazylev, B.; Huber, A.; et al. J. Nucl. Mater. 2007, 363–365, 96–100. 9. Sergienko, G.; Bazylev, B.; Hirai, T.; et al. Phys. Scripta 2007, T128, 81–86. 10. Neu, R.; Bobkov, V.; Dux, R.; et al. Phys. Scripta 2009, T138, 014038. 11. Maier, H.; Krieger, K.; Balden, M.; Roth, J. J. Nucl. Mater. 1999, 266–269, 1003–1008. 12. Maier, H.; Luthin, J.; Balden, M.; Linke, J.; Koch, F.; Bolt, H. Surf. Coat. Technol. 2001, 142–144, 733–737. 13. Maier, H.; Luthin, J.; Balden, M.; et al. J. Nucl. Mater. 2002, 307–311(1), 116–120. 14. Neu, R.; Rohde, V.; Geier, A.; et al. J. Nucl. Mater. 2001, 290–293, 206–210. 15. Neu, R.; Dux, R.; Geier, A.; et al. Fusion Eng. Des. 2003, 65(3), 367–374. 16. Neu, R.; Dux, R.; Geier, A.; et al. J. Nucl. Mater. 2003, 313–316, 116–126. 17. Neu, R.; Bobkov, V.; Dux, R.; et al. J. Nucl. Mater. 2007, 363–365, 52–59. 18. Roth, J.; Tsitrone, E.; Loarte, A.; et al. J. Nucl. Mater. 2009, 390–391, 1–9.
Tungsten as a Plasma-Facing Material 19. 20. 21. 22. 23. 24. 25. 26. 27. 28. 29. 30. 31. 32. 33. 34. 35. 36. 37. 38. 39. 40. 41. 42. 43. 44. 45. 46. 47. 48. 49. 50. 51. 52. 53.
Matthews, G. F.; Edwards, P.; Hirai, T.; et al. Phys. Scripta 2007, T128, 137–143. Mertens, Ph.; Altmann, H.; Hirai, T.; et al. Fusion Eng. Des. 2009, 84, 1289–1293. Neu, R.; Maier, H.; Gauthier, E.; et al. Phys. Scripta 2007, T128, 150–156. Federici, G.; Loarte, A.; Strohmayer, G. Plasma Phys. Control. Fusion 2003, 45, 1523–1547. Wilson, H. R.; Cowley, S. C. Phys. Rev. Lett. 2004, 92(17), 175006. Pamela, J.; Matthews, G. F.; Philipps, V.; Kamendje, R. J. Nucl. Mater. 2007, 363–365, 1–11. Pamela, J.; Becoulet, A.; Borba, D.; Boutard, J.-L.; Horton, L.; Maisonnier, D. Fusion Eng. Des. 2009, 84, 194–204. Bolt, H.; Barabash, V.; Krauss, W.; et al. J. Nucl. Mater. 2004, 329–333, 66–73. Maisonnier, D.; Cook, I.; Sardain, P.; et al. Fusion Eng. Des. 2006, 81, 1123–1130. Maisonnier, D. Fusion Eng. Des. 2008, 83, 858–864. Tobita, K.; Nishio, S.; Enoeda, M.; et al. Fusion Eng. Des. 2006, 81, 1151–1158. Mau, T. K.; Kaiser, T. B.; Grossman, A. A.; et al. Fusion Sci. Technol. 2008, 54(3), 771–786. Najmabadi, F.; Raffray, A. R. Fusion Sci. Technol. 2004, 46, 401–416. Holtkamp, N. Fusion Eng. Des. 2009, 84, 98–105. Merola, M.; Da¨nner, W.; Palmer, J.; Vielder, G.; Wu, C. H. Fusion Eng. Des. 2003, 66–68, 211–217. Merola, M.; Palmer, J. Fusion Eng. Des. 2006, 81, 105–112. Maisonnier, D.; Campbell, D.; Cook, I.; et al. Nucl. Fusion 2007, 47, 1524–1532. Cottrell, G. A. J. Nucl. Mater. 2004, 334, 166–168. Raffray, A. R.; El-Guebaly, L.; Malang, S.; et al. Fusion Sci. Technol. 2008, 54(3), 725–746. Ibrahim, A.; Henderson, D. L.; El-Guebaly, L. A.; Wilson, P. P. H.; Sawan, M. E. UWFDM-1331; Dec 2007. Abbott, F. R. J. Appl. Phys. 1942, 13, 384–389. Plansee Brochure: Tungsten, 663 DE 02.09, 2000; RWF, http://www.plansee.com/lib/Tungsten.pdf. Kang, H.-K. J. Nucl. Mater. 2004, 335, 1–4. Matejicek, J.; Koza, Y.; Weinzettl, V. Fusion Eng. Des. 2005, 75–79, 395–399. Deschka, S.; Garcia-Rosales, C.; Hohenauer, W.; et al. J. Nucl. Mater. 1996, 233–237, 645–649. Greuner, H.; Bolt, H.; Bo¨swirth, B.; et al. Fusion Eng. Des. 2005, 75–79, 333–338. Luo, G. N.; Umstadter, K.; Shu, W. M.; Wampler, W.; Lu, G.-H. Nucl. Instrum. Meth. Phys. Res. B 2009, 267, 3041–3045. Tokunaga, K.; Yoshida, N.; Noda, N.; Sogabe, T.; Kato, T. J. Nucl. Mater. 1998, 258–263, 998–1004. Tokunaga, K.; Yoshida, N.; Noda, N.; et al. J. Nucl. Mater. 1999, 266–269, 1224–1229. Liu, X.; Zhang, F.; Tao, S.; et al. J. Nucl. Mater. 2007, 363–365, 1299–1303. Ganne, T.; Crepin, J.; Serror, S.; Zaoui, A. Acta Mater. 2002, 50, 4149–4163. Cambe, A.; Gauthier, E.; Layet, J. M.; Bentivegna, S. Fusion Eng. Des. 2001, 56–57, 331–336. Murphy, J. D.; Giannattasio, A.; Yao, Z.; Hetherington, C. J. D.; Nellist, P. D.; Roberts, S. G. J. Nucl. Mater. 2009, 386–388, 583–586. Youchison, D. L.; Lutz, T. J.; Williams, B.; Nygren, R. E. Fusion Eng. Des. 2007, 82, 1854–1860. Hirai, T.; Maier, H.; Rubel, M.; et al. Fusion Eng. Des. 2007, 82, 1839–1845.
54. 55. 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. 82.
577
Maier, H.; Neu, R.; Greuner, H.; et al. J. Nucl. Mater. 2007, 363–365, 1246–1250. Ruset, C.; Grigore, E.; Maier, H.; et al. Phys. Scripta 2007, T128, 171–174. Anderl, R. A.; Pawelko, R. J.; Hankins, M. R.; Longhurst, G. R. J. Nucl. Mater. 1994, 212–215(2), 1416–1420. Kobayashi, A.; Sharafat, S.; Ghoniem, N. M. Surf. Coat. Technol. 2006, 200, 4630–4635. Matejicek, J.; Boldyryeva, H. Phys. Scripta 2009, T138, 014041. Riccardi, B.; Montanari, R.; Casadei, M.; Costanza, G.; Filacchioni, G.; Moriani, A. J. Nucl. Mater. 2006, 352, 29–35. Yahiro, Y.; Mitsuhara, M.; Tokunakga, K.; et al. J. Nucl. Mater. 2009, 386–388, 784–788. Gludovatz, B.; Wurster, S.; Hoffman, A.; Pippan, R. In Proceedings of the 17th Plansee Seminar, Reutte, Austria, May 25–29, 2009; Sigl, L. S., Ro¨dhammer, P., Wildner, H., Eds.; 21–1. Gumbsch, P. J. Nucl. Mater. 2003, 323, 304–312. Giannatassio, A.; Tanaka, M.; Joseph, T. D.; Roberts, S. G. Phys. Scripta 2007, T128, 87–90. Denissen, C. J. M.; Liebe, J.; van Rijswick, M. In Proceedings of the 16th Plansee Seminar, Reutte, Austria, May 30–June 3, 2005; Kneringer, G., Ro¨dhammer, P., Wildner, H., Eds.; RM15, pp 115–166. Barto, R. L.; Ebert, L. J. Metall. Trans. 1971, 2, 1643–1649. Humphreys, F. J.; Hatherly, M., Recrystallization and Related Annealing Phenomena, 2nd ed.; Elsevier: Amsterdam, 2004. Smid, I.; Akiba, M.; Vieider, G.; Plo¨chl, L. J. Nucl. Mater. 1998, 258–263, 160–172. Wesemann, I.; Spielmann, W.; Heel, P.; Hoffmann, A. In Proceedings of the 17th Plansee Seminar, Reutte, Austria, May 25–29, 2009; Sigl, L. S., Ro¨dhammer, P., Wildner, H., Eds.; RM50. Simpson, R. P.; Dooley, G. J.; Haas, T. W. Meter. Trans. 1974, 5, 585–591. Rupp, D.; Weygand, S. M. J. Nucl. Mater. 2009, 386–388, 591–593. Kurishita, H.; Matsuo, S.; Arakawa, H.; et al. J. Nucl. Mater. 2010, 398, 87–92. Cullip-Brennan, P.; Ordonez-Chu, A.; Villalobos, E.; et al. In Proceedings of the 16th Plansee Seminar, Reutte, Austria, May 30–June 3, 2005; Kneringer, G., Ro¨dhammer, P., Wildner, H., Eds.; RM77, pp 749–766. Davis, J. W.; Barabash, V. R.; Makhankov, A.; Plo¨chl, L.; Slatteri, K. T. J. Nucl. Mater. 1998, 258–263, 308–312. Fischer, B.; Vorberg, S.; Vo¨lkl, R.; Beschliesser, M.; Hoffman, A. In Proceedings of the 16th Plansee Seminar, Reutte, Austria, May 30–June 3, 2005; Kneringer, G., Ro¨dhammer, P., Wildner, H., Eds.; RM53, pp 513–517. Mabuchi, M.; Okamoto, K.; Saito, N.; et al. Mater. Sci. Eng. A 1996, 214, 174–176. Rieth, M.; Dafferner, B. J. Nucl. Mater. 2005, 342, 20–25. Rieth, M.; Hoffmann, A. Adv. Mater. Res. 2009, 59, 101–104. Rieth, M.; Hoffmann, A. Fusion Sci. Technol. 2009, 56, 1018–1022. Schmunk, R. E.; Korth, G. E. J. Nucl. Mater. 1981, 104, 943–947. Schmunk, R. E.; Korth, G. E. J. Nucl. Mater. 1984, 122–123, 850–854. Wie, Q.; Kecskes, L. J. Mater. Sci. Eng. A 2008, 491, 62–69. Pintsuk, G.; Prokhodtseva, A.; Uytdenhouwen, I. J. Nucl. Mater. 2011, in press.
578 83. 84. 85. 86. 87. 88. 89. 90. 91. 92. 93. 94. 95. 96. 97. 98. 99. 100. 101. 102. 103. 104. 105. 106. 107. 108. 109. 110.
111. 112. 113. 114. 115. 116.
Tungsten as a Plasma-Facing Material Zinkle, S. J.; Ghoniem, N. M. Fusion Eng. Des. 2000, 51–52, 55–71. Faleschini, M.; Kreuzer, H.; Kiener, D.; Pippan, R. J. Nucl. Mater. 2007, 367–370, 800–805. Wie, Q.; Zhang, H. T.; Schuster, B. E.; et al. Acta Mater. 2006, 54, 4079–4089. Wie, Q.; Jiao, T.; Ramesh, K. T.; et al. Acta Mater. 2006, 54, 77–87. Kurishita, H.; Kobayashi, S.; Nakai, K.; et al. J. Nucl. Mater. 2008, 377, 34–40. Kurishital, H.; Matsuo, S.; Arakawa, H.; et al. Mater. Sci. Eng. A 2008, 477, 162–167. Kurishita, H.; Matsuo, S.; Arakawa, H.; et al. J. Nucl. Mater. 2009, 386–388, 579–582. Makhankov, A.; Barabash, V.; Mazul, I.; Yohchison, D. J. Nucl. Mater. 2001, 290–293, 1117–1122. Ibbott, C.; Jakeman, R.; Ando, T.; et al. Fusion Eng. Des. 1998, 39–40, 409–417. Rigal, E.; Bucci, P.; Le Marois, G. Fusion Eng. Des. 2000, 49–50, 317–322. Watson, R. D.; Slattery, K. T.; Odegard, B. C., Jr.; et al. Fusion Technol. 1998, 34(3/2), 443–453. Giniyatulin, R. N.; Komarov, V. L.; Kuzmin, E. G.; et al. Fusion Eng. Des. 2002, 61–62, 185–190. Odegard, B. C.; Cadden, C. H.; Watson, R. D.; Slattery, K. T. J. Nucl. Mater. 1998, 258–263, 329–334. Barabash, V.; Akiba, M.; Mazul, I.; Ulrickson, M.; Vieider, G. J. Nucl. Mater. 1996, 233–237, 718–723. Chapa, J.; Reimanis, I. J. Nucl. Mater. 2002, 303, 131–136. Chong, F. L.; Chen, J. L.; Li, J. G. J. Nucl. Mater. 2007, 363–365, 1201–1205. Doering, J.-E.; Vaben, R.; Pintsuk, G.; Sto¨ver, D. Fusion Eng. Des. 2003, 66–68, 259–263. Pintsuk, G.; Bru¨ning, S. E.; Do¨ring, J.-E.; Linke, J.; Smid, I.; Xue, L. Fusion Eng. Des. 2003, 66–68, 237–240. Zhou, Z.; Song, S. X.; Du, J.; Ge, C. C. J. Nucl. Mater. 2007, 367–370, 1468–1471. Norajitra, P.; Giniyatulin, R.; Krauss, W.; et al. Fusion Sci. Technol. 2009, 56, 1013–1017. Ritz, G.; Hirai, T.; Linke, J.; Norajitra, P.; Giniyatulin, R.; Singheiser, L. Fusion Eng. Des. 2009, 84, 1623–1627. Ritz, G.; Hirai, T.; Norajitra, P.; et al. Phys. Scripta 2009, T138, 014064. Material Assessment Report (MAR); ITER Document No. G 74 MA 10 W 0.3; April, 2004. Makhankov, A.; Barabash, V.; Berkhov, N.; et al. Fusion Eng. Des. 2001, 56–57, 337–342. Keys, L. K.; Moteff, J. J. Nucl. Mater. 1970, 34(3), 260–280. Barabash, V.; Federici, G.; Matera, R.; Raffray, A. R. Phys. Scripta 1999, T81, 74–83. Sethian, J. D.; Raffray, A. R.; Latkowski, J.; et al. J. Nucl. Mater. 2005, 347, 161–177. Zinkle, S. J. Wiffen, F. W.; El-Genk, M. S., Eds. In Space Technology and Applications International Forum-STAIF 2004, AIP Conference Proceedings; American Institute of Physics: Melville, NY, 2004; Vol. 699, pp 733–740. Avrigeanu, V.; Chuvaev, S. V.; Eichin, R.; et al. Nucl. Phys. A 2006, 765, 1–28. Qaim, S. M.; Graca, C. Nucl. Phys. A 1975, 242, 317–322. Fischer, U.; Pereslavtsev, P.; Moeslang, A.; Rieth, M. J. Nucl. Mater. 2009, 386–388, 789–792. Seidel, K.; Eichin, R.; Forrest, R. A.; et al. J. Nucl. Mater. 2004, 329–333, 1629–1632. Seidel, K.; Eichin, R.; Fischer, U.; et al. Fusion Eng. Des. 2006, 81, 1211–1217. Cepraga, D. G.; Cambi, G.; Frisoni, M. Fusion Eng. Des. 2000, 51–52, 747–752.
117. 118. 119. 120. 121. 122. 123. 124. 125. 126. 127. 128.
129. 130. 131. 132. 133. 134. 135. 136. 137. 138. 139. 140. 141. 142. 143. 144.
145. 146.
Sublet, J. Ch.; Sawan, M. E. Fusion Eng. Des. 1999, 45, 65–73. Taylor, N. P.; Pampin, R. Fusion Eng. Des. 2006, 81, 1333–1338. Ho, S. K.; Lowenthal, M. D. Fusion Eng. Des. 1995, 29, 214–218. Shedd, K. B.; Tungsten, U. S. Geological Survey, Mineral Commodity Summaries; 2009, pp 178–179. Boonyongmaneerat, Y. J. Mater. Proc. Technol. 2009, 209, 4084–4087. Zhou, Z.; Pintsuk, G.; Linke, J.; Ro¨dig, M. Phys. Scripta 2009, T138, 014058. Mertens, Ph.; Altmann, H.; Hirai, T.; et al. J. Nucl. Mater. 2009, 390–391, 967–970. Roedig, M.; Kuehnlein, W.; Linke, J.; et al. Fusion Eng. Des. 2002, 61–62, 135–140. Hiraoka, Y.; Fujitsuka, M.; Fuji, T. J. Nucl. Mater. 1991, 179–181, 275–278. Liu, Y. L.; Zhou, H. B.; Zhang, Y.; Jin, S.; Lu, G. H. Nucl. Instrum. Meth. Phys. Res. B 2009, 267, 3282–3285. Brooks, J. N. Met. Powder Rep. 2006, 0026–0657, 22–24. Zeep, B.; Rath, S.; Ihli, T.; Piotter, V.; Ruprecht, R.; Hausselt, J. In Proceedings of the 16th Plansee Seminar, Reutte, Austria, May 30–June 3, 2005; Kneringer, G., Ro¨dhammer, P., Wildner, H., Eds.; RM9. Zeep, B.; Norajitra, P.; Piotter, V.; Boehm, J.; Ruprecht, R.; Hausselt, J. Fusion Eng. Des. 2007, 82, 2660–2665. Zhou, Z.; Linke, J.; Pintsuk, G.; Du, J.; Song, S.; Ge, C. J. Nucl. Mater. 2009, 386–388, 733–735. Zhou, Z.; Pintsuk, G.; Linke, J.; et al. Fusion Eng. Des 2009, 85, 115–121. Zhou, Z.; Ma, Y.; Du, J.; Linke, J. Mater. Sci. Eng. A 2009, 505, 131–135. Vorhauer, A.; Pippan, R. Mater. Sci. Forum 2003, 426–432, 2747–2752. Mathaudhu, S. N.; deRosset, A. J.; Hartwig, K. T.; Kecskes, L. J. Mater. Sci. Eng. A 2009, 503, 28–31. Kecskes, L. J.; Cho, K. C.; Dowding, R. J.; Schuster, B. E.; Valiev, R. Z.; Wei, Q. Mater. Sci. Eng. A 2007, 467, 33–43. Zhang, Y.; Ganeev, A. V.; Wang, J. T.; Liu, J. Q.; Alexandrov, I. V. Mater. Sci. Eng. A 2009, 503, 37–40. Liu, X.; Tamura, S.; Tokunaga, K.; et al. J. Nucl. Mater. 2004, 329–333, 687–691. Garcia-Rosales, C.; Franzen, P.; Plank, H.; Roth, J.; Gauthier, E. J. Nucl. Mater. 1996, 233–237, 803–808. You, J. H.; To¨schen, T.; Lindig, S. J. Nucl. Mater. 2006, 348, 94–101. Tamura, S.; Tokunaga, K.; Yoshida, N. J. Nucl. Mater. 2002, 307–311, 735–738. Taniguchi, M.; Sato, K.; Ezato, K.; Yokoyama, K.; Akiba, M. J. Nucl. Mater. 2002, 307–311, 719–722. Maier, H.; Neu, R.; Greuner, H.; et al. Phys. Scripta 2009, T138, 014031. Andersen, M.; Sharafat, S.; Ghoniem, N. Fusion Eng. Des. 2006, 81, 1639–1645. Rieth, M.; Hoffmann, A. In Proceedings of the 17th Plansee Seminar, Reutte, Austria, May 25–29, 2009; Sigl, L. S., Ro¨dhammer, P., Wildner, H., Eds.; RM22, 22–1. Veleva, L.; Oksiuta, Z.; Vogt, U.; Baluc, N. Fusion Eng. Des. 2009, 84, 1920–1924. Ghezzi, F.; Znai, M.; Magni, S.; Vanacore, G. M.; Tagliaferri, A. J. Nucl. Mater. 2009, 393, 522–526.
Tungsten as a Plasma-Facing Material 147. Taniguchi, M.; Nakamura, K.; Sato, K.; Ezato, K.; Yokoyama, K.; Akiba, M. Fusion Technol. 2001, 39, 890–893. 148. Wurster, S.; Gludovatz, B.; Pippan, R. In Proceedings of the 17th Plansee Seminar, Reutte, Austria, May 25–29, 2009; Sigl, L. S., Ro¨dhammer, P., Wildner, H., Eds.; RM24, 24–1. 149. Ishijima, Y.; Kurishita, H.; Yubuta, K.; et al. J. Nucl. Mater. 2004, 329–333, 775–779. 150. Ishijima, Y.; Kannari, S.; Kurishita, H.; et al. Mater. Sci. Eng. A 2008, 473, 7–15. 151. Kitsunai, Y.; Kurishita, H.; Kayano, H.; Hiraoka, Y.; Igarashi, T.; Takida, T. J. Nucl. Mater. 1999, 271–272, 423–428. 152. Kurishita, H.; Kobayashi, S.; Nakai, K.; et al. Phys. Scripta 2007, T128, 76–80. 153. Kurishita, H.; Amano, Y.; Kobayashi, S.; et al. J. Nucl. Mater. 2007, 367–370, 1453–1457. 154. Matsuo, S.; Kurishita, H.; Arakawa, H.; et al. Mater. Sci. Eng. A 2008, 492, 475–480. 155. Song, G. M.; Wang, Y. J.; Zhou, Y. Int. J. Ref. Met. Hard Mater. 2003, 21, 1–12. 156. Zhang, T.; Wang, Y.; Zhou, Y.; Song, G. Mater. Sci. Eng. A 2009, 512, 19–25. 157. Pintsuk, G.; Uytdenhouwen, I. Int. J. Ref. Met. Hard Mater. 2010, 28, 661–668. 158. Koch, F.; Bolt, H. Phys. Scripta 2007, T128, 100–105. 159. Koch, F.; Koeppl, S.; Bolt, H. J. Nucl. Mater. 2009, 386–388, 572–574. 160. Raffray, A. R.; Nygren, R.; Whyte, D. G.; et al. Fusion Eng. Des. 2010, 85, 93–108. 161. Linke, J.; Escourbiac, F.; Mazul, I. V.; et al. J. Nucl. Mater. 2007, 367–370, 1422–1431. 162. Pintsuk, G.; Ku¨hnlein, W.; Linke, J.; Ro¨dig, M. Fusion Eng. Des. 2007, 82, 1720–1729. 163. Hirai, T.; Pintsuk, G. Fusion Eng. Des. 2007, 82, 389–393. 164. Garkusha, I. E.; Arkhipov, N. I.; Klimov, N. S.; et al. Phys. Scripta 2009, T138, 014054. 165. Hirai, T.; Pintsuk, G.; Linke, J.; Batilliot, M. J. Nucl. Mater. 2009, 390–391, 751–754. 166. Parker, R.; Janeschitz, G.; Pacher, H. D.; et al. Nucl. Mater. 1997, 241–243, 1–26. 167. Bazylev, B. N.; Janeschitz, G.; Landman, I. S.; Pestchanyi, S. E. J. Nucl. Mater. 2005, 337–339, 766–770. 168. Bazylev, B. N.; Janeschitz, G.; Landman, I. S.; Loarte, A.; Pestchanyi, S. E. J. Nucl. Mater. 2007, 363–365, 1011–1015. 169. Linke, J.; Akiba, M.; Duwe, R.; et al. J. Nucl. Mater. 2001, 290–293, 1102–1106. 170. Filatov, V. J. Nucl. Mater. 2003, 313–316, 393–398. 171. Garkusha, I. E.; Bandura, A. N.; Byrka, O. V.; et al. J. Nucl. Mater. 2005, 337–339, 707–711. 172. Koza, Y.; Berthe, E.; Lehmann, E.; et al. J. Nucl. Mater. 2004, 329–333, 706–710. 173. Gervash, A.; Wallura, E.; Ovchinnikov, I.; Makhankov, A. N.; Linke, J.; Breitbach, G. Fusion Technol. 1997, 1, 499–502. 174. Hirooka, Y.; Bourham, M.; Brooks, J. N.; et al. J. Nucl. Mater. 1992, 196–198, 149–158. 175. Linke, J.; Duwe, R.; Gervash, A.; Qian, R. H.; Ro¨dig, M.; Schuster, A. J. Nucl. Mater. 1998, 258–263, 634–639. 176. Uytdenhouwen, I.; Decreton, M.; Hirai, T.; Linke, J.; Pintsuk, G.; Van Oost, G. J. Nucl. Mater. 2007, 363–365, 1099–1103. 177. Pintsuk, G.; Compan, J.; Hirai, T.; Linke, J.; Ro¨dig, M. In Proceedings of the 2007 IEEE 22nd Symposium on Fusion Engineering, Albuquerque, NM, June 17–21, 2007; IEEE: New York, NY, 2007; pp 1–4.
178. 179. 180. 181. 182. 183. 184. 185. 186. 187. 188. 189. 190. 191. 192. 193. 194. 195. 196. 197. 198. 199. 200. 201. 202. 203. 204. 205. 206. 207. 208. 209.
579
Linke, J.; Akiba, M.; Bolt, H.; et al. J. Nucl. Mater. 1997, 241–243, 1210–1216. Reichelt, W. Angew. Chon. J. 87, Jahrg. IY75/Nr. 7, 239–244. Barabash, V.; Baranov, A. G.; Burtseva, T. A.; et al. Fusion Eng. Des. 1991, 18, 145–150. Loarte, A.; Saibene, G.; Sartori, R.; et al. J. Nucl. Mater. 2003, 313–316, 962–966. Zohm, H. Plasma Phys. Control. Fusion 1996, 38, 1213–1223. Loarte, A.; Saibene, G.; Sartori, R.; et al. In Proceedings of the 22nd IAEA Fusion Energy Conference, Geneva, Switzerland, Oct 13–18, 2008; IT/P6-13. Garkusha, I. E.; Bazylev, B. N.; Bandura, A. N.; et al. J. Nucl. Mater. 2007, 363–365, 1021–1025. Garkusha, I. E.; Makhlaj, V. A.; Chebotarev, V. V.; et al. J. Nucl. Mater. 2009, 390–391, 814–817. Hirai, T.; Bekris, N.; Coad, J. P.; et al. J. Nucl. Mater. 2009, 392, 40–44. Klimov, N.; Podkovyrov, V.; Zhitlukhin, A.; et al. J. Nucl. Mater. 2009, 390–391, 721–726. Wittlich, K.; Hirai, T.; Compan, J.; et al. Fusion Eng. Des. 2009, 84, 1982–1986. Zhitlukhin, A.; Klimov, N.; Landman, I.; et al. J. Nucl. Mater. 2007, 363–365, 301–307. Bazylev, B. N.; Janeschitz, G.; Landman, I. S.; Pestchanyi, S. E. Fusion Eng. Des. 2005, 75–79, 407–411. Bazylev, B. N.; Janeschitz, G.; Landman, I. S.; et al. Fusion Eng. Des. 2008, 83, 1077–1081. Bazylev, B. N.; Janeschitz, G.; Landman, I. S.; et al. J. Nucl. Mater. 2009, 390–391, 810–813. Bazylev, B. N.; Janeschitz, G.; Landman, I. S.; et al. Fusion Eng. Des. 2009, 84, 441–445. Landman, I. S.; Bazylev, B. N.; Garkusha, I. E.; Loarte, A.; Pestchanyi, S. E.; Safronov, V. M. J. Nucl. Mater. 2005, 337–339, 761–765. Landman, I. S.; Janeschitz, G. Fusion Eng. Des. 2008, 83, 1797–1800. Wuerz, H.; Pestchanyi, S.; Bazylev, B.; Landman, I.; Kappler, F. J. Nucl. Mater. 2001, 290–293, 1138–1143. Makhlaj, V. A.; Garkusha, I. E.; Linke, J.; et al. Prob. Atom. Sci. Technol. Ser. Plasma Phys. 2009, 15, 58–60. Garkusha, I. E.; Bandura, A. N.; Byrka, O. V.; et al. J. Nucl. Mater. 2009, 386–388, 127–131. Ohno, N.; Kajita, S.; Nishijima, D.; Takamura, S. J. Nucl. Mater. 2007, 363–365, 1153–1159. Pestchanyi, S.; Linke, J. Fusion Eng. Des. 2007, 82, 1657–1663. Renk, T. J.; Olson, C. L.; Tanaka, T. J.; et al. Fusion Eng. Des. 2003, 65, 399–406. Renk, T. J.; Tanaka, T. J.; Olson, C. L.; Peterson, R. R.; Knowles, T. R. 2004, 329–333, 726–731. Umstadter, K. R.; Doerner, R.; Tynan, G. J. Nucl. Mater. 2009, 386–388, 751–755. Renk, T. J.; Provencio, P. P.; Tanaka, T. J.; et al. J. Nucl. Mater. 2005, 347, 266–288. Truszkowska, K. U.S. Patent 5,875,228, 1999. Ro¨dhammer, P.; Sprenger, D. EU Patent EP 0 874 385 B1, 1998. Tamura, S.; Liu, X.; Tokunaga, K.; et al. J. Nucl. Mater. 2004, 329–333, 711–716. Tokunaga, K.; Kubota, Y.; Noda, N.; et al. Fusion Eng. Des. 2006, 81, 133–138. Rapp, J.; Pintsuk, G.; Mertens, Ph.; Altmann, H.; Lomas, P. J.; Riccardo, V. Fusion Eng. Des. 2010, 85, 153–160.
580
Tungsten as a Plasma-Facing Material
210. Holstein, N.; Krauss, W.; Konys, J. Fusion Eng. Des. 2008, 83, 1512–1516. 211. Krauss, W.; Holstein, N.; Konys, J. Fusion Eng. Des. 2007, 82, 1799–1805. 212. Escourbiac, F.; Constans, S.; Vignal, N.; et al. Fusion Eng. Des. 2009, 84, 747–751. 213. Linke, J.; Lorenzetto, P.; Majerus, P.; Merola, M.; Pitzer, D.; Ro¨dig, M. Fusion Sci. Technol. 2005, 47, 678–685. 214. Merola, M.; Plo¨chl, L.; Chappuis, Ph.; et al. J. Nucl. Mater. 2000, 283–287, 1068–1072. 215. Missirlian, M.; Escourbiac, F.; Merola, M.; Durocher, A.; Bobin-Vastra, I.; Schedler, B. J. Nucl. Mater. 2007, 367–370, 1330–1336. 216. Roedig, M.; Duwe, R.; Kuehnlein, W.; et al. Fusion Eng. Des. 2001, 56–57, 417–420. 217. Roedig, M.; Kuehnlein, W.; Linke, J.; et al. J. Nucl. Mater. 2004, 329–333, 766–770. 218. Visca, E.; Escourbiac, F.; Libera, S.; et al. Fusion Eng. Des. 2009, 84, 309–313. 219. Missirlian, M.; Escourbiac, F.; Merola, M.; Durocher, A.; Bobin–Vastra, I.; Schedler, B. J. Nucl. Mater. 2007, 367–370, 1330–1336. 220. Li, H.; Chen, J. L.; Li, J. G.; Li, Z. X. J. Nucl. Mater. 2007, 363–365, 1226–1230. 221. Raffray, A. R.; Malang, S.; Wang, X. Fusion Eng. Des. 2009, 84, 1553–1557. 222. Greenwood, L. R.; Garner, F. A. J. Nucl. Mater. 1994, 212–215, 635–639. 223. Greenwood, L. R. J. Nucl. Mater. 1994, 216, 29–44. 224. Barabash, V.; Federici, G.; Ro¨dig, M.; Snead, L. L. J.; Wu, C. H. J. Nucl. Mater. 2000, 283–287, 138–146. 225. U¨beyli, M. J. Fusion Energy 2004, 22(4), 251–257. 226. Broeders, C. H. M.; Konobeyev, A. Yu.; Villagrasa, C. J. Nucl. Mater. 2005, 342, 68–76. 227. Kodeli, I. J. Nucl. Mater. 2004, 329–333, 717–720. 228. Gilbert, M. R.; Dudarev, S. L.; Derlet, P. M.; Pettifor, D. G. J. Phys. Condens. Mater. 2008, 20, 345214. 229. Park, N.-Y.; Kim, Y.-C.; Seok, H.-K.; Han, S.-H.; Cho, S.; Cha, P.-R. Nucl. Instr. Meth. Phys. Res. B 2007, 265, 547–552. 230. Bjo¨rkas, C.; Nordlund, K.; Dudarev, S. Nucl. Instr. Meth. Phys. Res. B 2009, 267, 3204–3208. 231. Fikar, J.; Scha¨ublin, R. Nucl. Instr. Meth. Phys. Res. B 2009, 267, 3218–3222. 232. Keys, L. K.; Smith, J. P.; Moteff, J. Phys. Rev. 1968, 176(3), 851–856. 233. Anand, M. S.; Pande, B. M.; Agarwala, R. P. Radiat. Eff. 1978, 39, 149–155. 234. Gorynin, I. V.; Ignatov, V. A.; Rybin, V. V.; et al. J. Nucl. Mater. 1992, 191–194, 421–425. 235. Sikka, V. K.; Moteff, J. J. Appl. Phys. 1972, 43(12), 4942–4944. 236. Shcherbak, V. I.; Zakharova, M. I.; Bykov, V. N. Fiz. Metal. Metalloved. 1975, 40(5), 1095–1099. 237. Matolich, J.; Nahm, H.; Moteff, J. Scripta Metall. 1974, 8, 837–842. 238. Steichen, J. M. J. Nucl. Mater. 1976, 60, 13–19. 239. Krautwasser, P.; Derz, H.; Kny, E. High Temp. – High Press. 1990, 22, 25–32. 240. Barabash, V.; Federici, G.; Linke, J.; Wu, C. H. J. Nucl. Mater. 2003, 313–316, 42–51. 241. Ro¨dig, M.; Conrad, R.; Derz, H.; et al. J. Nucl. Mater. 2000, 283–287, 1161–1165. 242. Katayama, K.; Imaoka, K.; Okamura, T.; Nishikawa, M. Fusion Eng. Des. 2007, 82, 1645–1650. 243. Ullmaier, H. Nucl. Fusion 1984, 24(8), 1039–1083. 244. Nishijima, D.; Ye, M. Y.; Ohne, N.; Takamura, S. J. Nucl. Mater. 2003, 313–316, 97–101.
245. 246. 247. 248. 249. 250. 251. 252. 253. 254. 255. 256. 257. 258. 259. 260. 261. 262. 263. 264. 265. 266. 267. 268. 269. 270. 271. 272. 273. 274. 275. 276. 277. 278.
Ye, M. Y.; Takamura, S.; Ohno, N. J. Nucl. Mater. 1997, 241–243, 1243–1247. Nishijima, D.; Ye, M. Y.; Ohno, N.; Takamura, S. J. Nucl. Mater. 2004, 329–333, 1029–1033. Fu, Z.; Yoshida, N.; Iwakiri, H.; Xu, Z. J. Nucl. Mater. 2004, 329–333, 692–696. Sakaguchi, W.; Kajita, S.; Ohno, N.; Takagi, M. J. Nucl. Mater. 2009, 390–391, 1149–1152. Tokunaga, K.; Doerner, R. P.; Seraydarian, R.; et al. J. Nucl. Mater. 2003, 313–316, 92–96. Gilliam, S. B.; Gidcumb, S. M.; Parikh, N. R.; et al. J. Nucl. Mater. 2005, 347, 289–297. Debelle, A.; Barthe, M. F.; Sauvage, T.; et al. J. Nucl. Mater. 2007, 362, 181–188. Debelle, A.; Barthe, M. F.; Sauvage, T. J. Nucl. Mater. 2008, 376, 216–221. Cipiti, B. B.; Kulcinski, G. L. J. Nucl. Mater. 2005, 347, 298–306. Subrahmanyam, V. S.; Nambissan, P. M. G.; Sen, P. Solid State Commun. 1994, 89(6), 523–527. Tokitani, M.; Miyamoto, M.; Tokunaga, K.; et al. J. Nucl. Mater. 2007, 363–365, 443–447. Baldwin, M. J.; Doerner, R. P. J. Nucl. Mater. 2010, 404, 165–173. Baldwin, M. J.; Lynch, T. C.; Doerner, R. P.; Yu, J. H. J. Nucl. Mater. 2010, in press. Sakaguchi, W.; Kajita, S.; Ohno, N.; Takagi, M.; Kurishita, H. Plasma Fusion Res 2010, 5, S10231–S10235. Hollis, K. J.; Castro, R. G.; Maggiore, C. J.; Ayala, A. J. Nucl. Mater. 2000, 283–287, 1085–1088. Ye, M. Y.; Kanehara, H.; Fukuta, S.; Ohno, N.; Takamura, S. J. Nucl. Mater. 2003, 313–316, 72–76. Alimov, V. Kh.; Shu, W. M.; Roth, J.; et al. Phys. Scripta 2009, T138, 014048. Luo, G. N.; Shu, W. M.; Nishi, M. Fusion Eng. Des. 2006, 81, 957–962. Shu, W. M.; Isobe, K.; Yamanishi, T. Fusion Eng. Des. 2008, 83, 1044–1048. Ogorodnikova, O. V.; Schwarz-Selinger, T.; Sugiyama, K.; et al. Phys. Scripta 2009, T138, 014053. Luo, G. N.; Shu, W. M.; Nishi, M. J. Nucl. Mater. 2005, 347, 111–117. Shu, W. M.; Kawasuso, A.; Miwa, Y.; Wakai, E.; Luo, G. N.; Yamanishi, T. Phys. Scripta 2007, T128, 96–99. Shu, W. M.; Kawasuso, A.; Yamanishi, T. J. Nucl. Mater. 2009, 386–388, 356–359. Causey, R. A.; Doerner, R.; Fraser, H.; et al. J. Nucl. Mater. 2009, 390–391, 717–720. Franzen, P.; Garcia-Rosales, C.; Plank, H.; Alimov, V. Kh. J. Nucl. Mater. 1997, 241–243, 1082–1086. Nakamura, H.; O’hira, S.; Shu, W.; et al. J. Nucl. Mater. 2000, 283–287, 1043–1047. Nakamura, H.; Hayashi, T.; Nishi, M.; Arita, M.; Okuno, K. Fusion Eng. Des. 2001, 55, 513–520. Nakamura, H.; Hayashi, T.; Kakuta, T.; Suzuki, T.; Nishi, M. J. Nucl. Mater. 2001, 297, 285–291. Nakamura, H.; Shu, W.; Hayashi, T.; Nishi, M. J. J. Nucl. Mater. 2003, 313–316, 679–684. Quastel, A. D.; Davis, J. W.; Haasz, A. A.; MacaulayNewcombe, R. G. J. Nucl. Mater. 2006, 359, 8–16. Tokunaga, K.; Baldwin, M. J.; Doerner, R. P.; et al. J. Nucl. Mater. 2005, 337–339, 887–891. Alimov, V. Kh.; Roth, J.; Causey, R. A.; et al. J. Nucl. Mater. 2008, 375, 192–201. Sakamoto, R.; Muroga, T.; Yoshida, N. J. Nucl. Mater. 1995, 220–222, 819–822. Ahlgren, T.; Heinola, K.; Vainonen-Ahlgren, E.; Likonen, J.; Keinonen, J. Nucl. Instrum. Meth. Phys. Res. B 2006, 249, 436–439.
Tungsten as a Plasma-Facing Material 279. Golubeva, A. V.; Mayer, M.; Roth, J.; Kurnaev, V. A.; Ogorodnikova, O. V. J. Nucl. Mater. 2007, 363–365, 893–897. 280. Poon, M.; Haasz, A. A.; Davis, J. W. J. Nucl. Mater. 2008, 374, 390–402. 281. Alimov, V. Kh.; Ertl, K.; Roth, J. J. Nucl. Mater. 2001, 290–293, 389–393. 282. Wright, G. M.; Kleyn, A. W.; Alves, E.; et al. J. Nucl. Mater. 2009, 390–391, 610–613. 283. Bolt, H.; Greuner, H.; Boeswirth, B.; Ku¨hnlein, W.; Huber, Th.; Sato, K. In Proceedings of the 16th Plansee Seminar, Reutte, Austria, May 30–June 3, 2005; Kneringer, G., Ro¨dhammer, P., Wildner, H., Eds. 284. Raunier, S.; Balat-Pichelin, M.; Sans, J. L. Nucl. Instrum. Meth. Phys. Res. B 2009, 267, 1841–1848. 285. Sze, F. C.; Doerner, R. P.; Luckhardt, S. J. Nucl. Mater. 1999, 264, 89–98. 286. O’hira, S.; Steiner, A.; Nakamura, H.; Causey, R.; Nishi, M.; Willms, S. J. Nucl. Mater. 1998, 258–263, 990–997. 287. Eleveld, H.; van Veen, A. J. Nucl. Mater. 1994, 212–215, 1421–1425. 288. Alimov, V. Kh.; Roth, J.; Mayer, M. J. Nucl. Mater. 2005, 337–339, 619–623. 289. Haasz, A. A.; Poon, M.; Macaulay-Newcombe, R. G.; et al. J. Nucl. Mater. 2001, 290–293, 85–88. 290. Poon, M.; Haasz, A. A.; Davis, J. W.; MacaulayNewcombe, R. G. J. Nucl. Mater. 2003, 313–316, 199–203. 291. Haasz, A. A.; Davis, J. W.; Poon, M.; et al. J. Nucl. Mater. 1998, 258–263, 889–895. 292. Miyamoto, M.; Nishijima, D.; Ueda, Y.; et al. Nucl. Fusion 2009, 49, 065035. 293. Ueda, Y.; Funabiki, T.; Shimada, T.; et al. J. Nucl. Mater. 2005, 337–339, 1010–1014. 294. Shu, W. M.; Luo, G.-N.; Yamanishi, T. J. Nucl. Mater. 2007, 367–370, 1463–1467. 295. Arkhipov, N.; Kanashenko, S. L.; Sharapovet, V. M.; Zalavutdinov, R. Kh.; Gorodetsky, A. E. J. Nucl. Mater. 2007, 363–365, 1168–1172.
296. 297. 298. 299. 300. 301. 302. 303. 304. 305. 306. 307. 308. 309. 310. 311. 312. 313. 314.
581
Hino, T.; Koyama, K.; Yamauchi, Y.; Hirohata, Y. Fusion Eng. Des. 1998, 39–40, 227–233. Iwakiri, H.; Morishita, K.; Yoshida, N. J. Nucl. Mater. 2002, 307–311, 135–138. Nishijima, D.; Ye, M. Y.; Ohno, N.; Takamura, S. J. Nucl. Mater. 2004, 329–333, 1029–1033. Nishijima, D.; Sugimoto, T.; Iwakiri, H.; et al. J. Nucl. Mater. 2005, 337–339, 927–931. Bizyukov, I.; Krieger, K.; Azarenkov, N.; Linsmeier, Ch.; Levchuk, S. J. Nucl. Mater. 2007, 363–365, 1184–1189. Fukumoto, M.; Ohtsuka, Y.; Ueda, Y.; et al. J. Nucl. Mater. 2008, 375, 224–228. Fukumoto, M.; Kashiwagi, H.; Ohtsuka, Y.; et al. J. Nucl. Mater. 2009, 386–388, 768–771. Fukumoto, M.; Kashiwagi, H.; Ohtsuka, Y.; et al. J. Nucl. Mater. 2009, 390–391, 572–575. Funabiki, T.; Shimada, T.; Ueda, Y.; Nishikawa, M. J. Nucl. Mater. 2004, 329–333, 780–784. Kobayashi, M.; Suzuki, S.; Wang, W.; et al. Phys. Scripta 2009, T138, 014050. Krieger, K.; Roth, J. J. Nucl. Mater. 2001, 290–293, 107–111. Ohya, K.; Tanabe, T.; Kirschner, A.; et al. J. Nucl. Mater. 2005, 337–339, 882–886. Ueda, Y.; Fukumoto, M.; Sawamura, I.; Sakizono, D.; Shimada, T.; Nishikawa, M. Fusion Eng. Des. 2006, 81, 233–239. Shimada, T.; Kikuchi, H.; Ueda, Y.; Sagara, A.; Nishikawa, M. J. Nucl. Mater. 2003, 313–316, 204–208. Oliver, B. M.; Causey, R. A.; Maloy, S. A. J. Nucl. Mater. 2004, 329–333, 977–981. Oya, Y.; Suzuki, S.; Wang, W.; et al. Phys. Scripta 2009, T138, 014501. Wampler, W. R.; Doerner, R. P. Phys. Scripta 2009, T138, 014037. Tamura, S.; Tokunaga, K.; Yoshida, N.; et al. J. Nucl. Mater. 2005, 337–339, 1043–1047. Tokunaga, K.; Yoshikawa, O.; Makise, K.; Yoshida, N. J. Nucl. Mater. 2002, 307–311, 130–134.
4.18
Carbon as a Fusion Plasma-Facing Material
L. L. Snead Oak Ridge National Laboratory, Oak Ridge, TN, USA
M. Ferraris Politecnico di Torino, Italy
ß 2012 Elsevier Ltd. All rights reserved.
4.18.1 4.18.1.1 4.18.1.2 4.18.1.3 4.18.2 4.18.2.1 4.18.2.2 4.18.2.3 4.18.3 4.18.3.1 4.18.3.2 4.18.3.3 4.18.3.3.1 4.18.3.3.2 4.18.3.3.3 4.18.4 4.18.4.1 4.18.4.2 4.18.4.3 4.18.4.4 4.18.4.5 4.18.5 4.18.6 4.18.6.1 4.18.6.2 4.18.7 References
Introduction Background Plasma-Facing Materials Particle–Matter Interactions The Advantages of Carbon as a PFC Plasma Impurities and the Need for Graphite Materials Thermomechanical Loading of PFMs Transient Loading of PFMs Irradiation Effects on Thermophysical Properties of Graphite and CFCs Graphite Irradiation Damage Surface Effects Properties and Property Evolution of Graphite Fiber Composite Irradiation-induced dimensional changes in CFCs Irradiation-induced changes in strength and modulus Thermal conductivity degradation Plasma–Particle Interactions Chemical Erosion Doping of Graphite to Suppress Erosion Physical Sputtering Radiation-Enhanced Sublimation Erosion of Graphite in Simulated Disruption Events Tritium Retention in Graphitic Materials HHF Component Technology Joining of CFC to Heat Sink Evaluation of HHF Joint Summary and Conclusions
Abbreviations ASTM ASTM International CFC Carbon(graphite) fiber composite CTE Coefficient of thermal expansion CVD Chemical vapor deposition DPA Displacement per atom EU European Union FoMd Disruption figure of merit FoMth Thermal figure of merit GMP Galvanic Metallization Process HIP Hot isostatic press
584 584 584 586 586 586 587 587 590 590 591 591 594 597 598 602 602 604 606 607 608 608 611 611 617 617 618
ITER
International Thermonuclear Experimental Reactor JET Joint European Torus LAM Low activation materials PAN Polyacrylonitrile PFC Plasma facing component PFM Plasma facing material(s) PVD Physical vapor deposition RES Radiation enhanced sublimation RT Room temperature SATIR Transient infrared thermography SEM Scanning electron microscopy XRD X-ray Diffraction
583
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Carbon as a Fusion Plasma-Facing Material
4.18.1 Introduction 4.18.1.1
Background
Graphite-moderated, gas-cooled reactors led the way into the nuclear age starting with the Chicago Pile-1 reactor, where the first controlled and sustained critical nuclear reaction was initiated in December 1942. The first commercial nuclear power plant, Calder Hall in the United Kingdom, went critical in 1956. As the graphite moderator was literally at the core of these early reactors, graphite became one of the first and most extensively studied nuclear materials. As discussed in Chapter 4.10, Radiation Effects in Graphite, the fission-born neutron results in significant thermophysical property changes in graphite. Moreover, depending on the type of fission reactor, other environmental factors such as graphite oxidation become extremely important. In addition to being the moderator of gas-cooled reactors, graphite has found a number of new nuclear power applications. As examples, pyrolytic graphite is a key functional element in TRi ISOtropic (TRISO) fuels, which continue to be developed and utilized for gas-cooled reactors; carbon fiber composites (CFCs) are now under development for core application in hightemperature gas-cooled reactors1 and have been widely used as plasma-facing components (PFCs) in fusion reactors.2 The latter application began in 1978, when the Princeton Large Torus made a transition from tungsten to graphite ‘limiters.’ This enabled the first thermonuclear temperatures, beginning the widespread application of graphite materials in fusion systems, the subject of this chapter. As will be discussed, the primary motivation for the use of graphite in fusion systems is not (as in fission reactors) for neutron moderation, but for reasons related to its exceptional high temperature performance and its relatively innocuous interaction with the plasma. However, the fusion reactor environmental effects on graphite, including irradiation-induced property evolution, are very similar to those of their fission reactor analogs. In contrast to the fission of heavy elements such as uranium or plutonium, which releases a large amount of energy in their fission fragments and a moderate amount of energy in the form of neutron kinetic energy (mean about 1 MeV), fusion can occur for a number of light elements, some of which have reactions that release very high-kinetic-energy neutrons. Several possible routes to fusion are shown below in eqn [1]:
þ1 H1 !1 D2 þ positron ¼ 1:4 MeV 2 1 3 1 H þ1 D !2 He ¼ 5:5 MeV 3 1 4 1 H þ1 T !2 He ¼ 19:9 MeV 2 2 3 1 D þ1 D !2 He þ neutron ¼ 3:3 MeV 2 2 3 1 1 D þ1 D !1 T þ1 H ¼ 4:0 MeV 2 3 4 1 D þ1 T !2 He þ neutron ¼ 17:6 MeV 2 3 4 1 1 D þ2 He !2 He þ1 H ¼ 18:2 MeV 1H
1
½1
For any of these reactions to take place, the ionized atoms must be brought together with sufficient force to overcome the coulombic barrier. In thermonuclear fusion, this is accomplished by heating the ‘plasma’ of these atoms to the point where the kinetic energy is sufficient to overcome that barrier. Currently, it is thought that DþT fusion is the most accessible route to fusion, though the gaseous temperature required for DþT reaction is more than 50 million Kelvin. Control and containment of high-temperature, high-density fusion plasmas is the primary challenge and obstacle to fusion power. Many reactor concepts have been studied in the past and attention is now focused on the ‘tokamak’ system. This toroidal confinement machine system was developed in the mid-1960s in Russia. In this design, a high-strength twisted helix of magnetic lines forms a magnetic bottle. Ions, which are trapped within a certain gyroradius, travel along these lines circulating around the helix in opposition to the plasma electrons. For noncollisional plasmas, the ions can be heated by magnetic induction or through various external means to the extreme temperature necessary for the fusion reaction to take place. This concept is the basis for the four largest present-day fusion machines (Table 1), and is the premise for the ITER machine currently under construction. To give an idea of scale, in all of the present-day machines listed in Table 1, the helical cavity is big enough in size for an adult to walk within, and the radius from the center of the machine to the middle of the helix is typically several meters. A depiction of the inside of the JET torus, complete with beryllium-coated CFC wall, is given in Figure 1. 4.18.1.2
Plasma-Facing Materials
Perfect containment of the high-density plasma needed for power production, where perfection means no interaction of the high-energy plasma and its surroundings, is not a practical reality. Whether through normal operation, or in off-normal incidents such as plasma ‘disruptions,’ plasma–material interaction (PMI) will occur in fusion devices. Components
Carbon as a Fusion Plasma-Facing Material
Table 1
585
Materials and heat loads for the major fusion machines worldwide
Fusion device
Location
Fuel system
First wall heat load (MW m2)
First wall material
Divertor or limiter heat load (MW m2)
Divertor or limiter material
ITER
Cadarache, France San Diego, USA
0.88 normal 1.75 off-normal 0.6
Beryllium
DIII-D
D/D D/T D/D
5–10 10 (ELM) 5.3
Sepcarb NB41, or CX 2002U CFCs Poco ATJ graphite
JT-60U
Naka, Japan
JET
Culham, England
D/D
18
Hitachi HCB-18S Ibiden EPT-10 Showa-Denko CC312 Dunlop DMS-704
TFTR
Princeton, USA
D/D D/T
Negligible
Figure 1 Inside the JET torus. Beryllium-coated carbon fiber composite.
in line of sight with the plasma, and therefore impacted by the hot gasses and particles, are referred to as plasma-facing components (PFCs) or materials (PFMs). The reactions between the fusion plasma and the PFMs are quite severe and typically cause melting or sublimation, component mechanical failure due to high thermal stress, and excessive surface erosion. The plasma ion flux and associated heat loading to the PFMs can be highly nonuniform and quite dependent on the tokamak design. The hot plasma gasses are made up of unburned hydrogen fuel, fusion byproducts such as helium, plasma electrons, and impurities, which include elements previously removed from PFCs. As can be seen in eqn [1], the types of particles that may strike the PFMs are dependent on the fusion fuel. For the DþT fuel system, the plasma will contain not only the
Poco ATJ graphite Dunlap DMS704 CFC Dunlop DMS-704 Sepcarb N11 Sepcarb N11-S (3D CFC) Poco AXF-5Q graphite
FMI 4D
DþT fuel, but also high-energy alpha particles (3.5 MeV He) and neutrons (14.1 MeV). The partitioning of the reaction energy between helium and the neutron is both an advantage and a disadvantage for the DþT fuel system. Because the energetic helium nucleus quickly collides with the surrounding gasses, most of its energy remains in the plasma and helps to sustain the high plasma temperature. Conversely, the neutron has very little chance of collision in the low-density plasma and loses its energy outside of the plasma, usually over meters of path length inside the structure of the reactor. Because less than 30% of the DþT reaction energy remains in the plasma, only this fraction is eventually dumped on the PFCs, thus reducing the heat load handling requirement and material erosion. However, as discussed in Section 4.18.3, the material damage associated with the 14.1 MeV neutron collisions is significant and perhaps offsets the advantages of reduced DþT heat loading. A characteristic classifying the fusion device type is the manner in which the plasma edge is defined and the plasma power handled. The classic approach is to define the plasma edge by placing a sacrificial component in contact with the plasma. This component, which intercepts the plasma edge particle flux, is known as a bumper or bumper limiter, and extends circumferentially around the torus. A second approach to defining the plasma edge is magnetically capturing and diverting the edge plasma onto a divertor plate well removed from the central plasma. Once the plasma gasses strike the surfaces and are thus cooled, they are pumped away. Unless mitigated, the energy
586
Carbon as a Fusion Plasma-Facing Material
deposited locally on the ‘divertor’ can be excessive. Many techniques, such as magnetic sweeping to spread the load and puffing of gas to ‘soften’ the ion impact, have been used to reduce the particle flux and energy. Regardless of whether the limiter or divertor design is employed, the majority of the particle and heat flux is intercepted by these components (Table 1). However, a significant flux also impacts the balance of the torus lining, generally referred to as the first wall. A convenient comparison for the heat loadings given in Table 1 is that the maximum output from a conventional propane torch is approximately 10 MW m2, or about the maximum seen in current fusion devices.
a steady state, and a cool down phase. In this case, the heat flux is approximately uniform around the circumference of the machine and scales with the machine power. However, a significant number of these plasma shots end in an abrupt and somewhat violent fashion referred to as disruption. When this occurs, the plasma rapidly becomes unstable and instantaneously ‘dumps’ its energy onto the PFC. This causes significantly larger heat loads than during normal operation, and in many cases, defines the design limits for these components.
4.18.1.3
4.18.2.1 Plasma Impurities and the Need for Graphite Materials
Particle–Matter Interactions
Of the flux of particles that will impact the PFMs, the highest particle flux will be the ionized fuel itself. The energy of the impacting fuel ions on the various plasma-facing areas depends on many variables. For the divertor, where most of the interactions will occur, the majority of particles will have energies in the eV range. On the first wall where the interaction intensity is less, the charge exchange particles will mostly be in the keV range. For larger fusion devices of the future where dense plasmas and higher magnetic fields result in the thermalization of the energetic helium (eqn [1]), those helium ions will have energies similar to the fuel ions. Electrons, which are in number density equilibrium with the plasma ions, also travel along the plasma field lines, albeit in the opposite direction. The high-energy neutrons present in the DþT reaction (14.1 MeV), or those for the DþD reaction (2.4 MeV) have mean free paths of several centimeters in graphite and so will not interact strongly with the first wall. However, these neutrons will be scattered and slowed down within and behind the first wall, resulting in a nearly isotropic flux of high-energy neutrons throughout the fusion device. The reaction of the plasma neutrons, ions, and electrons with graphite PFMs, which is discussed in some detail in the following sections, can have a wide range of effects. These effects include physical and chemical erosion of the first wall and thermomechanical property degradation of the bulk and surface material. The discussion thus far has been limited to the operation of tokamaks in the quasi-steady state (long pulse). All present-day large tokamaks are pulsed machines with pulse lengths of seconds, where the plasma discharge consists of a rapid heating phase,
4.18.2 The Advantages of Carbon as a PFC
The fusion plasma is maintained through a combination of internal heating, (i.e., the 3.5 MeV helium nucleus from the DþT reaction) and externally, by means of induction, radio frequency waves, or neutral particle injection. Plasma heating is balanced by plasma-cooling mechanisms among which electromagnetic radiation dominates. In fully ionized plasma, the radiative cooling comes from the Bremsstrahlung that occurs when the energetic ions interact with the plasma electrons. A fraction of the electromagnetic radiation released from this interaction is lost from the plasma. The energy lost in this manner is significantly increased by low concentrations of impurities. The plasma power loss in the Bremsstrahlung channel, Pbrem, is determined through: Pbrem ðMWm3 Þ 4:8 1043 Zi2 Ni Ne Te1=2 / Zi2 Ni
½2
where Zi, Ni, Ne, and T are the atomic number of the radiating species, their density, the electron density, and the plasma temperature, respectively. Clearly, from the linear dependence on the plasma impurity concentration, and the square dependence on the atomic mass of the impurity, the ideal PFMs comprise light elements that have a low tendency to erode and migrate into the plasma. Carbon and beryllium are two low atomic number elements commonly used in tokamaks. The next suitable element is aluminum, which would have almost a factor of five higher radiative loss on an atom-per-atom basis compared to carbon. On the same basis, molybdenum, which has been used in many tokamak experiments, has a radiative loss 49 times that of carbon, and tungsten 150 times the radiative loss of carbon. However,
Carbon as a Fusion Plasma-Facing Material
this is based on the assumption that the same number of impurity atoms find their way into the plasma (i.e., Ni), which, as discussed later, is not the case. 4.18.2.2 Thermomechanical Loading of PFMs As seen in Table 1, the first wall must handle high plasma surface heat fluxes under normal operation and volumetric heat loadings due to the penetrating neutron and electromagnetic radiation. Surface heat loading is dependent on line-of-sight distance from the plasma and can be as high as several MW m2. These surface and volumetric heat loadings will induce temperature gradients on the PFMs and corresponding thermal stress, and stresses at the interface between the PFM and the heat sink. For example, if one assumes the ideal case of a 2.5 cm thick, infinitely wide graphite plate that is perfectly bonded to a 50 C copper heat sink, the thermal stress at the graphite–copper interface for a heat flux of 5 MW m2 has been shown to be 200 MPa.3 The ability of the PFC to withstand this heat flux and thermal stress will depend both on the material properties and the component design. The two most obvious design parameters are the thickness of the PFM and how it is attached to the heat sink. The critical material property of thermal conductivity, which to a great extent can be engineered to optimize conduction to the heat sink, is a strong function of temperature. As discussed later in Section 4.18.3, this property and other performance properties such as elastic modulus and strength are also highly dependent on radiation-induced displacement damage. A typical design for a fusion reactor divertor is shown in Figure 2. In this design, the heat flux strikes the surface of CFC composite blocks and the heat flows into a water-cooled copper tube that has been brazed inside the block. The PFC is bolted to a stainless steel support structure. This configuration of PFC is called the monoblock structure, as compared to the flat plate and saddle types inset into Figure 2. To provide a quantitative comparison of candidate PFMs, a number of figures of merit (FoMs) have been derived, one of which may be written as follows: FoMth ¼
K sy aEð1 vÞ
½3
where K is the thermal conductivity, sy the yield strength, a the thermal expansion coefficient, E the Young’s modulus, and n the Poisson’s ratio. High values of FoMth provide guidance to superior
587
performing candidate materials. Figure 3 shows a comparison of the three primary candidate PFMs: graphite, beryllium, and tungsten. Graphite has been further broken down into fine and coarse-grained (Poco and H451 respectively) graphites, and a high-quality one-dimensional (1D) fiber architecture (MKC-1PH) and a balanced weave 3D fiber architecture (FMI-222) CFC. In Figure 3, it has been assumed that the high thermal conductivity direction for the 1D CFC is oriented at a normal angle to the surface of the PFC. From Figure 3, it is apparent that the graphites and graphite fiber composites, which possess higher strength and thermal conductivity, exhibit thermal FoMs considerably higher than either beryllium or tungsten. Thus, strictly from a thermal stress point of view, high-conductivity and high-strength graphite materials would be considered superior under normal operating conditions for fusion PFCs. 4.18.2.3
Transient Loading of PFMs
The disruption, or collapse, of the fusion plasma causes a potentially intense thermal load to the PFC of all large fusion devices. As discussed later in Section 4.18.4, such events will cause very high thermal stresses and significant material erosion. As these events are transient in nature, the ability of the PFC to withstand the disruption depends on the material’s ability to conduct, and its ability to absorb the deposited heat, before reaching a temperature or stress limit. For comparative purposes, a disruption figure of merit takes this into account: su ðCp rK Þ1=2 ½4 aE where su is the ultimate tensile strength, Cp the specific heat, and r the density. Figure 4 reports this disruption figure of merit for the materials in Figure 3. Consistent with the results of the thermal FoMth, high-quality, high-thermal conductivity composites and fine-grained graphites perform better than standard and larger grained graphites, and exhibit an order of magnitude better FoMd than beryllium and tungsten. As discussed later in Section 4.18.4, the erosion of graphite and beryllium is very high and dictates the use of thick tiles in high flux areas. This is in contrast to tungsten, which has a relatively low erosion yield, potentially allowing an armor thickness of only a few millimeters. Because the FoMs are essentially calculated on a per unit tile thickness, comparing tungsten with graphite can be somewhat misleading. However, because graphite FoMd ¼
588
Carbon as a Fusion Plasma-Facing Material
Ion
flux
Graphite tiles
Coolant tubes
Support structure
Coolant tube
Monoblock
Saddle type
Flat plate
Graphite
Graphite
Graphite
Figure 2 Schematic diagram of the proposed monoblock first wall structure for the ITER reactor. Redrawn from Kuroda, T. et al. ‘‘ITER Plasma Facing Components,’’ ITER Documentation Series, No. 30, International Atomic Energy Agency (1991).
and beryllium are erosion-limited, the FoMs and the melting temperatures are useful evaluation tools. While the sublimation temperature of graphite (3350 C) is comparable to the melting point of tungsten (3400 C), it is clear that beryllium, which has a melting point of 1300 C, is at a distinct disadvantage. Removal of beryllium, as well as other metallic PFCs, by melting has been seen in several large experiments. Performance calculations for graphite and CFCs have been conducted in both laboratory test stands and operating tokamaks. Some experimental data generated using electron beam simulation are given in Figure 5. Here, the power is deposited by a rastered electron beam for approximately one second up to surface heat loads of 11 MW m2. The samples were 2.5 2.5 cm tiles, 1 cm in thickness, facing the beam. Each sample had a large notch machined into
one edge (the highest stressed area) to serve as a stress intensifier. It was noted that, without the notch, the graphites did not crack. Figure 5 gives the maximum heat flux of which each material was tested and whether cracking of the tile occurred. The data indicate that CFC materials and graphites with a higher thermal conductivity and high density are superior. No cracking occurred in either the three composites studied, or the two FMI graphites, at the maximum power density applied. The superior performance of the composite materials agrees with the performance of CFCs in the large tokamaks such as TFTR and JT-60U. The reason for the superior performance of the CFCs and the graphites is most likely their low thermal expansion coefficient, high thermal conductivity, and high strength. In addition, the presence of the fibers in the CFCs may serve to
Carbon as a Fusion Plasma-Facing Material
589
106
Thermal stress figure of merit, FoMth
Mitsubishi Kasei MFC-1: 1D-C/C Fiber Materials Inc. FMI-222: 3D-C/C
105 Unocal, poco AXF-5Q: Graphite Sigri Great Lakes H451: graphite 316L stainless steel
104
Pure tungsten
Brush Wellman S65-C: Wrought be
1000 200
300
400
500
600
700
800
900
1000
Application temperature (⬚C) Figure 3 Thermal stress figure of merit for selected plasma-facing materials.
Thermal shock figure of merit, FoMd
108
Mitsubishi Kasei MKC-1PH: 1D-C/C Fiber Materials FMI-222: 3D-C/C
107
Unocal, poco AXF-5Q: graphite
316L stainless steel Sigri Great Lakes H451: Graphite
Pure tungsten
106 Brush Wellman S65-C: Wrought be
105 200
300
400
500
600
700
800
900
1000
Application temperature (⬚C) Figure 4 Thermal shock figure of merit for selected plasma-facing materials.
blunt and arrest cracks, thus increasing toughness. All monolithic graphites shown in Figure 5, with the exception of the two FMI-HDFG materials, cracked. It is interesting to note that this graphite possessed
the highest FoMd, even higher than that of the composites. However, strict correlation of improved performance with increased FoMd was not seen, although a loose correlation was noted. As pointed
590
Carbon as a Fusion Plasma-Facing Material
FMI 4-D C/C coarse FMI 4-D C/C fine BFG. 2-D staple KNIT FMI HDG FMI HDFG Union carbide CGW II Union carbide CGW Union carbide ATJS
Graphite or CFC type
Union carbide TS1792 Union carbide TS1909 Stackpole 2191 Stackpole 2204 SGL H-490 SGL H-489 Toyo carbon AX 280-K RINGS DORFF EK-98 IBIDEN ETP-10 Poco ZXF5Q Carbone Iorainne 589 SCHUNK&EBE FE219 Toyo Tanso IG-110 SGL H-478
Cracking range Failure range No cracking
Poco AXF5Q Stackpole 1225 Stackpole 1336 Isograph 880
0
2
4 6 8 Power density (kW cm−2)
10
12
Figure 5 The performance of several grades of graphite and graphite composites subject to thermal shock loading. Redrawn from Croessmann, C. D.; Gilbertson, N. B.; Watson, R. D.; Whitley, J. B. Fusion Technol. 1989, 127–135.
out by Watson,4 the CTE may be the most dominant property, with the lowest CTE graphites showing the best resistance to thermal shock. Finally, it should be noted that there are many issues regarding the selection of carbon materials as PFCs other than simply their thermal shock behavior. The issues of radiation damage, erosion, and hydrogen retention are the three leading issues/drawbacks to the use of graphite as a PFC and they are discussed in the following sections.
4.18.3 Irradiation Effects on Thermophysical Properties of Graphite and CFCs 4.18.3.1
Graphite Irradiation Damage
Gross physical property changes can occur in the graphite PFMs through two generic routes: (1) nearsurface damage caused by interaction with plasma ions and, to a lesser extent, electrons, and, (2) bulk displacements caused by neutrons emanating from the plasma or back scattered by the surrounding structure. Of the tokamaks, only TFTR had significant DþT fusion reactions and, therefore, experienced a
significant flux of fusion neutrons (see eqn [1]). Even so, the dose from that TFTR was not high enough for the structural materials to experience appreciable neutron effects. However, machines such as the ITER will see a significant neutron dose from both DþD and DþT reactions. As energetic particles travel through matter, they can interact with their surroundings, losing energy (per unit path length) in three ways: elastic collisions, electron excitations, and nuclear interactions. The interaction of primary interest from the materials property evolution point of view results from the particle elastic collisions with the graphite crystal. This has also been discussed in Chapter 4.10, Radiation Effects in Graphite (Section 4.10.4). If an ion or a neutron can provide sufficient energy to overcome an atom’s binding energy (Ed carbon20–30 eV), the carbon can be displaced from its original lattice position. If the energy transferred to the displaced atom is sufficient to displace further atoms, a series of displacement events or a ‘cascade’ occurs. In the simplest interpretation, the Kinchin–Pease5 model is used to calculate the total number of atoms displaced. For example, if a carbon atom were ejected by the plasma and
Carbon as a Fusion Plasma-Facing Material
reimpacted onto the carbon tile with a kinetic energy (Ecarbon) of 1 keV, the estimated number of atoms displaced (n) would be estimated as follows: Ecarbon ½5 Ed ¼ 20 atoms n¼ 2 The interaction of high-energy neutrons with matter is very similar to that of high-energy ions. The primary difference between the two is the amount of energy transferred in a single collision and the distance over which the interactions take place. An ion, which has a relatively large coulombic interaction radius, loses its energy over a short path length (typically less than a micron). In contrast, the comparatively small uncharged 14.1 MeV fusion neutrons undergo only simple elastic or ‘billiard ball’ collisions with a mean free path between collisions of 10 cm. So, on average, a fusion neutron will have an elastic collision with a carbon atom once in 10 cm of graphite. The amount of energy transferred to the carbon in this first collision (Ec) is calculated by simple elastic theory as: " # 4mc mn Eo cos2 a Ec ¼ ðmc þ mn Þ2 " # 4 12 1 ½6 14:1MeVcos2 a ¼ ð12 þ 1Þ2 where mc and mn are the carbon and neutron mass (in amu), respectively, Eo is the neutron energy, and a is the angle between the neutron path before and after the collision. For a totally back scattered neutron (the maximum imparted energy), the energy transferred to the displaced carbon is 4.7 MeV. Again, from eqn [5], the number of displaced carbon atoms in this 14.1 MeV neutron collision event is nearly 100 000. The vast majority of these atoms do not stay ‘displaced,’ but condense back into the graphitic structure within a few picoseconds. To assess the effects such collision events will have on a material, a convention has been adopted to compare irradiation doses. The displacement per atom (dpa) gives the average number of times an atom has been knocked from its original lattice position. The dpa is an integrated average quantity, and takes into account the atomic density, the interaction cross-section, and the neutron energy spectrum. For the next-generation fusion reactors such as ITER, peak end of life values for PFMs due to neutrons will be on the order of tenths of a dpa, while power fusion reactors could potentially be subjected to greater than 10 dpa year1.
4.18.3.2
591
Surface Effects
While fast neutrons will produce relatively uniform atomic displacements, ions will produce very high near-surface damage. This damage can be on the level of hundreds of dpa, even for the experimental machines in use today and certainly for machines such as ITER and beyond. However, the damage is typically limited to much less than a micron in depth. The effect of this high damage level will be the reduction of a well-graphitized structure into a structure that appears amorphous. However, these nearsurface regions are subjected to erosion either by physical sputtering (caused by elastic collisions), or by chemical interactions. Both these effects are addressed in Section 4.18.4. A second surface radiation damage issue, that is, the ability of the thin damaged surface layer to retain and transport hydrogen, is discussed in Section 4.18.5. 4.18.3.3 Properties and Property Evolution of Graphite Fiber Composite As mentioned earlier, the first wall materials in nextgeneration machines will receive many tens of dpa. At low doses (<0.01 dpa), there are essentially no mechanical property changes expected in graphite materials (see Chapter 4.10, Radiation Effects in Graphite). However, even at these low doses, thermal conductivity and stored energy are of concern, specially for low irradiation temperatures (<400 C). For displacement levels >0.01 dpa, significant property changes occur, including strength, elastic modulus, specific heat, coefficient of thermal expansion (CTE), Poisson’s ratio (n), and thermal conductivity. In addition, the dimensional stability under irradiation is important because the induced stresses may be significant and may need very tight tolerances at the plasma edge. It has been shown in fission neutron experiments that specific heat Cp6 and n7 are not greatly affected by irradiation. Moreover, only moderate changes in the CTE occur, but the magnitude and nature of the CTE change is highly dependent on the type of graphite.6,8–10 The irradiation-induced property changes for graphite and composites that have received the most study by the fusion community deal with dimensional stability, strength, elastic modulus, thermal conductivity, and hydrogen retention. A large body of data exists on the thermophysical changes in graphites, coming mainly from graphite-moderated nuclear reactor development programs. A smaller body of research
Carbon as a Fusion Plasma-Facing Material
exists on CFCs, mainly from the same source, but with some additional data from fusion research. These data suggest that CFCs have very similar irradiation behavior compared to graphite. In Chapter 4.10, Radiation Effects in Graphite, Burchell discusses radiation damage mechanisms in graphite, and some of the specific property changes that occur in fission reactor applications. Because they are of special significance to fusion energy, the radiation effects in CFCs in general and the radiation-induced degradation in thermal conductivity in graphite and CFCs in particular will be focused on in the remainder of this section. However, it is first important to contrast nuclear graphite (essentially a form of purified structural graphite) with that of graphite composites. For the purposes of discussing graphite materials for fusion applications, the term composites is applied specifically to continuous fiber composites, typically woven, and infiltrated with pitch or some other resin that is graphitized to form a highly crystalline graphite matrix. The fibers comprising these composites are, as compared with most forms of graphite, highly crystalline and of comparatively high strength, elastic modulus, and thermal conductivity. The fibers themselves are typically either polyacrylonitrile (PAN) or Pitch derived. In general, one would select the PAN-based fiber, which is somewhat less expensive, if the application required higher strength while the Pitch-based fibers would produce a product with superior elastic modulus and thermal conductivity. As observed in Sections 4.18.2.2 and 4.18.2.3, the composite materials, due to their typically higher strength and elastic modulus, have a superior performance in terms of thermal stress and thermal shock. Another key advantage of these materials stems from the fact that they tend to fail in a less abrupt manner than seen for graphite or ceramics in general due to the presence of the reinforcing fibers, which bridge evolving crack fronts. This can be seen by casual inspection of Figure 6, which compares the nuclear graphite Poco AXF-5Q (historically used in TFTR and for other nuclear applications) and the FMI-222 balanced weave, 3D CFC. From Figure 6, and by a comparison of the graphite and composite data of Table 2, it is clear that the FMI-222 CFC material has both higher bend strength and higher elastic modulus (greater slope) as compared to this Poco graphite. Moreover, it is clear from Table 2 that other engineering properties of importance, such as strength and thermal conductivity, are superior for the CFC. These superior properties are primarily attributable to the exceptional quality of graphite
250 2.3 ⫻ 6 ⫻ 30 mm bend bars 4-point bending, 6.45/19.05 mm load/support spans
200 Bend stress (MPa)
592
FMI-222 3-D CFC
150
Poco AXF-5Q graphite
100
50
0 0
0.2
0.4 0.6 Displacement (mm)
0.8
1
Figure 6 Comparison of the loading behavior of a typical graphite and carbon fiber composite.
fiber. Unlike nuclear graphite, which is on the order of 20% porosity with a relatively imperfect, heavily faulted, inhomogeneous amalgam of filler particles (such as coke) and graphitized binder (such as pitch; see Section 4.10.2 in Chapter 4.10, Radiation Effects in Graphite for a discussion of graphite manufacture), graphite fibers, while somewhat different depending on the starting material (PAN, Pitch, Rayon, etc.), are extremely uniform, and highly crystalline with density that can approach theoretical density. This leads to exceptional properties. For example, the PAN-based T-300 fiber has a tensile strength of 3.66 GPa, slightly higher than the 2.41 GPa strength of the P120 fiber of the FMI-222 composite of Table 2, or more than 40 times that of the Poco AXF-5Q graphite. Similarly, the elastic moduli of T-300 and P-120 fibers are 21 and 75 times the elastic modulus of the Poco AXF-5Q graphite. In the case of the P-120 fiber, which has been graphitized at a very high temperature, very long, defect-free basal planes oriented along the axis of the fiber result in exceptional 1D thermal conductivity (640 W m1 K1, twice that of copper). This property is the primary reason for the twofold increase in ambient thermal conductivity of the FMI-222 composite as compared to the Poco AXF-5Q graphite. Clearly from this example of thermal conductivity, the architecture (fiber weave or loading) will determine the composite properties. Examples of practical fusion CFCs are the materials chosen for consideration and application by the ITER project. Table 3 provides the nonirradiated thermophysical property data for selected CFCs
Carbon as a Fusion Plasma-Facing Material
Table 2
593
Comparison of thermophysical properties of a typical graphite and carbon fiber composite
Manufacturer Architecture Precursor Grain size/unit cell size (mm) Ambient thermal conductivity (W m1 K1) Apparent density (g cm3) Flexural strength (MPa) Elastic modulus (GPa)
Poco AXF-5Q nuclear graphite
FMI-222 3D carbon fiber composite
Poco specialty Near isotropic Pitch 9 95 1.78 86 11
Fiber Materials Inc. Balanced 3D weave Amoco P-120 fibers pitch matrix 900 200 1.96 175 52
http://www.poco.com/MaterialsandServices/Graphite/IndustrialGrades/GradeChart/tabid/95/Default.aspx.
Table 3
Thermophysical properties of CFCs of interest to fusion
Constituents
CX2002U
MFC-1
INOX Sepcarb NS31
Sepcarb NB31
Dunlop concept 1
Pitch fiber X: 18%, Y, Z: 6% HIP pitch matrix
K139 pitch fiber pitch matrix
Amoco P55 pitch fiber CVI pyrocarbon matrix SiC by liquid Si infiltrate 2800 C final heat treatment
X: Amoco P55 pitch fiber Y, Z PAN fiber X: 27%, Y, Z: 4% CVI and then pitch matrix 2800 C graphitization temperature 1.96 0.73
X: Amoco P120 pitch fiber Y, Z PAN fiber volume 30% CVI pyrocarbon matrix 2450 C graphitization temperature
Density Specific heat (J Kg1 K1) CTE 106
20 C 20 C
1.65–1.7
1.96 0.71
2.116 0.76
RT-400 C
X: 1.6, Y, Z: 5.2
X: 0.9, Y, Z: 12
X: 1.036, Y: 0.64, Z: 1.199
Thermal conductivity (W m1 K1, 20 C) Elastic modulus (GPa) Ultimate strength (MPa)
20 C
X: 368
X: 640
X: 265, Y: 124, Z: 109 X: 196, Y: 76, Z: 64 X: 146, Y: 58, Z: 49 X: 120, Y: 55, Z: 40
X: 0.339, Y: 1.376, Z: 0.018 X: 319, Y: 115, Z: 113 X: 196, Y: 72, Z: 68 X: 151, Y: 55, Z: 53 X: 107, Y: 15, Z: 12
X: 160, Y: 46, Z: 25
X: 130, Y: 30, Z: 19
X: 200, Y: 56, Z: 36 X: 230, Y: 67, Z: 40
X: 165, Y: 42, Z: 27 X: 185, Y: 50, Z: 30
X: 102, Y: 31
X: 102, Y: 31
Bend strength (MPa) Compressive strength (MPa) Shear strength (MPa) Poisson’s ratio
500 C 800 C 20 C
X: 100, Y, Z: 0.8 X: 400, Y, Z: 3
25 C 1000 C 1500 C 20 C
X: 39
20 C
X: 48
20 C
X: 480, Y, Z: >5 X: 216, Y, Z: >16
20 C
of the ITER project. A review of these properties emphasizes the anisotropic nature of the composite system, which is engineered through selection of the fiber type and route to matrix infiltration, fiber architecture, and final heat treatment of the system. All the materials for this application have been
1.88 0.7 X: 1.32, Y: 0.07, Z: 3.08 X: 413, Y: 102, Z: 78 X: 245, Y: 65, Z: 53 X: 78, Y: 52, Z: 38
XZ: 25, YZ: 15 XZ: 0.15, XY: 0.09, YZ: 0.15
XZ: 0.2, XY: 0.1, YZ: 0.1
engineered with a preferred thermal conductivity direction (the x direction in the table), and in order to maximize thermal conductivity, the composites will tend to have a higher volume fraction of fibers in that direction and the fibers will be of the higher conductivity pitch-based type. In the directions normal to
Carbon as a Fusion Plasma-Facing Material
this preferred thermal conductivity direction, for strength, cost, and fabricability reasons PAN-based fibers are typically chosen. The composite INOX Sepcarb NS31 underwent a final processing step of 10 2% liquid silicon infiltration. This silicon reacted with carbon-producing SiC, which is thought to mitigate chemical erosion and tritium retention while enhancing oxidation resistance. Also observed from Figure 6 is the clear difference in the shape of the load–displacement curves for the two materials. Clearly, the composite material has significant nonelastic behavior, which is attributed to the progressive load transfer from the composite to the high-strength fiber as the matrix becomes extensively microcracked. This contrasts with the graphite material, which undergoes abrupt failure when the load exceeds some critical stress adequate to propagate a crack through the test article. This added toughness of the composite is another key attribute to the systems that make it particularly attractive for fusion applications where disruption (shock) loading tends to produce interconnected cracking in materials leading to loss of material mass. 4.18.3.3.1 Irradiation-induced dimensional changes in CFCs
As discussed in Chapter 4.10, Radiation Effects in Graphite, irradiation-induced dimensional changes in graphite are highly anisotropic, and a strong function of irradiation temperature and neutron dose (dpa). The temperature range of interest for fusion applications varies from 100 C in areas well removed from the plasma of experimental devices, to over 1000 C for the surface of PFCs, which experience appreciable plasma flux, and for future power-producing machines. As described in detail in Chapter 4.10, Radiation Effects in Graphite, the mechanism of graphite irradiation-induced dimensional change is a combination of intra- and intercrystallite effects. Within the crystallites, displacement damage causes an hai-axis shrinkage (within the basal plane) and a hci-axis growth (perpendicular to the basal plane.) The upcoming ITER reactor will be the first fusion reactor to provide a flux of neutrons to produce measurable thermophysical effects to fusion structural materials. Even so, this will be a relatively modest fluence machine, with the maximum fast dose accumulating less than 1 1025 n m2 (E > 0.1 MeV), or less than a displacement per atom, over its lifetime. The work of Bonal provides data on the dimensional changes in CFCs, which are expected in this dose range. Specifically, his work11 irradiated 2D and 3D
0 Perpendicular −0.1 Irradiation-induced dimensional change (%)
594
−0.2 A05 CX 2002U DMS 678
−0.3
ITER range
0 Parallel
−0.1 −0.2 −0.3 −0.4 −0.5 −0.6 0
0.8 0.2 0.4 0.6 Neutron fluence (1025 n m–2), E > 0.1 MeV
1
Figure 7 Dimensional change at low (ITER-relevant) fluence for select two-dimensional (2D) (AO5 and DMS 678 polyacrylonitrile (PAN)-based composites) and CX-2002U 3-C composite. Parallel and perpendicular refer to parallel and perpendicular to fabric lay-up for the 2D composite and y–z plane of PAN fibers for the CX-2002U. Reproduced from Bonal, J. P.; Wu, C. H. J. Nucl. Mater. 1996, 228, 155–161.
composites to doses approaching 1 dpa in the temperature range of 610–1030 C. Figure 7 shows the dimensional instability that occurs in these materials in the sub-dpa region, specifically indicating a shrinkage. The work of Burchell12 in Figure 8 shows the dimensional change behavior of 1, 2, and 3 directional composites for doses somewhat in excess of the ITER lifetime. In this example, solid cylinders were irradiated at 600 C to doses ranging to 5 dpa and the resulting diameter and length measured. The behavior of each material can be explained by the accepted theory for dimensional change in graphite (Chapter 4.10, Radiation Effects in Graphite) after taking into account the individual fiber architectures, and by the observation that a model for fibers describes them as graphite fiber, filaments of circumferential or radial basal planes running parallel to the fiber axis. The irradiation-induced dimensional change of such a fiber is therefore a shrinkage in length and a growth in diameter. However, at doses
Carbon as a Fusion Plasma-Facing Material
1 Unidirectional fiber composite (UFC)
0
–1
–2 Fiber axis Fiber axis
–3
Axis parallel to fiber axes
Dimensional change (%)
–4 0.5
Random fiber (RFC) composite
0 –0.5 –1
Fiber axis Fiber axis
–1.5 Axis perpendicular to fibers axes
–2 –2.5 0
Three-dimensional balance weave
–0.5 –1 –1.5 –2 Composite diameter-PAN composite
–2.5
II Fiber axis - PAN composite II Fiber axis - pitch composite
–3
Axis parallel to a set of fiber axes
–3.5 –4
0
1
2 3 Neutron dose (dpa)
4
5
Figure 8 Dimensional change in carbon fiber composites at a moderately high neutron dose. Reproduced from Burchell, T. D. In Physical Processes of the Interaction of Fusion Plasmas with Solids, Plasma-Materials Interactions; Hofer, W. O., Roth, J., Eds.; Academic Press: New York, 1996; pp 341–382.
less than 1 dpa the dimensional change is relatively minor (Figure 7). As the dose is increased, the direction perpendicular to the fiber axis is more or less unchanged while a significant shrinkage along the direction parallel to the fiber axis occurs. At about 2–3 dpa, swelling in the composite occurs in the perpendicular direction. The random fiber composite
595
of Figure 8 has a random orientation of chopped PAN fibers in the plane of the composite. The specimen diameter shows practically no change perpendicular to the fiber axis to about 4.5 dpa, though it exhibits 2% shrinkage parallel to the fiber axis. The 3D balanced PAN weave fiber has essentially isotropic shrinkage to a dose of 2 dpa, at which point the diameter of the fibers, and hence the sample, begins to swell. Also given in the 3D composite plot in Figure 8 is the radiation-induced dimensional change parallel to the fiber axis of an Amoco P55 pitch fiber composite. This material was processed in an identical manner to the PAN fiber composite. From the plot, it appears that the pitch fibers, and thus the composite, undergo slightly less shrinkage, possibly due to the higher fiber crystallinity. This hypothesis is also supported by the observation that fibers with higher final heat treatment temperatures tend to exhibit less dimension change13 and it is also consistent with the observation that elevating the heat treatment temperature of graphite reduces the irradiation-induced shrinkage.14 The irradiation-induced dimensional changes are of fundamental importance to the design and performance of the fusion structure, and even more so of the PFCs. This is due to the need to precisely define the plasma edge. For this reason, it is instructive to look at the irradiation effects at the higher dose and temperature conditions representative of the nextgeneration fusion power devices. The data shown in Figures 9 and 10 provide higher temperature dimensional swelling data for the FMI-222 3D CFC and MKC-1PH 1D CFC, which were model, high thermal conductivity CFCs studied in the early phases of the ITER composite development program.15 In Figure 10, the dimensional change of the 1D composite yields substantial swelling perpendicular to the fiber axis and equally impressive shrinkage parallel to the fiber. The FMI-222 of Figure 10, a nearly isotropic orthogonal weave pitch-fiber composite with equivalent fiber volume fraction in the x, y, and z directions, undergoes a positive dimensional change (swelling) parallel to the cylindrical axis of the sample, which increased with increasing temperature. The magnitude of swelling was in excess of 10% at the highest temperatures studied at the 2 dpa dose level. This is in contrast to the FMI-222 swelling data reported by Burchell12 and Snead,16 also for HFIR irradiation, though at a lower irradiation temperature. Specifically, a contraction of 0.6% is interpolated from the data of Burchell for FMI-222 irradiated
596
Carbon as a Fusion Plasma-Facing Material
6 One-dimensional CFC MKC 1-PH 2 dpa irradiation
Diameter change (perpendicular axis)
Dimensional change (%)
4
2
0 33-N1
Length change (parallel to axis)
−2
−4 800
900
Figure 11 SEM image of the top surface of an FMI-222 composite following irradiation to 980 C, 2 dpa.
1000 1100 1200 1300 1400 1500 Irradiation temperature (⬚C)
Figure 9 Dimensional change at high irradiation dose and temperature for a balanced three-dimensional carbon fiber composite. Reproduced from Snead, L. L.; Burchell, T. D.; Katoh, Y. J. Nucl. Mater. 2008, 381, 55–61.
15
Dimensional change (%)
Three-dimensional CFC (FMI-222) 2 dpa irradiation
10
Length change (parallel to axis)
5
0
−5 600
800
1000
1200
1400
100 μm
1600
Irradiation temperature (⬚C) Figure 10 Dimensional change at high irradiation dose and temperature for a one-dimensional carbon fiber composite. Reproduced from Snead, L. L.; Burchell, T. D.; Katoh, Y. J. Nucl. Mater. 2008, 381, 55–61.
at 600 C to an equivalent fluence as the data of Figure 9. Snead16 reports on an 800 C irradiation to a substantially higher dose (7.7 1025 n m2) than the Figure 8 dose (2.4 1025 n m2). In this case, the material underwent a contraction of 3.6%
along the length of a bend-bar (2.3 6 30 mm). It was also noted in this work that the width and thickness direction exhibited swelling. Specifically, swelling parallel to the width direction (6 mm) was 1.4% and swelling parallel to the thickness direction (2.3 mm) was 5.9%. The overall dimensional effects were related to the effect of (measured) gross changes in the dimension of fiber bundles noting that gaps were evident on the surface of the bend bars. Figure 11 shows an example of the top surface of an FMI-222 composite irradiated in the present work to 980 C, 2.4 dpa. This composite underwent very low swelling. By inspection of the figure the contraction of the fiber tows below the free surface of the sample is evident. However, there is evidence from this micrograph that some of the fibers (particularly at the tow edge) have not withdrawn into the sample. This is evidence of shear within the fiber bundle as opposed to the tow–matrix interface. This observation is evidence of the large stresses that must be building in the composite under irradiation. The fact that the bundles are not failing at the tow–matrix interface also supports the previous finding that, at least in the initial period of gross dimensional change, the load-carrying capacity of the composite has not been degraded. In fact, previous measurement of FMI-222 irradiated to a dose of 7.7 1025 n m2 (E > 0.1 MeV) at 800 C described a 54% increase in strength.16 Figure 12 shows an scanning electron microscopy (SEM) image comparing the 2 dpa surface of the FMI-222 composite of Figure 11 with cylindrical samples of the same size, also irradiated near 1000 C, though at progressively higher doses. Clearly the dimensional instability continues with dose leading to gross changes in the composite.
Carbon as a Fusion Plasma-Facing Material
6 dpa, ~1025 ⬚C
2 dpa, ~980 ⬚C
22-A
1 mm
11N
1 mm
597
10 dpa, ~1065 ⬚C
29-J1
1 mm
Figure 12 Evolution of macrostructure for a three-dimensional composite under relatively high-dose, high-temperature irradiation.
Table 4
Properties before and after irradiation of IG-110 graphite and ITER-relevant composites
Young’s modulus (GPa) Bending strength (MPa) Bending fracture strain (%) Compressive strength (%) Compressive fracture strain (%)
Unirradiated Irradiated Unirradiated Irradiated Unirradiated Irradiated Unirradiated Irradiated Unirradiated Irradiated
IG-110 graphite
CC-312
8.83 11.5 35.2 38.4 0.532 0.364 85 82 2.67 1.49
34 31.3 90.5 110.8 0.324 0.394 65.1 93.7 0.79 0.393
MFC-1 X 74 98 103.9 98.4 0.332 0.174 59.8 55.9 0.165 0.156
CX-2002U Y
5.8 0.312 76.7 3.72
Z 87.6 87.2 99.2 88.9 0.294 0.325 59.6 51.0 0.71 0.11
14.9 18.4 36.3 46.7 0.44 0.435 33.3 41.2 1.94 1.82
Source: Eto, M.; Ishiyama, S.; Ugachi, H.; Fukaya, K.; Baba, S. J. Nucl. Mater. 1994, 212–215, 1223–1227.
4.18.3.3.2 Irradiation-induced changes in strength and modulus
Significant increases in both strength and elastic modulus occur in graphite at dose levels as low as 0.01 dpa.8 This increase continues to high displacement levels until volumetric expansion and extensive micro cracking occurs. For graphite, the reversal to property degradation typically occurs at tens of dpa depending on the graphite type and irradiation temperature. The increase in modulus is a result of dislocation pinning by lattice defects produced by neutron irradiation. The magnitude of the increase is dependent on the perfection of the graphite. For most graphite types, a maximum modulus increase of 2–2.5 times the nonirradiated value is typical for irradiation temperatures less than 300 C, with the change becoming less pronounced at higher irradiation temperatures. Irradiation-induced increase in strength occurs in a similar fashion as in the elastic modulus. Several authors17–23 report the effect of neutron irradiation on the elastic modulus of CFC. For example, Sato18 reports an increase of 42% and 30% in modulus following neutron irradiation to 1.2 1025 n m2,
E > 0.18 MeV in the temperature range of 750–810 C for a 2D pitch fiber and PAN fiber composite, respectively. Similar to the irradiation-induced increase in strength, the absolute increase and percent increase in elastic modulus is highly dependent on starting material and irradiation condition. The irradiated and nonirradiated mechanical properties of some candidate ITER PFC materials are shown in Table 4 for ITER relevant temperatures and doses somewhat higher than ITER neutron doses. Specifically, these materials were irradiated at approximately 1000 C to a dose of about 2 dpa.20 The change in properties is relatively small because of the high irradiation temperature and the relatively low dose. As with elastic modulus, reported data on the effect of irradiation on the strength of CFCs are somewhat sparse.16–22 Snead23 has reported the strength and elastic modulus of the 3D pitch fiber composite FMI-222 for doses higher than expected for ITER, or more consistent with a fusion power reactor. Figure 13 gives the modulus as a function of dose to 32 dpa at 800 C, exhibiting a marked increase to at least 10 dpa followed by a degradation by the 32 dpa
598
Carbon as a Fusion Plasma-Facing Material
4.18.3.3.3 Thermal conductivity degradation
is of primary importance. As with ceramics, graphite thermal conductivity is dominated by phonon transport and is therefore greatly affected by lattice defects, such as those caused by neutron irradiation. The extent of the thermal conductivity reduction is therefore directly related to the efficiency of creating and annealing lattice defects, and is therefore related to the irradiation temperature. The effect of neutron irradiation on the thermal conductivity of graphite has been widely studied. The majority of the literature10,13,23–31 in this area has been in support of the gas-cooled, graphite-moderated, fission reactor program in the United States and United Kingdom and has focused on ‘nuclear’ graphites as well as more fundamental work on pyrolytic graphite.8,27,32,33 In recent years, the emphasis of graphite radiation effects research has switched to its use in PFCs of graphite fusion reactors.10,11,34 Because of the significant advances in carbon– carbon composite (CFC) processing and fiber development, very high thermal conductivity materials have been recently demonstrated and they have become attractive for high heat flux applications. The highest thermal conductivity CFCs are made from highly crystalline graphite fibers having intrinsic conductivities approaching those of pyrolytic graphite. For example, vapor-grown carbon fibers35 have a thermal conductivity of 1950 W m1 K1. Along with advances in fiber properties, improvements have occurred in both monolithic graphite and the CFC matrix-processing areas, which also have enhanced thermal conductivities. The physical processes governing the thermal conductivity of graphites, as well as the mechanisms responsible for the radiation-induced degradation in conductivity, are well established.8 For all but the poorest grades of carbon, thermal conductivity is dominated by phonon transport along the graphite basal planes and is reduced by scattering ‘obstacles’ such as grain boundaries and lattice defects. For graphites with the largest crystallites, that is, pyrolytic graphite or natural flake, the in-plane room temperature thermal conductivity is approximately 2000 W m1 K1.36 The thermal conductivity of graphite-based materials can be written as a summation of the thermal resistance due to scattering obstacles: 1 1 1 1 þ þ ½7 K ðxÞ ¼ bðxÞ Ku Kgb Ki
CFCs for fusion applications are specifically designed to maximize thermal conductivity and for this reason irradiation-induced thermal conductivity degradation
where b(x) is a coefficient that includes terms due to orientation (with respect to the basal plane), porosity, and some other minor contributors. This coefficient
60 Sonic elastic modulus Strain gage
Elastic modulus (GPa)
55 50
Thermal control
45 40 35 30 0
10 20 Dose (dpa)
30
40
Figure 13 Effect of neutron irradiation on the elastic modulus of a balanced three-dimensional carbon fiber composite at high neutron dose. Reproduced from Snead, L. L.; Katoh, Y.; Ozawa, K. J. Nucl. Mater. 2010.
200
Ultimate flexural strength
Strength (MPa)
150 Proportional limit strength
100
50
0 0
5
10
15 20 Dose (dpa)
25
30
35
Figure 14 Effect of neutron irradiation on the strength of a balanced three-dimensional carbon fiber composite at intermediate and high neutron dose. Reproduced from Snead, L. L.; Katoh, Y.; Ozawa, K. J. Nucl. Mater. 2010.
value. The same samples, as seen in Figure 14, exhibit more than a 50% increase in strength, which is retained even at the 32 dpa value. This is particularly remarkable given that the composite had undergone significant dimensional change in this dose range.
Carbon as a Fusion Plasma-Facing Material
0.4 Normalized thermal conductivity Kirr / Kunirr
is assumed in most cases to be constant with temperature, with a value of around 0.6. The first two terms inside the parentheses are the contributions to the thermal conductivity due to umklapp scattering (Ku) and grain boundary scattering (Kgb). The grain boundary phonon scattering dominates the thermal resistance (1/Kgb) at low temperatures and is insignificant above a few hundred degrees Celsius, depending on the perfection of the graphite. The umklapp scattering, which defines the phonon– phonon scattering effect on the thermal conductivity, dominates at higher temperatures and scales nearly as T2.8 The umklapp scattering therefore defines the upper limit to the thermal conductivity for a ‘perfect’ graphite. Following Taylor’s analysis,37 the umklapplimited thermal conductivity of the graphite crystal would be 2200 W m1 K1 at room temperature, in close agreement with the best pyrolytic graphites or the vapor grown carbon fibers mentioned earlier. The third term in eqn [7], Ki, is the contribution to the basal plane thermal resistance due to defect scattering. Neutron irradiation causes various types of defects to be produced depending on the irradiation temperature. These defects are very effective in scattering phonons, even at flux levels that would be considered modest for most nuclear applications, and would quickly dominate the other terms in eqn [7]. Several types of irradiation-induced defects have been identified in graphite. For irradiation temperatures lower than 650 C, simple point defects in the form of vacancies or interstitials, along with small interstitial clusters, are the predominant defects. Moreover, at an irradiation temperature near 150 C,27 the defect that dominates thermal resistance is the lattice vacancy. Due to its sensitivity to the presence of defects, the temperature at which graphite is irradiated has a profound influence on the thermal conductivity degradation. As an example, Figure 15 shows one of the most complete sets of irradiation data on Pile Grade A nuclear graphite.38 This graphite is a mediumgrained, extruded, anisotropic material with a room temperature thermal conductivity of 172 W m1 K1 in the extrusion direction. Figure 15 presents the normalized room temperature thermal conductivity of this graphite of various irradiation temperatures. It is seen that as the irradiation temperature is decreased, the degradation in thermal conductivity becomes more pronounced. For example, following irradiation at 150 C, the thermal conductivity of this graphite appears to approach an asymptotic thermal conductivity of 1% of the original. The reason for
599
Pile grade A graphite Measured at ambient temperature
0.35 0.3
Tirr = 1150 ⬚C
0.25 600 ⬚C
0.2
920 ⬚C
450 ⬚C
0.15 300 ⬚C
0.1 250 ⬚C 200 ⬚C
0.05 150 ⬚C
0 0.1
1
10 dpa
Figure 15 Degradation in thermal conductivity as a function of irradiation dose and temperature. Reproduced from Kelly, B. T. Plot Constructed from Personally Communicated Data.
this is that as the irradiation temperature is decreased, the fraction of vacancies surviving a cascade event increases, and thus the number of vacancies available to scatter phonons is much higher for the lower temperature irradiation. Data have been published for CFCs whose thermal conductivities are similar to those of nuclear graphites, showing degradation similar to that expected from the graphite literature. For example, Burchell34 has shown that the saturation thermal conductivity for a 3-directional composite (FMI-222, Kunirr ¼ 200 W m1 K1 at RT) is reduced to 40% of the original room temperature conductivity following fast neutron irradiation at 600 C. Published data for the degradation of thermal conductivity in highly conductive CFCs have led to the conclusion that a higher initial conductivity composite results in higher absolute conductivity after irradiation.39,40 Figure 16 demonstrates this point. At the extremely damaging irradiation temperature of 150 C, it is observed that the absolute reduction (Kunirr Kirr) is substantially greater for the high thermal conductivity materials compared to the lower conductivity CFCs and graphite, as seen in Figure 16, although the normalized fraction (Kirr/Kunirr) is approximately the same for all the carbon materials in the figure. Moreover, saturation in thermal conductivity degradation occurs at a neutron dose of 1 dpa. Data for higher irradiation temperatures11,31 show that the higher thermal conductivity materials have a slightly larger fractional
600
Carbon as a Fusion Plasma-Facing Material
change in thermal conductivity (Kirr/Kunirr) compared to lower conductivity materials, although the absolute value of the irradiated thermal conductivity is still greater for the higher conductivity materials. A comparison of thermal conductivity degradation for a nuclear graphite (CH-45) with the composites FMI-222 and MFC-1 is given in Figure 17.31
Thermal conductivity (W m–1 K–1)
700 600 500
Tirr = 150 ⬚C, HFIR core Measurements at ambient temperature
FMI 1D
MFC-1
400
RGTi graphite Copper
300 Hercules 3D
200
FMI 222
100
H451 graphite
0 0.0
0.01 Damage level (dpa)
0.1
Figure 16 Comparison of absolute degradation in thermal conductivity for various graphite and carbon fiber composite materials irradiated at low temperature. Reproduced from Snead, L. L.; Burchell, T. D. J. Nucl. Mater. 1995, 224, 222–229.
For the low-dose regime relevant to machines such as ITER (less than 1 dpa or about 500 h in this figure), the conductivity is seen to decrease by a factor of two for the highest conductivity material (MFC-1) and by about 30% for the nuclear graphite. An algorithm has been developed to predict the thermal conductivity degradation in a high thermal conductivity composite (555 W m1 K1 at room temperature) as a function of radiation dose and temperature.41 The absence of irradiation data on CFCs of this type required the use of data from intermediate thermal conductivity materials and pyrolytic graphite to derive an empirical radiation damage term.24,27,28,39,42 An analysis of the effects of temperature and neutron dose on the thermal conductivity is shown in Figure 18. Specifically, the algorithm assumed the nonirradiated properties of the unidirectional fiber composite MFC-1 material compiled with an empirical radiation damage term. As with the experimental data of Figures 15 and 16, it is seen in Figure 18 that an enormous loss in thermal conductivity occurs at low irradiation temperatures. Presently, only a few data points exist that are relevant to the validation of this algorithm, and these are also plotted on the figure.39 The data do agree within the errors of irradiation temperature and thermal conductivity measurement, with the algorithm predictions. However, they
400 ~1 dpa
Thermal conductivity (W m–1 K–1)
350
~2 dpa
CFC : MFC-1; Tirr ~430 ⬚C
300 250 CFC : FMI-222; Tirr ~310 ⬚C
CFC : MFC-1; Tirr ~430 ⬚C
200 150 100
CFC : FMI-222; Tirr ~310 ⬚C H451 graphite; Tirr ~430 ⬚C H451 graphite; Tirr ~430 ⬚C
50 0 0.1
1
10
100
Time in reactor (h)
1000 0 200 400 600 800 1000 Time in reactor (h)
Figure 17 Comparison of the effect of neutron dose on the thermal conductivity degradation of a nuclear graphite and high conductivity carbon fiber composites. Reproduced from Snead, L. L. J. Nucl. Mater. 2008, 381, 76–82.
Carbon as a Fusion Plasma-Facing Material
601
400
Thermal conductivity (W m–1 K–1)
Conductivity at measurement temperature
Unirradiated
300
200 0.0 dpa 0.001 dpa 0.005 dpa 0.01 dpa 0.05 dpa 0.1 dpa 0.5 dpa 1 dpa
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100 Data from Bonal, et al. dpa = displacement per atom
0 300
400
500
600
700
800
900
1000
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Composite temperature (⬚C) Figure 18 Model for thermal conductivity degradation as a function of irradiation temperature and dose for high conductivity one-dimensional carbon fiber composite.
Flat plate design
Monoblock design Tmax = 1240 ⬚C – 700
– 870
– 1200
– 640 MFC-1PH 1D-C/C
Tmax = 920 ⬚C 8 mm
Tmax = 740 ⬚C
– 1100
– 570
– 700 – 520
– 940
– 390
– 500
– 810 – 300
– 390 – 570 – 440 – 360
Water coolant
33.6 mm
– 690
GlidCop DC copper
– 220 – 170
– 320 – 320
– 130
– 260
– 200
150 220 290 Unirradiated
140 1–3 dpa irradiation
1–3 dpa irradiation
Figure 19 Effect of neutron irradiation on thermal conductivity-driven temperature evolution in a monoblock and flat-plate divertor design.
are insufficient to validate the algorithm and the need clearly exists for additional data for this purpose. To illustrate the usefulness of such an algorithm, and the significance of the issue of thermal
conductivity degradation to the design and operation of PFCs, this algorithm has been used to construct Figure 19, which shows the isotherms for a monoblock divertor element in the nonirradiated and
Carbon as a Fusion Plasma-Facing Material
4.18.4 Plasma–Particle Interactions A range of particle types, fluxes, and energies strike the PFMs and interact in the near-surface region of the PFM. The most common interactions are with hydrogen fuel ions, ranging in energy from a
1 FMI-1D MFC-1 RGTi FMI-222 H451
0.8
Kirr/Kunirr
irradiated state and the ‘flat plate’ divertor element in the irradiated state. In constructing Figure 19, the thermal conductivity saturation level of the 1 dpa given in Figure 18 is assumed, and the flat plate and monoblock divertor shown are receiving a steady state flux of 15 MW m2. Both composite materials have been assumed to be in perfect contact with a copper coolant tube or plate. Figure 19 clearly shows two points. First, a very high conductivity composite is required to handle the extreme heat fluxes expected if the temperature is to be limited to <1200 C (Section 4.18.4). Second, the effect of neutron irradiation on the temperature is significant. In the case of the flat plate divertor, the temperature rise (DT) changes from 200 to 500 C following irradiation, while for the monoblock, it increases from 350 to 900 C. It should be noted that the larger temperature increase for the monoblock design is due not to the larger path length of graphite in that configuration, but rather to the larger amount of graphite material that is irradiated in the highly damaging low temperature regime (see Figures 15 or 18). The larger temperature increase for the monoblock design could be unacceptable from an erosion standpoint, as will be discussed in Section 4.18.4. Because of the serious thermal conductivity degradation in graphite, scenarios to limit the issue (such as baking the PFM) have been considered. Upon annealing above the irradiation temperature, some interstitial atoms become mobile and can recombine with vacancies, restoring the thermal conductivity of the lattice. It is therefore conceivable that intermittent annealing of the PFC could regain some of the irradiation-induced thermal conductivity degradation. Bake-outs are typically conducted between operating cycles of a fusion system for plasma impurity (usually oxygen) control. However, the wall-conditioning temperatures are typically limited to less than 300 C and for various reasons cannot be significantly increased. Inspection of data such as those given in Figure 2041 indicates that little recovery in thermal conductivity is possible unless bake-out temperatures approach 1000 C, and thus in situ annealing can be of only marginal benefit.
Normalized thermal conductivity
602
0.6 0.4 0.2 ~200 ⬚C irradiation 0.01 dpa
0 0
200
400 600 800 1000 1200 1400 1600 Annealing temperature (⬚C)
Figure 20 The effect of annealing on recovery in irradiation-degraded thermal conductivity of graphite materials. Reproduced from Snead, L. L.; Burchell, T. D. In Reduction in Thermal Conductivity due to Neutron Irradiation, 22nd, San Diego, CA, July 16–21, 1995; American Carbon Society: San Diego, CA, 1995; pp 774–775.
few eV to hundreds of eV. In addition to hydrogen ions, fuel by-product ions such as helium, and impurities from the first wall also impact the surface. Severe surface layer damage occurs because of such ion impacts, and significant erosion of surface material additionally occurs. Various mechanisms are responsible for erosion, dependent on the surface temperature of the graphite. The mechanisms can generally be characterized in order of increasing temperature phenomenon as physical sputtering, chemical erosion, and radiation enhanced sublimation (Figure 21).43 Above 2000 C, the vapor pressure of graphite dominates the erosion. In addition to the obvious issue of erosion-degraded component lifetime, removal and redistribution of carbon contributes to a high tritium inventory trapped in redeposited carbon.44 4.18.4.1
Chemical Erosion
When the eV or slightly more energetic ions of hydrogen atoms impact the graphite plasma-facing surface, they will chemically react with the surface. The form this reaction takes is quite dependent on the temperature of the graphite, and as the majority of the ions are reacting over nanometer lengths, the result is the production of an amorphous film of hydrogenated carbon suit at the near-surface layer. An analytical expression for this thermally activated erosion process has been put forward by Roth and Garcia-Rosales,45 with further development by Mech46 and Roth.47 It is noted that in the ion energies between 10 and 20 eV, there is still an active debate regarding the synergy of the ion impact and the chemical erosion.
603
Carbon as a Fusion Plasma-Facing Material
H⬚ onto soft amorphous C:H film
1 keV H 1 keV D 3 keV He
0.1
0.1
0.01
Physical sputtering
Chemical erosion
Radiation-enhanced sublimation
0.001 0
200 400 600 800 1000 1200 1400 1600 Temperature (⬚C)
Figure 21 Mechanisms of carbon removal from a graphite plasma-facing material as a function of temperature. Reproduced from Roth, J.; Bodhansky, J.; Wilson, K. L. J. Nucl. Mater. 1982, 111–112, 775–780.
Chemical erosion yield per incident ion
Sputtering yield (Y; atoms per ion)
1
(0.4 keV) H+ onto pyrolytic
0.01
H⬚ onto pyrolytic preirradiated by D+ -andhard amorphous C:H films H⬚ onto redeposited carbon
0.001 H⬚ onto pyrolytic
0.0001 0
The combination of energetic damage plus chemical reaction, which is sometimes referred to as ‘chemical sputtering,’ is discussed by Jacob and Roth48 and others. For low(RT) and intermediate temperatures, from 400 to 1000 C (Figure 21), the volatilization of carbon atoms by energetic plasma ions becomes important. As seen in the upper curve of Figure 21, helium does not have a chemical erosion component of its sputter yield. In the currently operating machines, the two major contributors to chemical erosion are the ions of hydrogen and oxygen. The typical chemical species that evolve from the surface as measured by residual gas analysis49 and optical emission50 are hydrocarbons, carbon monoxide, and carbon dioxide. The interaction of hydrogen with graphite appears to be highly dependent on the ion species, on material temperature, and on the perfection and type of the graphite. This is illustrated in Figure 22, which shows typical bell-shaped thermally activated erosion yield curves for hydrogen and deuterium ions on graphite. The shape of the yield curve is influenced by the competition for hydrogenation from the sp2 and sp3 hybridization states.51,52 Hydrogen ions incident to the surface are slowed down and preferentially attach to sp2 carbon atoms (such as graphite edge plane atoms) forming sp3 CH3 complexes. Above approximately 400 K, these CH3 complexes can be released, thus returning the structure to the sp2 state. It is important to note that this phenomenon
200
400 600 Temperature (⬚C)
800
Figure 22 Chemical erosion of graphitic materials as a function of temperature.
will only happen in the presence of simultaneous ion damage. It will not occur simply due to a thermal process. This step leads to chemical erosion products (a host of erosion species are possible). The ability of hydrogen to continue to be bonded to carbon drops as the temperature goes up. If there are no CH or CH2 precursors on the surface, then no volatile CH3 or CH4 complexes can be formed, and thus there is no chemical erosion. This balance yields a maximum erosion rate, which for undamaged pyrolytic graphite resides at 280–600 C.53 It is noted that more recent work by Balden54 has determined the maximum to be in the range of 872–1222 C. This mechanism was first elucidated by Horn55 and Wittmann.56 The rate of formation of CH2, CH3, and complex hydrocarbons from atomic hydrogen in well-graphitized material is fairly low unless the material is altered (damaged) in the near-surface layer. For preirradiated pyrolytic graphite (i.e., damaged graphite, meaning that a carbon atom has been removed from its lattice position, thus increasing the available sp2 sites) preirradiated by high-energy Dþ of Hþ ions, the total erosion yield following exposure to low-energy hydrogen increases dramatically. This is illustrated in the upper curves of Figure 22 that show more than an order of magnitude increase in erosion yield over
Carbon as a Fusion Plasma-Facing Material
the undamaged case. This increased carbon loss has been attributed to the creation of active sites for Ho attachment.57,58 This structurally dependent mechanism is supported by the data of Phillips et al.,59 which shows a factor of two difference in erosion yield between high- and low-quality pyrolytic graphite. 4.18.4.2 Doping of Graphite to Suppress Erosion Surface treatment of PFMs, while extremely effective for the current day short-pulse tokamaks (pulses typically less than a few seconds), are of limited value for the next-generation (quasi-steady state) machines because of the significant surface erosion expected. However, forming graphite or CFC homogenously with various erosion-mitigating elements is possible. The mechanisms for this mitigation are twofold: (1) geometric shielding by low erosion yield particulate, and (2) changes in local chemical reactivity due to the presence of doping atoms. Promising results have been obtained by doping of graphite with boron, which resides substitutionally in the graphite lattice, trapping migrating interstitials and altering the electronic structure of the material. Boron doping has been shown60 to both reduce the erosion due to oxygen, and to significantly reduce the sputtering yield due to methane formation. However, other factors, such as the drastic reduction in thermal conductivity that is unavoidable in boronized graphite, need to be factored into the overall picture. Boron is discussed in some detail here, mostly because it had received great early attention. However, newer dopant combinations61 have served to suppress erosion, and they have also not had such a negative impact on thermal conductivity, and are therefore considered superior for PFC application. As discussed by Garcia-Rosales,62 until the mid-1990s, boron was the primary ‘dopant’ of interest in fusion PFM graphites for chemical erosion mitigation. As seen in Figure 23,63 the inclusion of up to 15% boron in graphite can result in significant (an order of magnitude in the peak regions) reduction in the erosion yield. Studies to date indicate that for effective suppression a minimum of 3% boron in graphite is required. The mechanism behind this suppression may include the reduced chemical activity of the boronized material, as demonstrated by the increased oxidation resistance64 or the suppressed diffusion caused by the interstitial trapping at the boron sites. From the mid-1990s onward, many other metallic element additions to graphite were studied for
0.1 Chemical yield (CD4/D-ion)
604
Pyrolytic SEP-CFC (9% B) BE/C (JET) S2508 (3%B) USB 15(15% B) GB 120 (20% B) B4C
0.01
0.001 100
200
300
400
500
600
700
800
Temperature (⬚C) Figure 23 Effect of including the dopant atom boron on the suppression of chemical erosion of graphite. Reproduced from Roth, J.; et al. J. Nucl. Mater. 1992, 191–194, 45–49.
their possible beneficial effect on erosion mitigation. Specifically, elements such as silicon, titanium, tungsten, and vanadium have been studied with varying levels of success.62 These elements are somewhat less effective in erosion mitigation than boron, though a factor of two reduction is to be expected.65–67 In a recent review by Balden,61 considerably higher reductions in erosion are noted. More recently, emphasis has been on the use of multielement doping strategies.61 Because the removal of graphite is significant both in terms of gross material loss (possible consumption of the entire wall for power devices) and enhanced tritium retention for the resulting carbon dust, the effect of any additive to graphite in terms of physical properties or impact on plasma performance when eroded (Section 4.18.2.1) needs to be considered. With the exception of boron, additive elements discussed in the previous paragraph will all have a negative impact on plasma performance in comparison with carbon atoms, and therefore the balance of reduced mass loss compared to enhanced parasitic radiative plasma loss (eqn [2]) must be considered. As for physical properties, at levels well below the threshold at which they are effective for erosion suppression (3%), they are direct substitutional elements in the graphite lattice, effecting significant reduction in thermal conductivity (due to their mass-defect phonon scattering.) In contrast, titanium doping, as evidenced in materials such as the Russian RGTi material, serves to enhance the graphitization
Carbon as a Fusion Plasma-Facing Material
process, resulting in very well-crystallized materials of high (though somewhat anisotropic) thermal conductivity. A comparison of several element additions to graphite and their effects on the properties of graphite has been carried out by Paz68 and discussed by others.47,69 It is seen (Figure 2468) that all the metallic inclusions studied, with the exception of silicon,70 had, at these graphitization temperatures, the effect of enhancing the effective length (perfection) of the basal plane of the graphite crystals, which is directly linked to enhanced thermal conductivity. In comparison the basal crystal lengths of the Poco nuclear graphite see Table 2 and the Russian RGTi,71 which is processed at a similar graphitization temperature but with an applied electric current are shown in Table 2. The right hand side of Figure 24 shows the effect of varying the amount of titanium on the crystallite size, indicating that there is an increase in crystallite size with an increase of up to a few atomic percent of titanium. In addition to the thermal process for chemical erosion described in the previous section, a second route to erosion, which is limited to the regions very near the surface, is also of importance.62 Specifically, for ions of <100 eV, the formation of sp3 complexes occur at the surface of the graphite with very low binding energy (<2 eV compared to the 7.4 eV binding energy for carbon in bulk graphite). Because of this very low binding energy the complex can be easily physically sputtered from the surface. Doping of graphites is also somewhat effective in reducing this surface erosion, attributed to the build-up of higher mass ‘dopant’ elements as the carbon atoms
are preferentially sputtered from the surface, effectively armoring the surface.62 For machines that will run in steady state such as ITER, moisture and oxygen evolving from the surface may not be a significant issue. However, oxygen is the most damaging impurity to current tokamaks through its presence in the molecular form, or as water vapor, and its tendency to be strongly adsorbed by carbon PFMs. Consequently, this impurity has a large impact on the plasma performance and erosion. The release of oxygen from irradiated carbonaceous films has been reviewed by Haasz72,73 and others. It has been clearly demonstrated that the carbon flux away from the first wall is directly related to the evolving oxygen. Typically, the oxygen enters the plasma from the PFMs in the form of CO or CO2. Figure 25 shows the strong temperature dependence of the erosion yield of a variety of graphites and codeposited (near amorphous redeposited carbon) materials.74 It is noted that the data of Figure 25 are measurements of erosion yield by thermal oxidation in an O2 environment. Without special PFM surface treatment, such as plasma glow discharge and bake-out of the surface material, these fluxes dominate the surface erosion. For this reason, extensive research has been conducted into modification of graphite surfaces with impressive success in enhanced plasma performance.75 These improvements are due not so much to suppressed carbon erosion as to the decrease in the amount of oxygen released from the graphite. Toward this end, doped graphites have been modified to incorporate thermally and physically sputter-resistant carbides by doping wt% TiC
80
0 3 Graphitization temperature
Other materials
2100 ⬚C 2350 ⬚C
60
40
20
0
Pure C 5 at.% 5 at.% 5 at.% 5 at.% 5 at.% 5 at.% Poco RG-Ti
Si
W
B
Ti
Zr
V
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9
12 15
25 Basal crystal length, Lc (nm)
Basal crystal length, Lc (nm)
100
605
20
15 T = 2025 ⬚C 10
0
1
2 3 at. % Ti
4
Figure 24 Effectiveness of various dopants on increasing the basal crystal length of graphite. Reproduced from Burtseva, T. A.; Chugunov, O. K.; Dovguchits, E. F.; et al. J. Nucl. Mater. 1992, 191–194, 309.
606
Carbon as a Fusion Plasma-Facing Material
10−1
Temperature (K) 1000 800 700
2000 1500
10−2
Codeposits and implants TFTA N3-15 DIII-D ASDEX a-C:D Textor (estimate) DIARC film Ion implant (lower limit)
10−3 Erosion yield/(C/O2)
500
600
10−4 10−5 Pure graphites
10−6 10−7 10−8 10−9
10−10 10−11 0.4
HPG99 EK98 Okada and Ikegawa Duval Olander et al. basal plane, CO Olander et al. basal plane, CO2 Olander et al. prism plane, CO Olander et al. prism plane, CO2 Lang and Magnier Rodriguez-Reinoso et al. graphite M Rodriguez-Reinoso et al. graphite P Vietzke et al. CO Penzhorn et al. EK98 Penshorn et al. JET 2D C/C McKee C/C Bacos
0.6
0.8
1.0
Upper limit [2] (HPG99, EK98, AXF-5Q, ATJ, JET 2D C/C)
1.2
1.4
1.6
1.8
2.0
(1000/T) / (1 K–1) Figure 25 Temperature effect on erosion yield in various graphite grades. Reproduced from Davis, J. W.; Haasz, A. A. J. Nucl. Mater. 1999, 266–269, 478–484.
with titanium,75 boron,76,77 beryllium,78 and silicon.79 Comprehensive reviews can be found for the chemical erosion of graphite46,47,55,56,80 doped graphite by hydrogen,62 and also an article on the surface treatment of a graphite wall by Winter.75 4.18.4.3
Physical Sputtering
When a relatively high-energy impacting particle transfers energy to a near-surface carbon atom in an amount sufficient to overcome the lattice bond energy or surface binding energy, some carbon atoms may be displaced and these may move in a direction defined by the angle between their path and the initial path of the impacting atom. Analogous to striking a billiard ball, this angle must be between 0 and 90 . The energy imparted to the displaced atom follows that given in eqn [6]. For a relatively high-energy atom striking a surface normally, the recoiling atom cannot be sputtered from the surface. However, in off-normal impact or displacement cascade events from fusion neutrons that occur near the surface, some fraction of atoms will be emitted (physically sputtered) from the graphite surface. The amount of material lost from the surface is defined by the sputtering yield (Y ), which is the number of target PFM atoms emitted per plasma ion impacting the surface. From eqn [6], the energy transferred to a target PFM atom, which is directly
related to the erosion yield, is a strong function of the impacting particle mass and the mass of the material being sputtered. The impact angle also has a large effect on the number of atoms that receive adequate kinetic energy normal to the PFM surface to be physically sputtered. The plasma ions travel along the magnetic field lines that are at a shallow (grazing) angle with the PFM, typically 1–5 , though the ion impact angle will be modified by surface potentials and collisional processes.81 The quantitative effect of the mass, energy, and angle of impact on the sputter yield of impacting deuterium ions is shown in Figure 26(a) and 26(b). As the kinetic energy of deuterium increases, the total amount of energy transferred to the target atoms increases, and the average amount of energy per collision results in greater erosion. From Figure 26(a), it may be seen that the physical sputtering yield of light target atoms is considerably greater than that for the heavy atoms, primarily due to the reduced impact energy required to overcome the displacement energy of the higher target atoms. For example, in purely kinetic terms, approximately 20 eV is required to displace an atom of carbon from the surface, while 220 eV is required for an atom of tungsten. In the sub-keV energy range of plasma fuels, the high yield materials are therefore carbon and beryllium. As the impacting ion energy increases, the sputtering yield for all materials decreases as the depth of interaction of
Carbon as a Fusion Plasma-Facing Material
Sputtering yield (Y)
0.1 Be C Fe Mo W 0.01
0.001
Mo
Be C Fe
W
0.0001 10
(a)
100
104
1000
105
Impacting deuterium energy (eV)
Sputtering yield (Y)
10
Poco graphite Pyro : polished Trim.SP code Pyro : basal Pyro : edge
1
3 keV C+
(a)
0.1
Graphite
2 keV D+ 0.01 0
(b)
20
40
60
80
Angle (a)
Figure 26 Sputtering yields for plasma-facing material atoms of various candidate materials (a); and as a function of angle of incidence for various graphite materials.
the impacting ion becomes too great for displaced atoms to back scatter to the surface. In the case of graphite, the majority of the sputtered material comes from the top few atomic layers.82 With the correct combination of incident energy and target mass, it is possible for the sputtering yield to exceed unity, that is, more than one atom leaves the surface for every impacting particle. This quickly leads to what is called the catastrophic ‘carbon bloom,’ that is, the self-accelerating sputtering of carbon. As can be seen in Figure 26(b), this problem is the worst for carbon self-impacts at grazing angles to the surface. 4.18.4.4
Radiation-Enhanced Sublimation
The limiting temperature for graphite use in fusion systems is defined by thermal sublimation (1500– 2000 C). However, a process that is very similar to
607
thermal sublimation (in cause and in effect) appears to define the current temperature limit. This phenomenon, which is known as radiation-enhanced sublimation (RES), is not clearly understood yet, but dominates above a temperature of about 1000 C, and increases exponentially with increasing temperature. One theory says that the process responsible for initiating RES follows from the earlier discussion of radiation damage in graphite. Specifically, in a displacement event, a Frenkel pair is created. The interstitial has a low (0.5 eV) migration energy, is quite mobile between the basal planes, and thus diffuses readily. Some fraction of these interstitials are condensed at vacancy sites, which are essentially immobile below about 700 C (migration energy 4 eV). Other migrating interstitials can be trapped by microstructural defects or can coalesce into simple clusters, which limits their mobility. However, some fraction of the interstitials diffuse to the surface of the graphite and thermally sublime. The thermal sublimation of radiation-induced interstitials is RES, and must be distinguished from both physical and chemical sputtering. Time of flight measurements have shown that the thermal energy of RES ions has a Maxwellian energy distribution, which is directly coupled to the mean surface temperature.83 This clearly distinguishes RES atoms from physically sputtered atoms, which exhibit highly anisotropic energy distributions. RES atoms are also distinguished from thermally sublimed species in that only single carbon atoms are detected, whereas single atoms and atom complexes (C2, C3, . . .) are found during thermal sublimation. Another theory for the explanation of RES is simply that the bombarding hydrogen ions turn the very near-surface region into a low-density amorphous zone. A very large fraction of the carbon atoms in this zone are now edge atoms with weak bonding to the connecting atoms. These edge atoms are much more easily thermally volatized into the plasma. The effect of RES in the next generation of high surface particle flux fusion systems is presently unclear. Evidence suggests that the erosion yield does not scale linearly with flux, as physical sputtering does, but may in fact decrease significantly with increasing flux.84 Moreover, as with chemical erosion, the inclusion of interstitial boron into the crystal lattice can decrease RES and shift the threshold to higher temperatures. Boron will volatilize above 1500 C, thus limiting the PFM temperature to <1500 C.
608
Carbon as a Fusion Plasma-Facing Material
4.18.4.5 Erosion of Graphite in Simulated Disruption Events Finally, the effect of plasma disruptions needs to be considered. Section 4.18.2 discussed the thermomechanical response of the PFCs to the excessive plasma energy in a disruption. This large thermal energy dump can additionally cause enhanced erosion due to the increased particle flux, elevated surface temperature, or simply by exfoliation of the surface due to thermal shock. The latter two material losses are reduced for materials with high thermal conductivity. This has been demonstrated experimentally, and is shown in Figure 27,3 which gives weight loss as a function of thermal conductivity for a number of graphites and composites of varying thermal conductivities subjected to one electron beam pulse at 4.1 MW m2. As discussed in Section 4.18.2, and as seen in the data of Figure 27, high thermal conductivity materials reduce the surface temperature, and hence the overall erosion yield, during a disruption.
4.18.5 Tritium Retention in Graphitic Materials Tritium retention and transport is a critical phenomenon for graphite in fusion systems in general, and it is the subject of a chapter by Causey, San Marchi, and Karnesky in this series. In the previous section of this
4.5 4
Weight loss (mg)
3.5 3 2.5 2 1.5 1 0.5 50
100
150
200
250
300
350
Mean ambient thermal conductivity (W m–1 K–1) Figure 27 Weight loss as a function of graphite thermal conductivity. Reproduced from Akiba, M.; Madarame, H. J. Nucl. Mater. 1994, 212–215, 90–96.
chapter, the interaction of the plasma particle flux with the surface of graphite was discussed. However, the fate of the implanted particles, most importantly deuterium and tritium, following their impact with the graphite surface is also an important issue and is seen by some as the major impediment to the use of graphite as a PFM.85 Quantification of the problem and determination of possible mitigating steps is complicated by experimental data, which can vary by orders of magnitude,86–92 as reviewed by Wilson.93 The primary concern over retention of fuel in the PFC is the inventory of hydrogen adsorbed into the graphite and the subsequent release of near-surface hydrogen (due to physical or chemical sputtering, etc.) as plasma discharge begins. The hydrogen sputtered from the wall oversupplies the plasma edge with fuel, causing instabilities and making plasma control problematic. Tritium inventory concerns are generally safety-related but can have significant economic consequences because of the high cost of tritium. Tritium release to the environment in an accident situation had limited the allowed inventory in TFTR, and was a significant consideration for the sighting of the ITER. It has been estimated84 that as much as 1.5 kg of tritium would reside in the graphite PFM of ITER, corresponding to an additional fuel cost of 1.5–3 million dollars. A source of trapped hydrogen, not discussed in detail here, which may dominate the tritium inventory in ITER-like machines, is the ‘codeposited layer.’94 This layer is formed by the simultaneous deposition of carbon, which is eroded from the first wall, and hydrogen. Thick layers of carbon redeposited to low erosion areas are common, and have been seen in all large tokamaks utilizing graphite PFMs. As this layer grows, the hydrogen contained therein cannot be liberated by surface sputtering and becomes permanently trapped. This problem is unique to graphite and will require continual surface conditioning to minimize the total inventory of trapped species. It represents a vast sink for tritium and therefore must be managed in some way. Below is a discussion of the retention of tritium in bulk graphite. The physical process involved in the retention of hydrogen, as it corresponds to graphite PFMs, is fairly well understood. The energetic hydrogen isotopes are implanted to depths of less than a micron in the PFM surface. Once implanted, the hydrogen ions are either trapped, reemitted, or diffuse through the bulk. At temperatures less than 100 C,95,96
Carbon as a Fusion Plasma-Facing Material
the majority of ions are trapped near the end of their range. These trapped ions are not in solution in the graphite, but are held97 in the highly defected structure. The amount of hydrogen isotope that can be accommodated is largely dependent on the implantation temperature,96,98 the trap types and densities (defects), and, to a lesser extent, by the implantation depth.99,100 One model for bulk hydrogen trapping presented by Atsumi101 is shown in Figure 28. In his work, two distinct traps have been identified. The lower energy trap (2.6 eV) is associated with edge planes of the graphite crystals, with total trapping therefore depending on the effective size and accessibility of the crystals to diffusing hydrogen species. The second, higher energy trap (4.4 eV), is associated with dangling bonds, albeit at the edge of an interstitial loop. As the formation of small interstitial loops is one of the primary effects of neutron irradiation, the formation of these deeper traps is directly affected by neutron fluence. The total retained isotopic H can reach as much as 0.4–0.5 H/C in the implanted layer at room temperature.95,100,102 As the amount of implanted hydrogen increases toward its saturation value, a larger fraction of ions are released from the graphite surface. At intermediate and high temperatures (>250 C), diffusion of hydrogen in the graphite lattice occurs. This diffusion is most likely along internal surfaces such as micropores and microcracks, while transgranular diffusion has been seen above 750 C.103,104 This bulk diffusion, along with the associated trapping of hydrogen at defect sites, has been studied widely
H2
1–50 mm
with quite variable results. This variation can be seen in Figure 29, where the temperature dependence of the hydrogen diffusion coefficient for several carbon and graphite materials is shown. It is expected that the diffusion of hydrogen through graphite would be highly dependent on the graphite microstructure, which may explain the wide range of the data of Figure 29. In any event, the transport of hydrogen through the bulk graphite and associated solubility limits can significantly increase the hydrogen inventory for fusion devices. The effect of the perfection of graphitic structure on the solubility of hydrogen is shown by Atsumi’s data105 in Figure 30. The data in Figure 30 indicate that the more defect-free, highly graphitized materials have a lower solubility limit. Further evidence for the role of structural perfection comes from the observation that materials that have been disordered by neutron irradiation have significantly higher solubility for hydrogen.105,106 The effect of atomic displacements on the hydrogen retention of graphite was first shown by Wampler using 6 MeV ion beams.107 Wampler used four types of intermediate and high quality graphites, which were irradiated with a high-energy carbon beam at room temperature, and then exposed to deuterium gas. Wampler’s results indicated that the residual deuterium concentration increased by more than a factor of 30–600 appm for displacement doses appropriate to ITER. However, for reasons that are not entirely clear yet, neutron-irradiated high-quality CFCs retain significantly less tritium than expected
30–200 nm H atom
0.3354–0.365 nm
∼mm
ED = 1.3 eV Gas permeation through open pores
Molecular diffusion (with a sequence of dissociation and recombination)
Absorption
(Rate-determining step)
Desorption
609
Trap 2 (2.6 eV)
Trap 1 (4.4 eV)
Trapping at the edge surface of a crystallite
Trapping at the activated edge of an interstitial cluster loop
(>90%)
Detrapping (rate-determining step)
(<10%)
Detrapping (rate-determining step)
Figure 28 Schematic of the processing of hydrogen trapping and diffusion in graphite. Reproduced from Atsumi, H. J. Nucl. Mater. 2003, 313–316, 543–547.
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Carbon as a Fusion Plasma-Facing Material
10−9
2
Diffusion coefficient (cm2 s−1)
10−10 10−11 10−12 4
10−13
1
8
10−14
7
6 1 - Atsumi et al. 2 - Elleman 3 - Saeki 4 - Maika et al. 5 - Rohrig et al. 6 - Causey 7 - Morita et al. 8 - Tanabe and Watanabe
10−15 10−16 10−17 10−18 0.4
0.6
3 5
1
0.8 3/T
10
1.2
1.4
(K)
Figure 29 Variation in hydrogen diffusion coefficient as reported in the literature. Adapted from Causey, R. A. J. Nucl. Mater. 1989, 162–164, 151; Atsumi, H.; Tokura, S.; Miyake, M. J. Nucl. Mater. 1988, 155–157, 241; Saeki, M. J. Appl. Radiat. Isot. 1983, 43, 739; Malka, V.; Rorhig, H. D.; Hecker, R. In Tritium Technology in Fission, Fusion and Isotope Application; Dayton, OH, 1980; Rohrig, H. D.; Rischer, O. G.; Hecker, R. J. Am. Ceram. Soc. 1976, 59, 316; Morita, K.; Ohtsuka, K.; Hasebe, Y. J. Nucl. Mater. 1989, 162–164, 990; Tanabe, T.; Watanabe, Y. J. Nucl. Mater. 1990, 179–181, 231–234.
104
Tritium retention (appm)
Hydrogen solubility (appm)
104
1000
100
1000
N3M graphite FMI-222 CFC MFC-1 CFC
100
10
Unirradiated Neutron irradiated 10 0
20
40 60 80 Graphitic perfection (%)
100
1 0.001
0.01
0.1
1
10
Radiation dose (dpa)
Figure 30 Effect of perfection of graphite crystal on hydrogen solubility. Reproduced from Atsumi, H.; Iseki, M.; Shikama, T. J. Nucl. Mater. 1994, 212–215, 1478–1482.
Figure 31 Effect of graphite perfection on tritium retention as a function of neutron irradiation. Reproduced from Causey, R. A. Phys. Scripta 1996, T64.
from the earlier work. This was reported by Atsumi105 and clearly shown in the work of Causey108 (Figure 31). Causey irradiated high thermal conductivity MFC-1 unidirectional composite and FMI-222
3D composite at 150 C (a particularly damaging irradiation temperature regime), to a range of displacement doses up to 1 dpa. As seen in Figure 31, tritium retention is more than one order of
Carbon as a Fusion Plasma-Facing Material
4.0
611
20.0
Hydrogen retention (⫻ 10−3 H/C)
IG–430U 3.0
Trap 2 (crystallite edge surface)
15.0
Trap 1 (interstitial cluster loop) 2.0
10.0
1.0
5.0
0
0
Unirradiated 4.7 ⫻ 1022 9.4 ⫻ 1022 1.4 ⫻ 1023 3.9 ⫻ 1023 0.006 dpa 0.011 dpa 0.017 dpa 0.047 dpa
1.9 ⫻ 1024 5.4 ⫻ 1024 0.23 dpa 0.65 dpa
Figure 32 Loading of defects in irradiated graphite with hydrogen. Reproduced from Atsumi, H.; Muhaimin, A.; Tanabe, T.; Shikama, T. J. Nucl. Mater. 2009, 386–388, 379–382.
magnitude less than that expected from the earlier work on GraohNOL-N3M.109 In some more recent work by Atsumi,110 different neutron-irradiated graphites were irradiated and the amount of hydrogen that could be entrained in the crystal was measured. Specifically, Atsumi110 irradiated three isomolded grades of graphite (IG-110, IG-430U, and IG-880U) in the JMTR reactor below 200 C to various fluences and then baked the samples in an atmosphere of hydrogen. Figure 32 shows the relative abundance of hydrogen that can be loaded into the crystallite edge and interstitial cluster looptype defects (Type 1 and Type 2 defects of Figure 28) of the IG-430U graphite. Clearly, both defects are produced during irradiation and are accessible to postirradiation loading of hydrogen. In the same work, Atsumi carried out a series of preloading annealing of samples, which suggested that the edge-type defects would be preferentially annihilated. In this context, it is important to note that all work on hydrogen or tritium retention in irradiated graphite has followed the approach of irradiating the material at a relatively low temperature and then loading and unloading and measuring the released hydrogen from the sample at a comparatively higher temperature. This may be of significance in that, as inferred in the work of Atsumi and others,105,106,108 the relative crystalline perfection (amount of intrinsic defects) is strongly related to hydrogen retention. As discussed in Section 4.18.3 and Chapter 4.10, Radiation Effects in Graphite, irradiation at low temperature may result in a significantly different microstructure (an abundance of simple interstitial and vacancy clusters) as compared to the more fusion
relevant irradiation temperatures (formation of more perfect interstitial discs and collapsing of vacancy complexes). Moreover, the postirradiation annealing of a low-temperature-irradiated microstructure will likely not produce representative microstructures of irradiation at more relevant higher temperatures. For this reason, data generated to date should be considered as a guide for the trends likely to occur rather than as quantitative information on the actual tritium retention that will occur in fusion devices. Moreover, they are likely overly conservative.
4.18.6 HHF Component Technology 4.18.6.1
Joining of CFC to Heat Sink
CFCs, bonded either mechanically or otherwise to a metallic structure, are being used in most of the major fusion devices, including ITER,105,106 Tore Supra,111 Wendelstein 7-X,112 TFTR in United States,113 JT 60U, and JT60SU in Japan.110–113 Silicon-doped carbon is used as the first wall material for the Chinese reactor HT-7.114 See Table 3 for a description of ITER candidate composites (INOX Sepcarb NS-31, Sepcarb NB-31, and Dunlop Concept C1, for example). As mentioned, CFCs will be used in ITER (see Table 5 for performance specifications), and specifically a 3D CFC for the divertor in the initial phases of the ITER project. The diverter, which is among the most technically challenging ITER components, is located at the bottom of the plasma chamber where the CFCs (and tungsten) are bonded to a copper alloy (CuCrZr). CFCs have been selected
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Table 5 Operational conditions for CFC joined to copper alloy in the lower vertical target, component subjected to glancing incidence of heat flux Heat flux (MW m2)
10/20a
Number of cycles
3 103
Damage (dpa)
Tmax of the C/C – CuCrZr joint ( C)
0.2
Steady state 300
Transient 400
(a)
a 20 MW m2 transient events, duration 10 s, number 10% of normal shots. Reproduced from Ferraris, M., et al. In Ceramic Integration and Joining Technologies: From Macro to Nanoscale, 1st ed.; Singh, M., Ohji, T., Asthana, R., Mathur, S., Eds.; Wiley, 2011; Chapter 3, © 2010 The American Ceramic Society.
for the lower part of the vertical targets for the initial phase of ITER operation (without tritium), and they must be able to resist ‘steady state’ heat flux up to 10 MW m2 for at least three thousand 400 s pulses and up to 20 MW m2 in transient events. As discussed earlier, the use of graphite-based materials (particularly CFCs) in a divertor is believed to be an advantage for the first phase of ITER operations because CFCs have good demonstrated performance in the currently operational plants (e.g., Tore Supra). Their primary competitor, tungsten, also suffers from macroscopic cracking, melting, and possible melt layer loss, thus making the potential damage to the divertor components more serious than that of CFC. As conditions of additional heating and off-normal events will be very likely during initial operation of ITER, CFC is the present reference design solution for the lower part of the vertical targets for the ITER initial phase, when tritium retention-related issues are not relevant. In the ITER design, CFC must be joined to the copper alloy CuCrZr-IG (ITER Grade)116–119 in order to transfer the heat loads. Two different designs being considered for this component are the CFC-bonded flat-tile (Figure 33(a)) and monoblock (Figure 33(b)). In order to join the mating surface of the CFC to the copper alloy heat sink, a pure copper interlayer (1–2 mm thick, oxygen free high conductivity copper, 99.95%, CTE: 15.4 at RT, up 20.6 at 700 C and up to 21.6 at 800 C) is used to relieve, by plastic deformation, the thermal expansion derived between the CFC and copper alloy heat sink. The CTE of CFC can be found in Table 3, with that of the copper alloy stresses being (16–19) 106 K1 at 700 C).116 This joining design is shown schematically in Figure 34. Some alternatives to a pure copper interlayer have been proposed within the EU project ExtreMat. For example, an Mo interlayer (1–2 mm),
(b) Figure 33 (a) Flat tile design. Reproduced from Ferraris, M., et al. In Ceramic Integration and Joining Technologies: From Macro to Nanoscale, 1st ed.; Singh, M., Ohji, T., Asthana, R., Mathur, S., Eds.; Wiley, 2011; Chapter 3, © 2010 The American Ceramic Society. (b) Monoblock design. Courtesy of J. Linke, FZJ, Germany.
a Cu/W fiber interlayer, and CFC monoblocks with a Cu/W fiber interlayer (Figure 35) have been prepared and tested up to 10 MW m2. Results (unpublished) indicate that this method is less promising compared to the use of a pure Cu thin layer. Ti-doped CFC have also been prepared and tested for flat tiles.117,118 Among the several possible options, the flat tile and monoblock configurations (see Figure 33(a) and 33(b)) with a pure copper interlayer have yielded the most promising results. In particular, the monoblock gives a more robust solution in comparison with the flat tile for the vertical target and it is now considered as ITER reference geometry.119,120 The monoblock design requires drilled blocks of CFC into which a CuCrZr tube is inserted and joined; also necessary in the monoblock is a pure copper interlayer between the CFC and the copper alloy to relax interface stresses, which have been (modeled and) measured at 45 , 90 , 135 . If 0 is considered to be the flux direction,121,122 the monoblock is preferred to the flat tile design. This design is also much easier to manufacture because of its better heat flux performances and because of its intrinsic ability to attach even in the presence of cracks at the interface of CFC and CuCrZr, caused by its preparation process. Methods to join CFC to CuCrZr derive from techniques developed to join carbon-based components to metals. Brazing is a well-known joining process recommended for joining dissimilar
Carbon as a Fusion Plasma-Facing Material
613
Monoblock
Flat tile
CFC CFC Pure copper interlayer
Pure copper interlayer CuCrZr CuCrZr
(a)
(b)
Figure 34 Carbon fiber composite joined to CuCrZr in a (a) flat tile and (b) monoblock configuration (drawings show test components and not full-scale components). Reproduced from Ferraris, M., et al. In Ceramic Integration and Joining Technologies: From Macro to Nanoscale, 1st ed.; Singh, M., Ohji, T., Asthana, R., Mathur, S., Eds.; Wiley, 2011; Chapter 3, © 2010 The American Ceramic Society.
1500 mm
Figure 35 Cross-section of a carbon fiber composite monoblock with a Cu/W fiber interlayer. Courtesy of EU-Extremat Project and Pintsuk, G. JFZ, Germany.
materials. If properly done, it results in good mechanical strength, high fatigue performance, and minimal thermal resistance at the interface.123–125 In the case of joining of CFC for nuclear applications, some additional restrictions must be considered. A primary consideration is that the joining materials should be low activation materials (LAMs), even if the volume occupied by these materials is negligible in respect of the total volume of the structural materials in the reactor.126 Also important is that the materials be irradiation-stable at the anticipated conditions. Furthermore, the use of pressureless joining
techniques is preferred as the parts to be joined are relatively large. Some joining materials are not allowed, for example, elements with high vapor pressure (e.g., zinc or cadmium), or those giving dangerous transmutation reactions. The ITER project mandates thermodynamic and mechanical stability up to at least 800 C under vacuum for the joint, in order to satisfy requirements of Table 5; the joint must survive the thermal, mechanical, and neutron loads faced by the component, and it is expected to operate in a cyclic mode with an acceptable reliability and lifetime. The joining technology must also be compatible with the overall component manufacture process and in particular with the preservation of the thermomechanical properties of the precipitationhardened CuCrZr alloy.127 Wettability is a key factor in the joining of CFC to the heat sink. It is not within the scope of this chapter to review this subject, and an overview can be found elsewhere.124,128,129 However, to restate the most salient point, the Cu interlayer cannot be obtained by directly casting copper on the CFC surface, because Cu does not wet CFC at all, the contact angle of molten copper on carbon substrate being about 140 .128,129 The poor wettability of CFC is related to the nonmetallic character of its bonding, whereas the bonding electrons in copper are delocalized.123,124,128 Copper can be directly cast on CFC when the CFC surface is modified to form carbides by a solid state chemical reaction between the composite surface and elements such as Si, Al, Ti, Zr, Cr, Mo,
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Carbon as a Fusion Plasma-Facing Material
ABCD
Tungsten (upper part)
Plansee
Blocks 64
Blocks 33
Austenitic steel Blocks 32
CFC (lower part)
Copper/steel tube joint
1000 mm
or W; some metals can form carbides with ‘metal-like’ behavior128 and are usually well wetted by molten metals. Several patents refer to joints between carbon-based materials and copper130–135: for example, pure Cr and Ti react with C to form carbides. Cr and Ti wet carbon-modified surfaces very well with a contact angle of 35–40 at 1775 C and of 50–60 at 1740 C, in Ar, respectively.128 A joining technique based on CFC surface modification is the Active Metal Casting (AMC®) technique originally developed in the eighties by the Austrian company Plansee for nonnuclear purposes. ‘Active’ indicates the activation of the CFC surface to allow it to be wetted by Cu. Physical vapor deposition (PVD) or chemical vapor deposition (CVD) of Ti coating on the CFC surface is followed by a high temperature treatment to form TiC, which improves the wettability of CFC by molten copper. Active Metal Casting consists of casting a pure Cu layer onto a laser-textured and TiC-modified CFC surface.136,137 The laser texturing enhances Cu infiltration into the CFC, and the TiC-modified CFC surface improves the wetting. The special laser treatment of the CFC surface produces a large number of closed conical holes (diameter 50–500 mm, depth 100–750 mm), thus increasing the joined area and providing additional crack growth resistance. Due to the open porosity of the TiC-modified CFC and laser machining, the cast Cu penetrates into the CFC up to 2 mm. An example of a full-scale component produced by Plansee, Austria, is shown in Figure 36. AMC® was successfully applied for both flat tile and monoblock geometry. However, AMC® technology requires laser machining of CFC surfaces, which might not be economically attractive for large-scale production. Laser-induced stresses in the joined area and cracks induced during the joining process have recently been measured and modeled.121 However, Plansee has recently improved AMC® by using silicon and titanium to modify the CFC surface (TiSi-AMC) (Figure 37(c)).130 Ansaldo RicercheGenova, Italy, has proposed a joining technique based on a Cu–Ti-based (Cu ABA) commercial alloy, reinforced by 2D randomly oriented carbon fibers uniformly distributed in the brazing alloy. The joining is carried out at approximately 1000 C. The Ti reacts with carbon to form a thin TiC layer that promotes wetting.131 Carbon fibers are expected to mitigate thermal expansion mismatch between CFC and the braze (Figure 38) and to react with titanium in the brazing alloy, resulting in beneficial thermal fatigue strength of the joint. This technique was
Blocks 1
Figure 36 Vertical target full-scale prototype manufactured by Plansee (high heat flux units) and Ansaldo Ricerche (support structure and integration). Reproduced from Missirlian, M., et al. J. Nucl. Mater. 2007, 367–370(2), 1330–1336.
successfully tested on CFC NB31-Cu flat-tile and monoblock joints. Several other solutions have been investigated to modify CFC surface, for example, by TiN or TiC, within the EU project ExtreMat.118 A method has been proposed based on the CFC surface modification by reaction of Cr, Mo, and W. Both Cr and Mo have been extensively used as active elements in brazing alloys for copper active brazing and in patents referring to nonnuclear applications.130–135,138,139 As example W, Mo, and Cr powders were deposited on CFC (CFC NB31, Snecma Propulsion Solide, France) by the slurry technique: details related to the process can be found elsewhere.132–134 Cr-carbide-modified CFC appears to have yielded the best results (Figure 39). A 15-mm-thick carbide (Cr23C6, Cr7C3) layer has been identified by XRD on CFC; the CTE of the carbides lies between that of CFC and copper (reported above) (CTE of Cr7C3 is 10 106 K1).135 A commercial brazing alloy Gemco® (87.75 wt% Cu, 12 wt% Ge and 0.25 wt% Ni; Wesgo Metals) has been used to braze Cr-modified CFC to Cu and
Carbon as a Fusion Plasma-Facing Material
615
150 mm CFC
Braze CuCrZr
Cu interlayer
CFC
AMC
(a)
Cu
(b)
100 μm
CFC
0.5 mm
Figure 38 Carbon fiber composite/Cu active brazing with Cu–Ti alloy and dispersed carbon fibers. Reproduced from Bisio, M., et al., Fusion Eng. Des. 2005, 75–79, 277–283.
0.25 mm 0.2 mm
TiC
TiC
20 μm
Cu OFHC
Cu
TiC + SiC
Cr carbide 10 μm C/C (c) Figure 37 (a and b) Laser structuring of carbon fiber composite (CFC) in Active Metal Casting (AMCW) process and cross-section of the AMCW CFC–Cu joint. Courtesy of Chevet, G. Ph.D. Thesis 2010, University of Bordeaux, France.
to CuCrZr in a single step process,140 which is an advantage in comparison with other joining technologies that require two steps: first joining CFC to Cu, and then CFC–Cu to CuCrZr. Flat-tile (a) and monoblock (b) mock-ups have been obtained by this technique and tested (Figure 40). ENEA, Italy, has manufactured several actively cooled mock-ups of flat tile and monoblock type, by using different technologies; a new process (patented) for the production of monoblocks is based on prebrazed casting and hot radial pressing (PBCþHRP) (Figure 41). The CFC surface modification is obtained by a titanium–copper–nickel commercial brazing alloy, which is followed by a Cu casting, then a radial diffusion bonding between the cooling tube and the CFC by pressurizing only the internal tube and keeping the joining zone in vacuum at the required bonding temperature.120,141,142 Complete manufacture and testing of this vertical target medium-scale mock-up (Figure 41) can be
(a)
31 mm
38 mm
74 mm
(b) Figure 39 Carbon fiber composite–Cu active brazing with Cu–Ti alloy and dispersed carbon fibers in the braze (b). Reproduced from Schedler, B.; Huber, T.; Eidenberger, E.; Scheu, C.; Pippan, R.; Clemens, H. Fusion Eng. Des. 2007, 82(15–24), 1786–1792. Reproduced from Ferraris, M., et al. In Ceramic Integration and Joining Technologies: From Macro to Nanoscale, 1st ed.; Singh, M., Ohji, T., Asthana, R., Mathur, S., Eds.; Wiley, 2011; Chapter 3, © 2010 The American Ceramic Society.
considered as a success for both PBC and HRP processes, which can be an alternative to current techniques. Several joining techniques are based on active brazing, which do not require any manner of CFC surface
616
Carbon as a Fusion Plasma-Facing Material
CFC
Chromium carbide
Brazing alloy
Interface Cu
200 mm
C/C Cu CuCrZr
CuCrZr
Cu
Brazing alloy
(a)
(b)
Figure 40 Optical micrograph of the cross-section of a Cr-carbide modified carbon fiber composite (CFC)–Cu joint, (b) Cr-carbide modified CFC–Cu–CuCrZr mock-up. Reproduced from Ferraris, M., et al. In Ceramic Integration and Joining Technologies: From Macro to Nanoscale, 1st ed.; Singh, M., Ohji, T., Asthana, R., Mathur, S., Eds.; Wiley, 2011; Chapter 3, © 2010 The American Ceramic Society.
Figure 41 ENEA-manufactured actively cooled monoblocks.
modification. In this case, the active brazing alloys for CFC include elements such as Ti, Zr, Cr, and Si, which allow CFC wettability by the molten brazing alloy. A drawback of active brazing can be that active elements may form brittle intermetallics or compounds of low melting point. In one study,132 a TiCuNi brazing alloy produced by Wesgo Metals in the form of sheets (70Ti–15Cu–15Ni) was used to join CFC and silicon-doped CFC to pure copper, but the presence of Ni and Ti brittle intermetallics at the joint interface had a detrimental effect on thermal fatigue resistance tests of the joined component. In Japan, a brazing technology was developed for ITER143 to join CFC to Cu by a NiCrP brazing alloy, followed by the joining of the CFC–Cu to the CuCrZr by low temperature HIP (500 C). Recently,
Carbon as a Fusion Plasma-Facing Material
a new brazing process was developed, based on the NiCuMn alloy, after metallization of the CFC surface. Several other brazing alloys have been developed for CFC–Cu joints: Ag-based (63Ag–35Cu–2Ti, 59Ag–27Cu–13In–1Ti), or Cu-based brazing alloys (Cu–3Si–2Al–2Ti; Cu–Mn; Cu–Ti). Ag was discarded in view of nuclear transmutation-related issues.144 4.18.6.2
Evaluation of HHF Joint
Reliable mechanical tests on CFCs joined to heat sinks are still an issue. As ASTM tests to measure the shear strength of CFCs joined to metals are not available, several laboratories have independently developed tests for CFC to metal joints, making interlab comparison of results almost impossible. Joints obtained by AMC® have been extensively tested,146 in particular for shear strength, with data ranging from 20 to 60 MPa for prepared samples. At 600 C, the shear strength dramatically decreases to 20 MPa. The shear strength and tensile strength of the improved AMC® (TiSi-AMC) joint are in the range of 54–73 MPa and of 39–64 MPa, respectively.130,146 Monoblocks obtained by AMC® have been measured after HHF tests147: apparent shear strength has been measured in the range of 30–60 MPa. Some cracks have been found at 45 , 90 , and 135 , considering 0 as the flux direction, leading to the detachment of CFC from the Cu layer before testing.148 The Cr-modified CFC–Cu joined samples148–152 have been measured by single-lap test (adapted from ASTM C 1292, C 1425) and off-set single lap test (adapted from ASTM D 905) at room temperature. Independent of the CFC surface machining and different casting process, results obtained133,138 for Cr-modified CFC on more than 50 samples yielded average values of apparent shear strength ranging from 26 to 32 MPa. The average shear strength is in any case higher than the interlaminar shear strength of the CFC (15 MPa). The shear strength of the CFC–Cu joints (flat-tile geometry) obtained by using a commercial Gemco® brazing alloy to braze CFC to pure copper was 34 4 MPa, measured by single lap tests at room temperature. This is comparable to the values obtained by other joining processes and higher than the intrinsic CFC shear strength.140 Mechanical tests on the monoblock braze require specific designs: some of them are adapted by ASTM D 4562-01 ‘‘Standard Test Method for Shear Strength of Adhesives Using Pin-and-Collar Specimen’’, as in the compression test used by Plansee AG, but the joint is not stressed in uniform pure shear state.
617
Reliable nondestructive tests (NDTs) are also extremely important for nuclear fusion components, especially for high heat flux PFCs. NDTs on CFC–Cu joints are complex because of the different response of CFC and copper to the physical excitations used to test the component. The aim of NDT is to identify and localize defects in the joined components before submitting them to high heat flux tests or actual application. It is also important to identify the maximum acceptable defect size, as a function of its position, defined as the largest defect that is stable under specific loads in the fusion device.139,149 Several techniques are used for NDT150: X-ray microradiography, X-ray microtomography, ultrasonic inspection,151,153 lock-in thermography,152 and transient infrared thermography (SATIR). SATIR (Figure 42) is the French acronym for infra red acquisition and data processing device: it is a dedicated facility developed in Cadarache-France at CEA. SATIR consists of recording the surface temperature evolution of the component with an infrared device during the circulation of hot (95 C) and cold (5 C) water through the cooling channel of the component. The transient thermal response is compared to a ‘defect-free’ component; defects such as debonding of CFC tiles from heat sinks are detected by a slower temperature surface response (Figure 42). Ultrasonic inspection has been applied to the flat tile and the monoblock design. Defects on joints between materials having very different acoustic impedance (e.g., copper and CFC) result in the generation of a high reflected echo, making defect detectability more difficult.141 Lock-in thermography consists of applying a series of heat flashes on the CFC. The main advantage of this technique is that there is no need for an active cooling of the component and it can also be used as an inspection method during the manufacturing process.142,152 Several nondestructive tests have been performed on the Cr-modified CFC–Cu joined samples150 not only to test performance, but also to verify and compare the reliability of these tests on a CFC–Cu interface. However, the present conclusion is that nondestructive tests of joints should be validated by destructive tests such as morphological evidence of the detected defect and mechanical testing.
4.18.7 Summary and Conclusions Carbon and graphite materials have enjoyed considerable success as PFMs in current tokamaks because
618
Carbon as a Fusion Plasma-Facing Material
FE200 (final screening) Reference block per analyzed zone
D 537.3 759.5 967.2
A
(E = 0.9) B1–B6
B7–B11
B12–B16
(a) VTFS: CFC (straight) part 14.00
8A
12.00 10.00
1A
15D
3A 5D 6D
8.00 DTref (⬚C)
SATIR-2 (final examination)
2A
14D
7A
6.00 4.00 2.00 0.00 −2.00
1
2
3
4
5
6
7
8
9
10
11
−6.00
13
14
15
16
Unit A Unit D
−4.00
(b)
12
Block number
Figure 42 Scanning electron microscopy of the cross-section of CFC–Cu–CuCrZr sample brazed by Gemco alloy in a single-step process as in Salvo et al.155 (a) Flat-tile configuration and (b) monoblock configuration.
of their low atomic number, high thermal shock resistance, and favorable properties. However, their use is not without significant issues, and their application in next-generation fusion energy devices is by no means certain. Significant among the issues related to carbon and graphite PFMs are neutron irradiation damage, which causes significant dimensional change and degrades the thermal conductivity resulting in increased PFC surface temperatures; physical sputtering, chemical erosion, and radiation enhanced sublimation, which cause surface material loss to the plasma and redeposition of carbon with tritium; and tritium inventory, which constitutes both a safety problem and an economic impediment to the use of graphite. Joining of CFC to heat sinks has witnessed significant development in the past few decades, which has resulted in good performing designs for near-term test machines such as ITER. The high heat loads and surface temperatures that result after plasma disruptions are also problematic for carbons. However, the same high temperatures make the use of Be, which has a significantly lower melting temperature, very unlikely. Next-generation machines will impose increasingly greater thermal loads on their PFCs. High thermal conductivity CFC materials may offer a
solution for the high heat loads, but further research is needed to overcome the problems noted above and to secure the place of carbon materials in the future of fusion power reactors.
References 1. Snead, L. L.; Katoh, Y.; Windes, W. E.; Shinavski, R. J.; Burchell, T. D. In Ceramic Composites for Near Term Reactor Application, The 4th International Topical Meeting on High Temperature Reactor Technology (HTR-2008), Washington, DC, Sept 28–Oct 1, 2008; ASME: Washington, DC, 2008. 2. Snead, L. L. In Carbon Materials for Advanced Technologies; Burchell, T. D., Ed.; Elsevier Science: Kidlington, Oxford, 1999; pp 389–427. 3. Akiba, M.; Madarame, H. J. Nucl. Mater. 1994, 212–215, 90–96. 4. Croessmann, C. D.; Gilbertson, N. B.; Watson, R. D.; Whitley, J. B. Fusion Technology 1989, 127–135. 5. Kinchin, G. H.; Pease, R. S. Rep. Prog. Phys. 1955, 18(1), 1–1. 6. Simmons, J. W. H. Radiation Damage in Graphite; Pergamon: Oxford, 1965; Vol. 102. 7. Kennedy, C. R. In Extended Abstracts for 14th Biennial Conference on Carbon, Irvine, CA, 1979. 8. Kelly, B. T. Physics of Graphite; Applied Science: London, 1981. 9. Engle, G. B. Carbon 1974, 12, 291–306. 10. Burchell, T. D.; Eatherly, W. P. J. Nucl. Mater. 1991, 179–181, 205–208.
Carbon as a Fusion Plasma-Facing Material 11. 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. 40. 41.
Bonal, J. P.; Wu, C. H. J. Nucl. Mater. 1996, 228, 155–161. Burchell, T. D. In Physical Processes of the Interaction of Fusion Plasmas with Solids, Plasma-Materials Interactions; Hofer, W. O., Roth, J., Eds.; Academic Press: New York, 1996; pp 341–382. Burchell, T. D.; Oku, T. In Physical Processes of the Interaction of Fusion Plasmas with Solids; Hofer, W. O., Roth, J., Eds.; Academic Press, Inc.: San Diego, 1994; Vol. 5(Suppl.), pp 77–128. Yoshikawa, H. H.; et al. Radiation Damage in Reactor Materials; IAEA: Vienna, 1963. Snead, L. L.; Burchell, T. D.; Katoh, Y. J. Nucl. Mater. 2008, 381, 55–61. Snead, L. L.; Burchell, T. D.; Qualls, A. L. J. Nucl. Mater. 2003, 321, 165–169. Bullock, R. E.; McKague, E. L. Carbon 1973, 1973, 547–553. Sato, S.; Kuramada, A.; Kawamata, K. Fusion Eng. Des. 1990, 13, 159–176. Tanabe, Y.; Yasuda, E.; Kimura, S.; Iseki, T.; Maruyama, T.; Yano, T. Carbon 1991, 29(7), 905–908. Eto, M.; Ishiyama, S.; Ugachi, H.; Fukaya, K.; Baba, S. J. Nucl. Mater. 1994, 212–215, 1223–1227. Hamada, K.; Sato, S.; Kohyama, A. J. Nucl. Mater. 1994, 212–215, 1228–1233. Burchell, T. D.; Eatherly, W. P.; Strizak, J. P. In The Effect of Neutron Irradiation on the Structure and Properties of Carbon–Carbon Composite Materials, Effects of Radiation on Materials: 16th International Symposium; Kumar, A. S., Gelles, D. S., Nanstad, R. K., Eds.; ASTM: Philadelphia, PA, 1994; ASTM STP 1175. Ahlf, L.; et al. J. Nucl. Mater. 1990, 171, 31. Binkele, L. High Temp. High Press. 1972, 4, 401–409. Price, R. J. Thermal Conductivity of Neutron-Irradiated Reactor Graphites; GA-A13157; General Atomics: San Diego, CA, Oct 8, 1974. Kelly, B. T.; Martin, W. H.; Price, A. M.; Dolby, P.; Smith, K. J. Nucl. Mater. 1966, 20, 195–209. Taylor, R.; Gilchrist, K. E.; Poston, L. J. J. Phys. Chem. Solids 1969, 30, 2251–2267. Thiele, B. A.; Binkele, L.; Koizlik, K.; Nickel, H. In Proceedings of the 16th International Symposium on Effects of Radiation on Materials; ASTM: Philadelphia, USA, 1992. Kelly, B. T.; Schofield, P.; Brown, R. G. Carbon 1990, 28, 155–158. Kelly, B. T.; Brocklehurst, J. E. Carbon 1977, 15, 113. Snead, L. L. J. Nucl. Mater. 2008, 381, 76–82. Kelly, B. T. Carbon 1967, 5, 247–260. Kelly, B. T. Carbon 1971, 9, 783–789. Burchell, T. D.; Eatherly, W. P.; Strizak, J. P. In Effects of Radiation on Materials: 16th International Symposium; Kumar, A. S., Gelles, D. S., Nanstad, R. K., Little, E. A., Eds.; American Society for Testing and Materials: Philadelphia, PA, 1993; ASTM STP 1175, pp 1266–1282. Heremans, J.; Beetz, C. P. Phys. Rev. B 1985, 32, 1981. DeCombarieu, A. J. Phys. (France) 1968, 28, 931. Taylor, R. Philos. Mag. 1966, 13, 157–166. Kelly, B. T. Plot Constructed from Personally Communicated Data. Wu, C. H.; Bonal, J. P.; Thiele, B. J. Nucl. Mater. 1994, 212–215, 1168–1173. Snead, L. L.; Burchell, T. D. J. Nucl. Mater. 1995, 224, 222–229. Snead, L. L.; Burchell, T. D. In Reduction in Thermal Conductivity due to Neutron Irradiation, 22nd, San Diego, CA, July 16–21, 1995; American Carbon Society: San Diego, CA, 1995; pp 774–775.
42. 43. 44. 45. 46. 47. 48. 49. 50. 51. 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. 77. 78. 79. 80. 81. 82. 83.
619
Maruyama, T.; Harayama, M. J. Nucl. Mater. 1992, 195, 44–50. Roth, J.; Bodhansky, J.; Wilson, K. L. J. Nucl. Mater. 1982, 111–112, 775–780. Fererici, G.; Anderl, R. A.; Andrew, P.; et al. J. Nucl. Mater. 1999, 266–269, 14. Roth, J.; Garcia-Rosalez, C. Nucl. Fusion 1997, 36, 1647. Mech, B. V.; Haasz, A. A.; Davis, J. W. J. Appl. Phys. 1998, 84, 1655. Roth, J. J. Nucl. Mater. 1999, 266–269, 51. Jacob, W.; Roth, J. Top. Appl. Phys. 2007, 110, 1. Poschenrieder, W. J. Vac. Sci. Technol. 1987, A5, 2265. Pospieszcyk, A. Atomic and Plasma Material Interaction Processes in Controlled Thermonuclear Fusion; Elsevier: Amsterdam, 1993. Lutterloh, C.; Schenk, A.; Biener, J.; Winter, B.; Kuppers, J. Surf. Sci. 1994, 316, L1039–L1043. Schenk, A. Appl. Phys. Lett. 1992, 61, 2414. Davis, J. W.; Haasz, A. A.; Stangeby, P. C. J. Nucl. Mater. 1988, 155–157, 237. Balden, M.; Roth, J. J. Nucl. Mater. 2000, 280, 39–44. Horn, A.; et al. Chem. Phys. Lett. 1994, 61, 193. Wittmann, M.; Kuppers, J. J. Nucl. Mater. 1996, 227, 186. Roth, J.; Vietzke, E.; Haasz, A. A. Erosion of graphite due to particle impact. In Supplement of the Journal of Nuclear Fusion; IAEA: Vienna, 1991. Haasz, A. A. J. Nucl. Mater. 1984, 128–129, 593. Phillips, V.; Flashkamp, K.; Vietzke, J. J. Nucl. Mater. 1984, 122–123, 237. Hirooka, Y.; et al. J. Nucl. Mater. 1990, 176–177, 473–480. Balden, M.; Starke, P.; Garcia-Rosales, C.; et al. J. Nucl. Mater., in press. Garcia-Rosales, C.; Balden, M. J. Nucl. Mater. 2001, 290–293, 173–179. Roth, J.; et al. J. Nucl. Mater. 1992, 191–194, 45–49. Jones, L. E.; Thrower, P. A. Carbon 1991, 29, 251. Hino, T.; Yamashina, T.; Fukida, S.; Takasugi, Y. J. Nucl. Mater. 1991, 186, 54. Allen, Y. K.; Chen, A. A.; Haasz, A.; Davis, J. W. J. Nucl. Mater. 1995, 227, 66. Roth, J.; Plank, H.; Schworer, R. Phys. Scripta 1996, T64, 67. Paz, P.; Garcia-Rosales, C.; Echeberria, J.; Balden, M.; Roth, J.; Behrisch, R. Fusion Eng. Des. 2001, 56–57, 325–330. Balden, M. Phys. Scripta 1999, T81, 64. Oya, A.; Marsh, H. J. Mater. Sci. 1982, 17, 309. Burtseva, T. A.; Chugunov, O. K.; Dovguchits, E. F.; et al. J. Nucl. Mater. 1992, 191–194, 309. Haasz, A. A.; Chiu, S.; Pierre, J. E.; Gudimenko, Y. I. J. Vac. Sci. Technol. 1996, A14, 184–193. Haasz, A. A.; Stephens, J. A.; Vietzke, E.; Eckstein, W.; Davis, J. W.; Hirooka, Y. Nucl. Fusion Suppl. Atomic Plasma Mater. Interac. Data Fusion 1998, 7, 1. Davis, J. W.; Haasz, A. A. J. Nucl. Mater. 1999, 266–269, 478–484. Winter, J. Plasma Phys. Control. Fusion 1994, 36, B263. Winter, J.; et al. J. Nucl. Mater. 1989, 163–164, 236. Jackson, G. L. J. Nucl. Mater. 1992, 196–198, 236. Thomas, P. R.; et al. J. Nucl. Mater. 1990, 176–177, 3. Winter, J. Phys. Rev. Lett. 1993, 71, 1549. Vietzki, E.; Haasz, A. A. In Physical Processes of the Interaction of Fusion Plasmas with Solids; Hofer, W. O., Roth, J., Eds.; Academic Press: New York, 1996; pp 135–176. Schmid, K.; Mayer, M.; Adelhelm, C.; Balden, M.; Lindig, S. Nucl. Fusion 2010, 50(105004), 1–11. Sigmunc, P.; et al. Nucl. Instrum. Methods 1989, B36, 110. Vietzke, E.; Flaskamp, K.; Hennes, M.; Philips, V. Nucl. Instrum. Methods 1984, B2(1–3), 617–622.
620 84. 85. 86. 87. 88. 89. 90. 91. 92. 93. 94. 95. 96. 97. 98. 99. 100. 101. 102. 103. 104. 105. 106. 107. 108. 109. 110. 111. 112. 113. 114. 115. 116. 117. 118. 119. 120.
Carbon as a Fusion Plasma-Facing Material Eckstein, W.; Philipps, V. Physical Sputtering and Radiation-Enhanced Sublimation; Academic Press: New York, 1996; pp 93–143. Andrew, P. L.; Pict, M. A. J. Nucl. Mater. 1994, 212–215, 111–117. Causey, R. A. J. Nucl. Mater. 1989, 162–164, 151. Atsumi, H.; Tokura, S.; Miyake, M. J. Nucl. Mater. 1988, 155–157, 241. Saeki, M. J. Appl. Radiat. Isot. 1983, 43, 739. Malka, V.; Rorhig, H. D.; Hecker, R. In Proceedings Conference on Tritium Technology in Fission, Fusion and Isotope Applications 1980, p 739. Rohrig, H. D.; Rischer, O. G.; Hecker, R. J. Am. Ceram. Soc. 1976, 59, 316. Morita, K.; Ohtsuka, K.; Hasebe, Y. J. Nucl. Mater. 1989, 162–164, 990. Tanabe, T.; Watanabe, Y. J. Nucl. Mater. 1990, 179–181, 231–234. Wilson, K. Trapping, Detrapping and Release of Implanted Hydrogen Isotopes; IAEA: Vienna, 1991. Hsu, W. L.; Causey, R. A. J. Vac. Sci. Technol. 1987, A5, 1768. Langley, R. A.; Brewer, R. S.; Roth, J. J. Nucl. Mater. 1978, 76–77, 313. Braun, M.; Emmoth, B. J. J. Nucl. Mater. 1984, 128–129, 657. Moller, W.; Scherzer, B. M. U. J. Appl. Phys. 1988, 64(10), 4860–4866. Doyle, B. L.; Wampler, W. R.; Bryce, D. K. J. Nucl. Mater. 1981, 103–104, 513. Wampler, W. R.; Brice, D. K.; Magee, C. W. J. Nucl. Mater. 1981, 102, 304. Wampler, W. R.; Magee, C. W. J. Nucl. Mater. 1981, 103–104, 509. Atsumi, H. J. Nucl. Mater. 2003, 313–316, 543–547. Scherzer, B. M. U.; et al. J. Nucl. Mater. 1976, 63, 100. Wilson, K. L.; Hsu, W. L. J. Nucl. Mater. 1987, 145–147, 121. Causey, R. A.; Baskes, I.; Wilson, K. L. J. Vac. Sci. Technol. 1986, A4, 151. Atsumi, H.; Iseki, M.; Shikama, T. J. Nucl. Mater. 1994, 212–215, 1478–1482. Kwast, H. J. Nucl. Mater. 1994, 212–215, 1472–1477. Wampler, W. R. J. Nucl. Mater. 1990, 176–177, 983–986. Causey, R. A. Phys. Scripta 1996, T64, 32–35. Causey, R. A. Fusion Technol. 1991, 19, 1585. Atsumi, H.; Muhaimin, A.; Tanabe, T.; Shikama, T. J. Nucl. Mater. 2009, 386–388, 379–382. Mitteau, R.; Team, T. S. J. Nucl. Mater. 2005, 337–339, 795–801. Boscary, J.; Bo¨swirth, B.; Greuner, H.; et al. Fusion Eng. Des. 2007, 82(15), 1634–1638. Ying, A.; Abdou, M.; Wong, C.; et al. Fusion Eng. Des. 2006, 81, 433–441. Guo, Q. G.; Li, J. G.; Noda, N.; et al. J. Nucl. Mater. 2003, 313–316, 144–148. Schlosser, J.; Escourbiac, F.; Merola, M.; et al. Nucl. Fusion 2005, 45, 512–518. Visca, E.; Libera, S.; Mancini, A.; Mazzone, G.; Pizzuto, A.; Testani, C. Fusion Eng. Des. 2007, 82(15–24), 1651–1656. Garcı´a-Rosales, C.; et al. J. Nucl. Mater. 2009, 386–388, 801–804. Centeno, A.; et al. Fusion Eng. Des., in press. Tivey, R.; Akiba, M.; Driemeyer, D.; Mazul, I.; Merola, M.; Ulrickson, M. Fusion Eng. Des. 2001, 55, 219–229. Merola, M.; Vieider, G. J. Nucl. Mater. 1998, 258–263, 672–676.
121. 122. 123. 124. 125. 126. 127. 128. 129. 130. 131. 132. 133. 134. 135. 136. 137. 138. 139. 140. 141. 142. 143. 144. 145. 146. 147. 148. 149. 150. 151. 152. 153.
Chevet, G.; et al. Phys. Scripta 2009, T138, 014057. Schlosser, J.; et al. Phys. Scripta 2007, T128, 204–208. Nicholas, M. G.; Mortimer, D. A. Mater. Sci. Technol. 1985, 11, 657–665. Akselsen, O. J. Mater. Sci. 1992, 27, 1989–2000. Arroyave, R.; Eagar, T. W. Acta Mater. 2003, 51, 4871–4880. Wu, C. H.; Bonal, J.; Kwast, H.; et al. Fusion Eng. Des. 1998, 39–40, 263–273. Merola, M.; Orsini, A.; Visca, E.; et al. J. Nucl. Mater. 2002, 307–311, 677–680. Eustathopoulos, N.; Nicholas, M.; Drevet, B. Wettability at High Temperature; Pergamon: Oxford, 1999. Dezellus, O.; Eustathopoulos, N. Scripta Mater. 1999, 40(11), 1283–1288. Schedler, B.; Huber, T.; Friedrich, T.; et al. Phys. Scripta 2007, T128, 200–203. Bisio, M.; Branca, V.; Di Marco, M.; et al. Fusion Eng. Des. 2005, 75–79, 277–283. Appendino, P.; Casalegno, V.; Ferraris, M.; Grattarola, M.; Merola, M.; Salvo, M. Fusion Eng. Des. 2003, 66–68, 225–229. Appendino, P.; Ferraris, M.; Casalegno, V.; Salvo, M.; Merola, M.; Grattarola, M. J. Nucl. Mater. 2004, 329–333, 1563–1566. Appendino, P.; Ferraris, M.; Casalegno, V.; Salvo, M.; Merola, M.; Grattarola, M. J. Nucl. Mater. 2006, 348, 102–107. Pedzich, Z.; Haberko, K.; Babiarz, J.; Faryna, M. J. Eur. Ceram. Soc. 1998, 5, 1939–1943. Rainer, F.; Reheis, N. Method for Setting-Up a Cooling System. European Patent EP0663670, 1994. Schlosser, J.; Chappuis, P.; Durocher, A.; Moncel, L.; Garin, P. Phys. Scripta 2001, T91, 94–97. Casalegno, V. Ph.D. Thesis, Politecnico di Torino: Torino, Italy, 2006. Merola, M.; Chappuis, P.; Escourbiac, F.; et al. Fusion Eng. Des. 2002, 61–62, 141–146. Salvo, M.; Casalegno, V.; Rizzo, S.; Smeacetto, F.; Ferraris, M.; Merola, M. J. Nucl. Mater. 2008, 374, 69–74. Missirilian, M.; et al. J. Nucl. Mater. 2007, 367–370(Pt 2), 1330–1336. Escourbiac, F.; Constans, S.; Courtois, X.; Durocher, A. J. Nucl. Mater. 2007, 367–370, 1492–1496. Suzuki, S.; et al. Phys. Scripta 2009, T138, 014003. Akiba, M.; Suzuki, S. Fusion Eng. Des. 1998, 39–40, 219–225. Zhou, Z. J.; Zhong, Z. H.; Ge, C. C. Fusion Eng. Des. 2007, 82, 35–40. Schedler, B.; Huber, T.; Eidenberger, E.; Scheu, C.; Pippan, R.; Clemens, H. Fusion Eng. Des. 2007, 82(15–24), 1786–1792. Missirlian, M.; Escourbiac, F.; Merola, M.; Durocher, A.; Bobin-Vastra, I.; Schedler, B. J. Nucl. Mater. 2007, 367–370, 1330–1336. Chevet, G.; et al. Fusion Eng. Des. 2009, 84, 586–589. Ezato, K.; Dairaku, M.; Taniguchi, M.; Sato, K.; Akiba, M. J. Nucl. Mater. 2002, 307–311, 144–148. Casalegno, V.; Salvo, M.; Ferraris, M.; Smeacetto, F.; Merola, M.; Bettuzzi, M. Fusion Eng. Des. 2008, 83, 702–712. Krautkra¨mer, J.; Krautkra¨mer, H. Ultrasonic Testing of Materials, 4th ed.; Springer-Verlag: New York, 1990; pp 27–55. Wu Datong, B. G. Rev. Gen. Therm. 1998, 37, 693–703. Kar, A.; Palit Sagar, S.; KumarRay, A. Mater. Lett. 2007, 61, 4169–4172.
4.19 Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices G. Federici Fusion for Energy, Garching, Germany
R. Doerner University of California at San Diego, San Diego, CA, USA
P. Lorenzetto Fusion for Energy, Barcelona, Spain
V. Barabash ITER Organization, St Paul Lez Durance, France
ß 2012 Fusion for Energy (F4E). Published by Elsevier Ltd. All rights reserved.
4.19.1 4.19.2 4.19.2.1 4.19.2.2 4.19.2.3 4.19.3 4.19.3.1 4.19.3.1.1 4.19.3.1.2 4.19.3.1.3 4.19.3.1.4 4.19.3.2 4.19.3.2.1 4.19.3.2.2 4.19.3.3 4.19.3.3.1 4.19.3.3.2 4.19.4 4.19.4.1 4.19.4.1.1 4.19.4.1.2 4.19.4.2 4.19.4.3 4.19.4.4 4.19.4.4.1 4.19.4.4.2 4.19.4.4.3 4.19.4.4.4 4.19.4.4.5 4.19.5 4.19.5.1 4.19.5.1.1 4.19.5.1.2 4.19.5.2 4.19.6 4.19.6.1 4.19.6.1.1 4.19.6.1.2 4.19.6.2
Introduction Background Synopsis of PWIs in Tokamaks Brief History of Plasma-Facing Materials in Fusion Devices Experience with Beryllium in Tokamaks Beryllium PWI Relevant Properties Beryllium Erosion Properties Physical sputtering of beryllium Mixed-material erosion Chemically assisted sputtering of beryllium Enhanced erosion at elevated temperatures Hydrogen Retention and Release Characteristics Implantation Beryllium codeposition Mixed-Material Effects Be–C phenomena Be–W alloying Main Physical and Mechanical Properties General Considerations Physical properties Mechanical properties Selection of Beryllium Grades for ITER Applications Considerations on Plasma-Sprayed Beryllium Neutron-Irradiation Effects Thermal conductivity Swelling Mechanical properties Thermal shock effects Bulk tritium retention Fabrication Issues Joining Technologies and High Heat Flux Durability of the Be/Cu Joints Be/Cu alloy joining technology High heat flux durability of unirradiated Be/Cu joints Thermal Tests on Neutron-Irradiated Joints Tokamak PFC Design Issues and Predictions of Effects in ITER During Operation PFC Design Considerations Design of the beryllium ITER-like wall at JET Design of the beryllium ITER wall Predictions of Effects on the ITER Beryllium Wall During Operation
623 624 624 626 628 629 629 629 631 632 632 633 633 635 637 637 637 638 638 639 639 640 642 643 643 643 644 644 644 644 644 645 646 648 650 650 650 651 653
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4.19.6.2.1 4.19.6.2.2 4.19.6.3 4.19.7 References
Safety issues in ITER Erosion/damage of the ITER Be wall Prospect of Using Beryllium in Beyond-ITER Fusion Reactors Concluding Remarks
Abbreviations Alcator C-Mod
ANFIBE ASDEXUpgrade
ATC CFC CIP CP DIII-D
DIMES
DS-Cu EAST
The name Alcator was given to a class of tokamaks designed and built at the Massachusetts Institute of Technology; these machines are distinguished by high magnetic fields with relatively small diameters. The high magnetic field helps create plasmas with relatively high current and particle densities. The present incarnation is Alcator C-Mod Computer code for ANalysis of Fusion Irradiated BEryllium Axially Symmetric Divertor Experiment. The original ASDEX, located in Garching, Germany, and decommissioned in about 1990, would qualify today as a medium sized tokamak. It was designed for the study of impurities and their control by a magnetic divertor. Its successor, ASDEX-Upgrade (a completely new machine, not really an ‘upgrade’), is larger and more flexible. Adiabatic Toroidal Compressor Carbon-fiber composite Cold isostatic pressing Cold pressing A medium-sized tokamak, but the largest tokamak still operational in the United States. Operated by General Atomics in San Diego Divertor Material Evaluation Studies, a retractable probe that allows the insertion and retraction of test material samples to the DIIID divertor floor, for example, for erosion/deposition studies. Dispersion-strengthened copper Experimental advanced superconducting tokamak – an experimental superconducting
ELMs FISPACT FZJ HIP INEEL
ISX
ITER
JET
653 654 659 659 662
tokamak magnetic fusion energy reactor in Hefei, the capital city of Anhui Province, in eastern China Edge localized modes Inventory code included in the European Activation System Forschungszentrum Juelich, Germany Hot isostatic pressing Idaho National Engineering and Environmental Laboratory. Now Idaho National Laboratory (INL) Impurity study experiment (ISX-A and ISX-B where two tokamaks operated at Oak Ridge National Laboratory) ITER, the world’s largest tokamak experimental facility being constructed in the South of France to demonstrate the scientific and technical feasibility of fusion power. The project is being built on the basis of an international collaboration between the European Union, China, India, Japan, Russia, South Korea, and the United States. The international treaty was signed in November 2006 and the central organization established in Cadarache. Most of the components will be provided in kind by agencies set up for this purpose in the seven partners Joint European Torus – a large tokamak located at the Culham Laboratory in Oxfordshire, England, jointly owned by the European Community. First device to achieve >1 W of fusion power, in 1991, and the machine that has most closely approached Q ¼ 1 for DT operation (Q ¼ 0:95 in 1997)
Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices
JUDITH
Juelich Divertor Test Facility in Hot Cells KSTAR Korea Superconducting Tokamak Advanced Reactor – a long-pulse, superconducting tokamak built in South Korea to explore advanced tokamak regimes under steady state conditions LANL Los Alamos National Laboratory LCFS Last closed flux surface MAR ITER Materials Assessment Report MIT Massachusetts Institute of Technology MPH ITER Materials Properties Handbook NBI Neutral beam injection NRA Nuclear reaction analysis NRI Nuclear Research Institute in the Czech Republic PDX Poloidal divertor experiment PFCs Plasma-facing components PISCES Plasma Interaction with Surface Components Experimental Station. It is a plasma simulator located at the University of California San Diego in the United States (originally at University of California, Los Angeles) that is used to test materials and measure sputtering, retention, etc. expected in tokamaks PLT Princeton Large Torus PWIs Plasma–wall interactions RES Radiation enhanced sublimation RMP Resonance magnetic perturbation SNL Sandia National Laboratory SOL Scrape-off-layer ST Symmetric tokamak (in this chapter) STEMET 1108 Brazing alloy: Cu–Sn–In–Ni TFR Torus Fontenay-aux-Roses TPE Tritium plasma experiment TRIM Transport of ion in matter code UCSD University of California, San Diego UNITOR One of the first small tokamaks where beryllium was used UTIAS University of Toronto Institute for Aerospace Studies VDE Vertical displacement event VHP Vacuum hot pressing
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4.19.1 Introduction Beryllium, once called ‘the wonder metal of the future,’ 1 is a low-density metal that gained early prominence as a neutron reflector in weapons and fission research reactors. It then found a wide range of applications in the automotive, aerospace, defense, medical, and electronic industries. Also, because of its unique physical properties, and especially favorable plasma compatibility, it was considered and used in the past for protection of internal components in various magnetic fusion devices (e.g., UNITOR, ISX-B, JET). Most important future (near-term) applications in this field include (1) the installation of a completely new beryllium wall in the JET tokamak, which has been completed by mid of 2011 and consists of 1700 solid Be tiles machined from 4 t of beryllium; and (2) ITER, the world’s largest experimental facility to demonstrate the scientific and technical feasibility of fusion power, which is being built in Cadarache in the South of France. ITER will use beryllium to clad the first wall (700 m2 for a total weight of about 12 t of Be). Although beryllium has been considered for other applications in fusion (e.g., as neutron multiplier in the design of some types of thermonuclear breeding blankets of future fusion reactors and for hohlraums in inertial confinement fusion), this chapter will be limited to discussing the use of beryllium as a plasma-facing material in magnetic confinement devices, and in particular in the design, research, and development work currently underway for the JET and the ITER tokamaks. Considerations related to health and safety procedures for the use of beryllium relevant for construction and operation in tokamaks are not discussed here. Designing the interface between a thermonuclear plasma and the surrounding solid material environment has been arguably one of the greatest technical challenges of ITER and will continue to be a challenge for the development of future fusion power reactors. The interaction between the edge plasma and the surrounding surfaces profoundly influences conditions in the core plasma and can damage the surrounding material structures and lead to long machine downtimes for repair. Robust solutions for issues of plasma power handling and plasma–wall interactions (PWIs) are required for the realization of a commercially attractive fusion reactor. A mix of several plasma-facing materials is currently proposed in ITER to optimize the requirements of areas with different power and particle flux characteristics
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Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices
(i.e., Be for the first wall, carbon-fiber composite (CFC) for the divertor strike point tiles, and W elsewhere in the divertor). Inevitably, this is expected to lead to cross-material contamination and the formation of material mixtures, whose behavior is still uncertain and requires further investigation. The use of beryllium for plasma-facingcomponent (PFC) applications has been the subject of many reviews during the last two decades (see, e.g., Wilson et al.2 and Raffray et al.3 and references therein). Much of this fusion-related work has been summarized in a series of topical workshops on beryllium that were held in the past, bringing together leading researchers in the field of beryllium technology and disseminating information on recent progress in the field.4 Comprehensive reviews have also appeared recently in specialized journals5,6 containing state-of-the-art information on a number of topics such as manufacturing and development of coating techniques, component design, erosion/deposition, tritium retention, material mixing and compatibility problems, safety of beryllium handling, etc. This chapter reviews the properties of beryllium that are of primary relevance for plasma protection applications in near-term magnetic fusion devices (i.e., PWIs, thermal and mechanical properties, fabricability and ease of joining, chemical reactivity, etc.) together with the available knowledge on performance and operation in existing fusion machines. Special attention is given to beryllium’s erosion and deposition, the formation of mixed materials, and the hydrogen retention and release characteristics that play an important role in plasma performance, component lifetime, and operational safety. The status of the available techniques presently considered for joining the beryllium armor to the heat sink material of Cu alloys for the fabrication of beryllium-clad actively cooled components for the ITER first wall is briefly discussed together with the results of the performance and durability heat flux tests conducted in the framework of the ITER first-wall qualification programme. The effects of neutron irradiation on the degradation of the properties of beryllium itself and of the joints are also briefly analyzed. This chapter is organized as follows. Section 4.19.2 provides some background information for the reader and briefly reviews (1) the problem of PWIs in tokamaks; (2) the history of plasma-facing materials in fusion devices and the rationale for choosing beryllium as the material for the first wall of JET and ITER; and (3) the experience with the use of beryllium in tokamaks to date. Section 4.19.3
describes in detail the beryllium PWI-relevant properties such as erosion/deposition, hydrogen retention and release, and chemical effects such as material mixing, all of which influence the selection of beryllium as armor material for PFCs. Section 4.19.4 briefly reviews a limited number of selected physical and mechanical properties of relevance for the fabrication of heat exhaust components and the effects of neutron irradiation on material properties. Section 4.19.5 describes the fabrication issues and the progress of joining technology and high heat flux durability of beryllium-clad PFCs. Section 4.19.6 describes the main issues associated with the JET and ITER first-wall designs and discusses some constraints foreseen during operation. The prospects of beryllium for applications in fusion reactors beyond ITER are briefly discussed. Finally, a summary is provided in Section 4.19.7.
4.19.2 Background 4.19.2.1
Synopsis of PWIs in Tokamaks
A detailed discussion on this subject is beyond the scope of this review. The relevant PWIs are comprehensively reviewed by Federici et al.7,8 More recent interpretations of the underlying phenomena and impact on the ITER device can be found in Roth et al.9 Here we summarize some of the main points. PWIs critically affect tokamak operation in many ways. Erosion by the plasma determines the lifetime of PFCs, and creates a source of impurities, which cool and dilute the plasma. Deposition of material onto PFCs alters their surface composition and, depending on the material used, can lead to long-term accumulation of large in-vessel tritium inventories. This latter phenomenon is especially exacerbated for carbonbased materials but there are still some concerns with beryllium. Retention and recycling of hydrogen from PFCs affects fuelling efficiency, plasma density control, and the density of neutral hydrogen in the plasma boundary, which impacts particle and energy transport. The primary driver for the interactions between the core plasma, edge plasma, and wall is the power generated in the plasma core that must be handled by the surrounding structures. Fusion power is obtained by the reaction of two hydrogen isotopes, deuterium (D) and tritium (T), producing an a-particle and a fast neutron. Although the kinetic energy carried by the 14.1 MeV neutron escapes the plasma and could be converted in future reactors beyond ITER to thermal energy in a surrounding blanket system, the
Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices
kinetic energy of the a-particle is deposited in the plasma. The fraction of this power that is not radiated from the plasma core as bremsstrahlung or line radiation (and that on average is distributed uniformly on the surrounding structures) is transported across field lines to the edge plasma and intersects the material surfaces in specific areas leading to intense power loads. The edge plasma has a strong influence on the core plasma transport processes and thereby on the energy confinement time. A schematic representation of the regions of the plasma and boundary walls in a divertor tokamak is portrayed in Figure 1 taken from Federici et al.7 The outermost closed magnetic field surface forms an X-point of zero poloidal magnetic field within the vessel. This boundary is called the ‘last closed flux
Magnetic flux surfaces Separatrix (LCFS) Edge region Scrape-off layer
First wall
Plasma core
Separatrix (LCFS) X-point
Divertor region
Baffle
Vertical divertor target plate
Private flux region Separatrix strike point Pump
Figure 1 Poloidal cross-section of a tokamak plasma with a single magnetic null divertor configuration, illustrating the regions of the plasma and the boundary walls where important PWIs and atomic physics processes take place. The characteristic regions are (1) the plasma core, (2) the edge region just inside the separatrix, (3) the scrape-off-layer (SOL) plasma outside the separatrix, and (4) the divertor plasma region, which is an extension of the SOL plasma along field lines into the divertor chamber. The baffle structure is designed to prevent neutrals from leaving the divertor. In the private flux region below the X-point, the magnetic field surfaces are isolated from the rest of the plasma. Reproduced with permission from Federici, G.; Skinner, C. H.; Brooks, J. N.; et al. Plasma-material interactions in current tokamaks and their implications for next-step fusion reactors. Nucl. Fusion 2001, 41, 1967–2137 (review special issue), with permission from IAEA.
625
surface’ (LCFS) or ‘separatrix.’ Magnetic field surfaces inside the LCFS are closed, confining the plasma ions. The edge region, just inside the LCFS, contains significant levels of impurities not fully ionized, and perhaps neutral particles. Impurity line radiation and neutral particles transport some power from here to the wall. The remaining power enters the region outside the LCFS either by conduction or, depending on the degree to which neutrals penetrate the plasma, by convection. This region is known as the scrape-offlayer (SOL) as here power is rapidly ‘scraped off ’ by electron heat conduction along open field lines, which are diverted to intersect with target regions that are known as ‘divertors.’ Poloidal divertors have been very successful at localizing the interactions of plasma ions with the target plate material in a part of the machine geometrically distant from the main plasma where any impurities released are well screened from the main plasma and return to the target plate. The plasma density and temperature determine the flux density and energy of plasma ions striking the plasma-wetted surfaces. These, in turn, determine the rate of physical sputtering, chemical sputtering, ion implantation, and impurity generation. The interaction of the edge plasma with the surrounding solid material surfaces is most intense in the vicinity of the ‘strike point’ where the separatrix intersects the divertor target plate (see inset in Figure 1). In addition, the plasma conditions determine where eroded material is redeposited, and to what degree codeposition of tritium occurs at the wall. The plasma power flow also determines the level of active structural cooling required. Typical plasma loads and the effects expected during normal operation and off-normal operation in ITER are summarized in Table 1. Because of the very demanding power handling requirements (predicted peak value of the heat flux in the divertor near the strike-points is >10 MW m2) and the predicted short lifetime due to sputtering erosion arising from very intense particle fluxes (1023–1024 particles m2 s1) and damage during transient events, beryllium has been excluded from use in the ITER divertor and is instead the material selected for the main chamber wall of ITER. Recent observations in present divertor tokamaks have shown that plasma fluxes to the main wall are dominated by intermittent events leading to fast plasma particle transport that reaches the PFCs along the magnetic field (see Loarte et al.10 and references therein). The quasistationary heat fluxes to the main wall are thought to be dominated by convective
626 Table 1
Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices Major issues associated with operation of ITER PFCs
PFCs
Plasma loads
Candidate armor
Effects
Issue
Divertor – strike-point regions
Radiation and particle heat Large particle fluxes Disruptions ELM’s Slow-high power
CFCa
Chemical erosion evaporation brittle destruction and tritium codeposition
Erosion lifetime and
Divertor – baffle region Dome
Radiation heat Disruptions Radiation heat (MARFE’sb) 100 eV ions and CX
W
High sputtering evaporation/ melting High sputtering evaporation/ melting
Plasma contamination
Evaporation/melting
T retention in beryllium
transients
W
neutrals
First wall
Start-up limiters
Moderate power transients Plasma contact during VDEsc Disruptions and runaway electrons ELMs High start-up heat loads Plasma contact during VDEs Disruptions
Be
component replacement
High tritium inventory and safety
Erosion lifetime
codeposited layers
Chemical reactivity
especially with Be dust
Be
High sputtering evaporation/ melting
Erosion lifetime
a
W is also considered as an alternative. Multifaceted asymmetric radiation from the edge (MARFE). Vertical displacement event (VDE).
b c
transport,11 but still remain to be clarified. Although the steady-state parallel power fluxes associated with these particle fluxes will only be of the order of several MW m2 in the ITER QDT ¼ 10 reference scenario, local overheating of exposed edges of main wall PFCs can occur because of limitations in the achievable alignment tolerances. Similarly, transient events are expected to cause significant power fluxes to reach first-wall panels in ITER along the field lines. Edge localized modes (ELMs) deposit large amounts of energy in a short time, and in some cases in a toroidally localized fashion, which can lead to strong excursions in PFC surface temperatures. While the majority of ELM energy is deposited on divertor surfaces, a significant fraction is carried to surfaces outside the divertor. There are obvious concerns that ELMs will lead to damage of the divertor and the first wall.12 An additional concern is that even without erosion, thermal shock can lead to degradation of material thermomechanical properties, for example, loss of ductility leading to an enhanced probability of mechanical failure or spalling (erosion). Research efforts to characterize the ELMs in the SOL are described elsewhere.13–15 There are still large uncertainties in predicting the thermal loads of ELMs on the ITER beryllium first wall and the range of parallel
energy fluxes varies from 1.0 MJ m2 (controlled ELMs) to 20 MJ m2 (uncontrolled ELMs).16,17 Even for controlled ELMs, such energy fluxes are likely to cause melting of up to several tens of micrometers of beryllium at the exposed edges,18 which could cause undesirable impurity influxes at every ELM.10,11 4.19.2.2 Brief History of Plasma-Facing Materials in Fusion Devices PWIs have been recognized to be a key issue in the realization of practical fusion power since the beginning of magnetic fusion research. By the time of the first tokamaks in the 1960s in the USSR and subsequently elsewhere, means of reducing the level of carbon and oxygen were being employed.19,20 These included the use of stainless steel vacuum vessels and all-metal seals, vessel baking, and discharge cleaning. Ultimately, these improvements, along with improved plasma confinement, led to the first production of relatively hot and dense plasmas in the T3 tokamak (1 keVand 3 1019 m3).21,22 These plasmas, while being cleaner and with low-Z elements fully stripped in the core, still had unacceptable levels of carbon, oxygen, and metallic impurities. The metallic contamination inevitably consisted of wall and limiter materials.
Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices
Early in magnetic fusion research, it was recognized that localizing intense PWIs at some type of ‘sacrificial’ structure was desirable, if only to ensure that more fragile vacuum walls were not penetrated. This led to the birth of the ‘limiter,’ usually made to be very robust, from refractory material and positioned to ensure at least several centimeters gap between the plasma edge and more delicate structures like bellows, electrical breaks, vacuum walls, etc. Typical materials used for limiters in these early days included stainless steel in Adiabatic Toroidal Compressor (ATC)23 and ISX-A24 and many others, molybdenum in Alcator A25 and Torus Fontenay-aux-Roses (TFR),26 tungsten in symmetric tokamak (ST)27 and Princeton Large Torus (PLT),28 and titanium in poloidal divertor experiment (PDX).29 Poloidal divertors have been very successful at localizing the interactions of plasma ions with the target plate material in a part of the machine geometrically distant from the main plasma where any impurities released are well screened from the main plasma and return to the target plate.30 By the early 1980s, it was also recognized that in addition to these functions, the divertor should make it easier to reduce the plasma temperature immediately adjacent to the ‘limiting’ surface, thus reducing the energies of incident ions and the physical sputtering rate. Complementary to this, high divertor plasma and neutral densities were found. The high plasma density has several beneficial effects in dispersing the incident power, while the high neutral density makes for efficient pumping. Pumping helps with plasma density control, divertor retention of impurities and, ultimately, in a reactor, helium exhaust. By the late 1970s, various tokamaks were starting to employ auxiliary heating systems, primarily neutral beam injection (NBI). Experiments with NBI on PLT resulted in the first thermonuclear class temperatures to be achieved.28,31,32 PLT at the time used tungsten limiters, and at high powers and relatively low plasma densities, very high edge plasma temperatures and power fluxes were achieved. This resulted in tungsten sputtering and subsequent core radiation from partially stripped tungsten ions. For this reason, PLT switched limiter material to nuclear grade graphite. Graphite has the advantage that eroded carbon atoms are fully stripped in the plasma core, thus reducing core radiation. In addition, the surface does not melt if overheated – it simply sublimes. This move to carbon by PLT turned out to be very successful, alleviating the central radiation problem. For these reasons, carbon has tended to be the favored limiter/divertor material in magnetic fusion research ever since.
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By the mid-1980s, many tokamaks were operating with graphite limiters and/or divertor plates. In addition, extensive laboratory tests and simulations on graphite had begun, primarily aimed at understanding the chemical reactivity of graphite with hydrogenic plasmas, that is, chemical erosion. Early laboratory results suggested that carbon would be eroded by hydrogenic ions with a chemical erosion yield of Y 0.1 C/Dþ, a yield several times higher than the maximum physical sputtering yield. Another process, radiation-enhanced sublimation (RES), was discovered at elevated temperatures, which further suggested high erosion rates for carbon. Carbon’s ability to trap hydrogenic species in codeposited layers was recognized. These problems, along with graphite’s poor mechanical properties in a neutron environment (which had previously been known for many years from fission research33), led to the consideration of beryllium as a plasma-facing material. This was primarily promoted at JET.34 A description of the operation experience to date with Be in tokamak devices is provided in Section 4.19.2.3. At present, among divertor tokamaks, carbon is the dominant material only in DIII-D. Alcator C-Mod at Massachusetts Institute of Technology (MIT), USA35 uses molybdenum. ASDEX-Upgrade (Axially Symmetric Divertor Experiment) is fully clad with tungsten,36 and JET has completed in 2011 a large enhancement programme37 that includes the installation of a beryllium wall and a tungsten divertor. New superconducting tokamaks, such as Korea Superconducting Tokamak Advanced Reactor (KSTAR) in Korea38 and experimental advanced superconducting tokamak (EAST) in China,39 employ carbon as material for the in-vessel components, but with provisions to exchange the material later on in operation. The current selection of plasma-facing materials in ITER has been made by compromising among a series of physics and operational requirements, (1) minimum effect of impurity contamination on plasma performance and operation, (2) maximum operational flexibility at the start of operation, and (3) minimum fuel retention for operation in the DT phase. This compromise is met by a choice of three plasma-facing materials at the beginning of operations (Be, C, and W). It is planned to reduce the choices to two (Be and W) before DT operations in order to avoid long-term tritium retention in carbon codeposits during the burning plasma phase. Beryllium has been chosen for the first-wall PFCs to minimize fuel dilution caused by impurities released from these surfaces, which are expected to
628
Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices
have the largest contamination efficiency.40–44 Moreover, the consequences of beryllium contamination on fusion performance and plasma operations are relatively mild. This has been demonstrated by experiments in tokamaks (see Section 4.19.2.3). The main issues related to the use of beryllium in ITER are (1) the possible damage (melting) during transients such as ELMs, disruptions, and runaway electron impact, and its implications for operations and (2) the codeposition of tritium with beryllium which is eroded from the first wall and deposited at the divertor targets (and possibly also locally redeposited into shadowed areas of the shaped ITER first wall). Both issues are part of ongoing research, the initial results of which are being taken into account in the ITER design so that the influence of these two factors on ITER operation and mission is minimized. This includes ELM control systems based on pellets and resonance magnetic perturbation (RMP) coils, disruption mitigation systems, and increased temperature baking of the divertor to release tritium from the beryllium codeposited layers. Carbon is selected for the high power flux area of the divertor strike points because of its compatibility with operation over a large range of plasma conditions and the absence of melting under transient loads. Both of these characteristics are considered to be essential during the initial phase of ITER exploitation in which plasma operational scenarios will require development and transient load control and mitigation systems will need to be demonstrated. 4.19.2.3 Experience with Beryllium in Tokamaks Only three tokamaks have operated with beryllium as the limiter or first-wall material. The first experiments were performed by UNITOR,45 which were then followed by ISX-B.46 Both tokamaks investigated the effects of small beryllium limiters on plasma behavior (UNITOR had side limiters at two toroidal locations and ISX-B had one top limiter) in support of the more ambitious beryllium experiment in JET (see below). The motivation to use beryllium came from the problem of controlling the plasma density and impurities when graphite was used. Both UNITOR and ISX-B showed that once beryllium is evaporated from the limiter and coated over a large segment of the first wall, oxygen gettering leads to significant reduction of impurities. When the heat load on the beryllium limiter was increased to the point of evaporating beryllium, the oxygen
concentration was decreased dramatically. Although the concentration of beryllium in the plasma was increased, its contribution to Zeff (the ion effective charge of the plasma Zeff provides a measure for impurity concentration) was more than compensated by the reduction of oxygen, carbon, and metal impurities.45 The plasma Zeff was observed to be reduced from 2.4 to near unity with beryllium. It must be noted that there was a negative aspect associated with beryllium operation during the ISX-B campaign. The reduction in plasma impurities was not observed until the limiter surface was partially melted causing beryllium to be evaporated and coated on the first wall. Once melting did occur, the droplets made subsequent evaporation more likely but hard to control. The consequent strong reduction in plasma impurities associated with gettering then made discharge reproducibility hard to obtain. However, if a much larger plasma contact area is already covered with Be, one does not need to rely on limiter melting to obtain the beneficial effect of beryllium. This effect could be achieved by using large area beryllium limiter, or coating the inside wall with beryllium which was the approach taken by JET when it introduced beryllium in 1989. Large tokamak devices such as JET had found it very difficult to control the plasma density with graphite walls as the discharge pulse length got longer. Motivated by the frequent occurrence of a phenomenon that plagued the earlier campaigns – the so-called carbon blooms due to the overheating of poorly designed divertor tiles and the subsequent influx of carbon impurities in the plasma due to evaporation – JET decided to use beryllium as a plasma-facing material. Thin evaporated beryllium layers on the graphite walls were used initially (100 A˚ average thickness per deposition) on the plasma-facing surface of the device. Subsequently, beryllium tiles were installed on the toroidal belt limiter. The main experimental results with beryllium can be summarized as follows: 1. The concentration of carbon and oxygen in the plasma were 4–7% and 0.5–2%, respectively, when graphite was used as belt limiter. With a beryllium belt limiter, the carbon content was reduced to 0.5% and oxygen became negligible, because of oxygen gettering by beryllium. During ohmically heated discharges, the concentration of beryllium remained negligible even though beryllium was the dominant impurity.
Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices
2. While the value of Zeff was 3 using the graphite limiter and auxiliary heating power of 10 MW, Zeff was 1.5 even with additional heating powers of up to 30 MW with a beryllium limiter. 3. The fuel density control was greatly improved with the beryllium limiter and beryllium evaporated wall. Gas puffing to maintain a given plasma density was typically 10 times larger when using beryllium than graphite. Following the beryllium limiter experience, divertor beryllium targets were installed in JET for two configurations. An extensive set of experiments with toroidally continuous X-point divertor plates was carried out in JET in the period 1990–1996 to characterize beryllium from the point of view of its thermomechanical performance, as well as its compatibility with various plasma operation regimes.47–50 In the JET Mk I experiments, melting of the beryllium tiles was reached by increasing (in a progressive way) the power flux to a restricted area of the divertor target in fuelled, medium density ELMy H-mode discharges (Pinp 12 MW). Large beryllium influxes were observed when the divertor target temperature reached 1300 C. In these conditions, it became difficult to run low-density ELMy H-mode discharges (Pinp 12 MW) without fast strike point movement (to achieve lower average power load) and the discharges either had very poor performance or were disrupted. However, no substantial plasma performance degradation was observed for medium density H-modes with fixed strike point position, or if fast strike point movement was applied in lowdensity H-modes, despite the large scale distortion of the target surface caused by the melt layer displacement and splashing due to the previous 25 high power discharges48,51 (see Figure 252). This demonstrated that it was possible to use the damaged Be divertor target as the main power handling PFC if the
Figure 2 Melting of the Joint European Torus Mk I beryllium target plate tiles after plasma operation. Image courtesy of EFDA-JET.
629
average power load was decreased, either by increasing plasma density and radiative losses, or by strike point sweeping. The damage did not prohibit subsequent plasma operation in JET, but would seriously limit the lifetime of Be PFCs in long-pulse ITER-like devices. The latest results of the operation of JET with beryllium have been reviewed recently by Loarte et al.10
4.19.3 Beryllium PWI Relevant Properties This section describes the present understanding of PWIs for beryllium-containing surfaces. First, it focuses on the erosion properties of ‘clean’ beryllium surfaces at different temperatures. Retention of plasma fuel species in both bulk and codeposited layers of beryllium is then described. As beryllium will not be used as the exclusive plasma-facing material in future confinement devices, issues associated with mixed, beryllium-containing surfaces are also addressed. 4.19.3.1
Beryllium Erosion Properties
The term erosion is used to describe a group of processes that remove material from a surface subjected to energetic particle bombardment. Included under the general classification of erosion are processes such as physical sputtering, chemically assisted physical sputtering, chemical sputtering, and thermally activated release from surfaces. Of these processes, only chemical sputtering, where volatile molecular species are formed on the surface, appears to be inactive in beryllium. 4.19.3.1.1 Physical sputtering of beryllium
Physical sputtering results from the elastic transfer of energy from incoming projectiles to atoms on the surface of the target material. Target atoms can be sputtered when the energy they receive from the collisional cascade of the projectile exceeds the binding energy of the atom to the surface. The physical sputtering rate is usually referred to as the sputtering yield, Y, which is defined as the ratio of the number of atoms lost from a surface to the number of incident energetic particles striking the surface. The lower the binding energy of surface atoms, the larger the physical sputtering yield. As physical sputtering can be approximated using a series of binary collisions within the surface, it is relatively easy to estimate
Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices
the physical sputtering yield of given projectile-target scenarios. Monte-Carlo based simulation codes (such as transport of ions in matter (TRIM))53 have been used to generate extensive databases of sputtering yields based on incident particle angle, energy, and mass, for a variety of targets54 including beryllium. Measurement of the physical sputtering yield from a beryllium surface is complicated by the natural affinity of beryllium for oxygen. A beryllium surface will quickly form a thin, stable, passivating oxide surface layer when exposed to atmosphere. In ion beam devices, it is possible to clean any oxides from the beryllium surface before a measurement and with careful control of the residual gas pressure, make the measurements before the oxide reforms on the surface and alters the measurement.55 It has also been shown that it is possible to deplete the beryllium surface of oxide by heating the sample to temperatures exceeding 500 C, where the beryllium below the oxide can diffuse through the oxide to the surface,56 thereby allowing measurements on a clean beryllium surface. The comparison between the calculated sputtering yield and measurements made using mass-selected, monoenergetic ion-beams devices impinging on clean beryllium surfaces is fairly good.57 Measurements of sputtering yields in plasma devices, however, are complicated by several factors. In plasma devices, the incident ions usually have a temperature distribution and may contain different charge state ions. Each different charge state ion will be accelerated to a different energy by the electrostatic sheath in the vicinity of the surface. When hydrogenic plasma interacts with a surface, one must also account for a distribution of molecular ions striking the surface. In the case of a deuterium plasma, for example, the distribution of molecular þ ions (Dþ, Dþ 2 , D3 ) must be taken into account as the incident molecule disassociates on impact with the surface and a Dþ 2 ion becomes equivalent to the bombardment of two deuterium particles with one-half the incident energy of the original Dþ 2 ion. Figure 3 shows the change to the calculated sputtering yield when one includes the effects of molecular ions in a plasma, compared to the calculated sputtering yield from pure Dþ ion bombardment. The trajectory of the incoming ions can also be altered by the presence of electrostatic and magnetic sheaths. Plasmas also contain varying amounts of impurity ions, originating either from PWIs in other locations of the device, or ionization of residual background gas present in the device and these impurity ions, or simply neutral gas atoms, may interact with
the surface. Finally, the incident flux from the plasma is usually so large that the surface being investigated, and its morphology, becomes altered by the incident flux and a new surface exhibiting unique characteristics may result. Nevertheless, the physical sputtering yield from beryllium surfaces exposed to plasma ion bombardment has been measured in several devices. Unfortunately, there is little consensus on the correct value of the physical sputtering yield. In JET, the largest confinement device to ever employ beryllium as a PFC sputtering yield measurements range from values far exceeding47 to values less58 than one would expect from the predictions of TRIM. In the Plasma Interaction with Surface Components Experimental Station B (PISCES-B) device, systematic experiments to measure the physical sputtering yield routinely show values less59–61 than those expected from TRIM. This difference is shown in Figure 3, where the energy dependence of the calculated yield is compared to experimental measurements. Another primary difference between the conditions in an ion beam device and those encountered in a plasma device has to do with the neutral density near the surface being investigated. In an ion beam experiment, the background pressure is kept very low 0.1
Physical sputtering yield
630
0.01
0.001
0.0001
10-5 0
20
40
60 80 100 120 140 160 Ion energy (eV)
Measured yield in D plasma Calculated yield (D+ ions only) Calculated yield (D+, D+2 , D+3 ions) Figure 3 Calculated sputtering yields from pure Dþ bombardment at normal incidence compared to that þ calculated for a (0.25, 0.47, 0.28) mix of Dþ, Dþ 2 , and D3 ; also shown is the measured yield from such a plasma.
Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices
so that the surface being probed maintains its clean properties. On the other hand, the incident flux in a plasma device is usually several orders of magnitude larger than in an ion beam device, ensuring that the surface remains clean because of the large incident flux. However, this plasma-facing surface undergoes not only energetic ion bombardment, but also bombardment by neutral atoms and molecules. The neutral density in plasma generators is typically on the order of 1020 m3 (a few millitorr) which is necessary for breakdown of the plasma. The estimated neutral atom flux is approximately equal to the incident ion flux to the surface61 and it is often not possible to alter significantly this flux ratio. In the case of a beryllium surface which can form a hydride (see Section 4.19.3.1.3), the presence of adsorbed deuterium on the surface could affect the measured sputtering yield by decreasing the beryllium concentration at the surface and altering the binding energy of surface beryllium atoms. Some evidence of this effect may be discerned in data from JET measurements of the beryllium sputtering yield. Two sets of sputtering yield measurements have been reported from JET; one from beryllium divertor plate measurements and the other from beryllium limiter measurements. In the divertor region, one expects a neutral density similar to that encountered in plasma generators (1020 m3 or more) and the measured sputtering yield is lower than that predicted by TRIM calculations.58 When sputtering measurements are made on the limiter, where the neutral density is typically lower, the sputtering yield agrees with, or exceeds, the calculated value.47 Of course, other issues such as impurity layers on the divertor plate and angle of incidence questions tend to confuse the results. However, the data sets from JET are consistent with the impact of neutral absorption on the beryllium plasma-facing surface. Effects associated with plasma operation will need to be taken into account when predicting sputtering yields from different areas of confinement devices. In addition to the low-energy neutral atom flux and higher-energy charge exchange neutral flux, the impact of small impurity concentrations in the incident plasma flux will also have a large impact on the expected sputtering yield. Some of the implications of the formation of a mixed-material surface are discussed in the next section and in Section 4.19.3.3. 4.19.3.1.2 Mixed-material erosion
As was pointed out in the previous section, it is important to have accurate knowledge of a target’s surface
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composition to predict its erosion rate. A small impurity concentration contained within the incident plasma can drastically alter the surface composition of a target subjected to bombardment by the impure plasma. Oxygen impurities in the plasma, either from ionization of the residual gas, or due to erosion from some other surface, will readily lead to the formation of beryllium oxide on the surface of a beryllium target. Depending on the arrival rate of oxygen to the surface compared to the erosion rate of oxygen off the surface, one can end up measuring the sputtering rate of a clean beryllium surface or a beryllium oxide surface. Careful control of the residual gas pressure in ion beam sputtering experiments55 has documented this effect. Unfortunately, it is not always so easy to control the impurity content of an incident plasma. In the case of a magnetic confinement device composed of groups of different plasma-facing material surfaces, erosion from a surface in one location of the device can result in the transport of impurities to other surfaces throughout the device. Mixedmaterial surfaces are the result. To first order, a mixed-material surface will affect the sputtering of the original surface material in two ways. The first is rather straightforward, and is true even for materials which do not form chemical bonds, in that the surface concentration of the original material is reduced thereby reducing its sputtering rate. The second effect changing the sputtering from the surface results from changes in the chemical bonding on the surface, which can either increase, or decrease the binding energy of the original material. If the chemical bonds increase the binding energy, the sputtering rate will decrease. If the bonding acts to reduce the surface binding energy, the sputtering rate will increase (assuming the change in surface concentration does not dominate this effect). A recent review of mixed materials62 provides some background information on the fundamental aspects of general mixed-material behavior. If a plasma incident on a beryllium target contains sufficient condensable, nonrecycling impurities (such as carbon), it will affect the sputtering rate of the beryllium. This effect was first referred to as ‘carbon poisoning.’ 5,9,63 A simple particle balance model has been used to adequately explain the results for formation of mixed carbon-containing layers on beryllium at low surface temperature.64 However, as the target temperature increases, additional chemical effects, such as carbide formation, have to be included in the model. An interesting change occurs when the bombarding species is a mixture of carbon and oxygen.
632
Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices
Measurements of the chemical composition of a beryllium surface bombarded with a COþ ion beam showed almost exclusive bonding of the oxygen to the beryllium in the implantation zone.65,66 The formation of BeO on the surface left the carbon atoms easily chemically eroded. The amount of oxygen present in the incident particle flux plays a strong role in the final chemical state of the surface atoms and their erosion behavior. The inverse experiment, beryllium-containing plasma incident on a carbon surface, has also been investigated.67–69 In the case of beryllium impurities in the plasma, a much more accurate measurement of the impurity concentration was possible. Contrary to the carbon in beryllium experiments, a simple particle balance model could not account for the amount of beryllium remaining on the surface after the plasma exposure. Clearly, the inclusion of chemical effects on the surface needs to be taken into account to interpret the results. Beryllium carbide (Be2C) was observed to form on the surface of carbon samples exposed to berylliumcontaining deuterium plasma even during bombardment at low surface temperature. Carbide formation will also act to increase the binding of beryllium atoms to the surface and decrease the binding of carbon atoms. This effect will result in an increase in the concentration of beryllium on the surface compared to a simple particle balance equation and must be included to understand the evolution of the surface. In addition, the formation of the carbide was correlated with the decrease of carbon chemical erosion70 (see Section 4.19.3.3.1 for more discussion of the chemical erosion of the beryllium–carbon system). 4.19.3.1.3 Chemically assisted sputtering of beryllium
The term chemically assisted physical sputtering refers to the transfer of energy from an incident particle to a molecule on the surface. The energy gained is sufficient to break any remaining bonds of the molecule to other atoms on the surface resulting in the release of the molecule, or a fragment of the molecule, from the surface. In the case of beryllium bombarded by deuterium plasma, the sputtering of beryllium deuteride was first recorded in JET71 during operation with a beryllium divertor plate. Since that time, a series of systematic investigations were performed in PISCES-B to quantify the magnitude of this erosion term.72,73 The results from PISCES-B show a surface temperature dependence of the sputtering rate72 of BeD
molecules. The maximum in the BeD sputtering rate (at 175 C) corresponds with the onset of thermal decomposition of BeD2 molecules73 from a standardized sample of BeD2 powder. Even at the maximum loss rate, the chemical sputtering remains smaller than the physical sputtering rate of beryllium atoms from the surface over the incident energy range examined (50–100 eV). Molecular dynamics simulations have predicted,74 and subsequent measurements have validated the prediction, that chemical sputtering can dominate physical sputtering of beryllium as the incident deuterium ion energy decreases below 50 eV. A distinction should be made between chemical sputtering and chemically assisted physical sputtering. Chemical sputtering involves the formation and loss of volatile molecules from a surface. In the case of beryllium deuteride, the molecule decomposes into a deuterium molecule and a beryllium atom before it becomes volatile, so at least to date there is no evidence for chemical sputtering of beryllium during deuterium particle bombardment. Documentation of the chemically assisted physical sputtering of beryllium may be important for determining material migration patterns in confinement devices and the identification of beryllium deuteride molecular formation in plasmaexposed surfaces may also help explain the hydrogenic retention properties of beryllium. 4.19.3.1.4 Enhanced erosion at elevated temperatures
In addition to the temperature-dependent chemical sputtering of beryllium when exposed to deuterium plasma, another temperature-dependent loss term is present in beryllium exposed to plasma bombardment at elevated temperature. The classical picture of the temperature dependence of erosion from chemically inert surfaces exposed to energetic particle bombardment is composed of the superposition of a constant physical sputtering yield with an exponentially varying thermal sublimation curve. The classical picture is contradicted, however, by experiments that show an exponential increase in erosion at lower temperature that cannot be explained by classical thermodynamic sublimation. First observed by Nelson75 for a variety of metal surfaces, similar results have been measured for Be,76,77 W,78 and C79,80 surfaces. In the case of carbon, this mechanism has been called RES. In the case of beryllium, two explanations have been proposed and both rely on the large flux of ions incident during plasma bombardment to modify the
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Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices
4.19.3.2 Hydrogen Retention and Release Characteristics 4.19.3.2.1 Implantation
The use of beryllium as a plasma-facing material in tokamaks has prompted many experimental studies of retention and emission of hydrogen implanted into beryllium-like metals from ion-beams or plasmas. References and discussions of these studies can be found in reviews.82–85 Here, we review those studies which are relevant to H retention in Be in a fusion plasma environment. This section is mainly excerpted from Federici et al.7 Two basic parameters for understanding H retention are the hydrogen diffusivity and
Solubility (H/metal atom Pa-1/2)
solubility. Studies of solubility and diffusivity are reviewed in Causey and Venhaus85 and Serra et al.86 Figures 487–90 and 587,91,92 show experimental values for hydrogen solubility and diffusivity in W and Be. For Be there are significant differences between results from various studies. These differences may be due to effects of traps and surface oxide layers. The presence of bulk traps tends to increase the measured values of solubility and to decrease the measured values of diffusivity (see Federici et al.7), especially under conditions where the concentration
10-7 Be
1
2 10-8 3
W
10-9
0.5
1.5
1.0 1000/T(K)
Figure 4 Measured solubility of hydrogen in tungsten (dashed line87) and beryllium (solid lines 1,88 2,89 and 390). Reproduced with permission from Federici, G.; Skinner, C. H.; Brooks, J. N.; et al. Plasma-material interactions in current tokamaks and their implications for next-step fusion reactors. Nucl. Fusion 2001, 41, 1967–2137 (review special issue), with permission from IAEA.
1´10-7 1´10-8
Diffusivity (m2 s−1)
plasma-facing material surface. In the first, the incident plasma ions, in addition to creating sputtered atoms from the surface, also create a population of surface adatoms. An adatom is an atom from a lattice site on the surface that has gained sufficient energy to leave its lattice location, yet does not have sufficient energy to escape from the surface as a sputtered atom. The atom then occupies a site on top of the regular lattice sites. Because an adatom does not have the same number of adjacent atoms as those in the lattice, it is less strongly bound to the surface and can therefore sublime at a lower temperature than one associates with equilibrium thermodynamic sublimation. In the second explanation, incident plasma ions that have thermalized somewhere below the surface of the lattice exert a stress on the surface atoms of the target again resulting in a lower binding energy of the surface atoms to the bulk of the material. Measurements show atoms are being lost from the surface at thermal energies,77 rather than the energy associated with sputtered particles (i.e., on the order of electron volts). This seems to verify the loss mechanism that occurs because of the thermal release of an ensemble of particles with a lower surface binding energy than that of bulk atoms of that element. Additional measurements at elevated temperature have documented the variation in Be atom surface loss rate with changes of the incident flux of energetic particles.81 The larger the incident flux, the lower the onset temperature for the enhanced erosion. The implication of this enhanced loss rate at elevated temperature is a reduction of the permissible operating temperature of any plasma-facing material, or alternatively that the lifetime of a component operating at extreme temperature may be less than that expected based on the predictions from classical surface loss terms.
W
1´10-9 1´10-10
1a
10-11
2
Be
1b
10-12 3 10-13
0.5
1.0 1000/T(K)
1.5
Figure 5 Measured diffusivity of hydrogen in tungsten (dashed line87) and beryllium (solid lines 1a&b,91 2,89 and 392). Reproduced with permission from Federici, G.; Skinner, C. H.; Brooks, J. N.; et al. Plasma-material interactions in current tokamaks and their implications for next-step fusion reactors. Nucl. Fusion 2001, 41, 1967–2137 (review special issue), with permission from IAEA.
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Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices
of hydrogen in solution is smaller than the concentration of traps. For this reason, studies done on materials of higher purity and crystalline perfection, and at higher temperatures and with higher concentrations of hydrogen in solution, tend to give more reliable results. The porosity and oxide inclusions present in beryllium produced by powder metallurgy are also likely to lead to inconsistent results in measurements of hydrogen solubility and diffusivity. In the Be experiments, the effects of traps were not characterized and may be dominant. One firm conclusion is that the solubility of hydrogen is very low in both Be and W. Many studies have been done on the retention and emission of H implanted into materials to provide data needed to predict H retention in fusion reactor environments. Figure 6 shows the retention of 1 keV deuterium implanted into Be at 300 K versus incident fluence, measured by thermal desorption.93 D retention in Be was close to 100% at low fluences but saturated at high fluences. Earlier nuclear reaction analysis (NRA) measurements of D retained in Be within 1 mm of the surface gave very similar results.94 This saturation behavior indicates that D implanted into Be at 300 K does not diffuse, but accumulates until it reaches a limiting concentration of 0.3–0.4 D/Be within the implantation zone. At high fluences, the implanted zone becomes porous allowing additional implanted D to escape. This
saturation mechanism is confirmed by electron microscopy, which shows bubbles and porosity in the implantation zone after high fluence H implantation.95 Saturation of retention by the same mechanism is observed for D implanted into stainless steel at 150 K where the D is not mobile.96 H retention in Be increases with increasing ion energy and decreases with increasing sample temperature.84,97 The retention of 1 keV deuterium implanted into W and Mo at 300 K98 is also shown for comparison in Figure 6. Figure 784 shows retention of deuterium and tritium as a function of incident particle fluence from a set of high fluence experiments in which Be specimens were exposed to laboratory ion-beams (Idaho National Engineering and Environmental Laboratory (INEEL), University of Toronto Institute for Aerospace Studies (UTIAS)), linear plasma devices (Sandia National Laboratory (SNL)/Los Alamos National LaboratoryTritium plasma experiment (LANL-TPE), University of California, San Diego-Plasma Interaction with Surface Components Experimental Station B (UCSDPISCES-B)), a tokamak divertor plasma (DIIID-DIMES), and a neutral beam (NB-JET). In some of these studies carbon deposition or formation of carbide or oxide surface layers occurred, which is likely to affect D retention. The figure shows the D retention in Be observed under a wide range of exposure conditions. The high fluence saturated concentration tends to be lower at higher temperatures.
1022
100% retention
Retained fluence (D m−2)
3 2
Be
1021
W
3 2
1020
Mo, slope = 0.345
3 2 19
10 20 10
3 keV D+3 (1020 D m–2 s) 2
21
10
2
Be, Mo, or W (300 K) 10
22
2
Incident fluence
1023 2
1024 2
1025
(D m−2)
Figure 6 Retention of 1 keV deuterium implanted into Be and W, at 300 K versus incident fluence, measured by thermal desorption. Adapted and reproduced with permission from Haasz, A. A.; Davis, J. W. J. Nucl. Mater. 1997, 241–243, 1076–1081, Federici, G.; Skinner, C. H.; Brooks, J. N.; et al. Plasma-material interactions in current tokamaks and their implications for next-step fusion reactors. Nucl. Fusion 2001, 41, 1967–2137 (review special issue), with permission from IAEA.
Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices
635
1023 SNL/LANL-TPE
Retained quantity (D (T) m–2)
PISCES-B-1
1022
PISCES-B-2 PISCES-B-3 INEEL(200–400) UTIAS(27)
1021
100
DIII-D(DiMES)(130)
250
100 NB-JET(120)
200 250 250
1020
200
200
200 500
500 500
40 200 300
300 500
500
500 540 500 700
1019 21 10
1022
1023 1024 1025 Particle fluence (D (T ) m–2)
1026
1027
Figure 7 Retention of deuterium and tritium in Be as a function of incident particle fluence. For purposes of comparing results from different experiments using different ion energies, the data have been scaled to correspond to an equivalent 100 eV deuterium ion energy. Numerical values next to the symbols and in the legend are specimen exposure temperatures, in degrees Celsius. Reproduced with permission from Anderl, R. A.; et al. J. Nucl. Mater. 1999, 273, 1–26.
It must be noted that this phenomenon is very important because it implies that tritium inventories and permeation due to implantation in beryllium for ITER PFC applications should be significantly lower than was previously estimated using classical recombination-limited release at the plasma surface. A first attempt to model this saturation by allowing the recombination coefficient to become exponentially large as the mobile atom concentration near the plasma-facing surface approaches a critical value was made by Longhurst et al.99 For Be, calculations suggest that the critical concentration is related to the yield strength using Sieverts’ law of solubility. On the basis of the results of these calculations, it can be concluded that the inventory of tritium in the beryllium first wall of a device such as ITER, because of implantation, diffusion, trapping, and neutroninduced transmutation, will be of the order of 100 g rather than the kilogram quantities estimated previously,100,101 and most of that will result from neutroninduced transmutations in the Be itself and from trapping in neutron-induced traps. Current predictions of tritium inventory in ITER are briefly discussed in Section 4.19.6.2.1. Fusion neutrons will create vacancies and interstitials in plasma-facing materials. For metals at reactor wall temperatures, these defects will be mobile and will annihilate at sinks (e.g., surfaces or grain boundaries), recombine, or agglomerate into defect clusters. Vacancy agglomeration may also lead to the formation
of voids. In beryllium, neutron-induced nuclear reactions produce helium and tritium, which may be trapped at defects or precipitate as gas bubbles. These defects, resulting from neutron irradiation, will increase the retention of hydrogen, by increasing the concentration of sites where diffusing hydrogen can precipitate as gas or become trapped as atoms. The effect of neutron irradiation on hydrogen retention in metals is complex, but, in principle, this can be modeled, provided the material parameters are known, such as hydrogen diffusivity, solubility, trap binding energy, and defect microstructure produced by the neutron irradiation. For many metals, most of these parameters are known well enough to attempt such modeling. For beryllium, however, uncertainties in solubility, diffusivity, and trapping of hydrogen make such modeling of hydrogen retention difficult. The problem of production of helium and tritium by nuclear transmutation in beryllium itself is discussed in Section 4.19.4.4.5. 4.19.3.2.2 Beryllium codeposition
As deuterium retention in plasma-exposed beryllium targets saturates after a given ion fluence (see Section 4.19.3.2.1), it is apparent that retention in codeposits will eventually be the dominant accumulation mechanism with respect to beryllium PFCs. This is primarily due to the fact that the thickness of a codeposit will continue to grow linearly with time. It is, therefore, critical to understand both the
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Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices
1
1 Mayer et al. Mayer et al.
0.1
0.1 Causey et al. TPE
0.01
Present data PISCES 400
O/Be
D/Be
Present data PISCES
Causey and Walsh TPE
Causey and Walsh TPE 0.01
600 800 Temperature (K)
400
800
Figure 8 Comparison of D/Be levels in beryllium codeposits with the O/Be levels in the same codeposits. Reproduced with permission from Baldwin, M. J.; Schmid, K.; Doerner, R. P.; Wiltner, A.; Seraydarian, R.; Linsmeier, Ch. J. Nucl. Mater. 2005, 337–339, 590–594.
retention amounts and the release behavior of hydrogen isotopes from beryllium codeposits. In this section, a ‘codeposit’ includes both the codeposition (where a BeD or BeD2 molecule is deposited on a surface) and co-implantation (where deposited layers of beryllium are bombarded with energetic hydrogen isotopes) processes. Initial interpretation of studies of beryllium codeposits were made difficult by relatively high oxygen impurity content within the codepositing surface.102,103 Subsequent measurements104 with lower oxygen content seemed to indicate that the oxygen level within the codeposit was correlated to the level of hydrogen isotope retention in the codeposit. The other variable that was identified to impact the retention level in these studies was the temperature of the codepositing surface. Measurements seriously questioning the importance of oxygen on the retention level in beryllium codeposits were made by Baldwin et al.105 In this data set, the oxygen content throughout the codeposit was measured by depth profiled X-ray photoelectron spectroscopy and the oxygen content did not correlate with the deuterium retention level (Figure 8), although the temperature of the codepositing surface was still a dominating term in determining the deuterium retention level. Later, more detailed measurements confirmed that the presence of a beryllium
oxide surface layer was not correlated with an increase in retention in beryllium.106 A systematic study of beryllium codeposition followed,107 identifying three experimental parameters that seemed to impact the retention level in a codeposit. Along with the surface temperature, the incident deuterium energy and the beryllium deposition rate were determined to be influential scaling parameters. The previously reported data in the literature was also evaluated using the derived scaling and found to agree with the predictions of the retention levels measured under the various experimental conditions present in the different machines. Later the derived scaling was revised108 to use the ratio of the fluxes of the codepositing species, rather than the deposition rate to permit more accurate extrapolation to conditions expected in the edge of confinement devices. The ability to predict the level of tritium retention in beryllium codeposits is an important aspect of a safety program; however, developing techniques to remove the trapped tritium from codeposits is a more important issue. The deuterium release behavior during thermal heating of beryllium codeposits has been investigated.109 The results show that the maximum temperature achieved during a bakeout is the figure of merit for determining the amount of deuterium release from beryllium. Increasing
Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices
the time spent at lower baking temperatures did not increase the amount of deuterium released from the beryllium codeposits. These results, along with the retention level predictions, should make it possible to design baking systems for different areas of a confinement device to control the accumulation rate of tritium to a desired level. 4.19.3.3
Mixed-Material Effects
A recent review of mixed-material effects in ITER62 provides background information on mixed-material formation mechanisms and plasma–surface interaction effects. Here, the focus is on berylliumcontaining mixed-material surfaces (i.e., Be/C and Be/W) and the conditions when one might expect these surfaces to dominate the observed plasma– surface interactions. In addition to plasma interactions with mixed-material surfaces, which will be discussed here, other aspects such as changes to thermal conductivity, material strength, and ductility, the impact of impurities on material joints, etc., must also be carefully evaluated. 4.19.3.3.1 Be–C phenomena
Beryllium and carbon have been observed to begin thermally interdiffusing at a temperature of around 500 C,56 resulting in the formation of a beryllium carbide layer. However, beryllium carbide has also been observed to form during energetic carbon ion bombardment of beryllium surfaces at room temperature.110 As mentioned in Section 4.19.3.1.2, the change in the binding energy of the carbide molecule affects the sputtering yield of both the beryllium and carbon atoms. In addition, the formation of beryllium carbide also has a dramatic effect on the chemical erosion properties of a carbon surface bombarded with energetic beryllium ions.67,68,111 The presence of beryllium carbide on the surface of a carbon sample exposed to deuterium plasma has been shown to correlate with the reduction of chemical erosion of the carbon surface.70 The speculation for the cause of this effect is that the carbide enhances the recombination of deuterium in the surface, thereby lessening the amount of deuterium available to interact with carbon atoms on the surface. This is similar to the impact of small amounts of boron carbide in a graphite surface affecting chemical erosion.112 However, the difference here is that instead of obtaining the carbide through an expensive production technique, the carbide forms naturally as beryllium ions in the plasma interact with the carbon surface.
637
A systematic study of the time necessary to suppress chemical erosion of a graphite surface due to the interaction with beryllium-containing plasma has been carried out.69 Increasing the surface temperature of the graphite was seen to have the biggest impact on reducing the suppression time. Increasing the beryllium content of the plasma also reduced the suppression time in a nonlinear fashion. An increase of the incident particle energy was observed to increase the time necessary to suppress the chemical erosion of the surface, presumably due to an increase in the removal of the carbide-containing surface layer. A subsequent study showed that applying heat pulses to a graphite surface interacting with beryllium-containing plasma, to simulate surface heating due to intermittent events, acted to reduce the time necessary for the carbide surface to form and suppress the chemical erosion of the surface.113 4.19.3.3.2 Be–W alloying
The existence of tungsten beryllide alloys (i.e., Be2W, Be12W, and Be22W) is an excellent example of the importance of mixed-material surface formation in plasma-facing components.114 Figure 9 shows the tungsten–beryllium phase diagram. Each of the beryllides shown in the figure exhibits a lower melting temperature than one would expect from a tungsten plasma-facing surface. If plasma containing beryllium impurities interacts with a tungsten surface, there is a possibility of these lower melting temperature beryllide alloys being formed. In thermodynamic equilibrium, various beryllide alloys of tungsten have been observed to form,115 and their reaction rates have been measured,116 at temperatures in excess of 800 C. However, as was seen with beryllium carbide forming during plasma bombardment at lower temperature than expected thermodynamically, the concern exists that tungsten beryllide could form at temperatures below 800 C as well. Well controlled laboratory measurements in vacuum117 and in plasma simulators118 have shown that although thin, nanometer scale, Be2W layers form at the interface between beryllium and tungsten surfaces, their growth below 800 C is negligible. In addition, above 800 C, rapid beryllium sublimation from surfaces can act to limit the amount of beryllium available for reacting with tungsten and thereby also limit the growth rate of the alloys. In the present low wall temperature confinement devices, modeling shows that the divertor strike point locations are the only areas where significant beryllide growth might be expected and in these regions there does not
638
Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices
Weight percent tungsten 90 95
0 4060 70 80
100
3500 3422 ºC
3000
L
<2250 ºC 1289 ºC
?
2000
2100 ± 50 ºC -60
-95
W
<1750 ºC ?
? 1500
Be2W
Temperature ( ºC)
2500
? (bBe) (aBe) Be12W
Be22W
1000
500 0 Be
10
20
50 60 70 30 40 Atomic percent tungsten
80
90
100 W
Figure 9 Phase diagram for the Be–W system. Reproduced with permission from Doerner, R. P.; Baldwin, M. J.; Causey, R. A. J. Nucl. Mater. 2005, 342, 63–67.
appear to be enough beryllium deposition to raise significant concerns.119 One caveat to these predictions would be the existence of intermittent events that raise the temperature of surfaces where significant beryllium deposits are located, thereby possibly allowing the optimized beryllide growth conditions. Another concern with regard to thin Be2W surface layers on plasma-exposed tungsten is the impact of these layers on tritium retention. While a thin Be2W surface layer is not likely to retain much tritium itself, the thin beryllide surface layer could alter the recombination characteristics of the bulk material and change the accumulation rate of tritium within the device. To date, there is little or no data available to address this issue. While it appears likely that the most serious issues of tungsten beryllide formation may be avoided in present confinement devices, the issues associated with these alloys highlight the uncertainties and importance of understanding and predicting mixed-material
formation in plasma environments. Mixed materials often interact with plasma in much different ways than their elemental components. In the case of the beryllium–carbon system (Section 4.19.3.3.1), the mixed material appears to offer the potential for beneficial effects, whereas in the case of the beryllium– tungsten system, the mixed material appears likely to be detrimental to the operation of the device. Each mixed-material system must, therefore, be individually evaluated to determine its potential impact on all aspects of operating surfaces in contact with plasma.
4.19.4 Main Physical and Mechanical Properties 4.19.4.1
General Considerations
A comprehensive, although not recent, review of the science and technology of beryllium can be found in Beryllium Science and Technology.120
Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices
Several reviews have been published recently related to use of beryllium in tokamaks and the status of the investigations of the Be properties for the fusion application.3,121–126 Various production and processing methods of beryllium metal fabrication have been reviewed in Dombrowski.127 The majority of methods are based on powder metallurgy and include powder preparation from cast product by grinding (i.e., attrition milling, impact grinding, ball mill grinding); further powder consolidation (i.e., by cold pressing (CP), cold isostatic pressing (CIP), vacuum hot pressing (VHP), hot isostatic pressing (HIP)); and possible additional mechanical treatment (e.g., extrusion, rolling, forging). Beryllium protective armor can also be produced by plasma spray (see Section 4.19.4.3) and vapor deposition. Several proposals were made at the beginning of the ITER Research Programme during the ITER Engineering Design Phase to develop a fusion grade beryllium with high ductility, high resistance to heat flux, and high radiation resistance. However, it was recognized that this development would require significant efforts and could not be supported only by requests from the fusion community. There are various beryllium grades, which have been developed for different applications. These grades differ by chemical composition (BeO content, impurities), by method of powder preparation, by method of consolidation, etc. The nonexhaustive list of various beryllium grades from the US and the Russian Federation is presented in ITER Materials Properties Handbook (MPH).128 Grades with similar composition are under production in Kazakhstan and in China. We briefly discuss below some of the most relevant physical and mechanical properties of beryllium, in relation to its application as armor for PFCs. 4.19.4.1.1 Physical properties
The physical properties of beryllium are summarized in Table 2, which is taken from ITER MPH.128 These properties have been used for design and performance assessments. In addition to its low atomic number, beryllium has several excellent thermal properties that make it well-suited for heat removal components. The thermal conductivity is comparable with that of graphite or CFC at low and high temperatures but, in contrast to C-based materials, is not significantly degraded as a result of neutron-irradiation. The specific heat of beryllium exceeds that of C-based materials typically by a factor of 2 over the temperature range of interest for operation. However, Be has poor refractory
Table 2
Physical properties of beryllium
Atomic number Atomic weight Crystal structure Density (kg m3) Melting temperature ( C) Thermal conductivity (W m1 C1) Specific heat (J kg1 C1) Latent heat of fusion (kJ kg1) Latent heat of vaporization (kJ kg1) Electrical resistivity (mO cm) Thermal expansion coefficient 106 C1 Emissivity
639
4 9.013 Hexagonal closepacked 1830–1850 1283–1287 200 (RT) 82 (800 C) 1900 1300 3.66 104 4.4 (RT) 11.6 (RT) 14.96 (400 C) 0.1–0.5* (at 300 C)
Source: ITER MPH, ITER Final Design Report 2001 (internal project document distributed to the ITER participants). RT, room temperature. * Depending on quality of surface
properties, such as low melting temperature and high vapor pressure. The high heat capacity and good thermal conductivity of Be can be used to maintain low surface temperatures in PFCs during normal operation, but its low melting temperature and high vapor pressure cause great design difficulties from the standpoint of survivability from offnormal events such as vertical displacement event (VDE), ELMs, disruptions, and runaway electron impact (see Section 4.19.6.2). For the beryllium hexagonal close packed crystal structure, the main physical properties, such as the coefficient of thermal expansion, elastic modulus etc. have some anisotropy. However, for the polycrystalline grades these properties could be, in the first approximation, considered as isotropic. Some anisotropy is also typical for the highly deformed grades. The physical properties (thermal conductivity, specific heat, elastic modulus, etc.) in first approximation are the same for beryllium grades with similar BeO and other impurity content and they are produced by the same fabrication method. 4.19.4.1.2 Mechanical properties
Beryllium is known to be a brittle material, with a typical elongation to failure in room temperature tensile tests of roughly 0.8–6%. For material with strong anisotropy (e.g., rolled plate or sheet), elongation in the rolling direction could be higher, but in the transverse direction the elongation is
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Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices
typically significantly lower than 1%. Recently, the mechanical properties of beryllium have been summarized in ITER MPH128 and ITER Materials Assessment Report (MAR).129 The mechanical properties of beryllium depend on the production method used and they are sensitive to a variety of factors including BeO and impurity content (which varies from less than 1% to 2–3% for various grades), method of powder preparation (impact grinding, attrition grinding), method of consolidation, and further treatments. The main problem in using beryllium is its low ductility related to the hexagonal-close-packed structure. There is limited slip in directions not parallel to the basal planes, resulting in very small ductility perpendicular to the basal direction. Depending on the production method, ductility of beryllium can be severely anisotropic. The grain size is an important factor in determining the ductility of various beryllium components. Much of the fine grain size present in the starting powder is retained during hot pressing at 1060 C. Without an oxide network, grain growth occurs at a much lower temperature, about 800 C. Among various beryllium grades, it was found that grade S-65C VHP (production of Brush Wellman, US) has the highest guaranteed fracture elongation at room temperature (minimum 3%; typical is more than 4–5%). This grade is produced using impact grinding powder and has a guaranteed BeO content <1%. The level of impurities is also controlled. The high ductility of the grade is one of the advantages of this material. Because of the VHP production method, there is some anisotropy of properties in relation to hot pressing direction, but the differences are not significant. As typical for all metals, the tensile properties of beryllium depend on the testing temperature. As the testing temperature increases, a decrease of the ultimate tensile and yield strength are observed. However, rupture elongation increases with increasing test temperature and could reach a value higher by 40–50% for temperatures around 300–350 C (see as example data for grade S-65C VHP in the ITER MPH128). A further increase in the test temperature leads to a decrease of the elongation. At temperatures above 600 C, the ductility depends on the impurity content, mainly aluminum, which tends to segregate at grain boundaries, impairing the mechanical properties. By heat treatment in the temperature range 650–800 C, aluminum can be combined with other elements, mainly iron and beryllium itself, to form a stable beryllide as AlFeBe4. However, the stable
beryllide dissolves progressively when heated at temperatures >850 C. This last feature is important for the selection of the joining technology for manufacturing of the PFCs. Further details on mechanical properties, such as creep and fracture toughness, can be found elsewhere (see, e.g., ITER MAR129). 4.19.4.2 Selection of Beryllium Grades for ITER Applications For ITER PFC applications, various commercially available beryllium grades from the United States (Brush Wellman Inc.) and from the Russian Federation, listed in Table 3, were evaluated more than a decade ago as potential candidates during the ITER Engineering Design Activity (EDA). The selection of the optimum grade for ITER PFC applications is driven mainly by the requirements of ITER operation for structural integrity and stability against various thermal loads, and in particular, the absence or minimization of macrodamage. It is believed that ion-induced and thermal erosion at elevated temperatures is very similar for various grades of Be. However, performance under high heat fluxes, especially under transient thermal loads such as disruptions, VDE, and ELMs resulted in different behavior and damage mechanisms. It is considered that the ease of joining beryllium to copper alloys (see Section 4.19.5) is not so sensitive to BeO content, impurity levels, and method of consolidation, which are the parameters defining the grade of beryllium material. It should be noted that for tokamak applications (see Section 4.19.6) beryllium is used in the form of tiles. Some surface cracking of the tiles could be acceptable, if there is no macrodamage or delamination along the surface of tiles, which leads to the loss of macroscopic particles. The resistance to thermal fatigue is the most important factor that affects the material selection Table 3
Candidate grades of beryllium
Producer
Grades
Brush Wellman, US
S-65C VHP, S-65C HIP, S-65 CIP, SR-200 VHP S-200F HIP, S-200F VHP, I-400 VHP DShG-200, TShG-56, TR-30, TGP-56 TShGT, DIP-30, TShG-200
Russian Federation
VHP, vacuum hot pressing; HIP, hot isostatic pressing; CIP, cold isostatic pressing.
Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices
because cracking could lead not only to enhanced armor erosion, delamination, and loss of particles, but also potentially to crack propagation to the heat sink structure. Neutron irradiation resistance is another factor to be taken into account because it may affect the thermal performance and structural integrity. Because of some of the uncertainties in the ITER thermal loads, especially during transient events, preference is given to beryllium grade(s) with potentially higher resistance to transient thermal loads. The selection of the reference grades was made on the basis of comprehensive assessment of the results of various tests carried out during the ITER EDA. The detailed analysis is presented in ITER MAR.129 Among the various studies, the following shall be mentioned:
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the molten zone, whereas for some grades the cracks propagated to the bulk of the sample. Results of VDE simulation tests have been reported in Linke et al.132,133 Severe melting of Be was observed for energy densities of 60 MJ m2 (1 s pulse duration); however, no cracks were observed between molten and unmolten material and in the bulk of unmolten parts for S-65C VHP grade. On the basis of the available data, Be S-65C VHP (Brush Wellman, US) was selected as the reference material on the basis of excellent thermal fatigue and thermal shock behavior, and for the good available database on materials properties, including neutron irradiation effects. DShG-200 (produced in the Russian Federation) was proposed as a backup, but this grade is no longer commercially available. Recently, China and the Russian Federation, that are two of the seven International Parties engaged in the construction of ITER, have proposed the fabrication of additional first-wall grades as part of their ITER contribution. The Russian Federation proposes to use beryllium grade TGP-56-FW. This grade is produced by VHP in almost the final form of the tiles foreseen for the first wall. The recent results on development of this grade have been reported in Kupriyanov et al.134 China proposes instead to use a grade called CN-G01135 that is produced from
Screening low cycle fatigue test of 21 different beryllium grades was performed in the past.130 It was shown that the grades with the best thermal fatigue resistance are S-65C VHP, DShG-200, TShG-56, and TShGT. Figure 10 shows the results of the comparative low cyclic thermal fatigue study of different grades of beryllium. Various grades of beryllium were also tested in conditions simulating the disruption heat loads.131 The tests show that crack formation and behavior after surface layer melting in different grades are quite different. For Be S-65C, all cracks stopped in
2500 S-65C (L)
Grades with best fatigue performance
DShG-200 (T) Cycles to crack initiation
2000 S-65C (T) S-65-H TShGT(T) TShG-56 (T)
1500
S-200F (T)
98% S-65 1000 Extruded (T) S-200F-H SR-200
500
Be/60% BeO 0
0
94% S-65
S-200F (L) TGP-56 Extruded (L) I-400 (T)
Be/30% BeO 0.5 1 1.5 Side crack propagation depth (mm)
2
Figure 10 Results of low cycle thermal fatigue tests of different Be grades: number of cycles to crack initiation versus crack propagation depth. Reproduced with permission from Watson, R.; et al. Low cyclic fatigue of beryllium. In Proceedings of the 2nd IEA International Workshop on Beryllium Technology for Fusion, Wyoming, Sept 6–8, 1995; p 7.
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impact ground powder (similar to powder used for S-65C grade) by VHP. The grade is produced by Ningxia Orient Non-Ferrous Metals Co. Ltd. In order to accept these newly proposed beryllium grades a specific qualification program is underway. 4.19.4.3 Considerations on Plasma-Sprayed Beryllium In the past, plasma spraying was considered as a high deposition rate coating method, which could offer the potential for in situ repair of eroded or damaged Be surfaces. Development work was launched during the early phase of the ITER R&D Program in the mid-1990s.136 In the plasma spray process, a powder of the material to be deposited is fed into a small arc-driven plasma jet, and the resulting molten droplets are sprayed onto the target surface. Upon impact, the droplets flow out and quickly solidify to form the coating. With recent process improvements, high quality beryllium coatings ranging up to more than 1 cm in thickness have been successfully produced. Beryllium deposition rates up to 450 g h1 have been demonstrated with 98% of the theoretical density in the as-deposited material. Several papers on the subject have been published.136–138 A summary of the main achievements can be found in Table 4. However, based on the results available, the initial idea of using plasma-sprayed beryllium for in situ (in tokamak) repair was abandoned for several reasons. First was the complexity of the process and requirements to control a large number of parameters, which affect the quality of the plasma sprayed
Table 4 together)
coatings. Some of the most important parameters include plasma spray parameters such as (1) power, gas composition, gas flow-rate, nozzle geometry, feed, and spray distance; (2) characteristics of the feedstock materials, namely, particle size distribution, morphology, and flow characteristics; (3) deposit formation dynamics, that is, wetting and spreading behavior, cooling and solidification rates, heat transfer coefficient, and degree of undercooling; (4) substrate conditions, where parameters such as roughness, temperature and thermal conductivity, and cleanliness play a strong role; (5) microstructure and properties of the deposit, namely, splat characteristics, grain morphology and texture, porosity, phase distribution, adhesion/cohesion, and physical and mechanical properties; and (6) process control, that is, particle velocity, gas velocity, particle and gas temperatures, and particle trajectories. Second, plasma-sprayed beryllium needs (1) inert gas pressure, (2) reclamation of the oversprayed powder (more than 10%), and (3) strict control of the substrate temperature. The higher the temperature the higher the quality of the plasma-sprayed coating, but unfortunately, an easy and reliable method to heat the first wall to allow in situ deposition was not found. Finally, tools to reliably measure the quality of the coating and its thickness are not available today and a strict control of the coating parameters is difficult to achieve. Thus, it was concluded that plasma-sprayed beryllium for in situ repair is too speculative for ITER without further significant developments. Nevertheless, this method still remains attractive and could be used for refurbishment of damaged components in
Main achievements of ITER-relevant plasma-sprayed technology (summary of best results, not always achieved
Parameter
Value/results
Comments
Residual porosity (%) Thermal conductivity (W mK1)
2 Up to 160 at RT
Bond strength (MPa) Substrate temperature ( C)
100–200 >450
Substrate preparation
Negative transfer arc 4.5 >10 >90 5/680; 1/3000
Could be more than 5% Depends on temperature of substrate, maximum achieved at T 600–800 C with addition of H Reasonable Very important for good strength, adhesion, and thermal conductivity. Keep in mind that CuCrZr temperature should not be higher than 500 C for several hours due to overageing of CuCrZr Needed, but very difficult to do in situ
Deposition rate (kg h1) Thickness (mm) Deposition efficiency (%) Thermal fatigue (MW m2/ number of cycles)
Reasonable Reasonable It means that more than 10% of powder will be lost in chamber For first-wall conditions tested
Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices
hot cell, albeit it may be cheaper to replace a damaged component with a new one. 4.19.4.4
Neutron-Irradiation Effects
Several authors have reviewed the properties of neutron-irradiated beryllium for fusion applications in the past.139–141 Neutron irradiation leads to complex changes in the microstructure, such as the radiation-induced change of volume in beryllium, which is dominated by the nucleation and growth of He bubbles. There are two important pathways for gas production. One is the (n, 2n) reaction in which the 9Be is reduced to 8Be, which then splits into two 4He atoms. The second is the (n,a) reaction where the 9Be absorbs a neutron and then splits to form a 4He and a 6He. The 6He rapidly undergoes a b decay to become 6Li. The 6Li then reacts with a thermal neutron to produce 4He and 3H. These processes have been incorporated into the inventory code FISPACT,142 which is used (see, e.g., Forty et al.143) to estimate the generation rates of gas and other reaction products in a tokamak. Helium generation has significant effects on the properties of materials, especially at elevated temperatures. Helium is initially trapped within the beryllium lattice in submicroscopic clusters. At higher neutron fluence massive helium-bubble-induced swelling occurs, especially at elevated irradiation or postanneal temperatures. Because of the atomistic nature of the helium bubble nucleation and growth, porous beryllium microstructures, such as from powder metallurgy or plasma spray technology, were not found to be effective in releasing significant amounts of helium under fusion reactor conditions.2 The maximum neutron-induced damage and helium production expected in Be for ITER firstwall applications (fluence of 0.5 MWam2) are 1.4–1.7 dpa and 1500 appm, respectively and the expected irradiation temperatures are in the range of 200–600 C. The maximum temperature is on the surface of beryllium tile and depends on thickness and heat flux. Tritium production in beryllium is expected to be about 16 appm. Recently, Barabash et al.144 have analyzed the specific effects of neutroninduced material property changes on ITER PFCs foreseen during ITER operation. Typically, property changes induced by neutron irradiation are investigated by exposing samples/ mock-ups in fission reactors. However, the differences between the fission and fusion neutron spectra
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are important to interpret and predict the effects. The key difference is transmutation production, which needs to be considered for the correct prediction of the material performance.145 During irradiation in fission reactors, for example, the typical value of the ratio (appm He per dpa) is 100–250, whereas for a fusion neutron spectrum this value is 1000. Depending on operational temperature, the dpa or He transmutation must be used as a reference neutron damage parameter. For beryllium, during lowtemperature irradiation (<300 C) the dpa value must be considered. For high-temperature irradiation (more than 500 C), the He generation must be taken as the reference parameter. A detailed discussion on this subject is beyond the scope of this review. We summarize only some of the main findings with emphasis on results for ITER relevant grades. Considerations of the effects of neutron irradiation of duplex Be/Cu alloy mock-ups are provided in Section 4.19.5. 4.19.4.4.1 Thermal conductivity
For S-65C Be grade irradiated up to 1025 n m2 (0.74 dpa) at 300 C, the thermal conductivity was found to be similar, within experimental error, to that of the unirradiated material.146 Similarly, no effect was seen for Be S-65C after irradiation at 350 and 700 C to a damage dose 0.35 dpa.147 Significant changes in the thermal conductivity were observed only for conditions that lead to significant changes of the beryllium structure, such as the formation of a high density of radiation defects (especially at low irradiation temperature and high dose) or high (more than tens of percent) swelling.144 Other physical properties (elastic modulus, coefficient of thermal expansion, etc.) are not influenced by neutron irradiation (at least at the fluence and temperature ranges relevant for the beryllium armor for the ITER PFCs). 4.19.4.4.2 Swelling
It is well known that beryllium swells when irradiated by neutrons, especially during high temperature irradiation. Reviews of the available swelling data for different Be grades can be found elsewhere (see, e.g., ITER MAR,129 Billone,139 and Barabash et al.141). The computer code ANFIBE (ANalysis of Fusion Irradiated BEryllium), has been developed and applied in the past as an interpretative and predictive tool148 for the prediction of beryllium swelling. The driving force for the swelling is the presence of He, which forms He bubbles, especially during
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high-temperature irradiation (more than 400 C) or after high-temperature annealing. The maximum values of swelling could reach approximately tens of percent at temperatures more than 600 C and helium content more than several thousand atomic parts per million. Swelling depends on the structure of the beryllium: beryllium grades with small grain size (8–10 mm) and high BeO content (3–4 wt%) have a higher resistance to swelling than conventional Be grades.141 As concluded in ITER MAR,129 for an irradiation temperature <400 C, swelling of beryllium containing 1500 appm He is <1%. At higher temperature, swelling could reach the value of a few percent at the end of life. 4.19.4.4.3 Mechanical properties
Beryllium undergoes hardening from dislocation pinning and grain-boundary decohesion from the helium bubbles at the interfaces. The general qualitative trends in the strength and ductility behavior of neutron-irradiated Be were summarized in Gelles et al.140: (1) at low and moderate irradiation temperatures (20–500 C) the strength typically increases, while ductility decreases, in some cases to zero value; (2) at high temperature (more than 600 C) the ductility decreases without an increase in the strength; and (3) an increase of the dose leads to saturation of the changes of the strength and ductility. The main effect is the embrittlement of Be at low irradiation temperature (<200 C) due to accumulation of the radiation defects in the form of dislocation loops, and at high irradiation temperature (more than 400 C) due to He bubble formation at grain boundaries. For the range of conditions expected at the ITER first wall (e.g., Be temperature range 200–600 C and damage level at the end of life of 1 dpa), it is expected that the ductility of Be will be at the level <1%. Recent irradiated Be mechanical property studies (at ITER relevant low temperatures) demonstrated that no catastrophic embrittlement or thermal conductivity degradation occurred.149 4.19.4.4.4 Thermal shock effects
First-wall beryllium protection tiles in ITER will be subjected to thermal shocks (disruptions, VDEs). The damage during these events is a complex function of the heat loads and material properties, which, as described above, are neutron irradiation dependent. Only limited studies to investigate the combination of these effects have been performed,150,151 but more work in this area is necessary. Several Be grades (S-65C, plasma sprayed Be, TShG-56, etc.) have
been irradiated at 350 and 700 C to a fluence of 0.35 dpa and corresponding He content of 55 appm. An increase of erosion after neutron irradiation (up to 100%) was observed for different grades of beryllium for a thermal shock disruption load of 15 MJ m2. It was concluded that in this case thermal erosion is not caused by simple evaporation, but by the loss of some particles due to the brittle destruction of the surface. 4.19.4.4.5 Bulk tritium retention
Several studies have been conducted in the past to investigate tritium trapping in irradiated beryllium152,153 and it has been found that the tritium retention can increase with neutron fluence by a factor 3–10. Tritium loading in these experiments was performed by exposure to tritium in the gaseous phase. Tritium atoms are indeed trapped by neutroninduced defects such as dislocation loops and helium bubbles. Bulk tritium retention due to implantation, and, as a result, the effect of the damaged microstructure on this retention, is expected to be less serious than previously anticipated.154
4.19.5 Fabrication Issues 4.19.5.1 Joining Technologies and High Heat Flux Durability of the Be/Cu Joints One of the most critical aspects of the design of the ITER first wall is the attachment of the Be tiles to the actively cooled copper alloy substrate. The primary threat to this attachment comes from the heat flux. However, secondary effects also arise from the fact that there will be thermal gradients as well as mechanical loads from disruptions, which will cause distortion. These effects cannot easily be simulated experimentally. It should, therefore, be the responsibility of the design to minimize these effects. At the time of writing this chapter, the first wall of ITER was undergoing a major redesign. Because of the lack of more detailed information on the new design, most of the considerations presented here, albeit general, are on the basis of the design of actively cooled ITER first-wall panels, which consist of stainless steel tubes in a copper heat sink with 10 mm thick beryllium tiles bonded to the plasmafacing surface. This technology was deemed to be adequate for handling the previously assumed surface heat flux of 0.5 MW m2, and could potentially withstand 3 MW m2(155) for a limited number of cycles. However, because of fatigue it would be incompatible
Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices
with areas subjected to higher power densities during each pulse. In such cases, other options for the heat exhaust technology are being considered,156 using thinner Be tiles. 4.19.5.1.1 Be/Cu alloy joining technology 4.19.5.1.1.1 Background information
The main problem of bonding Be to Cu alloys is that Be reacts with almost all possible metals (except Al, Si, Ag, and Ge) and forms brittle intermetallic phases.157–159 Such bond joints have poor mechanical integrity. More robust joints use metal interlayers to act as either diffusion barriers and/or strain accommodating compliant layers to avoid the formation of deleterious phases and to assist in the accommodation of thermal cycling-induced strains.160 R&D has been performed over a number of years to develop the design and manufacturing techniques required to meet the demanding design requirements. Significant experience has been gained with these manufacturing techniques and the associated inspection techniques. It must be noted that in the 1990s the best joining technology developed for manufacturing the Be/Cu actively cooled components was brazing with Ag-base alloys (e.g., InCuSil with 41.75% Ag) which was successfully used in JET. However, the use of Ag base brazing alloy was not allowed in ITER mainly because of the transmutation by neutron irradiation to Cd (5 wt% Cd will be produced in Ag–Cu eutectic alloy at 1 MWam2) whose presence would (1) reduce the melting temperature of the braze; (2) lead to the formation of highly radioactive isotopes; and (3) affect the pumping system in case of Cd release to the vacuum chamber and codeposition in the cryopumps panels. During the early stage of the ITER first-wall design development, dispersion-strengthened copper (DS-Cu) alloys (e.g., Glidcop Al25) were considered as the first option because (1) the stresses were within the design allowable, and (2) they had better thermal stability under the manufacturing route, which required a first wall to be integrated with a 4 t shield. The main developments for fabrication of joints between Be and DS-Cu alloys are reported in ITER MAR129 and Lorenzetto et al.161 However, as a result of a design change that took place from an integrated first-wall panel to a separated first-wall panel design, a precipitation-hardened copper–chromium–zirconium alloy (CuCrZr), was subsequently chosen. This was because the fracture toughness of DS-Cu is very low above 200 C even for unirradiated material. Fracture toughness of the
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unirradiated and irradiated CuCrZr alloy decreases with increasing temperature, but it remains at a relatively high level in the ITER working temperature range and it is significantly higher than fracture toughness of DS-Cu. The use of separable first-wall panels makes it possible to perform heat treatments with fast cooling rates, which are mandatory to adequately retain the mechanical properties of precipitationhardened materials. Thus, extensive studies were then performed during the last 10 years to develop reliable silver free Be/CuCrZr alloy joining techniques and to modify the joining conditions to minimize the mechanical strength loss of the CuCrZr alloy. Different methods have been considered and investigated. Some of these were eliminated because of bad results (e.g., explosive bonding, inertial welding, joint rolling, and some types of brazing). Two methods gave good results and were kept for further investigations: HIPping and brazing. Good results were achieved with the HIP joining technique by lowering the HIP temperature as close as possible to the CuCrZr alloy ageing temperature (about 480 C) and with the brazing technique in the development of a fast brazing method to minimize the holding time at high temperature. The latter was achieved by induction brazing in Europe and by fast heating and cooling using an e-beam test facility in the Russian Federation. 4.19.5.1.1.2
HIP joining technique
An extensive development programme performed especially in Europe has enabled the production of very good Be/CuCrZr alloy joints by HIPping. The progress on the fabrication of Be/CuCrZr joints in Europe is described in Lorenzetto et al.155,161,162 and Sherock et al.163 The HIP joining temperatures ranged from 540 to 580 C. Different interlayers such as Cr, Ti, and Cu were tested and reported. The selection of the joining conditions used for the fabrication of representative first-wall mock-ups was done on the basis of mechanical test results performed with guillotine shear test specimens. The best results were obtained with Ti and Cu interlayers at 580 C, in which the shear strength exceeded the yield strength of the parent materials as creep of the CuCrZr part or rupture of the Beryllium part was observed. Performance achieved with representative first-wall mock-ups exceed the present ITER design requirements, namely, 30 000 cycles at 0.5 MW m2 peak heat flux plus transient events up to 1.4 MW m2 for about 1000 cycles (see Section 4.19.5.1.2).
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Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices
A neutron irradiation programme is still in progress to complete the full characterization with irradiated mock-ups (see Section 4.19.5.2). In the frame of the ITER Programme, in addition to Europe, the countries that in the past were interested to supply the ITER first wall were China, Korea, the Russian Federation, and the United States. The underlying development work which has recently been performed in these countries to establish the necessary fabrication capability is summarized elsewhere (see, e.g., Sherock et al.,163 Nishi et al.,164 Hong et al.,165 Youchison et al.,166 Lee et al.,167 Park et al.,168 Liu,169 and Chen170). 4.19.5.1.1.3 Fast brazing techniques
A development programme has been launched to develop a fast brazing technique to minimize the holding time at high temperature and consequently retain adequate mechanical properties of the CuCrZr alloy. This was achieved by induction brazing in Europe and by fast heating and cooling using an e-beam test facility in the Russian Federation. Induction brazing tests were done using the only appropriate silver free braze alloy available in the market, the STEMET 1108 procured from the Russian Federation. It was found that this braze alloy had poor wetting properties and the quality of the product was variable. Difficulties were met for brazing Be tiles of dimensions representative of the Be tiles of first-wall panels. A few first-wall mock-ups were fabricated with inductively brazed Be tiles but showed thermal fatigue performance well below HIPped mock-ups, with detachment of Be tiles between 1.5 and 2 MW m2.162,171 This result has been considered unsatisfactory. The development work on fast brazing equipment has been stopped in Europe and the effort is being concentrated on the development of a new silverfree braze alloy. The fast brazing development in the Russian Federation has also been done using the STEMET 1108 alloy but fast heating was performed using an e-beam test facility. First-wall mock-ups were heated on the beryllium side by the e-beam and temperatures as high as 780 C were achieved at the Be/CuCrZr joints for a very short time, followed by fast active cooling of the mock-ups, minimizing therefore the production of brittle intermetallic compounds.172 Good results were achieved on hypervapotron type mock-ups, developed at the early stage of the ITER design, for Be coated divertor components.129
4.19.5.1.2 High heat flux durability of unirradiated Be/Cu joints
Actively cooled first-wall mock-ups have been tested under relevant heat fluxes in electron beam facilities (accelerated fatigue tests) or in low-heat flux facilities, which are capable of delivering moderate heat flux (<1 MW m2) via electrical heaters for a very large number of heat pulses (>10 000) of larger duration to investigate the thermomechanical performance (including integrity and possible creep phenomena) of the bonded interfaces and fatigue lifetime. Because of intrinsic limitations, cost, or the limited availability of testing facilities, electron beam facilities are used with the shortest duration of the heating cycle, compatibly with the conditions of near thermal steadystate. In many cases, to reduce the testing time, the tests are performed at a heat flux well above the nominal one. These tests are often used to select the best joining conditions and then confirmation and final selection are done including tests at representative moderate heat fluxes but for the specified large number of cycles. Because of the health hazards linked to the use of beryllium material, only a limited number of test facilities were available or built for testing beryllium components. As far as Be-compatible electron beam facilities are concerned, two are available in Europe, the Juelich divertor test equipments in hot cells ( JUDITH 1 and JUDITH 2) at the Forschungszentrum Juelich (FZJ) (Germany), one in the Russian Federation at the Efremov Institute in St. Petersburg and one in the United States at Sandia National Laboratory in Albuquerque (New Mexico). For the performance of thermal fatigue tests at moderate heat fluxes and large number of cycles, three facilities were built in Europe, one at the Ispra Joint Research Center in Italy, one at the ENEA Research Center of Brasimone in Italy, and one at the Nuclear Research Institute (NRI) of Reˇz close to Prague in the Czech Republic. From these three facilities, only the last one is still under operation. The best high heat flux test results from representative first-wall mock-ups, namely, 10-mm beryllium tiles joined onto 10 mm CuCrZr heat sink layer with embedded 10/12 mm diameter stainless steel pipes and the assembly joined onto a stainless steel backing plate (see Figure 11), were achieved with Ti and Cu interlayers and HIP temperatures of about 580 C as described in Section 4.19.5.1.1.2. The performance limit of this assembly is presently at about 3 MW m2.162 As far as the thermal fatigue tests are concerned, representative first-wall mock-ups with Be/CuCrZr
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Figure 12 First-wall mock-up for thermal fatigue tests. Reproduced with permission from Lorenzetto, P.; et al. Fusion Eng. Des. 2008, 83, 1015–1019. Figure 11 First-wall mock-up for high heat flux tests.
joints made as above have been successfully tested up to 30 000 cycles at about 0.6 MW m2 (see Figure 12). Ultrasonic testing of the Be/CuCrZr joints after testing did not show any indication of defects. These mock-ups will be tested at higher heat fluxes representative of transient and off-normal events to check the available performance margins. Originally, six ITER Domestic Agencies were candidates for the procurement of the ITER first wall: China, European Union (EU), Japan, South Korea, Russian Federation, and United States. (The seven members of the international ITER project have all created Domestic Agencies to act as the liaison between national governments and the ITER Organization. The Domestic Agencies’ role is to handle the procurement of each member’s in-kind contributions to ITER.) However, as stated in the Final Report of Negotiations on ITER Joint Implementation of 1 April 2006, a prequalification ‘‘. . . will be needed for the critical procurement packages shared by multiParties. . .’’, such as the blanket first wall. Well in advance of the assumed start of the procurement, each ITER Domestic Agency shall first demonstrate its technical capability to carry out the procurement with the required quality, and in an efficient and timely manner. For the first wall system, this is achieved via a two-stage qualification process: a mock-up qualification stage and a semiprototype qualification stage. Each stage is also split into two phases: a manufacturing acceptance phase and a heat flux testing acceptance phase. The successful manufacturing and testing of two first wall mockups for stage I (see Figure 13) demonstrating in particular the know-how to assemble beryllium (Be)
Figure 13 ITER first-wall qualification mock-up, EU mock-up before testing.
tiles on a CuCrZr alloy and stainless steel bimetallic structure is the prerequisite to be eligible for stage II. The qualification tests for stage I have been split between the United States and EU in three test facilities: at the SNL for the US and at the NRI in the Czech Republic and the FZJ in Germany for the EU. At the SNL facility, the qualification test programme consists of the performance of 12 000 cycles at 0.88 MW m2 for 1.6 min followed by 1000 cycles at 1.4 MW m2, while in the EU test facilities it consists of the performance of 12 000 cycles at 0.62 MW m2 for 5 min (at NRI) followed by 1000 cycles at 1.75 MW m2 (at FZJ). To be qualified, a Domestic Agency shall fabricate two mockups which pass both tests. The first wall mock-ups fabricated by the EU Domestic Agency have successfully achieved the above test programme conditions without any indication of failure. Additional tests were also performed on these mock-ups to assess the limit and tests were performed for 200 cycles at 1.7 MW m2 at the SNL and for 100 cycles at 2.25 MW m2 plus 100 cycles at 2.75 MW m2 at FZJ without any indication of failure. Tests were stopped so as not to exceed the maximum acceptable Be temperature in the test facilities. The progress on the fabrication and thermal tests is described elsewhere.163–171
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4.19.5.2 Thermal Tests on Neutron-Irradiated Joints Thermal fatigue durability is a necessary but not sufficient prerequisite for the Be/CuCrZr alloy joint. Neutron irradiation is also expected to affect the high heat flux durability of these joints. This is particularly true for the first-wall components, where the joint will experience rather high neutron fluence. To investigate the effects of neutron irradiation on the joint properties there are two possible ways. The first method is postirradiation testing, that is, irradiation of small scale mock-ups in fission reactors and postirradiation testing in heat flux test facilities. The second method is to try to reproduce more faithfully the situation in ITER and to apply a cyclic heat flux during irradiation. This consists of in-pile heat flux testing of small mock-ups in fission reactors. Given the technical difficulty of achieving simultaneously representative values of heat flux and neutron irradiation and to determine exactly what the heat fluxes are within a reactor, it is suggested that the effect of neutron irradiation on the mock-ups be determined by pre- and postirradiation heat flux tests on mock-ups. This method is probably conservative as the postirradiation tests are performed on materials and joints having accumulated damage corresponding to the total neutron irradiation dose. However, this should be supported by analysis and material testing. Most of the tests conducted in the past were done for DS-Cu as a heat sink. Studies of neutron irradiation effects on the durability of the Be/Cu-alloy joints have been performed in at least two of the ITER parties: Europe,150,162 and the Russian Federation.173 For the Russian experiment, the irradiation conditions were 0.3 dpa at 350 C. The irradiated Be/Cu mock-ups were then tested in the JUDITH facility. The results of the postirradiation high heat flux testing of the different Be/Cu mock-ups are presented elsewhere.174 On the basis of the results of these tests, it was concluded that the effect of the neutron irradiation is not critical for the joints. Metallographic inspections did not show any significant changes in the braze joint after neutron irradiation. For the CuMnSnCe, a small intermetallic phase was observed in the middle of the braze layer for unirradiated and irradiated samples. No crack formation was found in the intermetallic. In addition to thermal fatigue tests, shear tests were conducted and it was found that for CuMnSnCe the shear strength decreased after neutron irradiation from 200 to
155 MPa, whereas for the InCuSil braze no irradiation influence was observed (300 MPa). Nevertheless, the thermal performance of the joints during high heat flux tests was very similar to the performance of unirradiated mock-ups. For the European experiments, the first irradiation campaign (named Paride 1 and Paride 2) of small scale Be/Cu mock-ups and Be/Cu joints took place in 1996–1999. Mock-ups were fabricated from a single 10-mm thick Be tile (grade S-65C) of dimensions 22 mm 60 mm, HIPped on a 20-mm thick CuAl25 substrate (grade IG1) with a drilled 10-mm diameter cooling channel (Figure 14). HIPping was done at 830 C for 2 h with the use of a 50 mm Ti interlayer. One mock-up was neutron irradiated in the test reactor high flux reactor (HFR) and then high-heat flux tested in JUDITH. The neutron irradiation was done at about 200 C up to a neutron dose of about 0.6 dpa in the Be material. The neutron dose expected at the end of life in the Be armor of the first-wall panels is about 1 dpa (Be). The irradiated mock-up was tested for 1000 cycles at 1.6 MW m2 plus 100 cycles at 1.9 MW m2 plus 100 cycles at 2.4 MW m2 plus 1000 cycles at 2.8 MW m2 plus 100 cycles at 3.3 MW m2 without any sign of failure. It failed during the first cycle at 4.25 MW m2 with a partial detachment of the Be tiles on one end (Figure 14). An unirradiated mock-up was tested for 1000 cycles at 1.5 MW m2 plus 1000 additional cycles at 3 MW m2 without any sign of failure but a detachment of the Be tile occurred during the first cycle at 4.5 MW m2. It was therefore concluded that no significant degradation of the Be/CuAl25 joint was observed up to a neutron dose of about 0.6 dpa. Neutron irradiation test experiments are ongoing or in preparation with Be coated first-wall mock-ups made from CuCrZr alloy to confirm the above result with this Cu alloy.
Figure 14 Be/CuAl25 mock-up after postirradiation high-heat flux test at 4.25 MW m2. Reproduced with permission from Lorenzetto, P.; et al. Fusion Eng. Des. 2006, 81, 355–360.
Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices
The first experiment was a joint European/Russian irradiation test campaign. It was prepared by the Efremov Institute of St. Petersburg. The original objective was to perform thermal fatigue testing of two first-wall mock-ups at about 0.5 MW m2 simultaneously to neutron irradiation. A failure of the surface heating system made from graphite after about 5000 cycles resulted in the discontinuation of the thermal fatigue testing and a continuation of the campaign with only neutron irradiation. The irradiation campaign has been stopped with the achievement of a neutron dose of 0.75 dpa. The first wall mock-ups, one made with Be tiles HIPped at 580 C and a Cu interlayer (Figure 15) and another with brazed Be tiles with STEMET 1108 braze alloy, will then be high heat flux tested together with unirradiated reference first wall mock-ups. Two other test campaigns are in preparation at the NRI of Reˇz (Czech Republic) and at Petten (The Netherlands) with the objective of thermal fatigue testing three first wall mock-ups in parallel to neutron irradiation. The question as to whether the correlation between fusion and fission neutron spectra assumed in many of the above measurements is valid or not needs to be discussed. Comparison of changes in the mechanical properties, especially at low temperature, needs to be made with the same He to dpa ratio to ensure that the results will be valid for ITER. Be/Cu alloy mock-ups have been tested in an electron beam for 1.5 s under a deposited energy density of 60 MJ m2(132) to simulate Be damage during a VDE. For a 6 mm thick Be tile, the melt layer was 1.5 mm, while that calculated for the same
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condition is 1.25 mm. A cross section of the Be tile after the simulated VDE event is shown in Figure 16. The embrittlement of the Be due to neutron irradiation increases the loss of material particles, especially at low irradiation temperature. A clear pore formation (which is expected to be He filled) has been observed in the melt layer of all neutron irradiated specimens after thermal shock loading (see Figure 16). An area of possible concern is the small surface cracks that form when molten metals resolidify. These resolidification cracks could serve as thermal fatigue crack initiation sites and accelerate this type of damage. While this effect has not been extensively studied because of the difficulty of simulating disruptions in the laboratory, it may not be a critical issue as thermal fatigue cracks form after a few hundred cycles in most materials and they grow only to depths where the thermal stress level is above the yield stress.175 High heat flux tests of neutron irradiated mockups conducted in the past did not reveal any damage in Be and in Be/Cu joints,176 although the irradiation conditions were not fully ITER relevant (damage dose of 0.3 dpa instead 1 dpa for the end of life and also lower He per dpa). An increase in crack formation and erosion rate has been observed in the surface of irradiated Be at 350 and 700 C.177 The S-65C grade presented the lowest damage after
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Figure 15 First-wall mock-up for irradiation test experiments. Reproduced with permission from Lorenzetto, P.; et al. Fusion Eng. Des. 2008, 83, 1015–1019.
Figure 16 Micrograph of S-65C armor on CuCrZr (CuMnSnCe braze) after vertical displacement event simulation. The actively cooled modules have been loaded with energy densities of 60 MJ m2 (effective pulse duration: 1 s). Reproduced with permission from Linke, J.; Duwe, R.; Gervash, A.; Qian, R. H.; Roedig, M.; Schuster, A. J. Nucl. Mater. 1998, 258–263, 634–639.
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irradiation. High heat flux tests (at more severe conditions than needed for the first wall) of the cracked unirradiated Be did not reveal any detrimental behavior and loss of material due to cracking.178 From an engineering point of view, to avoid possible crack formation and delamination of the brittle Be, it is recommended to use Be tiles without any stress concentrations.
4.19.6 Tokamak PFC Design Issues and Predictions of Effects in ITER During Operation 4.19.6.1
PFC Design Considerations
Robust engineering solutions are needed for the components that directly surround the plasma in order to withstand the thermal and mechanical loads during normal and off-normal operation and contribute to protect the outer components from the effect of neutrons, especially the vacuum vessel and the magnets. Considerations here are limited to the portion of the PFC surface that protects the main chamber (the so-called first wall), which for both JET and ITER is made of beryllium. The strategy and the criteria adopted to design the PFCs of JET and ITER are substantially different. One of the main differences is the longer plasma pulse duration foreseen in ITER that has required the use of water-cooled structures to handle the power in steady-state conditions. This translates into design solutions that consist of tiles bonded to an actively cooled copper-alloy substrate that must withstand the thermomechanical loads with a good engineering margin and achieve the desired fatigue lifetime. An additional problem in ITER is the degradation, albeit limited, of thermal and mechanical properties and the properties of the joints due to neutron irradiation (see Section 4.19.4.4). In contrast to ITER, technical constraints prevented the use of actively cooled first-wall structures in JET and the first-wall protection relies on a series of discrete poloidal limiters. The tiles are inertially cooled, and their power handling performance is driven by the need to (1) avoid surface melting and (2) reduce thermally induced stresses to give an adequate fatigue lifetime. At 40 mm typical thickness, the tiles are thermally thick for a typical 10 s pulse and will handle up to about 60 MJ m2 without melting. Another important driver for designing PFCs is the magnitude of the electromagnetic loads
associated with plasma disruptions. In tokamaks, disruptions produce large changes in magnetic field, dB/dt, which induce eddy currents in the conducting materials. The currents interact with the local magnetic field, B, to produce a torque, which is strongly dependent on the geometry. The design of the PFCs has to manage this torques via a combination of the castellations of the tiles along with cuts, which will interrupt the eddy current loops. The tile assembly must also withstand electromagnetic forces due to halo currents which, during disruptions, pass between plasma and vacuum vessel via the tiles. Enormous attention has been paid at JET and ITER during the design phase to address this problem. The problems associated with the design of the first wall at JET and ITER are briefly described in Section 4.19.6.1.1 and 4.19.6.1.2, respectively. 4.19.6.1.1 Design of the beryllium ITER-like wall at JET
JET has completed in 2011 a large enhancement programme that includes, among other things, the installation of a beryllium wall and a tungsten divertor. An overview of the status of the JET ITER-like wall project is presented in Matthews et al.179 The material combination chosen for the wall and the divertor is that chosen for the DT phase of ITER and experiments in JET with the new wall configuration will provide the first fully representative test of material migration, material mixing, and consequent tritium retention under ITER relevant conditions.180 Equally important is the opportunity to develop fully integrated scenarios and control schemes for protecting the wall. The project will therefore provide essential information for interpreting material behavior in ITER and a sound technical basis for guiding the development of ITER scenarios. The design layout, the main engineering challenges, and the operational limits of the JET ITER-like wall are discussed elsewhere (see, e.g., Nunes et al.,181 Thompson et al.,182 Riccardo,183 and Riccardo et al.184). Figure 17 describes the design layout and the planned material layout. It must be noted that the existing JET wall relies on a series of discrete poloidal limiters whereas at the moment ITER relies on a plasma conforming wall. The electrical resistivity of Be 0.08 mO m is more than a 100 times lower than that of CFC (10 mO m). Therefore, after replacing the CFC tiles in JET, the mechanical loads due to eddy currents associated with disruptions, which were negligible in the case of carbon tiles, have become dominant for Be tiles and this
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Figure 17 (a) View of the present Joint European Torus main chamber with indications of how the carbon-fiber composite tiles will be replaced (reproduced with permission from Riccardo, V. J. Nucl. Mater. 2009, 390–391, 895–899). (b) Cross-section with allocations of materials. Image courtesy of EFDA-JET.
has posed a significant engineering challenge. Calculations for the main limiter tile types clearly show that a tile of the size of the existing CFC tiles would give unacceptable eddy torques, leading to the investigation of several slicing options.181 The chosen design has vertical slices with a large central block and one to three side slices, depending on the toroidal extent of the tile assembly, supported on a carrier via pins (see Figure 18). The design is defined by the balance between conflicting requirements of eddy currents (avoidance of large low resistance loops) and power handling (minimum number of vertical cuts to be shadowed). The problems associated with the design of the JET beryllium tiles (power handling capacity and disruption induced eddy currents) are discussed in detail elsewhere (see, e.g., Thompson et al.182). The installation of the new ITER-like wall and the NB enhancement has been completed by the mid of 2011 and operation is now restarting to provide important information for ITER. 4.19.6.1.2 Design of the beryllium ITER wall
At the present time, the ITER first wall and shielding blanket is still undergoing a major redesign to overcome some of the main design shortcomings that were identified in the context of design review
conducted in 2007; for example, the thermal load requirements were updated, in light of experimental experience.16 Most important in this respect was the recognition that the upper X-point region would see much higher loads during burn than 0.5 MW m2; long transients (approximately up to 5–10 s) of plasma contact with the wall would have to be withstood. In addition, NB shine-through at low densities would necessitate high heat flux first-wall protection, and a new requirement has been introduced to provide remote maintainability of the first-wall panel to be done in situ and independently of the shield module (which would also have to be maintainable). The rationale for the ongoing effort is described by Lowry et al.156 Proposed design modifications are being developed while trying to avoid and minimize changes to other components which are on the critical fabrication path, especially the vacuum vessel, which is under fabrication. The main features of the proposed design are the following: (1) to abandon the port-limiters and to exploit the first wall for plasma startup by relying on more benign plasma start-up scenarios, including an early X-point formation; (2) to use suitably shaped plasma-facing surfaces to hide edges such that there is no illumination of component surfaces by
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‘Toast rack’ carrier
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Figure 18 Inner wall guard limiter tile (exploded view, top, and prototype, bottom). The five castellated Be slices have interslice and outer slice internal toroidal edges ski-slope shadowed. The slices are held on an inconel carrier by pins which allow bowing under thermal load. The RH bolts are designed to be shadowed by the next installed tile. Reproduced with permission from Riccardo, V. J. Nucl. Mater. 2009, 390–391, 895–899.
normal or near-normal field lines, emanating in the near SOL; (3) provide power load capability of 4–5 MW m2, in order to be able to use the first wall as a limiter for startup and termination; and (4) withstanding transients is still subject to discussion. In particular, it must be noted that the vulnerability to damage induced by thermal transients is recognized and linked to the feasibility and efficiency of all processes required for full remote maintenance of the first-wall panels, which is yet to be demonstrated. In practical terms, the approach adopted is to provide a shadowed poloidal band in the center of the first-wall panel, the two sides being shaped in a form typical for limiters both to provide the shadowing of the band and to ensure that the toroidally facing edge of the first-wall panel is shadowed. Because of the port regions on the low-field side, which contain a variety of structures with varying power handling capabilities, and because the toroidal field ripple is variable with the toroidal field, it is not possible to exploit the entire
wall surface in this location. For this reason, the firstwall panels on the low-field side have the poloidal bands between the ports advanced with respect to those in line with the ports (see Figure 19). The amount of set back required at the edges of the first wall is determined by the penetration angle of the field lines and the power scrape-off length, with the optimization taking into account the differing power handling capability of the front face and the edge of the first wall. Considerations discussed here are limited to some problems associated with the design of the beryllium tiles and prediction of PWI effects during operation in ITER. An important design driver for the first wall in the past was the specification of the thermal load during off-normal transient events.3 In particular, the thickness of the beryllium tiles had to be such as to prevent overheating of the joints and possible damage of the coolant pipes (see Section 4.19.6.2.2). Also, the thickness of the tile determines the temperature
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Figure 19 (a) View of the low field side first-wall surface showing how the panels in line with the port openings are recessed with respect to those between. It also shows the shielded central section of the panels allowing for access to the mechanical and hydraulic connections. Reproduced from Hawryluk, R. J.; et al. Nucl. Fusion 2009, 49, 15, 065012; with permission from IAEA. (b) Allocation of armor materials. Reproduced from Hawryluk, R. J.; et al. Nucl. Fusion 2009, 49, 15, 065012; Federici, G.; Loarte, A.; Strohmayer, G. Plasma Phys. Contr. Fusion 2003, 45, 1523–1547, with permission from IOPP.
gradient and the thermal stress under a prescribed thermal load during steady-state. Limits on the tile temperature during operation arise as a result of many processes including melting, excessive vaporization, thermal fatigue, reduced mechanical integrity, and chemical reactions during accidental exposure of armor or structure to air or steam. The last one of the above processes is important as explosion of hydrogen liberated from the steam–Be reaction is a major concern. In the past, a tile thickness of 10 mm was adopted. This corresponded to a Be maximum temperature limit of 650–750 C, roughly the level at which the relevant Be material properties (including mechanical, embrittlement, thermal fatigue, and swelling effects) start to degrade considerably. Because of the differences in the product of the elastic modulus and the coefficient of thermal expansion (E) between beryllium and copper or copper alloys (EBe/ECuCrZr ¼ 2.4), large thermal stresses are set up around the bond between the beryllium tile and copper allot heat-sink. The difficulty to successfully join low thermal expansion armor materials such as beryllium and tungsten to high thermal expansion heat sink materials has been a major problem and has been discussed in
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Section 4.19.5. Thermomechanical modeling has shown the desirability of using very small tiles of brush like structure for PFC armor because of the reduction of the stress at the armor–heat sink interface. The proper selection of the size of beryllium tile is an important issue which impacts all aspects of component manufacturing such as increased cost of machining, nondestructive examination features, reliability and repair of unbonded tiles, etc. In general, the fatigue life issue is difficult to quantify because of a number of factors. The thermal stresses depend on the temperature profile and the degree of constraint in the tiles. Tile castellations must be introduced to further relieve the constraints, and these have been sized following an extensive program of coupled thermal and mechanical analyses using finite elements codes such as ANSYS185 and ABAQUS.186 4.19.6.2 Predictions of Effects on the ITER Beryllium Wall During Operation 4.19.6.2.1 Safety issues in ITER 4.19.6.2.1.1
In-vessel tritium inventory
Estimates of the tritium inventory and of permeation in the PFCs of a magnetic fusion device are important for assessing the radiological hazards from routine operation and from potential accidents, for the design of the water detritiation system, and for predicting the tritium supply requirements. In addition, these estimates have contributed to the decisions involving the choice of different armor materials in ITER options, which have a strong impact on tritium retention. In spite of the experimental and modeling progress which has taken place in the recent past, understanding of the subject of tritium–wall interactions is still far from complete and quantification of the tritium inventory in ITER is highly uncertain. The retention and permeation of implanted tritium in ITER PFCs have been widely studied in the past (see Section 4.19.3 and the example of calculations found elsewhere).9,187–189 On the basis of the results of these calculations, it can be concluded that the inventory of tritium in the beryllium first wall of a device like ITER, due to implantation, diffusion, trapping, and neutron-induced transmutation, will be on the order of 100 g rather than the kilogram quantities estimated previously70,100 and most of that will come from neutron-induced transmutations in the Be itself. The dominant process for long-term retention of tritium in beryllium for ITER is expected to be
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codeposition (see Section 4.19.3.2.2) with eroded wall material (i.e., the incorporation of tritium in the deposited layers where impurity atoms or molecules are deposited together with eroded material and a flux of energetic or thermal atoms). The inventory of this potentially volatile tritium must be kept as low as reasonably achievable (1 kg tritium), in order to minimize the impact on the environment in case of an accidental release, in particular to avoid the evacuation of the neighboring population. The rate of formation of the codeposited material depends on the energy of the incident particles and on the substrate temperature during the deposition. In ITER, the total amount of tritium trapped in the codeposited layers will strongly depend on whether carbon is retained in the divertor during DT operation. But even in a full metal ITER configuration (e.g., with Be wall and W divertor) there is evidence for potential tritium accumulation for ITER in deposited Be layers.9 In contrast to carbon, tritium codeposition in beryllium layers is expected to be released at relatively low temperature and there are provisions to periodically bake the divertor in ITER at 350 C to release tritium trapped in the codeposited Be layers (see Section 4.19.3.2.2). 4.19.6.2.1.2 Chemical reactivity of beryllium dust with steam in ITER
Although not a concern in present day tokamaks, in-vessel dust and tritium inventories have been recognized as a safety and operational issue for next step devices such as ITER.190–193 In particular, accident scenarios that result in water or steam exposure of hot plasma-facing materials are one of the greatest concerns for ITER, because steam interacts with hot beryllium leading to the production of hydrogen, and hydrogen in the presence of air can lead to an explosion. The steam-chemical reactivity of different grades of Be has been studied extensively in the past.194–200 The amount of hydrogen produced depends on the specific material, temperature, exposure time, and especially the effective surface area. Because of the large surface area of dust, its chemical reactivity is an issue. Dust is expected to be produced inside the vacuum vessel of a tokamak by interaction of the plasma with the components of the first wall and the divertor. A detailed discussion of the mechanisms of dust production and of the influence of parameter variations is beyond the scope of this contribution, but it should be noted that the processes and the production rate of dust are not fully understood and the extrapolation
of knowledge from existing tokamaks to ITER is difficult. Research into dust production mechanisms and rates, the appropriate dosimetric limits for personnel exposure, and methods of removal has only recently begun.201,202 The location where the dust settles will determine its temperature, and consequently, its chemical reactivity. At the moment about 6 kg of C, 6 kg of W, and 6 kg of Be dust are allowed ‘on hot surfaces’ in ITER, with these limits set by the H production risk. This corresponds to the maximum allowable quantity of H (2.5 kg) for the vessel integrity to be guaranteed in case of explosion. A complete oxidation of Be at 400 C and C at 600 C is assumed for the calculation. If no C is present in the machine, the limits are relaxed to 11 kg for Be, or 230 kg for W. These quantities are set such that the overall hydrogen combustion limit is not exceeded.9 It must be recognized that a limit on the order of 10 kg for beryllium dust on ‘hot-surfaces’ is very restrictive, and in particular, the development of diagnostics techniques that can determine from local measurements the global inventory in the machine could prove to be very challenging.203 However, it is also likely that dust in ITER produced by Be eroded from the wall and deposited on the divertor will not survive on plasma-facing surfaces exposed to heat fluxes and will tend to accumulate in grooves or castellations in the armors of PFCs. They are an essential feature of the design of PFCs to relieve stresses during cyclic high heat flux loading, thus maximizing the fatigue lifetime of the armor to heat-sink joint. Some reduction in reaction rates is expected because the steam supply is not unlimited and steam must diffuse through the dust in the grooves. Experiments have been carried out in the Russian Federation, both in the Bochvar Institute of Moscow and the Efremov Institute of St. Petersburg.204 Although not conclusive, the main results summarized in Figure 20, show a reduction of the measured Be steam reactivity, particularly at high temperatures (more than a factor of 20). However, further experimental and modeling work is needed to clarify if the observed slower kinetics at high temperatures (800–900 C) eliminates the risk of explosion in the event of an accident. 4.19.6.2.2 Erosion/damage of the ITER Be wall
The erosion mechanisms that affect the erosion/ damage of the first wall in ITER are (1) sputtering erosion by D–T ions and charge-exchange neutrals
Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices
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Figure 20 Initial reactivity of Be powder in grooves with steam in recent experiments carried out at Efremov and Bochvar Institute in the Russian Federation. (Be powder: BET-0.38 m2 g1, average partial size ¼ 15 mm, free (nonpressed) dust density ¼ 0.7 g cm3).
during normal operation; and (2) evaporation and loss of melt layers during off-normal transient events such as thermal quench disruptions, ELMs, VDEs, and runaway electrons impact. There are additional localized erosion phenomena such as arcing, overheating with evaporation, and, possibly, loss of melt layer on exposed edges, but it is very difficult to make predictions of these effects for ITER. Special design attention has been given to avoid the misalignments of PFCs and avoid thermal overloading with possible localized damage. 4.19.6.2.2.1 Erosion of Be wall during normal operation
Calculations have been done to compute erosion of the first wall (due to fuel charge-exchange neutrals and ions, and impurity ions).205,206 It was found that about 20–40 g of Be per 400 s discharge are eroded from the wall with a beryllium peak erosion rate of the order of 0.1 nm s1. These predictions are confirmed by extrapolation of experimental data from JET.180 This erosion rate would be acceptable from a component lifetime standpoint, especially during the low duty-factor operation of ITER. However, the total amount of eroded material may be significant. This material will most likely go to the divertor, and this will affect the composition of the divertor surface; therefore, it will affect the divertor performance
and contribute to tritium codeposition and dust inventories. Modeling of the influx of the eroded beryllium on the divertor is in progress to extrapolate from present machines and, in particular, to account for effects arising from material mixing including codeposition as expected in ITER. Several studies have been recently published on this subject (see, e.g., Kirschner et al.207,208). 4.19.6.2.2.2 Erosion of the beryllium wall during ELMs
Depending on the actual energy flux on the Be PFCs in ITER during ELMs, melt damage may or may not occur. For Type I ELMs, which are compatible with the ITER divertor lifetime (10 MJ ELMs16,18), the expected energy flux on the main chamber in ITER will be in the range of 2–3 MJ. The area of the wall over which this flux will be distributed is 30–60 m2, for a toroidally symmetric energy deposition. This leads to ELM energy fluxes 0.02–0.08 MJ m2 on the main chamber wall, which will cause no Be melting at all. If toroidal asymmetries and/or poloidal structures dominate the ELM energy deposition on the first wall, a substantial reduction of the first-wall effective area for energy deposition is expected (by a factor of 5). In this case, the ELM energy fluxes on the first wall would be 0.1–0.4 MJ m2, which can cause up to 18 mm of
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melting, lasting 300 ms.209 Figure 21 shows the results of an analysis carried out with the code described in Raffray and Federici.210,211 The erosion lifetime, expressed in number of ELMs or corresponding ITER full power pulses (approximately 700 ELMs/pulses for a Be target initially 10 mm thick) is found to sharply decrease above a certain ELM energy threshold. Depending on the duration of the ELM event, the threshold energy density varies between 0.2 and 0.7 MJ m2. For comparison, the results of a W wall are also shown. More recently, analysis has been carried out using more sophisticated modeling tools and the results are described elsewhere.212 From the JET Be divertor experience, we expect that only a very small part of the melt layer produced during each ELM will be mobilized (typically <5%) and may lead to a Be influx into the plasma. Larger ELM energy fluxes onto the Be wall in ITER are indeed possible and would lead to serious problems for the use of Be as main plasma PFC in ITER, both because of lifetime issues and because of plasma contamination. However, for the arguments explained above, a regime with repetitive ELM energy loads which are not compatible with the lifetime of the Be main chamber wall in ITER is not compatible with the ITER divertor lifetime either and will not be the reference regime of ITER operation. The development of techniques that can either eliminate or greatly reduce ELM energy losses
without significantly degrading confinement have therefore been recognized to be critically important for successful operation of ITER and have stimulated further worldwide research on ELMs mitigation.213 4.19.6.2.2.3 Erosion of the beryllium wall during thermal quench disruptions
Thermal quench of a full-performance ITER plasma, with 350 MJ of thermal energy will result in significant transient heat loads causing vaporization and melting especially of divertor material. The erosion lifetime due to these events will depend on mitigating effects resulting from vapor shielding, redeposition of the eroded materials, and melt layer behavior. Disruptions could also result in significant Be erosion due to vaporization and possible loss of the melt layer. The evaporated and melt layer thicknesses are of the order of 10 mm and 50 mm, respectively, for the radiation energy density of 1 MJ m2 over 1 ms expected for the first wall under the assumption of no vapor shielding.214 Figure 22 compares the predicted melted layer thickness for a thermal quench time of 0.1 and 1 ms for beryllium and tungsten. For a disruption energy density of 1 MJ m2, we see that about 50 mm of Be are melted, as compared to 60 mm of tungsten. This result occurs even though beryllium melts at 1283 C, whereas the melting point of tungsten is 3410 C. The explanation for this result is that, under very intense energy deposition, a nearly instantaneous thermal balance is established between
107
1 : 0.1 ms 2 : 0.3 ms 3 : 0.5 ms 4 : 1.0 ms
105
104
103
100 1
3
4
2
Be W 10
3
Number of pulses
Number of fELMs
4
2 1
103
1
0.5
1.0 1.5 Energy density (MJ m-2)
2.0
Figure 21 Erosion lifetime, expressed in number of edge localized modes (ELMs) or corresponding ITER full power pulses (approximately several hundreds ELMs/pulse) for a Be target (initial thickness 10 mm) and for a W wall (initial thickness 10 mm) as a function of the ELM energy density.
Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices
120 Melt thickness (mm)
100 Be(2) 80 W(1) 60 40
Be(1) W(2)
20 0 0
(a)
1 1.5 0.5 Deposited energy (MJ m-2)
2
Evaporated thickness (mm)
20
15
10
5
Be(1)
W(1)
Be(2) W(2) ´ 10
0 (b)
0
0.5 1 1.5 Deposited energy (MJ m-2)
2
Figure 22 Calculated amount of material (a) melted and (b) evaporated during (1) a 0.1 ms and (2) a 1 ms plasma disruption for beryllium and tungsten.
the energy deposited by the plasma and cooling by vaporization of beryllium. The vaporization temperature of beryllium is variously reported as 2480– 2979 C, as compared to over 5630 C for tungsten. Similarly, the latent heats of melting and vaporization of Be are also lower than the corresponding tungsten values. This explanation is consistent with the results in Figure 22(b), which shows the amount of material that is vaporized for a thermal quench time of (1) 0.1 ms and (2) 1 ms. At an energy density of 1 MJ m2 and for a time of 0.1 ms the thickness of vaporized material is 10 mm for beryllium and 2.5 mm for tungsten. It must be noticed that vapor shielding is not included in these calculations and that the results therefore should be considered conservative. High-pressure noble-gas-jet injection, for example, of neon and argon, has shown to be a simple and robust method to mitigate the deleterious effects of disruptions in tokamaks.215 The gas jet penetrates the central plasma at its sonic velocity. The deposited species dissipate >95% of the plasma energy by radiation and substantially reduce mechanical stress on the vessel caused by poloidal halo currents. Nevertheless, there remains some concern that even mitigated disruptions could damage the Be wall
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in ITER. Preliminary calculations show that even during a mitigated disruption in which the plasma energy is intentionally dissipated by radiation in 1 ms by disruption mitigation techniques, the entire first wall of beryllium can melt to a depth of roughly 20–50 mm.212,216 The fate of this melted layer is uncertain. If the melt layer resolidifies, it provides a means of removing the oxide layer and creating a clean Be layer for oxygen gettering. On the other hand, if significant j B forces associated with the plasma termination mobilize the melt layer within the vessel, it will likely lead to operational difficulties. Another area of possible concern is the small surface cracks that form when molten metals resolidify. These resolidification cracks could serve as thermal fatigue crack initiation sites and accelerate this type of damage. While this effect has not been extensively studied because of the difficulty of simulating disruptions in the laboratory, it may not be a critical issue as thermal fatigue cracks form after a few hundred cycles in most materials and they grow to depths only where the thermal stress level is above the yield stress. 4.19.6.2.2.4 Erosion of the beryllium wall during VDEs
At some locations (mostly at the upper inboard region and the lower region of the first wall), the Be armored PFCs must withstand a certain number of ‘slow’ thermal transients resulting from loss-ofcontrol of plasma position during VDEs. Typical parameters for these events are 60 MJ m2 over 0.3 s. In contrast to thermal quench disruptions, VDEs lead not only to significant erosion or melting, but also to high heat fluxes and a subsequent temperature increase at the armor/heat sink interfaces that can result in a failure of the armor/heat sink joints.217 As a matter of fact, because of their short duration (<10 ms), ELMs and the thermal-quench phase of a disruption have no significant thermal effects on structural materials and coolant channels. In contrast, plasma instabilities such as VDEs (duration 100–300 ms), and runaway electron impact, in addition to causing severe surface melting and erosion, can result in substantial bulk damage to these components. Elevated temperatures and high thermal stresses in the structure can seriously degrade the integrity of the interface bonding and burnout the coolant channels. Runaway electrons (up to many megaelectron volt) penetrate many centimeters of beryllium and directly heat underlying metal structures, potentially damaging coolant channels.218–220
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Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices
The erosion due to VDEs in a device like ITER has been modeled by various authors,221 whereas the results of analysis carried out to quantify the effects on PFCs resulting from runaway electrons can be found in Raffray et al.222 As an example, Figure 23 shows the surface temperature of a 5 mm copper 1400 Tmett
10 mm W, Be, or CFC on a 5 mm Cu substrate
Copper surface temperature (K)
W 1200
VDE 60 MJ m-2 0.3 s
CFC 1000
800 Be 600
400
0
2
4
6
8
10
Time (s)
substrate at its interface with a tungsten, beryllium or carbon tile of 10 mm thickness during a typical VDE releasing about 60 MJ m2 to the surface in 300 ms.7 Tungsten and carbon armors of similar thickness usually result in a similar and higher copper surface temperature than that of beryllium armor of the same thickness. This is because most of the incident plasma energy is removed by the beryllium’s higher surface vaporization rate, which leaves little energy to be conducted through the structural material. In order to reduce the temperature at the copper interface, thicker tiles would be required. Only beryllium tiles of reasonable thickness (<5–10 mm) or very thick carbon or W tiles (>20 mm) can withstand the acceptable temperature rise in the copper structure for the conditions shown. The coolant flux and, consequently, the Be/Cu interface temperature increase with decreasing Be thickness. The evaporated and melting thickness and temperature at the Be/Cu alloy interface during each VDE is shown in Table 5 for Be tiles (5 and 10 mm thick); for two values of the VDE energy density (30 and MJ m2) and for two VDE durations (10 and 100 ms).
Plasma-facing component design
4.19.6.2.2.5 Erosion of the beryllium wall during runaway impact
Graphite, Be, or W
5–10 mm
5 mm
Copper substrate
(Carbon, beryllium, or tungsten coating on copper structure)
Figure 23 Interface copper surface temperature rise during a vertical displacement event for different surface coating materials. Reproduced with permission from Federici, G.; Skinner, C. H.; Brooks, J. N.; et al. Plasma-material interactions in current tokamaks and their implications for next-step fusion reactors. Nucl. Fusion 2001, 41, 1967–2137 (review special issue), with permission from IAEA.
Compared to disruptions, the thermal effects from runaway electrons are confined to a much smaller area, but the localized damage is expected to be more severe and can cause severe melting/vaporization in virtually all materials and can lead to surface spallation. These events have been observed to cause severe damage to graphite tiles in present day tokamaks. While the beryllium in the strike region will probably be severely melted, the most critical issues for runaway electron damage and VDE are damage of coolant pipes with resulting risk of water spillage. Because of the deep
Table 5 Vaporization and melting thickness (mm) and temperature at the Be/Cu alloy during each VDE (w/o any vapor shielding effect) Material/thick (mm) Energy density (MJ m2) Time (ms) 10 100 a
Beryllium/5 mm
Beryllium/10 mm
30 Vap. (mm) 4.03 3.03
60 Melta (mm) 309.9 821.6
Tjoimb ( C) 163 217.1
Vap. (mm) 8.44 7.5
Only stationary melt thickness (without splashing). Temperature at the interface armor/ heat sink.
b
30 Melta (mm) 648.2 812.3
Tjoimb ( C) 163 233.1
Vap. (mm) 4.04 3.08
60 Melta (mm) 318.6 861.5
Tjoimb ( C) 169.6 170.2
Vap. (mm) 8.46 7.56
Melta (mm) 652 870.2
Tjoimb ( C) 169.6 170.5
Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices
penetration and large spatial dispersion of the highenergy electrons, a thick armor may be required to avoid overheating of the coolant channels with subsequent coolant leakage. As thicker armor implies higher surface temperatures, the best solution may be local regions that are either uncooled or with thick armor that receives low heat flux during normal operations. Typically a runaway electron energy deposition transient of 50 MJ m2 over 0.3 s on the Be first-wall modules results in a maximum heat flux to the coolant of 7.4 MW m2, a maximum Cu alloy temperature 640 C, and a Be melt layer thickness 1.8 mm.222 A reduction of the armor thickness will lead to an increase in the maximum Cu alloy temperature and could lead to the damage of joints. 4.19.6.3 Prospect of Using Beryllium in Beyond-ITER Fusion Reactors The main differences between a future power reactor and ITER are the much longer operation time (e.g., >108 s vs. >107 s), high duty cycle, and the higher temperature of the fluid to cool the PFCs to maintain a high plant energy conversion efficiency. The higher surface temperature of the PFC will affect particle recycling, tritium uptake, chemical erosion, and material-mixing effects. Erosion rates at the divertor target are very difficult to predict in conventional fusion power plant concepts with solid high-Z targets because the net erosion or deposition is strongly dependent on plasma parameters. The fraction of ions arriving above the sputtering threshold is crucial, as is the efficiency of the prompt local redeposition. ELMs have not really been considered in this context but we can see from the analysis in Section 4.19.6.2.2.2 that ELMs in power plant systems will have to be extremely small – much smaller than will be allowable in ITER, which still has a relatively low duty cycle. The ideal would in fact be a quiescent ELM free high density steady state edge plasma. Calculations of the minimum erosion rate for the main wall are somewhat more robust as there has to be a hot plasma in the main chamber and the rate of leakage of neutrals into the main chamber from the divertor can be calculated using Monte-Carlo codes. In a recent study, Be, C, Fe, Cu, Mo, and W walls were compared,206 with the conclusion that in all cases the erosion rate was 1–2 t per year of continuous operation. A Be or CFC wall will erode too rapidly in a reactor and the large amount of eroded material might give rise to deleterious problems as far as
659
control of the tritium and dust inventories are concerned. A medium-Z material, such as Fe, does not seem to be acceptable purely from the standpoint of erosion lifetime. As molybdenum is unfavorable for long-term activation problems, W is the best and only solution we have available for a reactor. Effects of plasma contamination from Mo and W at the wall of tokamaks are being addressed in Alcator C-Mod, ASDEX-Upgrade and in the near-future at JET. There is considerable gross erosion by sputtering for all materials. The contributions of ions and neutrals from the plasma to this erosion are of the same order of magnitude. The integrated total erosion due to ions and the energetic neutrals for the different wall materials (Figure 24) show that because of the larger sputtering yields for the low-Z materials, the number of atoms eroded for these materials is a factor of 10–20 larger than for high-Z materials such as W. However, the total mass loss is similar for all materials, up to several kilograms per day or about 1 t per year. The maximum wall thinning for the low-Z materials is about 3.5 mm year1, while for high-Z materials, such as W, it is 0.22 mm year1, that is, lower by about a factor of 15. These values are in reasonable agreement with erosion measurements at the JET vessel walls.223 With respect to wall thinning, W is favorable for the use at the vessel walls because it has the longest ‘erosion lifetime’ (Figure 24(b)). With respect to plasma contamination, the probability of the eroded atoms entering into the plasma core, their lifetime in the plasma core, and the tolerable concentration of these ions in a burning fusion plasma all have to be taken into account.206 The tolerable concentration of W in the plasma is nearly three orders of magnitude lower than for low-Z atoms, such as Be and C. However, recent observations have shown that W can be effectively removed from the plasma center by central heating.224 As this central heating is natural for burning plasmas, W may be a possible plasmafacing material, even from the viewpoint of plasma contamination. The ion and neutral flux densities on the vessel walls are of the order of 1020 m2 s1, which may be critical with respect to the tritium implantation, accumulation in and permeation through the vessel walls.
4.19.7 Concluding Remarks Beryllium is a low-density metal that is used in a number of industries, including the nuclear, automotive, aerospace, defense, medical, and electronics
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Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices
25
(a)
Time (year) to erode 5 mm
Wall erosion (atoms s-1)
1022 Be Ctot Cchem
1021
Cphys
Cu Fe Mo
W 1020
20
W
15 10 5
Mo Be C
Cu
0 0
10 20 30 40 50 60 70 Atomic number of the material
80
0
(b)
Fe
10 20 30 40 50 60 70 Atomic number of the material
80
Figure 24 (a) Integrated gross erosion due to ions and the energetic neutrals. (b) Upper estimate for the time until a thickness of 5 mm is eroded by sputtering at the area of largest erosion, that is, at the reference distance of about 11 m. Reproduced with permission from Behrisch, R.; Federici, G.; Kukushkin, A.; Reiter, D. J. Nucl. Mater. 2003, 313–316, 388–392.
industries, for various applications because it is exceptionally strong, is light in weight compared with other metals, has high heat-absorbing capability, and has dimensional stability in a wide range of temperatures. Beryllium has been considered for many years as a primary candidate for protection of PFCs in tokamaks because it offers distinct advantages when compared with alternative materials such as carbon and tungsten. It has a low atomic number and is an excellent oxygen getter. The interaction of beryllium with tritium is also significantly weaker than that of carbon, leading to potentially reduced tritium inventory. Beryllium does not form stable hydrides above 300 C, so there should be very little trapping expected in codeposited layers formed at such temperatures in the divertor after sputtering, although work is still underway to clarify this problem. However, beryllium has a relatively high physical sputtering rate and a relatively low melting temperature and as such is more susceptible to melting damage that may occur in a tokamak during thermal transients. In addition, because of its toxicity, special precautions are needed for working with beryllium, either for manufacturing or research investigation purposes. Beryllium has been used with success in various tokamaks in the past mainly because of its ability to getter oxygen and to improve plasma performance. In particular, its successful deployment in JET that started in 1989 and is continuing today with the installation of a completely new beryllium wall is the main rationale for the selection of beryllium as a plasma-facing material for the first wall of ITER, on the basis of a combination of plasma compatibility and design considerations.
This paper reviewed the properties of beryllium that are of primary relevance for plasma protection applications in magnetic fusion devices (i.e., PWIs, thermal and mechanical properties for power handling, fabricability and ease of joining, chemical reactivity, etc.) together with the available knowledge on performance and operation in existing fusion machines. Special attention was given to beryllium’s erosion and deposition, formation of mixed materials, and its hydrogen retention and release characteristics. These phenomena have a profound impact on component design, machine operation, and safety. Extensive data on the behavior of Be with plasmas have been collected from existing tokamaks and simulators during the last two decades and this has enabled great strides to be made in our understanding of the PWI processes involved. However, there are many issues for which there are still uncertainties and we will only learn from operating the next two major experiments that foresee the use of large amounts of Be ( JET and ITER). Much work remains to be done in this area and more machine operational time and diagnostics dedicated to PWIs are required. Initiatives on these fronts, together with modeling of the results, are essential to advance the understanding of PWIs. This includes (1) the possible surface damage (melting) during transients such as ELMs and disruptions and its implications for operations and (2) the problem of beryllium mixing with other armor materials and in particular the issue of codeposition of tritium with Be, which is eroded from the first wall and deposited at the divertor targets. Such material may also be locally redeposited into shadowed areas of the shaped ITER first wall. Both issues are part of
Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices
ongoing research, the initial results of which are being taken into account in the ITER design so that the influence of these two factors on ITER operation and mission are minimized. For example, ITER will very likely employ, ELM control systems based on pellets and RMP coils, disruption mitigation systems, and increased temperature baking of the divertor to release T from Be-codeposits. Dust generation is still a process which requires more attention. Conversion from gross or net erosion to dust and the assessment of dust on hot surfaces need to be investigated. At the time of writing this paper, the ITER first wall and shielding blanket is undergoing a major redesign effort to overcome some of the main shortcomings that were identified in the context of a recent design review scrutinizing the internal components. Complex and interrelated materials, manufacturing, and design issues were briefly reviewed in this paper together with the progress of the manufacturing technologies being used and tested to demonstrate the durability of the joints. A critical feature of the ITER first-wall design is the beryllium to copper alloy bond. The joints must withstand the thermal, mechanical, and neutron loads and the cyclic mode of operation, and operate under vacuum, while providing an acceptable design for lifetime performance and reliability. The availability of reliable joining technologies has a large impact on the design of the PFCs and on the overall cost of these components. The status of the available techniques presently considered to join the Be armor to the heat sink material of Cu alloys for the fabrication of Be-clad actively cooled components for the ITER first wall was discussed. During earlier ITER design phases, the feasibility of manufacturing reliable Be–CuCrZr joints was demonstrated. The results of the performance and durability heat flux tests conducted in the framework of the further ITER first-wall qualification program were described. This program has been launched and is in progress in the ITER parties in order to qualify the design and manufacturing routes. The integrity of this bond must be assured for reliable ITER performance whatever process is used to fabricate joints. The original procurement sharing that assigned the fabrication of first-wall panels up to six parties was seen as a risk and the number of parties supplying these critical components has now been reduced to three, Europe, the Russian Federation, and China. The selection of specific grades of specific beryllium for the ITER first wall was described. The
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effects of neutron irradiation on the degradation of the properties of beryllium itself and on the joints were also analyzed. Some of the changes are important while others are not significant for the ITER conditions. Change of thermal conductivity and swelling are not important because of the low fluence. The bulk tritium retention in neutron irradiated Be is expected to be significantly less than tritium retention in the codeposited layers. The most critical consequence of neutron irradiation under ITER conditions is embrittlement. This is typical of all grades of beryllium. The structural integrity of neutron irradiated brittle Be is a key issue. Embrittlement of neutron-irradiated Be could lead to increased thermal erosion and crack formation, which is also observed to occur for unirradiated beryllium under severe transient heat loads. These cracks could serve as thermal fatigue crack initiation sites and accelerate this type of damage. While this effect has not been extensively studied because of the difficulty of simulating disruptions in the laboratory, it may not be a critical issue as thermal fatigue cracks form after a few hundred cycles in most materials and they grow only to depths where the thermal stress level is above the yield stress. On the basis of the information available from existing fusion machines, we discussed the problems that are still at issue in the design and operation of ITER. This includes, in particular, the problem of erosion/ damage and the problem of up-take and control of tritium in the beryllium-based codeposited films. Finally, on the basis of these results some tentative and speculative consideration of the limited prospects that beryllium has in future reactors was offered. The worldwide fusion energy research over the last four decades has developed a tremendous amount of knowledge on plasma physics and related technologies. From this point of view, collecting the latest information from a wide range of studies is important in order to help the fusion community to recognize the critical issues and the status. That has been the intent of this chapter. (See also Chapter 4.17, Tungsten as a Plasma-Facing Material and Chapter 4.18, Carbon as a Fusion Plasma-Facing Material).
Acknowledgments The views expressed in this publication are the sole responsibility of the authors and do not necessarily reflect the views of Fusion for Energy. Neither Fusion for Energy nor any person acting on behalf
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Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices
of Fusion for Energy is responsible for the use which might be made of the information in this publication. The authors from the ITER Organization wish to acknowledge that this paper was prepared as an account of work by or for the ITER Organization. The Members of the Organization are the People’s Republic of China, the European Atomic Energy Community, Republic of India, Japan, Republic of Korea, the Russian Federation, and the United States. The views and opinions expressed herein do not necessarily reflect those of the Members or any agency thereof. Dissemination of the information in this paper is governed by the applicable terms of the ITER Joint Implementation Agreement.
25. 26. 27. 28.
29.
30. 31.
References 1. 2. 3. 4. 5. 6. 7.
8. 9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19. 20. 21. 22. 23. 24.
Smith, E. A. Financ. Anal. J. 1960, 16(6), 51–54. Wilson, K. L.; Causey, R. A.; Hsu, W. L.; Mills, B. E.; Smith, M. F.; Whitley, J. B. J. Vac. Sci. Technol. 1990, A8(3), 1750–1759. Raffray, A. R.; Federici, G.; Barabash, V.; et al. Fusion Eng. Des. 1997, 37, 261–286. Proceedings of the IEA International Workshop on Beryllium Technology for Fusion (1990–2000). Hirooka, Y. (guest ed.) Beryllium studies for fusion – Part 1 (special issue). Fusion Eng. Des. 1997, 37, 225. Hirooka, Y. (guest ed.) Beryllium studies for fusion – Part 1 (special issue). Fusion Eng. Des. 1997, 37, 451. Federici, G.; Skinner, C. H.; Brooks, J. N.; et al. Plasma-material interactions in current tokamaks and their implications for next-step fusion reactors. Nucl. Fusion 2001, 41, 1967–2137 (review special issue). Federici, G. Phys. Scripta 2006, T124, 1–8. Roth, J.; et al. J. Nucl. Mater. 2009, 390–391, 1–9. Loarte, A.; et al. Nucl. Fusion 2007, 47, S203–S263. Lipschultz, B.; et al. Nucl. Fusion 2007, 47, 1189–1205. Federici, G.; Loarte, A.; Strohmayer, G. Plasma Phys. Contr. Fusion 2003, 45, 1523–1547. Eich, T.; et al. Plasma Phys. Contr. Fusion 2005, 47, 815–842. Fenstermacher, M. E.; et al. Nucl. Fusion 2005, 45, 1493–1502. Maingi, R.; et al. Phys. Plasma. 2006, 13(9), 092510–092513. Loarte, A.; et al. In Proceedings of the 22nd International Conference on Fusion Energy, Geneva, Switzerland, 2008; IAEA: Vienna, Austria, 2008; IT/P6-13. Hawryluk, R. J.; et al. Nucl. Fusion 2009, 49, 15, 065012. Loarte, A.; et al. J. Nucl. Mater. 2006, 337–339, 816–820. Behrisch, R. Nucl. Fusion 1972, 12, 695–713. McCracken, G. M.; Stott, P. E. Nucl. Fusion 1979, 19, 889–981. Gorbunov, E. P.; et al. In 4th European Conference on Controlled Fusion and Plasma Physics, Rome, Italy, 1970; Nat. Nucl. Energy Committee 19. Peacock, N. J.; Robinson, D. C.; Forrest, M. J.; Wilcock, P. D.; Sannikov, V. V. Nature 1969, 224, 488–490. Ellis, R. A. Nucl. Fusion 1985, 25, 1145. Zuhr, R. A.; Clausing, R. E.; Emerson, L. C.; Heatherly, L. J. Nucl. Mater. 1979, 85–86, 979.
32.
33. 34. 35. 36. 37. 38. 39. 40. 41. 42. 43. 44. 45. 46. 47. 48. 49.
50. 51.
52.
53.
Ascoli-Bartoli, U.; et al. In 5th International Conference on Plasma Physics Controlled Nuclear Fusion Research; IAEA: Vienna, Austria, 1975; Part I 19. TFR Group. In Plasma–Wall Interactions, Proceedings of the International Symposium, Julich, 1976; Pergamon: Oxford, 1977; Vol. 3. Dimock, D.; et al. In 4th Conference on Plasma Physics and Controlled Thermonuclear Fusion; IAEA: Vienna, Austria, 1971; Part 1 45. Bol, K.; et al. In Plasma Physics Controlled Nuclear Fusion Research, Proceedings of the 7th International Conference, Innsbruck, 1978; IAEA: Vienna, Austria, 1979; Vol. 1, p 11. Meade, D.; et al. In Plasma Physics Controlled Nuclear Fusion Research, Proceedings of the 8th International Conference, Brussels, 1980; IAEA: Vienna, Austria, 1981; Vol. 1, p 665. Stangeby, P. The Plasma Boundary of Magnetic Fusion Devices; Chapters 5 and 6, Book ISBN 0 07503 0559 2. Meservey, E. B.; et al. J. Nucl. Mater. 1980, 93–94, 267–271. Eubank, H.; et al. In Proceedings of the 7th Conference on Plasma Physics Controlled Nuclear Fusion Research Held by the IAEA, Innsbruck, Austria, 1978; IAEA: Vienna, Austria, 1979; Vol. 1, p 167. Burchell, T. D. Carbon 1996, 34, 297–316. The JET Team. J. Nucl. Mater. 1990, 176–177, 3–13. Hutchinson, I. H.; et al. Phys. Plasma. 1994, 1, 1511–1518. Neu, R.; et al. Plasma Phys. Contr. Fusion 2007, 49(12B), B59–B70. Pamela, J. J. Nucl. Mater. 2007, 363–365, 1–11. Bak, J. S. In Proceedings of the 22nd International Conference on Fusion Energy, Geneva, Switzerland, 2008, IAEA: Vienna, Austria, 2008; IT/FT/1-1. Wan, B. In Proceedings of the 22nd International Conference on Fusion Energy, Geneva, Switzerland, 2008, IAEA: Vienna, Austria, 2008; OV/3-4. McCracken, G.; et al. Phys. Plasma. 1997, 4, 1681–1689. Lipschultz, B.; Pappas, D. A.; LaBombard, B.; Rice, J. E.; Smith, D.; Wukitch, S. J. Nucl. Fusion 2001, 41, 585–596. Strachan, J.; et al. Nucl. Fusion 2003, 43, 922–941. Kallenbach, A.; et al. Plasma Phys. Contr. Fusion 2005, 47(12B), B207–B222. Dux, R.; et al. J. Nucl. Mater. 2009, 390–391, 858–863. Mioduszewski, P. K. Fusion 1986, 26, 117. Hackmann, J.; Uhlenbusch, J. J. Nucl. Mater. 1984, 128–129, 418–421. Thomas, P. R. JET Team. J. Nucl. Mater. 1991, 176–177, 3–13. Deksnis, E. B.; Peacock, A. T.; Altmann, H.; Ibbot, C.; Falter, H. D. Fusion Eng. Des. 1997, 37, 515–530. Lomas, P. J. The JET Team. In Plasma Physics Controlled Fusion Research, Proceedings of the 14th International IAEA Conference, Wuerzburg, Germany, 1992; IAEA: Vienna, Austria, 1993; Vol. 1, p 181. Campbell, D. J. The JET Team. J. Nucl. Mater. 1997, 241–243, 379–384. Tubbing, B. J. D.; et al. In Proceedings of the 22nd European Physical Society Conference on Controlled Fusion and Plasma Physics, Bournemouth, 1995; European Physical Society: Geneva, 1995; Vol. 19C, p 453, Part 3. Wesson, J. The Science of JET. The achievements of the scientists and engineers who worked on the Joint European Torus. 1973–1999. Published March 2000. http://www.jet.efda.org/documents/books/wesson.pdf. Biersack, J. P.; Eckstein, W. Appl. Phys. 1984, A34, 73.
Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices 54. 55. 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. 82. 83. 84.
Eckstein, W. Calculated Sputtering, Reflection and Range Values; IPP Report, IPP 9/132; Max-Plank-Institut fur Plasmaphysik: Garching, 2002. Roth, J.; Eckstein, W.; Guseva, M. Fusion Eng. Des. 1997, 37, 465–480. Roth, J.; Wampler, W. R.; Jacob, W. J. Nucl. Mater. 1997, 250, 23–28. Doerner, R. P. In Nuclear Fusion Research – Understanding Plasma-Surface Interactions; Clark, R. E. H., Reiter, D. H., Eds.; Springer: 2005; Chapter 14, pp 335–355, Book ISBN 3-540-23038-6. Guo, H. Y.; et al. J. Nucl. Mater. 1997, 241–243, 385–389. Won, J.; Spada, F. E.; Boivin, R.; Doerner, R. P.; Luckhardt, S.; Sze, F. C.; Conn, R. W. J. Nucl. Mater. 1997, 241–243, 1110–1116. Doerner, R. P.; Grossman, A.; Luckhardt, S.; et al. J. Nucl. Mater. 1998, 257, 51–58. Nishijima, D.; Doerner, R. P.; Baldwin, M. J.; De Temmerman, G. J. Nucl. Mater. 2009, 390–391, 132–135. Doerner, R. P. J. Nucl. Mater. 2007, 363–365, 32–40. Hirooka, Y.; Won, J.; Boivin, R.; Sze, D.; Neumoin, V. J. Nucl. Mater. 1996, 228, 148–153. Doerner, R. P.; Grossman, A. A.; Luckhardt, S.; Seraydarian, R.; Sze, F. C.; Whyte, D. G. J. Nucl. Mater. 1999, 266–269, 392–398. Goldstrauss, P.; Linsmeier, Ch. J. Nucl. Mater. 2001, 290–293, 71–75. Goldstrauss, P.; Klages, K. U.; Linsmeier, Ch. J. Nucl. Mater. 2001, 290–293, 76–79. Schmid, K.; Baldwin, M.; Doerner, R. P.; Wilthner, A. Nucl. Fusion 2004, 44, 815–819. Baldwin, M. J.; Doerner, R. P.; Nishijima, D.; et al. J. Nucl. Mater. 2006, 358, 96–105. Nishijima, D.; Baldwin, M.; Doerner, R. P.; Seraydarian, R. J. Nucl. Mater. 2007, 363–365, 1261–1265. Baldwin, M. J.; Doerner, R. P. Nucl. Fusion 2006, 46, 444–450. Duxbury, G.; Stamp, M. F.; Summers, H. P. Plasma Phys. Contr. Fusion 1998, 40, 361–370. Nishijima, D.; Doerner, R. P.; Baldwin, M. J.; De Temmerman, G.; Hollmann, E. M. Plasma Phys. Contr. Fusion 2008, 50, 125007. Doerner, R. P.; Baldwin, M. J.; Buchenauer, D.; De Temmerman, G.; Nishijima, D. J. Nucl. Mater. 2009, 390–391, 681–684. Bjorkas, C.; Vortler, K.; Nordlund, K.; Nishijima, D.; Doerner, R. P. New J. Phys. 2009, 11, 123017. Nelson, R. S. Philos. Mag. 1965, 11(110), 291–302, 1478-6443. Doerner, R. P.; Krasheninnikovv, S. I.; Schmid, K. J. Appl. Phys. 2004, 95, 4471–4475. Doerner, R. P.; Baldwin, M. J.; Conn, R. W.; et al. J. Nucl. Mater. 2001, 290–293, 166–172. Vaulin, E. P.; Georgieva, N. E.; Martynenko, T. P.; et al. Sov. J. Plasma Phys. 1981, 2, 437. Roth, J.; Moeller, W. Nucl. Instrum. Methods Phys. Res. B Beam Interactions Mater. Atoms 1985, 7–8(Pt 2), 788–792. Philipps, V.; Vietzke, E.; Trinkaus, H. J. Nucl. Mater. 1991, 179–181, 25–33. Doerner, R. P.; Baldwin, M. J.; Krasheninnikov, S. I.; Schmid, K. J. Nucl. Mater. 2005, 337–339, 877–881. Wilson, K. L.; et al. In Atomic and Plasma-Material Interaction Data for Fusion; Supplement to Nuclear Fusion; IAEA: Vienna, Austria, 1991; Vol. 1, p 31. Zakharov, A. P.; Gorodetsky, A. E.; Alimov, V. Kh.; Kanashenko, S. L.; Markin, A. V. J. Nucl. Mater. 1997, 241–243, 52–67. Anderl, R. A.; et al. J. Nucl. Mater. 1999, 273, 1–26.
85. 86. 87. 88. 89. 90. 91. 92.
93. 94. 95. 96. 97.
98. 99. 100. 101. 102. 103.
104. 105. 106. 107. 108. 109. 110. 111. 112. 113. 114. 115. 116. 117.
663
Causey, R. A.; Venhaus, T. J. Phys. Scripta 2001, T94, 9–15. Serra, E.; Benanati, G.; Ogorodnikova, O. V. J. Nucl. Mater. 1998, 255, 105–115. Frauenfelder, R. J. Vac. Sci. Technol. 1969, 6, 388–397. Shapovalov, V. I.; Dukel’ski, Yu. M. Russ. Metall. 1988, 5, 210. Jones, P. M. S.; Gibson, R. J. Nucl. Mater. 1967, 21, 353–354. Swansiger, W. A. J. Vac. Sci. Technol. 1988, A4, 1216–1217. Abramov, E.; Riehm, M. P.; Thompson, D. A.; Smeltzer, W. W. J. Nucl. Mater. 1990, 175, 90–95. Tazhibaeva, I. L.; et al. In Proceedings of the 18th Symposium on Fusion Technology, Karlsruhe, Germany, 1994; Elsevier Science: B.V. Amsterdam, 1994; Vol. 1, p 427. Haasz, A. A.; Davis, J. W. J. Nucl. Mater. 1997, 241–243, 1076–1081. Wampler, W. R. J. Nucl. Mater. 1984, 122–123, 1598–1602. Chernikov, V. N.; Alimov, K. Kh.; Markin, A. V.; Zakharov, A. P. J. Nucl. Mater. 1996, 228, 47–60. Altstetter, C. J.; Behrisch, R.; Scherzer, B. M. U. J. Vac. Sci. Technol. 1978, 15, 706–709. Moeller, W.; Scherzer, B. M. U.; Bohdansky, J. Retention and Release of Deuterium Implanted into Beryllium; Final Report – IPP-JET-Report-No.26; Max-Planck-Institut fuer Plasmaphysik: Garching, 1985. Haasz, A. A.; Davis, J. W.; Poon, M.; MacaulayNewcombe, R. G. J. Nucl. Mater. 1998, 258–263, 889–895. Longhurst, G. R.; et al. J. Nucl. Mater. 1998, 258–263, 640–644. Longhurst, G. R.; Anderl, R. A.; Dolan, T. J.; Mulock, M. J. Fusion Technol. 1995, 28, 1217–1222. Federici, G.; Holland, D.; Janeschitz, G.; Wu, C. H. J. Nucl. Mater. 1997, 241–243, 260–267. Mayer, M. J. Nucl. Mater. 1997, 240, 164–167. Causey, R. A.; et al. In Proceedings of International Workshop on Present Status and Prospect of Tritium Material Interaction Studies, Toyama, Japan, July 18–19, 1996. Causey, R. A.; Walsh, D. S. J. Nucl. Mater. 1998, 254, 84–86. Baldwin, M. J.; Schmid, K.; Doerner, R. P.; Wiltner, A.; Seraydarian, R.; Linsmeier, Ch. J. Nucl. Mater. 2005, 337–339, 590–594. Reinelt, M.; Allouche, A.; Oberkofler, M.; Linsmeier, Ch. New J. Phys. 2009, 11, 19, 043023. De Temmerman, G.; et al. Nucl. Fusion 2008, 48, 7, 075008. De Temmerman, R. P. Nucl. Fusion 2009, 49, 3, 042002. Doerner, R. P.; Baldwin, M. J.; De Temmerman, G.; et al. Nucl. Fusion 2009, 49, 6, 035002. Doerner, R. P.; Luckhardt, S. C.; Seraydarian, R.; Sze, F. C.; Whyte, D. G. Phys. Scripta 1999, T81, 35–39. Nishijima, D.; Baldwin, M. J.; Doerner, R. P.; Seraydarian, R. J. Nucl. Mater. 2007, 363–365, 1261–1265. Roth, J. J. Nucl. Mater. 1999, 266–269, 51–57. Hanna, J.; Baldwin, M. J.; Doerner, R. P.; Nishijima, D.; Seraydarian, R. J. Nucl. Mater. 2009, 386–388, 756–759. Doerner, R. P.; Baldwin, M. J.; Causey, R. A. J. Nucl. Mater. 2005, 342, 63–67. Watts, R. Int. J. Powder Metall. 1968, 4, 49. Vasina, E. A.; Panov, A. S. Russ. Metall. 1974, 1, 119. Linsmeier, Ch.; Ertl, K.; Roth, J.; et al. J. Nucl. Mater. 2007, 363–365, 1129–1137.
664
Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices
118. Baldwin, M. J.; Doerner, R. P.; Nishijima, D.; et al. J. Nucl. Mater. 2007, 363–365, 1179–1183. 119. Brooks, J. N.; Allain, J. P.; Doerner, R. P.; et al. Nucl. Fusion 2009, 49, 6, 035007. 120. Beryllium Science and Technology, Vol. 1, Webster, D., London, G. J., Eds.; Vol. 2, Floyd, D. R., Lowe, J. N., Eds.; Plenum: New York, 1979. 121. Barabash, V.; Tanaka, S.; Matera, R. Beryllium assessment and recommendation for application in ITER plasma-facing components. In Proceedings of the 3rd IEA International Workshop on Beryllium Technology for Fusion, Mito City, Japan, Oct 22–24, 1997; pp 1–13, JAERI-Conf 98-001. 122. Barabash, V.; Federici, G.; Matera, R.; Raffray, A. R. Phys. Scripta 1999, T81, 74–83. 123. Cardella, A.; Barabash, V.; Ioki, K.; Yamada, M. Application of beryllium as first wall armour for the ITER primary baffle and limiter blanket modules. In Proceedings of the 4th IAE International Workshop on Beryllium Technology for Fusion, Sept 15–17; 1999; pp 192–198, FZK Report 6462. 124. Scaffidi-Argentina, F.; Longhurst, G. R.; Shestakov, V.; Kawamura, H. J. Nucl. Mater. 2000, 283–287, 43–51. 125. Scaffidi-Argentina, F.; Longhurst, G. R.; Shestakov, V.; Kawamura, H. Fusion Eng. Des. 2000, 51–52, 23–41. 126. Dombrowski, D. E.; Deksnis, E.; Pick, M. A. In Thermomechanical Properties of Beryllium; Supplement to Nuclear Fusion; IAEA: Vienna, Austria, 1995; Vol. 5, pp 19–75. 127. Dombrowski, D. E. Fusion Eng. Des. 1997, 37, 229–242. 128. ITER Materials Properties Handbook (MPH) (internal project document distributed to the ITER participants). 129. ITER Materials Assessment Report (MAR). ITER Final Design Report 2001, (internal project document distributed to the ITER participants). 130. Watson, R.; et al. Low cyclic fatigue of beryllium. In Proceedings of the 2nd IEA International Workshop on Beryllium Technology for Fusion, Wyoming, Sept 6–8; 1995; p 7. 131. Roedig, M.; Duwe, R.; Gervash, A.; et al. Thermal shock tests with beryllium coupons in electron beam facility JUDITH. In Proceedings of the 2nd IEA International Workshop on Beryllium Technology for Fusion, Wyoming, Sept 6–8; 1995; pp 39–57. 132. Linke, J.; Duwe, R.; Merola, M.; Qian, R. H.; Roedig, M.; Schuster, A. Response of beryllium to severe thermal shocks – simulation of disruptions and vertical displacement events in further thermonuclear devices. In Proceedings of the 3rd IEA International Workshop on Beryllium Technology for Fusion, Mito, Japan, 1997; pp 189–199. 133. Linke, J.; Duwe, R.; Gervash, A.; Qian, R. H.; Roedig, M.; Schuster, A. J. Nucl. Mater. 1998, 258–263, 634–639. 134. Kupriyanov, I.; et al. Status of RF beryllium characterization for ITER fist wall. In Proceedings of International Conference on Fusion Reactor Materials, Sapporo, Japan, Oct 14, 2009 (in press). 135. Jiming, C. China ITER Party, Communication with ITER Organization; 2005–2009. 136. Castro, R. G.; Stanek, P. W.; Elliott, K. E.; Youchison, D. L.; Watson, R. D.; Walsh, D. S. Phys. Scripta 1996, T64, 77–83. 137. Castro, R. G.; Vaidya, R. U.; Hollis, K. J. Characterisation of plasma sprayed beryllium ITER first wall mock-ups. In Proceedings of the 3rd IEA International Workshop on Beryllium Technology for Fusion, Mito, Japan, Oct 22–24; 1997; p 137.
138. 139. 140. 141. 142. 143.
144. 145. 146.
147.
148.
149. 150. 151. 152. 153. 154. 155. 156.
157. 158. 159.
160.
161. 162. 163.
Nygren, R. E.; Youchison, D. L.; Hollis, K. J. J. Nucl. Mater. 2007, 367–370, 1325–1329. Billone, M. Irradiation Effects on Beryllium Properties; Report, ITER/US/95/IV BL-06; 1995. Gelles, D.; et al. J. Nucl. Mater. 1994, 212–215, 29–38. Barabash, V.; Federici, G.; Roedig, M.; Snead, L. L.; Wu, C. H. J. Nucl. Mater. 2000, 283–287, 138–146. Forrest, R. A.; Endacott, D. A. J. Fispact User Manual; Harwell Report, AERE-M-3654 (Rev); Harwell Laboratory: Oxfordshire, 1990. Forty, C. B. A.; Forrest, R. A.; Compton, J. J.; Rayner, C. Handbook of Fusion Activation Data Part 2: Elements Niobium to Bismuth; AEA UFS 232, AEA Technology Fusion; Euratom/UKAEA Fusion Association, May 1993. Barabash, V.; Federici, G.; Linke, J.; Wu, C. H. J. Nucl. Mater. 2003, 313–316, 42–51. Behrisch, R.; Khripunov, V.; Santoro, R. T.; Yesil, J. M. J. Nucl. Mater. 1998, 258–263, 686–693. Snead, L. L. Low temperature low dose neutron irradiatin effects on Brush Wellman S65-C and Kawecki Berylco P0 beryllium; Fusion Materials Semiannual Progress Report Period Ending Jun 30, 1998; DOE/ER-0313/24, pp 215–226. Roedig, M. Neutron Irradiation and Post Irradiation Testing of Beryllium Samples and Be/Cu Mock-Ups; Final Report ITER Task T221/1 (GB5), Report of FZ Juelich, IWV2-TN-3/2000; 2000. Dalle Donne, M.; Scaffidi-Argentina, F.; Ferrero, C.; Rocci, C. Modeling He induced swelling in beryllium during fast neutron irradiation. In Proceedings of the 1st Workshop on Beryllium Technology for Fusion, 1993; p 105, KFZ Report 52711. Snead, L. L. J. Nucl. Mater. 2004, 326, 114–124. Roedig, M.; et al. J. Nucl. Mater. 2000, 283–287, 1161–1165. Linke, J.; et al. J. Nucl. Mater. 2001, 290–293, 1102–1106. Kwast, K.; Werle, H.; Wu, C. H. Phys. Scripta 1996, T64, 41–47. Wu, C. H.; et al. Fusion Eng. Des. 1998, 39–40, 263–273. Federici, G.; et al. J. Nucl. Mater. 2000, 283–287, 110–119. Lorenzetto, P.; et al. Fusion Eng. Des. 2008, 83, 1015–1019. Lowry, C. G.; et al. Progress in design and R&D on ITER plasma facing components. In Proceedings of the 22nd International Conference on Fusion Energy, Geneva, Switzerland, 2008; IAEA IT/1-4. Youchison, D. L.; et al. Fusion Technol. 1996, 29, 599–614. Odegard, B. C., Jr.; Cadden, C. H.; Yang, N. Y. C.; Watson, R. D.; Youchison, D. L. Fusion Eng. Des. 2000, 49–50, 309–316. Manly, W. D.; et al. Report of a Technical Evaluation Panel on the Use of Beryllium for ITER Plasma Facing Material and Blanket Breeder Material; Sandia Report, SAND95-1693; Aug 1995. Puskar, J. D.; Goods, S. H.; Cadden, C. H. Diffusion bonding of beryllium to CuCrZr for ITER applications, trends in welding research. In Proceedings of the 8th International Conference, Pine Mountain, GA, Jun 1–6, 2008. Lorenzetto, P.; Cardella, A.; Daenner, W.; et al. Fusion Eng. Des. 2002, 61–62, 643–648. Lorenzetto, P.; et al. Fusion Eng. Des. 2006, 81, 355–360. Sherock, P.; Lorenzetto, P.; Walton, N.; Keil, S. Fusion Eng. Des. 2009, 84, 1759–1762.
Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices 164. Nishi, H.; et al. Study on characterisation of dissimilar material joints for ITER first wall. In 22nd International Conference on Fusion Energy, Geneva, Switzerland, 2008; IAEA: Vienna, Austria, 2008; IT/P7-11. 165. Hong, B. G.; et al. Progress on the development of the fabrication technology for the ITER first wall in Korea. In 22nd International Conference on Fusion Energy, Geneva, Switzerland, 2008; IAEA: Vienna, Austria, 2008; IT/P7-12. 166. Youchison, D. L.; et al. Fusion Eng. Des. 2009, 84, 2008–2014. 167. Lee, D. W.; et al. Fusion Eng. Des. 2009, 84, 1160–1163. 168. Park, J. Y.; et al. Fusion Eng. Des. 2009, 84, 1468–1471. 169. Liu, X. High heat flux tests of small-scale Be/Cu mock-ups for ITER. In 22nd International Conference on Fusion Energy, Geneva, Switzerland, 2008; IAEA: Vienna, Austria, 2008; IT/P2-6. 170. Chen, J. In ITER first wall fabrication technology in China. In 22nd International Conference on Fusion Energy, Geneva, Switzerland, 2008; IAEA: Vienna, Austria, 2008; IT/P7-10. 171. Boudot, C.; Boireau, B.; Lorenzetto, P.; Macelc, D. Fusion Eng. Des. 2007, 82, 1639–1644. 172. Kalin, B.; Fedotov, V.; Sevryukov, O.; et al. J. Nucl. Mater. 1999, 271–272, 410–414. 173. Mazul, I.; Giniyatulin, R.; Melder, R. In-pile thermocycling testing and post test analysis of beryllium divertor mock-ups. In Proceedings of the 3rd IEA Workshop on Beryllium Technology for Fusion, Mito, Japan, Oct 1997; p 221. 174. ITER Material Assessment, Final Design Report 200. Section 3.3, Armour/Heat Sink Joining Technologies (Internal Project Document Distributed to the ITER Participants). 175. Watson, R. D.; Whitley, J. B. Nucl. Eng. Des. Fusion 1986, 4, 49–60. 176. Mazul, I.; Giniyatulin, R.; Melder, R. Post irradiation examination of Be/Cu divertor mock-up after in-pile thermocycling experiment with active cooling. In Proceedings of the Third IEA Workshop on Beryllium Technology for Fusion, Mito, Japan, Oct 1997; p 281. 177. Lodato, A.; Roedig, M.; Duwe, R.; et al. High heat flux performance of beryllium before and after neutron irradiation. In Proceedings of the Fourth IAE Workshop on Beryllium Technology for Fusion, Karlsruhe, Germany, Sept 1999; p 206. 178. Gervash, A.; Giniyatulin, R.; Mazul, I. Fusion Eng. Des. 1999, 46, 229–235. 179. Matthews, G. F.; et al. Phys. Scripta 2007, T128, 137–143. 180. Matthews, G. F. J. Nucl. Mater. 2005, 337–339, 1–9. 181. Nunes, I.; de Vries, P.; Lomas, P. J. Fusion Eng. Des. 2007, 82, 1846–1853. 182. Thompson, V.; Krivchenkov, Y.; Riccardo, V.; Vizvary, Z. Fusion Eng. Des. 2007, 82, 1706–1712. 183. Riccardo, V. J. Nucl. Mater. 2009, 390–391, 895–899. 184. Riccardo, V. Phys. Scr. 2009, T138, 014033. 185. ANSYS. Finite Element Analysis Code, http://www. ansys.com/. 186. ABAQUS. Analysis Users Manual Version 6.6, http://www.abaqus.com/. 187. Federici, G.; et al. J. Nucl. Mater. 1999, 266–269, 14–29. 188. Federici, G.; et al. Preliminary assessment of the tritium inventory and permeation in the PFC’s of ITER. In Proceedings of the 16th IEEE/NPSS Symposium on Fusion Engineering, Champaign, IL, 1995; IEEE: Piscataway, NJ, 1996; Vol. 1, pp 418–423.
189. 190. 191. 192. 193. 194. 195. 196. 197.
198. 199. 200. 201. 202. 203. 204. 205. 206. 207. 208. 209. 210. 211. 212. 213. 214. 215. 216. 217. 218. 219. 220.
665
Federici, G.; Brooks, J. N.; Iseli, M.; Wu, C. H. Phys. Scripta 2001, T91, 76–83. Gaeta, M. J. Tokamak Dust Explosion Hazard Examination and Oxidation Database Review for ITER; ITER Report, ITER/US/94/TE/SA-11; 1994. Piet, S. J.; Federici, G. ITER White Paper on Integrated Picture of In-Vessel Tritium and Dust; ITER Report, S 81 RI 13 96-06-28 W 1.4; 1996. Technical Basis for the ITER Interim Design Report, Cost Review and Safety Analysis; ITER EDA Documentation Series No. 7; IAEA: Vienna, Austria, 1996. Rosanvallon, S.; et al. J. Nucl. Mater. 2009, 386–388, 882–883. Smolik, G. R.; Merril, B. J.; Wallace, R. S. J. Nucl. Mater. 1992, 191–194, 153–157. Anderl, R. A.; et al. J. Fusion Energ. 1997, 16, 101–108. Smolik, G. R.; Merrill, B. J.; Wallace, R. S. Reaction of Porous Be in Steam; INEL Report, EGG-FSP-10346; INEL: Idaho Falls, ID, 1992. Anderl, R. A.; Pawelko, R. J.; Oates, M. A.; Smolik, G. R. Steam Chemical-Reactivity Experiments for LANL Plasma Sprayed Beryllium; US Home Team Report, ITER/US/97/TE/SA-23; 1997. Sharpe, J. P.; Petti, D. A.; Bartels, H. W. Fusion Eng. Des. 2002, 63–64, 153–163. Petti, D. A.; Smolik, G. R.; Anderl, R. A. J. Nucl. Mater. 2000, 283–287, 1390–1395. McCarthy, K. A.; Petti, D. A.; Carmack, W. J.; Smolik, G. R. Fusion Eng. Des. 1998, 42, 45–52. Winter, J. Plasma Phys. Contr. Fusion 1998, 40, 1201–1210. Winter, J.; Gebauser, G. J. Nucl. Mater. 1999, 266–269, 228–233. Razdobarin, G. T.; Federici, G.; Kozhevin, V. M.; Mukhin, E. E.; Semenov, V. V.; Tolstyakov, S. Yu. Fusion Sci. Technol. 2002, 41, 32–43. Ovchinnikov, I.; Komarov, A.; Kuznetsov, V.; Titov, V. Fusion Eng. Des. 2006, 81, 2073–2084. Federici, G.; Wu¨rz, H.; Janeschitz, G.; Tivey, R. Fusion Eng. Des. 2002, 61–62, 81–94. Behrisch, R.; Federici, G.; Kukushkin, A.; Reiter, D. J. Nucl. Mater. 2003, 313–316, 388–392. Kirschner, A.; et al. J. Nucl. Mater. 2007, 363–365, 91–95. Kirschner, A.; et al. J. Nucl. Mater. 2009, 390–391, 152–155. Federici, G. ITER-IT. Private communication; 2004. Raffray, A. R.; Federici, G. J. Nucl. Mater. 1997, 244, 85–100. Federici, G.; Raffray, A. R. J. Nucl. Mater. 1997, 244, 101–130. Bazylev, B.; Janeschitz, G.; Landman, I.; Pestchanyi, S.; Loarte, A. J. Nucl. Mater. 2009, 386–389, 919–921. Thomas, P. R.; et al. In Proceedings of the 22nd International Conference on Fusion Energy, Geneva, Switzerland, 2008; IAEA: Vienna, Austria, 2008; IT/1-5. Federici, G. Unpublished results; 2009. Whyte, D. G.; et al. Phys. Rev. Lett. 2002, 89, 055001. Whyte, D. G.; Davis, J. W. J. Nucl. Mater. 2005, 337–339, 560–564. Pacher, H. D.; Smid, I. Net Internal Note N/I/3340/1/A; Garching, 1996. Kunugi, T. Fusion Eng. Des. 1994, 23, 329–339. Merola, M. J. Nucl. Mater. 1993, 202, 29–36. Bartels, H. W. In Proceedings of the 17th Symposium on Fusion Technology, Rome, Italy, 1992; Ferro, C., Gasparotto, M., Knoepfel, H., Eds.; Elsevier: Amsterdam, 1993; Vol. 1, p 181.
666
Beryllium as a Plasma-Facing Material for Near-Term Fusion Devices
221. Federici, G.; et al. In Proceedings of the 16th IEEE/NPSS Symposium on Fusion Engineering, Champaign, IL, 1995; Ed.; IEEE: Piscataway, NJ, 1996; Vol. 1, p 430, IEEE 95CH35852. 222. Raffray, A. R.; et al. In Proceedings of the 20th Symposium on Fusion Technology, Marseille, France, 1998; Beaumont, B., et al. Ed.; 1998; Vol. 1, p 211.
223. 224.
Mayer, M.; Behrisch, R.; Plamann, K.; Andrew, P.; Coad, J. P.; Peacock, A. T. J. Nucl. Mater. 1999, 266–269, 604–610. Dux, R.; et al. J. Nucl. Mater. 2005, 337–339, 852–856.
4.20 Physical and Mechanical Properties of Copper and Copper Alloys M. Li Argonne National Laboratory, Argonne, IL, USA
S. J. Zinkle Oak Ridge National Laboratory, Oak Ridge, TN, USA
Published by Elsevier Ltd.
4.20.1 4.20.2 4.20.2.1 4.20.2.2 4.20.2.2.1 4.20.2.2.2 4.20.2.2.3 4.20.2.3 4.20.3 4.20.4 4.20.4.1 4.20.4.2 4.20.4.3 4.20.4.4 4.20.5 4.20.5.1 4.20.5.2 4.20.5.2.1 4.20.5.2.2 4.20.5.2.3 4.20.5.2.4 4.20.5.3 4.20.5.3.1 4.20.5.3.2 4.20.6 4.20.7 References
Introduction Copper and High-Strength, High-Conductivity Copper Alloys Pure Copper PH Copper Alloys CuCrZr alloy CuNiBe alloy CuNiSi DS Copper Alloys Physical Properties of Copper and Copper Alloys Mechanical Properties of Copper and Copper Alloys Tensile Properties Fracture Toughness Creep Fatigue and Creep–Fatigue Irradiation Effects in Copper and Copper Alloys Effect of Irradiation on Physical Properties of Copper and Copper Alloys Effect of Irradiation on Mechanical Properties of Copper and Copper Alloys Tensile properties Fracture toughness Fatigue and creep–fatigue Irradiation creep and void swelling Effect of Irradiation on Microstructure of Copper and Copper Alloys Defect structure in irradiated copper and copper alloys Dislocation channeling Joining Summary
Abbreviations CW DS FFTF G-P HIP IACS JET MOTA OFHC PH SAA
Cold worked Dispersion strengthened Fast Flux Test Facility Guinier–Preston Hot isostatic pressing International Annealed Copper Standard Joint European Torus Materials Open Test Assembly Oxygen-free, high conductivity Precipitation hardened Solution annealed, and aged condition
SFT TCH
667 668 668 668 669 670 670 670 671 671 671 673 674 674 675 676 676 676 678 678 678 681 681 684 685 687 688
Stacking fault tetrahedral Tension and compression hold
4.20.1 Introduction Copper alloys are prime candidates for high heat flux applications in fusion energy systems. High heat flux is a major challenge for various fusion devices because of the extremely high energy density required in controlled thermonuclear fusion. The removal of a large amount of heat generated in the plasma through
667
668
Physical and Mechanical Properties of Copper and Copper Alloys
the first wall structure imposes a major constraint on the component design life. Materials with high conductivity are needed to assist heat transfer to the coolant and to reduce the thermal stress for pulsed mode of operation. A number of issues must be considered in the selection of materials for high heat flux applications in fusion reactors. While high conductivity is the key property for such applications, high strength and radiation resistance are also essential for the effective performance of materials in a high heat flux, high irradiation environment. In addition, fatigue behavior is a major concern for many high heat flux applications because of planned or inadvertent changes in the thermal loading. Pure copper has high thermal conductivity but rather low strength, and therefore its application as heat sinks is limited. The strength of copper can be improved by various strengthening mechanisms. Among them, precipitation hardening and dispersion strengthening are the two most viable mechanisms for improving the strength of copper while retaining its high electrical and thermal conductivities. A number of precipitation-hardened (PH) and dispersion-strengthened (DS) copper alloys are commercially available, and have been evaluated for fusion applications, for example, PH CuCrZr, CuNiBe, CuNiSi, and DS GlidCop® Al15, Al25, Al60, MAGT-0.2, etc. Two copper alloys that are most appealing are PH CuCrZr and DS CuAl25. Surveys of copper alloys for fusion applications were conducted by Butterworth and Forty1 and Zinkle and Fabritsiev.2 In this chapter, a brief description of pure copper and several copper alloys of interest to fusion applications is presented, followed by a summary of their physical and mechanical properties. The radiation effects on the physical and mechanical properties of copper and copper alloys as well as their irradiated microstructure are then discussed. Joining techniques for plasma facing components in fusion reactors are also discussed.
4.20.2 Copper and High-Strength, High-Conductivity Copper Alloys 4.20.2.1
Pure Copper
Copper is widely used where high electrical or thermal conductivity is required. Pure copper is defined as having a minimum copper content of 99.3%. Copper with oxygen content below 10 ppm is called ‘oxygenfree.’ ‘Oxygen-free, high conductivity’ (OFHC) grade
copper has room temperature electrical conductivities equal to or greater than 100% International Annealed Copper Standard (IACS), where 100% IACS ¼ 17.241 nO m at 20 C.3 Copper grades with the ASTM/SAE unified number system (UNS) designation C10100, C10200, C10400, C10500, and C10700 are classified as OFHC copper. Grades C10400, C10500, and C10700 have significant silver content, which creates activation hazards. Only C10100 and C10200 are considered for fusion systems. The use of unalloyed copper is often limited by its low strength. Copper can be strengthened by various processes, for example, cold working, grain refinement, solid solution hardening, precipitation hardening, dispersion strengthening, etc. While these approaches can significantly increase the strength, they can also lead to a pronounced reduction in conductivity. The challenge is to design a material with the best combination of strength and conductivity. Cold work can significantly increase the strength of pure copper and has a relatively moderate effect on conductivity.4 However, cold-worked copper can be softened at relatively low temperatures (200 C) because of its low recrystallization temperature.5 A recent study has shown that ultrahigh-strength and high-conductivity copper can be produced by introducing a high density of nanoscale twin boundaries.6 The tensile strength of the nano-grained copper can be increased by a factor of 10 compared to conventional coarse-grained copper, while retaining a comparable conductivity. The potential of highstrength, high-conductivity bulk nano-grained copper in nuclear energy systems, however, has not been widely explored. Alloying in copper can significantly improve mechanical strengths and raise the softening temperatures. However, additions of alloying elements also reduce electrical and thermal conductivity. Among the three alloying strengthening mechanisms, namely, solid solution hardening, precipitation hardening, and dispersion strengthening, solid solution hardening has the most detrimental effects on the conductivity4 and is the least favored mechanism to obtain highconductivity, high-strength copper alloys. 4.20.2.2
PH Copper Alloys
PH copper alloys are heat-treatable alloys. The high strength of PH copper alloys is attributed to the uniform distribution of fine precipitates of secondphase particles in the copper matrix. PH copper alloys are produced by conventional solution treatment
Physical and Mechanical Properties of Copper and Copper Alloys
and aging treatment. Solution treatment produces a homogeneous solid solution by the heating of an alloy to a sufficiently high temperature to dissolve all solutes. The alloy is then quenched to a lower temperature to create a supersaturated condition. A subsequent aging treatment heats the alloy to an intermediate temperature below the solvus temperature, to precipitate fine second-phase particles. Precipitates not only give rise to high strength, but also reduce the solute content in the matrix, maintaining good conductivity. The strength of a PH alloy depends on particle size, particle shape, volume fraction, particle distribution, and the nature of the interphase boundary.7 Despite their ability to develop significant strength, PH copper alloys may be softened substantially as a result of precipitation coarsening (overaging) at intermediate to high service temperatures or because of recrystallization during brazing or diffusion bonding. Therefore, heat treatment and thermal processing histories can have a large influence on the strength and conductivity of this class of alloys. A number of commercial PH copper alloys have been investigated for applications in fusion design, for example, CuCrZr, CuNiBe, and CuNiSi. 4.20.2.2.1 CuCrZr alloy
PH CuCrZr alloy is commercially available under several trade names, for example, Elbrodur® CuCrZr from KME Germany AG, Outokumpu Oy CuCrZr, Zollen CuCrZr, C18150®, Trefimetaux CuCrZr, MATTHEY 328® from Johnson Matthey Metals, and YZC® from Yamaha Co, Ltd. The chemical compositions of these alloys differ by a small amount, with Cr varying from 0.4 to 1.5% and Zr 0.03–0.25%. Low Cr content is to prevent the formation of coarse Cr precipitates. The element, Zr,
improves the hardening by the enhancement of fine homogeneous precipitation and improves the ductility of the alloy by inhibiting intergranular fracture.8–10 CuCrZr-IG is the ITER grade with tighter specification for composition and heat treatment. CuCrZr alloys are available in different forms, for example, bars, tubes, wires, foils, sheets, and plates. Hot forming, brazing, and inert gas welding are applicable for component manufacturing. CuCrZr alloys are used in the conventional aged condition. The reference ITER heat treatment includes solution annealing at 980–1000 C for 1 h, water quench, and aging at 450–480 C for 2–4 h.11 Typical microstructure of the prime-aged CuCrZr is shown in Figure 1(a). The alloy contains an equiaxed grain structure and uniformly distributed fine Guinier–Preston (GP) zones exhibiting primarily black dot contrasts and a small number of precipitates with lobe–lobe contrast. The number density of precipitates is on the order 1022 m3, with a mean diameter of 3 nm. A low density of micron-size Cr particles and grain boundary precipitate-free zones were also observed.12–18 CuCrZr is susceptible to overaging and recrystallization during prolonged exposure at elevated temperatures. Overaging of CuCrZr causes significant coarsening of grain structure and fine precipitates. Li et al.14 reported a lower number density (1.9 1022 m3) of larger (9 nm in diameter) precipitates with a mixture of coherent and incoherent particles after CuCrZr was hot isostatic pressing (HIP) treated at 1040 C for 2 h at 140 MPa followed by solutionizing at 980 C for 0.5 h with a slow cooling rate of 50–80 C min1 between 980 and 500 C, and final aging at 560 C for 2 h (Figure 1(b)). The average grain size was >500 mm in comparison with 27 mm grain size in the prime-aged alloy.
50 nm
50 nm (a)
669
(b)
Figure 1 Representative weak-beam dark-field images showing precipitates in unirradiated CuCrZr (a) solutionized, water quenched, and aged, and (b) hot isostatic pressed, solutionized, slow-cooled, and aged.
670
Physical and Mechanical Properties of Copper and Copper Alloys
4.20.2.2.2 CuNiBe alloy
Copper–beryllium (<1 wt% Be) binary alloys provide a good combination of strength and conductivity. The precipitation of Cu–Be binary alloys occurs in both continuous and discontinuous modes. Continuous precipitation creates uniformly distributed fine particles in the copper matrix, as a result of the following precipitation process19: a0 ðsupersaturatedÞ ! GP zones ! g00 ! g0 ! gðCuBeÞ The sequence and morphology of precipitation depends mainly on aging temperature. The first phase to nucleate from a supersaturated Cu–Be solid solution is coherent Cu-rich GP zones. Following the GP zones formation is the precipitation of so-called transition phases, g00 and g0 . The equilibrium phase, g, forms after the transition phases, and its appearance indicates overaging of the alloy. Discontinuous precipitation in Cu–Be binary alloys leads to nonuniform precipitation of long, lamellar precipitates, resulting in cell structure at grain boundaries, which increases the tendency to intergranular fracture in the alloy. High-conductivity Cu–Be alloys generally contain a third element. The addition of a small amount of nickel to Cu–Be binary alloys further increases the strength of the alloys without degrading electrical and thermal conductivities. The addition of nickel increases the precipitate solvus temperatures of Cu–Be binary alloys.20 A higher solute supersaturation condition can be reached in the solution treatment which provides a larger driving force for precipitation during the aging treatment. The strength of ternary Cu–Ni–Be alloys, therefore, is significantly increased from enhanced precipitation hardening. The electrical and thermal conductivities of Cu–Ni–Be alloys are also increased because of the depletion of the alloying elements from the solid solution during aging, resulting in high strength and high conductivity. CuNiBe exhibits very high strength with respect to other PH copper alloys. The drawback of this alloy is its very low ductility and low fracture toughness after low-dose irradiation. 4.20.2.2.3 CuNiSi
CuNiSi is another PH copper alloy that has been considered for fusion applications. CuNiSi has a nominal composition of 2.5% Ni and 0.6% Si. When heat treated properly, CuNiSi can have a much higher yield strength and higher electrical
resistivity than CuCrZr. It has been extensively used for the Joint European Torus (JET) components, for example, the divertor cryopump, the water-cooled baffles, and the Lower Current Hybrid Drive cryopump.21 4.20.2.3
DS Copper Alloys
DS copper alloys contain a fine dispersion of nanometer-sized oxide particles such as alumina, zirconia, hafnia, or chromia in the copper matrix, giving rise to high-strength and thermal stability of the alloys. This class of copper alloys can be manufactured by either conventional powder metallurgy or internal oxidation. Their properties strongly depend on the type, dimension, and volume fraction of the dispersed phase and processing techniques. Unlike PH copper alloys, the addition of finely dispersed oxide particles into the copper matrix can prevent recrystallization of the matrix and consequent softening even after exposure to temperatures approaching the melting point of the copper matrix. In addition, the oxide particles are insoluble in the solid state, and are essentially immune to coarsening because of their high melting point and high thermodynamic stability. This extends the useful temperature range of a DS alloy far beyond that possible for conventional PH alloys. Several DS copper alloys have been evaluated for fusion applications, for example, GlidCop® Al15, Al25, Al60, and MAGT 0.2. Both GlidCop® and MAGT class alloys are strengthened by Al2O3 particles, produced by internal oxidation. GlidCop® Al25 and MAGT-0.2 have been studied extensively because of their balanced strength, thermal conductivity, and ductility. GlidCop® Al25 (0.25 wt% Al) is produced by OMG America. CuAl25-IG is the ITER grade with the optimized fabrication process for improved ductility and reduced anisotropy. The microstructure of the CuAl25 alloy is characterized by elongated grain structure along the extrusion or rolling direction and a high density (average of 3.27 1022 m3) of dispersed Al2O3 particles with a mean diameter of 6–9 nm. The distribution of alumina particles can be highly heterogeneous, with some grains free of strengthening particles. A low number density of micron-size a-Al2O3 particles exists at grain boundaries. The density of dislocations in the as-wrought condition can be as high as 1.5 1015 m2.15–18,22–24 Most of the oxide particles in GlidCop alloys are triangular platelets with the remainder in the form of circular or irregular-shaped disks.25
Physical and Mechanical Properties of Copper and Copper Alloys
MAGT 0.2 is a Russian alloy produced by SPEZSPLAV Company. It contains 0.17% Al, 0.05% Hf, and 0.09% Ti in the form of oxide particles.25,26 GlidCop contains Al-oxide particles only, while in MAGT alloy, there are Al-, Ti-, and Hf-oxide particles, and mixed Al- and Ti-oxide particles. A majority of the oxide particles in MAGT 0.2 are spherical in shape with a small fraction in the form of circular disks, with an average particle size of 6 nm.25,26
4.20.3 Physical Properties of Copper and Copper Alloys
resistivity in copper is dr/dT ¼ 6.710–11 O m K1.30 Severe cold work can reduce the electrical conductivity of copper by only 2–3% IACS. All alloying elements in copper reduce the electrical conductivity, and the amount of degradation depends on the type of element, the concentration, and microstructural form (e.g., solid solution, precipitation, or dispersion). Figure 2 compares the strength and conductivity of copper and several types of copper alloys.31
4.20.4 Mechanical Properties of Copper and Copper Alloys
Physical properties of pure copper and copper alloys are quite similar in terms of the melting point, the density, the Young’s modulus, and the thermal expansion coefficient. Table 1 compares the room temperature physical properties of pure copper, PH CuCrZr, and DS CuAl25.2,27–29 Because PH copper alloys and DS copper alloys contain only a small amount of fine second-phase particles, the physical properties of these copper alloys closely resemble those of pure copper. The conductivity of copper and copper alloys is the most important physical property for their applications. The electrical conductivity of copper can be reduced by thermal vibration of atoms and crystal imperfections, for example, solute atoms, vacancies, dislocations, and grain boundaries. These different mechanisms have additive contributions to the increase in resistivity. As with other metals, the thermal conductivity of copper, kth, is proportional to the electrical conductivity, l, described by the Wiedemann–Franz law, that is, kth ¼ lLT
½1
where T is the absolute temperature and L is the Lorentz number. The electrical conductivity of pure copper is sensitive to temperature, and less sensitive to the amount of cold work and the grain size. The linear temperature coefficient for electrical Table 1 Physical properties of pure copper, PH CuCrZr, and DS CuAl25
Melting point ( C) Density (g cm3) Thermal conductivity (W m-K1) Elastic modulus (GPa)
671
Cu
CuCrZr
CuAl25
1083 8.95 391
1075 8.90 314–335
1083 8.86 364
117
123
130
4.20.4.1
Tensile Properties
The influence of test temperature, strain rate, and thermal–mechanical treatments on the tensile properties of copper and copper alloys has been studied extensively. Figure 3 illustrates the effect of test temperature on the yield strength of pure copper (in the annealed condition), PH CuCrZr and CuNiBe alloys, and DS CuAl25.15–18,28,32–39 The strength of copper alloys decreases with increasing test temperature. The decrease in strength is moderate up to 200 C. Significant drops in strength occur at higher temperatures, except that the CuNiBe ATalloy shows a relatively small reduction in strength even up to 400 C. Pure copper has the lowest yield strength. The tensile properties of pure copper strongly depend on the thermal–mechanical treatment and the impurity content.15–18,32,33 CuNiBe alloy has the highest strength over the entire temperature range.34 The tensile properties of PH copper alloys are sensitive to the thermal–mechanical treatments. CuCrZr in the solution-annealed, cold-worked, and aged condition (SA þ CW þA) has superior yield strength at low temperatures relative to CuCrZr in the solutionannealed, and aged condition (SAA). However, the strength of CuCrZr SA þ CW þA alloy drops more rapidly with increasing temperature.29,34–39 The yield strength of CuNiBe can be quite different, depending on the processing techniques. The tensile ductility of copper alloys also shows strong temperature dependence. The uniform elongation of the CuAl25 alloy decreases considerably as the test temperature increases, but increases with increasing test temperature above 400 C. The CuNiBe AT alloy shows a moderate drop of uniform elongation below 200 C, but a sharp drop in ductility at higher temperature.34 The uniform elongation of the CuCrZr alloy shows the smallest sensitivity to test temperature. Among
Physical and Mechanical Properties of Copper and Copper Alloys
1200
175
0.2% yield strength (MPa)
1000
150
Cu–2% Be (cold worked and aged)
Cu–Ni–Be (thermomechanical treated)
125 800 Cu–Ni–Be (cold worked and aged)
Cu–2% Be (cast and aged)
600
Cu–Ni–Be (solutionized and aged)
100 Cu–Al2O3 (cold worked)
75
Cu–Cr–Zr (cold worked and aged)
400
50
0.2% yield strength (ksi)
672
Cu (cold worked) Cu–Al2O3 (wrought)
200
25
Cu–Cr–Zr (solutionized and aged) Cu (annealed)
0 0
100
200 300 Thermal conductivity (W m-K–1)
0 500
400
Figure 2 Strength and conductivity of copper and copper alloys. After Li, G.; Thomas, B. G.; Stubbins, J. F. Metall. Mater. Trans. A 2000, 31A, 2491.
1000 CuCrZr, SAA (Zinkle and Eatherly,34 Zinkle,38 Singh et al. 39)
900
CuCrZr, SA+CW+A (Piatti and Boerman,29 Fabritsiev et al.35 Fabritsiev and Pokrovsky36,37) OFHC Cu, (Singh et al.,32, Singh et al.,15–18 Singh and Toft33) CuAL25, (Zinkle and Eatherly34) CuNiBe, HT1, HT2, AT (Zinkle and Eatherly34)
800
Yield strength (MPa)
700 600 500 400 300 200 100 0
0
100
200
300
400 500 Temperature (⬚C)
600
700
800
Figure 3 The yield strength of copper alloys as a function of temperature.
the three copper alloys, the CuCrZr alloy has the best ductility over the temperature range, and the ductility of the CuNiBe alloy is the lowest. Because of the sensitivity of mechanical properties to thermal–mechanical treatments in PH copper
alloys, the strength of large components made of these alloys can be significantly lower. For example, during component manufacturing, CuCrZr often experiences additional thermal cycles, such as brazing, welding, or HIPing. While solution annealing
Physical and Mechanical Properties of Copper and Copper Alloys
can be conducted during or after a brazing or HIPing process, rapid quenching is not feasible for large components, and a much slower cooling rate (e.g., furnace cooled or gas cooled) is applied in the manufacturing cycle. Significant reduction in strength due to slow cooling rates has been reported in CuCrZr.30,40–42 A slow cooling rate (50–80 C min1) and overaging at 560 C/2 h significantly reduced the yield stress and the ultimate tensile strength, and tensile elongations of CuCrZr relative to prime-aged CuCrZr.14 Cooling rates >1200 C min1 are required to fully quench the Cu–Cr solid solution.43–45 The effect of strain rate on tensile properties for pure copper and PH CuCrZr and CuNiBe alloys as well as DS CuAl25 alloy was studied at temperatures of 20 and 300 C.14,34,46 All three copper alloys are relatively insensitive to strain rate at room temperature. The strain rate sensitivity parameter of m (where sy ¼ Ce_ m and C is a constant) is 0.01 for the CuAl25 alloy at room temperature. The strain rate sensitivity of this alloy increases significantly with increasing temperature as reflected by a strain rate sensitivity parameter of m 0.07 at 300 C. Stephens et al.47 reported a strain rate sensitivity parameter of m 0.1 in the temperature range of 400–650 C for CuAl25. A similar effect of strain rate on ultimate tensile strength was also observed on these materials.34,46 Edwards46 investigated the strain rate effect of copper alloys in air and vacuum, and found that
testing in air or vacuum did not appear to change the strain rate dependence of the CuAl25 alloy, but that testing the CuNiBe alloy in air shifted the embrittlement to a lower temperature. 4.20.4.2
Fracture Toughness
Fracture toughness data for PH copper alloys, CuCrZr and CuNiBe, and DS copper alloys, CuAl15 and CuAl25, are summarized in Figure 4.14,48–50 CuCrZr has the highest toughness, and CuNiBe the lowest among these alloys. The large scatter in measured fracture toughness values for CuCrZr in different studies is likely due to different heat treatments, specimen geometry and dimensions, and testing methods. The temperature dependence of the fracture toughness in CuCrZr, while difficult to accurately define, shows an initial decrease with increasing temperature, and then a slight recovery at temperatures above 250 C. The effect of thermal–mechanical treatment on fracture toughness of CuCrZr is insignificant in comparison with its effect on tensile properties.14 The minimum value of the JQ for unirradiated CuCrZr is as high as 100 kJ m2. The fracture toughness of DS CuAl15 and CuAl25 is significantly lower than that of CuCrZr, and shows a strong directional dependence. The toughness is higher in the L-T orientation than in the T-L orientation. The fracture toughness decreases
500 Black = CuAl15 or CuAl25 Red = CuCrZr Green = CuNiBe
450 400
T-L,
L-T
----------------------------------------------
: Tahtinen et al.50 : Alexander et al.48 : Alexander et al.48 : Alexander et al.48 : Li et al.14 : Li et al.14
JQ (kJ m−2)
350 300 250
------------------------------------------------
: Suzuki et al.49
200 150 100 50 0 0
100
673
200 Temperature (⬚C)
300
Figure 4 Fracture toughness data of PH CuCrZr, CuNiBe and DS CuAl15, CuAl25.
400
674
Physical and Mechanical Properties of Copper and Copper Alloys
rapidly with increasing temperature. The JQ value for CuAl25 is only 7 kJ m2 at 250 C in the T-L orientation.48 4.20.4.3
rates of copper alloys strongly depend on the applied stress and the temperature, and can be described by the Norton power law relation; that is, e_ ¼ Asn expðQ =RT Þ where e_is creep rate, s is the applied stress, n is the stress exponent, Q is the activation energy, R is the gas constant, and T is the temperature. DS copper alloys exhibit unusually high values of the stress exponent, for example, 10–21 in the temperature range of 472–721 C for GlidCop Al15.52 Because of the time-dependent nature of creep deformation, softening behavior due to overaging and recrystallization must be considered during the creep analysis for PH copper alloys. The creep properties of this class of alloys could be significantly changed during prolonged exposure at elevated temperature.
Creep
Thermal creep of copper and copper alloys can be significant at relatively low temperatures, because of copper’s low melting point (0.3Tm ¼ 134 C, Tm is the melting point). Nadkarni51 and Zinkle and Fabritsiev2 compared the 100-h creep rupture strength of copper and several PH and DS copper alloys at elevated temperatures. Copper alloys have significantly higher creep rupture strength than pure copper. Creep rupture strength decreases drastically as temperature increases in PH alloys such as CuCrZr, as well as in pure copper, between 200 and 450 C. DS alloys such as CuAl25 have superior creep rupture strength even above 400 C because of their thermal stability at high temperatures. Li et al.31 summarized steady-state thermal creep data for pure copper and several copper alloys, as shown in Figure 5. Pure copper can suffer significant creep deformation at high temperature even with a very low applied stress. The creep rate of pure copper can be as high as 10–4 s1 at 100 MPa at 400 C. The creep resistance of copper alloys is considerably higher than that of pure copper. The creep
4.20.4.4
Fatigue and Creep–Fatigue
Copper alloys are subjected to severe thermal cycles in high heat flux applications in fusion systems, and so, fatigue as well as creep–fatigue performance is a primary concern. Figure 6 shows the fatigue performance of OFHC Cu, PH CuCrZr and CuNiBe, and DS CuAl25.53 All three copper alloys show significantly better fatigue performance than OFHC copper. Among the three alloys, CuNiBe has the best
Applied stress (ksi) 0
8
16
24
32
40
48
56
GlidCop Al-25 at 350 ⬚C (Solomon et al., 1995)
0.01 Pure copper at 400 ⬚C (Nix et al., 1985)
0.01 GlidCop Al15 at 400 ⬚C47
Creep rate (1 s–1)
10−4
10−4
10−6 GlidCop Al15 at 472 ⬚C52
Cu–Cr–Zr at 300 ⬚C (Gorynin et al., 1992)
10−8
Cu–Cr–Zr at 300 ⬚C5 Cu–Cr–Zr at 216 ⬚C (Thomas, 1993)
10−10
10−12
10−6
Ag–Cu at 193.3 ⬚C (Thomas, 1993)
0
50
100
Cu–Ni–Be at 229 ⬚C (Thomas, 1993)
150 200 250 Applied stress (MPa)
300
350
10−8
10−10
10−12 400
Figure 5 Steady-state thermal creep laws for copper alloys. After Li, G.; Thomas, B. G.; Stubbins, J. F. Metall. Mater. Trans. A 2000, 31A, 2491.
Physical and Mechanical Properties of Copper and Copper Alloys
675
1
OFHC Cu, no hold OFHC Cu, TCH 10 s CuAl25, no hold CuAl25, TCH 2 s CuAl25, TCH 10 s CuCrZr PA, no hold CuCrZr PA, TCH 10 s CuCrZr HT1, no hold CuCrZr HT1, TCH 10 s CuCrZr HT2, no hold CuCrZr HT2, TCH 10 s
0.1 0.1 102
103 104 105 Number of cycles to failure (Nf)
Ttest = 22 ⬚C
1
Cu, 25 ⬚C CuCrZr, 25 ⬚C CuCrZr, 350 ⬚C CuNiBe, 25 ⬚C CuNiBe, 350 ⬚C CuAl25, 25 ⬚C CuAl25, 350 ⬚C
Strain amplitude (%)
Strain range (%)
5
106
Figure 6 Fatigue performance of OFHC copper, precipitation-hardened CuCrZr and CuNiBe, and dispersion-strengthened CuAl25 in the temperature range of 25–350 C.
fatigue response. The temperature dependence of fatigue behavior is stronger in CuAl25 and CuNiBe than in CuCrZr at temperatures between 25 and 350 C. Heat treatments have an insignificant effect on fatigue life in CuCrZr.54 The fatigue life of copper and copper alloys can be significantly reduced when a hold time is applied at peak tensile and/or compressive strains during fatigue cycling. The hold time effect is evident even at room temperature and with a hold time as short as a few seconds.53,55,56 As shown in Figure 7, the fatigue life of OFHC copper is reduced significantly by the introduction of a hold time of 10 s at both tensile and compressive peak strains. The reduction in fatigue life is more severe in the high-cycle, longlife regime than in the low-cycle, short-life fatigue regime. A similar effect of the hold time was observed in copper alloys. The hold time effect appears to be more severe in CuAl25 than in CuCrZr. The effect of hold time is stronger in overaged CuCrZr (e.g., HT2 in Figure 7) than in prime-aged CuCrZr. Stress relaxation was observed during the hold periods even at room temperature where thermally activated creep processes are not expected. The reduction in fatigue life is apparently due to a change in the crack initiation mode from transgranular with no hold period to intergranular with a hold period.56,57 The fatigue life reduction under creep–fatigue loading could be more severe at high temperatures, particularly in PH copper alloys. Their softening behavior at elevated temperature due to overaging
1000
10 000 Cycles to failure (Nf)
100 000
Figure 7 Hold time effect on the fatigue life of OFHC copper, DS CuAl25, and PH CuCrZr with three different heat treatments (prime aged (PA): solution annealed at 1233 K for 3 h, water quenched, and then heat treated at 733 K for 3 h; heat treatment 1 (HT1): PA plus an additional anneal in vacuum at 873 K for 1 h and water quenched; and heat treatment 2 (HT2): PA plus an additional anneal in vacuum at 873 K for 4 h (and water quenched) tested at room temperature. TCH, tension and compression hold.
and recrystallization could have significant impact on the fatigue life with a very long hold time. Few studies have been performed to characterize the fatigue propagation rates of copper alloys. The fatigue crack growth rate of CuAl25 was found to be higher than that of CuCrZr at a lower stress intensity range, DK, at room temperature.58 Crack growth rates of CuCrZr and CuAl25 alloys increase with increasing temperature.49,59
4.20.5 Irradiation Effects in Copper and Copper Alloys The irradiation behavior of copper and copper alloys has been extensively studied up to high doses (>100 dpa) for irradiation temperatures of 400– 500 C.60 Most of the irradiation experiments of copper and copper alloys have been done in mixed spectrum or fast reactors, such as HFIR, Fast Flux Test Facility (FFTF), or EBR-II. It should be noted that the accumulation rate of helium in copper in fusion reactors is significantly higher than in fission reactors (10 appm dpa1 in fusion reactors vs. 0.2 appm dpa1 in fast reactors).22 Attention must be paid to transmutation effects such as helium when the irradiation data of copper and copper alloys from fission reactors are applied for fusion reactor design.
676
Physical and Mechanical Properties of Copper and Copper Alloys
4.20.5.1 Effect of Irradiation on Physical Properties of Copper and Copper Alloys
reactions. The data from fission reactor irradiation experiments must be treated with care when they are applied for fusion design.
Neutron irradiation leads to the formation of transmutation products and of irradiation defects, dislocation loops, stacking fault tetrahedra (SFT), and voids. All these features result in reduction of electrical and thermal conductivities.36,37,61–63 At irradiation temperatures between 80 and 200 C, the electrical resistivity is controlled by the formation of dislocation loops and stacking fault tetrahedra and transmutation products. The resistivity increase from radiation defects increases linearly with increasing dose up to 0.1 dpa and saturates. The maximum measured resistivity increase at room temperature is about 6%. At irradiation temperatures above 200 C, the conductivity change from extended radiation defects becomes less significant, and void swelling becomes important to the degradation of the electrical conductivity. Fusion neutrons produce a significant amount of gaseous and solid transmutation products in copper. The major solid transmutation products include Ni, Zn, and Co. The calculated transmutation rates for copper in fusion first wall at 1 MW-year m2 are 190 appm dpa1 Ni, 90 appm dpa1 Zn, and 7 appm dpa1 Co.2 Ni is the main transmutation element that affects the thermal conductivity of copper. It should be noted that water-cooled fission reactors would produce significantly higher transmutation rates of copper to Ni and Zn (up to 5000 and 2000 appm dpa1, respectively) because of thermal neutron
4.20.5.2 Effect of Irradiation on Mechanical Properties of Copper and Copper Alloys 4.20.5.2.1 Tensile properties
Irradiation causes large changes in tensile properties of copper and copper alloys. Copper and copper alloys can be hardened or softened by irradiation, depending on the irradiation temperature and the amount of the cold work prior to irradiation. Irradiation hardening of copper and copper alloys due to defect cluster formation is significant at irradiation temperatures <300 C. Irradiation softening occurs at irradiation temperatures >300 C because of radiation-enhanced recrystallization and precipitate coarsening in PH copper alloys. Low-temperature neutron irradiation of pure copper leads to development of a yield drop and significant hardening. Typical stress–strain behavior of pure copper and copper alloys irradiated to low doses at low temperatures is illustrated in Figure 8. The data of irradiated copper are from the work of Edwards et al.,64 and the data of irradiated CuCrZr from Li et al.14 Irradiation significantly changes the work hardening behavior of pure copper. Work hardening capability is progressively reduced with increasing doses. Appreciable work hardening still exists at the dose of 0.1 dpa. The effect of irradiation on the tensile behavior of copper alloys can be quite different. A complete loss of work hardening capability and
600
300 0.3 dpa
Ttest = 373 K 0.2 dpa
250
CuCrZr SAA
Tirr = 373 K
500
0.01 dpa
Stress (MPa)
Stress (MPa)
0.1 dpa
200
Unirradiated
150
400 300
100
200
50
100
0.14 dpa
Unirradiated
1.5 dpa
OHFC Cu
0 0
10
20
30
40
Strain (%)
50
60
70
0
0
5
10
15
20
25
30
35
Strain (%)
Figure 8 Engineering stress–strain curves for OFHC copper (left) neutron irradiated at 100 C and for precipitationhardened CuCrZr (right) neutron irradiated at 80 C. The plot for copper is from the reference. Reproduced from Edwards, D. J.; Singh, B. N.; Bilde-Sørensen, J. B. J. Nucl. Mater. 2005, 342, 164.
Physical and Mechanical Properties of Copper and Copper Alloys
uniform elongation occurs at 0.14 dpa in neutronirradiated CuCrZr in the prime-aged condition. Irradiation to 1.5 dpa further reduces the yield strength, and recovers some total elongation in CuCrZr. The dose dependence of radiation hardening in copper at irradiation temperatures of 30–200 C is summarized by Zinkle et al., and shown in Figure 9.65,66 Radiation hardening in copper can be observed at a dose as low as 0.0001 dpa. The yield stress increases dramatically with increasing dose and saturates at 0.1 dpa. Significant radiation hardening is accompanied by loss of strain hardening capabilities, resulting in prompt necking upon yielding. The temperature dependence of radiation hardening of pure copper at different irradiation temperatures was summarized and discussed by Fabritsiev and Pokrovsky.67 The radiation hardening decreases with increasing irradiation temperature in copper. The magnitude of radiation hardening is 200 MPa at 80 C, while only 40 MPa at 300 C at a dose of 0.1 dpa. Annealing at temperatures higher than 0.4 Tm can effectively reduce the defect cluster density in copper. Annealing at 300 C for 50 h after irradiation of copper to 0.01–0.3 dpa at 100 C and annealing at 350 C for 10 h after irradiation of CuCrZr IG and GlidCop Al25 IG to 0.4 dpa at 150 C can essentially recover the ductility of the copper and copper alloys.68,69 However, postirradiation
350 Tirr = 30–200 ⬚C
annealing also reduces the critical stress for flow localization in pure copper.70 Irradiation creates a large increase in strength and decrease in ductility in copper alloys for irradiation temperatures below 300 C. The strengthening effect decreases with increasing temperature. The crossover to radiation softening occurs at approximately 300 C. The radiation softening effect in CuAl25 alloy is not as strong as for CuCrZr alloy where precipitate stability may be an issue. Neutron-irradiated copper alloys exhibit low uniform elongation after low-dose, low-temperature irradiation. The uniform elongation is recovered to near unirradiated values at 300 C. Figure 10 compiles the yield strength data for PH CuCrZr and DS copper alloys (CuAl 25, CuAl15, MAGT 0.2) as a function of dose for the irradiation temperature of 100 C.14,71 Both alloys show significant radiation hardening at low doses and an apparent saturation at 0.1 dpa. Irradiation-induced hardening is accompanied by the loss of strain hardening capability and a complete loss of uniform elongation, while the total elongation remains on the level of 10% for doses up to 2.5 dpa for CuCrZr. The strain rate dependence of tensile properties in neutron-irradiated CuCrZr was investigated at room temperature by Li et al.14 The strain rate sensitivity is small at room temperature in unirradiated CuCrZr. The measured strain rate sensitivity parameter, m, is <0.01 for CuCrZr. The strain rate sensitivity parameter increased to 0.02 in CuCrZr after neutron irradiation to 1.5 dpa. Zinkle et al.65 observed a small strain rate dependence of tensile strength in GlidCop Al15 and MAGT 0.2 neutron irradiated to 13 dpa at 200 C with m 0.02 for GlidCop
250
800 Ttest ~ = Tirr = 60–100 ⬚C
200
600 Kruglov et al. (1969) EI-Shanshoury 1972) Mohamed et al. (1982) Vandermeulen (1986) Heinisch (1988) Fabritsiev et al. (1994) Singh et al.23 Zinkle and Gibson65 Singh et al.75
150
100
50 0.0001
Yield stress (MPa)
0.2% yield strength (MPa)
300
677
0.001
0.01
0.1
1
10
Damage level (dpa)
Figure 9 Radiation hardening in copper. Reproduced from Zinkle, S. J.; Gibson, L. T. Fusion Materials Semi-annual Progress Report; DOE/ER-0313/27; Oak Ridge National Laboratory, 1999; p 163.
400
200 CuCrZr DS Cu 100
0
0
1E-4
1E-3
0.01
0.1
1
Dose (dpa)
Figure 10 Dose dependence of the yield strength in CuCrZr and DS copper alloys irradiated at low temperatures.
10
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Physical and Mechanical Properties of Copper and Copper Alloys
Fracture toughness data for irradiated copper alloys are scarce. The effect of neutron irradiation on fracture toughness has been studied in two alloys, CuCrZr and CuAl25.14,50,72 Fracture toughness data on neutron-irradiated CuAl25 are available to a dose of 0.3 dpa, and for CuCrZr, the data are available up to 1.5 dpa (Figure 11). Neutron irradiation to 0.3 dpa significantly reduced the fracture toughness of CuAl25 in the temperature range of 20–350 C. The toughness of irradiated CuAl25 is two to three times lower than that of the unirradiated alloy. The effect of neutron irradiation on fracture toughness of CuCrZr was less pronounced, despite the significant effect on the tensile properties even at relatively low doses (0.14–0.15 dpa). Reduction of fracture toughness in irradiated CuCrZr was small, and the JQ value was still >200 kJ m2 up to 1.5 dpa (Figure 11).14
at 250 and 350 C because of radiation exposure. The fatigue life of the CuCrZr alloy was reduced following irradiation at 250 and 350 C, similar to CuAl25. The degradation in the fatigue performance of these two alloys from irradiation exposure was not as severe as that in the tensile properties. Creep–fatigue behavior of neutron-irradiated CuCrZr was investigated at a dose level of 0.2–0.3 dpa at 22 and 300 C by Singh et al.54 Hold times of 10 and 100 s were applied during fatigue cycling. Radiation hardening at low temperatures (e.g., 60 C) is beneficial to the fatigue performance, while irradiation at high temperatures (e.g., 300 C) has no significant effect on the creep–fatigue life of irradiated CuCrZr. A number of in-reactor creep–fatigue experiments were performed on a CuCrZr alloy in the BR-2 reactor at Mol (Belgium) by Singh et al.77 The irradiation experiments were carried out at 70 and 90 C at the strain amplitude of 0.5% with hold times of 10 and 100 s. The key finding was that neither the irradiation nor the hold time has any significant effect on the fatigue life of CuCrZr during the in-reactor tests.
4.20.5.2.3 Fatigue and creep–fatigue
4.20.5.2.4 Irradiation creep and void swelling
The effect of irradiation on fatigue performance has been evaluated for PH CuCrZr and DS CuAl25.73 The fatigue data for unirradiated and irradiated CuAl25 and CuCrZr in the temperature range of 20–350 C are compiled and compared in Figure 12.24,53,74–76 The effect of irradiation on the fatigue response of CuAl25 is small at low temperature. However, the fatigue life is reduced significantly
There is limited literature on irradiation creep of copper and copper alloys.78–82 A study by Witzig82 showed no enhancement of creep rates in copper relative to thermal creep at 260 C and 69 MPa under light ion irradiation. Jung79 studied irradiation creep of 20% cold-worked copper foils at temperatures of 100–200 C and the applied tensile stress of 20–70 MPa under 6.2 MeV proton irradiation with displacement rates of 0.7–3.5 10–6 dpa s1. The irradiation creep rate showed a linear stress dependence with the irradiation creep compliance of 6.2 10–11 Pa1 dpa1 at stresses <50 MPa at 150 C, comparable to that of other fcc metals such as Ni and austenitic stainless steels. At higher stresses (>50 MPa), the creep rate showed a power law relation with the stress exponent of 4. Ibragimov et al.78 investigated in-reactor creep of copper in the WWR-K water-cooled reactor at a neutron flux of 2.5 1015 m2 s1 (E > 0.1 MeV) at 150–500 C and 20–65 MPa. The in-reactor creep rate of copper was significantly higher than the thermal creep rate at temperatures below 0.4 Tm (Tm is the melting point). The stress dependence of the in-reactor creep rate showed a power law relation with the stress exponent of 3. Pokrovsky et al.80 reported irradiation creep data for DS MAGT 0.2. The irradiation creep experiments were performed using pressurized tubes
Al15 and m < 0.01 for MAGT 0.2. In general, the strain rate and temperature dependence of flow stresses is small in fcc metals. 4.20.5.2.2 Fracture toughness
500 Tirr = 80 ⬚C; Ttest = 22 ⬚C
JQ (kJ m−2)
400
CuCrZr SAA
300 CuCrZr SCA 200
100
0
Solid symbols: JQ Open symbols: Jmax < JQ 0
0.01
0.1
CuCrZr SCA CuCrZr SAA Tahtinen et al. Singh et al. Suzuki et al. Gillian et al. Rowcliffe 1
Dose (dpa)
Figure 11 Fracture toughness of CuCrZr with two heat treatments as a function of dose. The heat treatment, SCA, was to simulate the manufacturing cycle for ITER large components. Reproduced from Li, M.; Sokolov, M. A.; Zinkle, S. J. J. Nucl. Mater. 2009, 393, 36.
Physical and Mechanical Properties of Copper and Copper Alloys
679
Unirr RT, small size, UIUC Unirr RT, standard size, UIUC Tirr = 47 ⬚C, RT, RISO
3
Total strain range (%)
Unirr RT, HT at 650 ⬚C, longitudinal, srivatsan Unirr RT, HT at 650 ⬚C, transverse, srivatsan Unirr RT, HT at 650 ⬚C, longitudinal, srivatsan Unirr RT, HT at 650 ⬚C, transverse, srivatsan Unirr 200 ⬚C in air, UIUC Unirr 250 ⬚C in vac, RISO Tirr = Ttest = 250 ⬚C, 0.1 dpa, RISO Tirr = Ttest = 250 ⬚C, 0.3 dpa, RISO Unirr 350 ⬚C in air, UIUC Unirr 350 ⬚C in vac, UIUC Unirr 350 ⬚C in vac, RISO Tirr = Ttest = 350 ⬚C, 0.1 dpa, RISO
1
GlidCopTM CuAl25 unirradiated and irradiated 0.2
100
1000
10 000
100 000
Cycles to failure (Nf) 3
Total strain range (%)
CuCrZr alloy, unirradiated and irradiated
1
Unirr RT, small size, UIUC Unirr RT, standard size, UIUC Unirr, 200 ⬚C in air, UIUC Unirr, 250 ⬚C in vac, RISO
Tirr = Ttest = 250 ⬚C 0.3 dpa, RISO Unirr, 350 ⬚C in air, UIUC Unirr, 350 ⬚C in vac, RISO
Tirr = Ttest = 350 ⬚C 0.3 dpa, RISO
0.1
100
1000
10 000
100 000
Cycles to failure (Nf) Figure 12 Effect of irradiation on fatigue life of CuAl25 (top) and CuCrZr (bottom) between room temperature and 350 C.
irradiated in coolant water in the core position of the SM-2 reactor to 3–5 dpa at temperatures of 60–90 C. A creep rate as high as 2 10–9 s1 was observed at a hoop stress of 117 MPa. Radiation-induced void swelling in copper and copper alloys has been studied extensively. Zinkle and Farrell83,84 measured the temperaturedependence of void swelling in pure copper and a dilute Cu–B alloy neutron irradiated to 1.1–1.3 dpa at a damage rate of 2 10–7 dpa s1 at temperatures of 180–500 C (Figure 13). Maximum swelling occurs at 300–325 C in pure copper under fission neutron irradiation conditions. The lower
temperature limit for void swelling is 180 C, and the higher temperature limit 500 C. Low-dose irradiation (<0.2 dpa) often leads to inhomogeneous void formation and nonlinear swelling behavior.60 A steady-state swelling rate of 0.5%/dpa is observed in copper at high doses, and the swelling level can be as high as 60%.60,85 Variations in displacement damage rate can shift the peak swelling temperature. An order of magnitude decrease in neutron flux can lower the peak swelling temperature by 20 C. The peak swelling temperature shift can be as high as 165 C between neutron irradiation (10–7 dpa s1) and ion irradiation (10–3 dpa s1).
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Physical and Mechanical Properties of Copper and Copper Alloys
0.7 Cu-100 appm 10B 0.6
Density change (%)
0.5
0.4 Pure Cu 0.3
0.2 Stage V 0.1
0 150
200
250
300
350
400
450
500
550
Irradiation temperature (⬚C) Figure 13 Swelling in pure copper and Cu–B alloy. Reproduced from Zinkle, S. J.; Farrell, K. J. Nucl. Mater. 1989, 168, 262; Zinkle, S. J.; Farrell, K.; Kanazawa, H. J. Nucl. Mater. 1991, 179–181, 994.
Residual impurity oxygen can have a significant effect on void swelling in copper. A number of neutron, ion, and electron irradiation studies have shown that voids are not formed in high-purity, low-oxygen copper over the wide range of irradiation temperatures.60,86 The oxygen content should be maintained below 10 wt ppm to minimize void swelling in copper. The effect of helium production on void formation and swelling in copper is a significant concern for its fusion applications.87 Helium effects have been studied by either dual-beam ion irradiation88,89 or neutron irradiation of Cu–B alloys.89 Significant enhancement of void formation and swelling was observed in copper under ion irradiation with simultaneous helium implantation. Neutron irradiation of copper containing 18 wppm 10B to 1.2 dpa for the irradiation temperatures of 182–500 C showed that the peak swelling temperature and the lower swelling temperature limit shifted to lower values (Figure 13). A recent study by Xu et al.90 of materials enriched in the copper isotopes, 63Cu, 63þ65Cu, and 65Cu neutron irradiated in the Materials Open Test Assembly (MOTA) in the FFTF at irradiation temperatures of 373–410 C to doses up to 15.4 dpa found that both H and He enhanced void swelling in copper. The H effect is important at lower temperatures when the H production is considerably higher than the
He production. At 410 C the hydrogen effect decreases dramatically and void swelling is affected by the helium concentration. PH and DS copper alloys have superior void swelling resistance compared to pure copper under fission neutron irradiation.2,71 Both PH CuCrZr and DS CuAl25 showed <2% swelling after irradiation to 150 dpa at 415 C.85,91 When irradiated to 98 dpa at 450 C, only 2% swelling was observed in CuAl25. The CuAl25 alloy appears to have the best resistance to void swelling among the copper alloys.92 However, the swelling resistance of DS copper alloys can be significantly reduced when there is a high generation rate of helium. While CuAl25 showed negligible swelling after irradiation to 103 dpa at 415 C in the FFTF, boron-doped CuAl15 showed 11% swelling under the same irradiation condition.93 Fabritsiev et al.22 reported a swelling rate of 1%/dpa for CuAl25 þ B alloy even at a low dose of 0.5 dpa at 300 C, because of high helium accumulation. The boron-free MAGT 0.2 alloy did not show swelling in the same experiment. Simultaneous heavy ion irradiation and helium implantation in GlidCop Al60 at 350 C showed an increase of the swelling rate from 0.01%/dpa (single-beam irradiation) to 0.05%/dpa (dual-beam irradiation).94 The initial thermal–mechanical treatment of PH copper alloys can have a significant impact on their
Physical and Mechanical Properties of Copper and Copper Alloys
swelling resistance. CuNiBe in the cold-worked and aged condition showed 28% swelling, while CuNiBe in the annealed and aged condition swelled only 13% after fission neutron irradiation to 98 dpa at 450 C.95 The susceptibility to radiation-enhanced recrystallization is more severe in a cold-worked alloy, leading to the swelling instability. 4.20.5.3 Effect of Irradiation on Microstructure of Copper and Copper Alloys 4.20.5.3.1 Defect structure in irradiated copper and copper alloys
Copper is among the most extensively studied metals in terms of fundamental radiation damage. Several reviews on the effect of irradiation on the
681
microstructure of copper and copper alloys are available in the literature.60,96,97 Neutron irradiation of copper at low temperatures produces small defect clusters, dislocation loops, and SFTs. At temperatures above 150–180 C, the density of defect clusters starts to decrease with increasing temperature, accompanied by the formation of voids. This temperature-dependent formation of defect structures is shown in Figure 14.60 Low-temperature neutron irradiation produces a high number density of SFTs and a low number density of dislocation loops in copper. Edwards et al.64 reported a number density of SFTs, 2–4 1023 m3 and a number density of dislocation loops, 5 1021 m3 in OFHC copper neutron irradiated to 0.01 dpa at 100 C. Dislocation loops are believed to be of interstitial type.
300 ⬚C
10 nm
100 nm
(a)
(b)
Normalized units
1.0
Loops, SFT (irradiation hardening) 0.5
0
(c)
Void swelling
0
100
200
300
400
500
600
Temperature (⬚C)
Figure 14 (a) Stacking fault tetrahedra and defect clusters produced in OFHC copper during irradiation to 1.9 dpa at 180 C (reproduced from Zinkle, S. J.; Matsukawa, Y. J. Nucl. Mater. 2004, 329–333, 88), (b) voids in copper irradiated at 300 C (reproduced from Zinkle, S. J.; Farrell, K. J. Nucl. Mater. 1989, 168, 262). (c) Schematic drawing showing the temperature dependence of defect cluster formation and void swelling (reproduced from Zinkle, S. J. In Effects of Radiation on Materials, ASTM STF 1125, 15th International Symposium); Stoller, R. E., et al., Eds.; American Society for Testing and Materials: Philadelphia, 1992; p 813.
682
Physical and Mechanical Properties of Copper and Copper Alloys
The size of SFTs is small, 2–3 nm. As doses increased, the density of SFTs increased to a saturation level at 0.1 dpa, while the size of SFT is independent of the dose and temperature. In general, the dislocation loop density is low, and a significant dislocation network is not formed in irradiated copper.96 Radiation hardening in copper can be adequately described by Seeger’s dispersed barrier model, and the yield strength increase is due to the formation of defect clusters.98 Singh and Zinkle96 summarized the dose dependence of the TEM-visible defect cluster density in copper irradiated near room temperature with fission neutrons, 14 MeV neutrons, spallation neutrons, and 800 MeV protons (Figure 15)96 TEMvisible defect clusters were observed at a very low dose (10–5 dpa). The defect cluster density showed a linear dependence on irradiation dose at low doses. The dose dependence of the defect cluster density shifts to either a linear or a square root relation at intermediate doses (>0.0002 dpa). The cluster density reaches an apparent saturation (1 1024 m3) at 0.1 dpa. The dislocation loops range in size from 1 to 25 nm.99 Differences in the type of irradiation (fission, fusion, spallation, etc.) have no significant effect on the defect cluster accumulation behavior in copper. The density of defect clusters in irradiated copper shows strong temperature
dependence (Figure 16).100 The defect cluster density is essentially independent of the irradiation temperature between 20 and 180 C (upper temperature limit is dependent on dose rate). At higher temperature, the cluster density decreases rapidly with increasing irradiation temperature. At irradiation temperatures between 182 and 450 C, the density of defect clusters was reduced by over three orders of magnitude.83,84 The saturation dose of the defect cluster density is similar, 0.1 dpa, for all irradiation temperatures.96 The size distribution of visible defect clusters can be described by an exponential function101: N(d ) ¼ N0 exp(d/d0), where N(d ) is the number of defects of diameter d, N0, and d0 are constants, and their values depend on irradiation conditions and material purity. As the irradiation temperature decreases, a fraction of small clusters increases relative to large clusters. Void formation occurs above 180 C in neutronirradiated copper.60 The peak void swelling temperature in copper is about 320 C at a dose rate of 2 10–7 dpa s1. Singh and Zinkle96 summarized the dose dependence of void density measured by TEM in copper irradiated with fission and fusion neutrons at 250–300 C from several studies. The data showed a large variation (up to two orders of magnitude differences) of void density between
1025 800 MeV protons
Cluster density (m−3)
1024
n = 1/2 1023
1022
Yoshida et al. (1985) Zinkle89 Satoh et al. (1988) Horsewell et al. (1990) Makin et al. (1962) Shimomura et al. (1985) Brager et al. (1981)
n=1 1021
1020 19 10
1020
1021
1022
1023
1024
1025
ft (n m−2) Figure 15 Dose dependence of defect cluster density in copper irradiated near room temperature. Reproduced from Singh, B. N.; Zinkle, S. J. J. Nucl. Mater. 1993, 206, 212.
Physical and Mechanical Properties of Copper and Copper Alloys
experiments. One possible source could be residual gas atoms in copper that can have a dramatic effect on void swelling in copper. Zinkle and Lee86 discussed in detail the effect of oxygen and helium on the formation of voids in copper. The stacking fault tetrahedron is predicted to be the most stable configuration of vacancy clusters in copper. A small amount of oxygen (10 appm) or helium ( 1 appm) in copper is needed to stabilize voids. High-purity copper with low oxygen concentration (<5 wppm) showed no significant 16 10 dpa 2 ⫻ 10−3 dpa s-1
Cluster density (1022 m−3)
14 12
Total 10 8 6
SFT
4 2 0
100
200 300 Irradiation temperature
400
Figure 16 Measured defect cluster density in 14-MeV Cu3þ ion-irradiated copper as a function of irradiation temperature. Reproduced from Zinkle, S. J.; Kulcinski, G. L.; Knoll, R. W. J. Nucl. Mater. 1986, 138, 46.
683
void formation after 14 MeV Cu ion irradiation to 40 dpa at temperatures of 100–500 C.100 The defect microstructure (SFTs and dislocation loops) in irradiated copper alloys is essentially the same as in irradiated pure copper.22,25,64 Neutron irradiation can affect precipitate microstructure in copper alloys. When irradiated at 100 C, the precipitate density in CuCrZr was slightly reduced, and the mean size of the precipitates increased.13,64 Zinkle et al.25,26 reported that when GlidCop Al25 and MAGT 0.2 were ion irradiated to 30 dpa at 180 C, a high number density (5 1023 m3) of defect clusters (primarily SFTs) with a mean size of 2 nm was produced. The geometry of oxide particles in GlidCop Al25 was transformed from triangular platelets to nearly circular platelets, and the particle size was reduced from 10 to 6 nm after irradiation (Figure 17).25,26 The geometry and size of oxide particles in MAGT 0.2 were essentially unchanged by irradiation. In general, DS copper alloys showed superior particle stability under irradiation. Limited data are available in terms of the effect of solution additions on the irradiated microstructure of copper. A study by Zinkle25 showed that solute additions (e.g., Al, Mn, Ni) to 5 at% in copper do not have significant effect on the total density of small defect clusters at low irradiation temperatures (130 C). However, solute additions reduce the formation of SFTs and enhance the formation of dislocation loops. The loop density and mean size in Cu–5% Mn irradiated to 1.6 dpa at 160 C were 3 1021 m3 and 23 nm, and 1.8 1022 m3 and 18 nm in Cu–5% Ni irradiated to 0.7 dpa at 90 C
20 nm
100 nm
Figure 17 Defect structure (left) and Al2O3 particle morphology (right) in 50% cold-worked GlidCop Al25 irradiated with 3 MeV Arþ ions to 30 dpa at 180 C. Reproduced from Zinkle, S. J.; Horsewell, A.; Singh, B. N.; Sommer, W. F. J. Nucl. Mater. 1994, 212-215, 132; Zinkle S. J.; Nesterova, E. V.; Barabash, V. R.; Rybin, V. V.; Naberenkov, A. V. J. Nucl. Mater. 1994, 208, 119.
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Physical and Mechanical Properties of Copper and Copper Alloys
0.1 µm Figure 18 Comparison of the dislocation loop microstructure in irradiated pure copper (left), Cu–5% Mn (center) and Cu–5% Ni (right) alloys. The irradiation conditions were 0.7 dpa at 90 C (Cu), 1.6 dpa at 160 C (Cu–5% Mn), and 0.7 dpa at 90 C (Cu–5%Ni). Reproduced from Zinkle, S. J.; Horsewell, A.; Singh, B. N.; Sommer, W. F. J. Nucl. Mater. 1994, 212-215, 132; Zinkle S. J.; Nesterova, E. V.; Barabash, V. R.; Rybin, V. V.; Naberenkov, A. V. J. Nucl. Mater. 1994, 208, 119.
(Figure 18).25,26 These loop densities are more than an order of magnitude larger than the highest loop density observed in pure copper. The effect of the stacking fault energy on void formation in copper alloys was also investigated. Generally speaking, the lower the stacking fault energy, the less favorable for the formation of 3D voids. For example, swelling occurred in Cu–1–2.5% Ge alloys irradiated at 250 C, while no measurable swelling occurred in Cu–3–5% Ge that has lower stacking fault energies.97 4.20.5.3.2 Dislocation channeling
Dislocation channels are frequently observed during postirradiation deformation of copper and copper alloys.102,103 Greenfield and Wilsdorf104 were the first who observed an area free of irradiation defects in the middle of a slip-line cluster by TEM in a neutron-irradiated copper single crystal. Extensive studies were conducted to establish the correlation between the deformation behavior and the slip-line structure in neutron-irradiated copper single crystals.104–107 Sharp108–110 studied the deformation and dislocation channels in neutron-irradiated copper single crystals in detail, and established a direct correlation between the surface slip steps and dislocation channels. The channels are nearly free of irradiationproduced defects, and operate parallel to the primary {111} slip plane. The cleared channels are formed by cooperative localized motion of glide dislocations that interact with and annihilate the preexisting radiation defect clusters. The channel characteristics
have strong dependence on irradiation dose and test temperatures. The channel width decreases and the slip step height increases with increasing irradiation dose, and the channel width and the slip step height decrease with decreasing deformation temperature. Howe111 confirmed that the channel width, spacing, the slip step height, and the average shear per slip band increase with increasing test temperature in the temperature range of 4–473 K. The reduction in channel width was considered to be a consequence of impeded cross-slip.108,111 Dislocation channels were also observed in neutron-irradiated copper single crystals under cyclic straining.112,113 The width and average spacing of channels changed with the number of cycles, in contrast to channels formed during tensile straining where the width and spacing of channels were constant over a large range of strains.108 Dislocation channels are formed in neutronirradiated copper alloys as well. Sharp114 observed the channeling effect in three different copper alloys neutron irradiated at ambient temperature, that is, Cu–0.8% Co, Cu–Al2O3, and Cu–4% Al single crystals. The channel spacing in the copper alloys were 1.2–1.5 mm, about half that observed in neutron-irradiated copper single crystals (2.3 mm). The channel width in Cu–0.8% Co alloy is similar to that for irradiated copper crystal (0.16 mm), and the channels have the uniform width along the length. The presence of the second-phase particles in Cu–0.8% Co alloy has little effect on channeling. In the DS Cu–Al2O3 alloy, the channels are wider (0.24 mm) and
Physical and Mechanical Properties of Copper and Copper Alloys
more irregular in width. The channel width can vary by a factor of 2 within a few microns along the length of a channel. A high density of dislocations surrounding the particles within the channel was observed in Cu– Al2O3, indicating great difficulty of dislocations in bypassing the (nondeforming) second-phase particles. In the single-phase Cu–4% Al alloy, however, no dislocation channels were observed. Edwards13,40,64,115 studied thoroughly the deformed microstructure in neutron-irradiated CuCrZr alloys, and compared with the deformation microstructure in neutron-irradiated OFHC-Cu (Figure 19). Dislocation channels were observed during postirradiation deformation of the CuCrZr alloy neutron irradiated to 0.2–0.3 dpa at 100 C. Channels were formed even before the upper yield point, and continued throughout the tensile deformation process. Some channels are completely free of defect clusters, and others contain a sizeable population of defect clusters. The width of cleared channels varied between about 100 and 250 nm. The channel formation is more pronounced in a higher-dose specimen than in a lowerdose specimen. In comparison with OFHC-Cu, CuCrZr showed little difference in deformation mode and channel characteristics in terms of width and size. While the channels in the OFHC-Cu were free of defects and dislocation debris, the channels in the CuCrZr alloy contained a small fraction of defects and precipitates. When the irradiated CuCrZr was annealed and deformed, deformation occurs in a much more homogeneous fashion, and no well-defined channels were observed. The formation of dislocation channels in pure copper was investigated by in situ straining experiments on ion-irradiated copper in an electron microscope.116,117 Postirradiation straining of the thin foils of polycrystalline copper irradiated with 200 keV Kr
685
ions to about 2 10–4 to 0.02 dpa at room temperature showed that defect-free channels nucleate at grain boundaries, or in the vicinity of cracks, suggesting that grain boundaries and crack tips are nucleation sites for channels.117 Cross-slips were found to be responsible for channel widening and defect removal within the channel. Edwards et al.64 studied the initiation and propagation of dislocation channels in neutron-irradiated OFHC-Cu (Figure 20) and CuCrZr alloy in an interrupted tensile test. TEM observations suggested that channels are initiated at boundaries, large inclusions, or existing channels. Channels are formed by interactions of newly formed dislocations with irradiation defects on the glide plane. Once formed, the channels propagate rapidly in the grain interior until they intercept another boundary, interface, or channel. Despite significant efforts, the exact mechanism of channel formation and evolution still remains unresolved, and a clear connection between the slip processes, dislocation channeling, and localized flow in neutron-irradiated metals is still lacking.
4.20.6 Joining Copper and copper alloys can be joined by a variety of techniques, including mechanical coupling, welding, brazing, and diffusion bonding. A comprehensive overview of joining techniques for copper and copper alloys can be found in the reference.118 The welding techniques commonly used for copper and copper alloys include arc welding, resistance welding, oxyfuel welding, and electron beam welding. Welding is generally not recommended for joining high-strength copper alloys. PH copper alloys lose their mechanical strength because of the dissolution of precipitates
Figure 19 Dislocation channels observed in OFHC-Cu (left) and CuCrZr (right) irradiated to 0.3 dpa at 100 C. Edwards, D. J.; Singh, B. N.; Xu, Q.; Toft, P. J. Nucl. Mater. 2002, 307–311, 439; Edwards, D. J.; Singh, B. N.; Bilde-Sørensen, J. B. J. Nucl. Mater. 2005, 342, 164.
686
Physical and Mechanical Properties of Copper and Copper Alloys
(a)
(c)
(b)
Strained to 1.5%
(d)
Strained to 14.5%
Figure 20 Examples of cleared channels formed in the OFHC-Cu irradiated (to 0.3 dpa) and tested at 323 K to different strain levels: (a) before yield, (b) before yield, (c) 1.5%, and (d) 14.5%. Note that at 14.5% strain level the grain is subdivided by numerous channels formed on different slip planes. All images shown in this figure were taken in the STEM bright field mode. Reproduced from Edwards, D. J.; Singh, B. N.; Bilde-Sørensen, J. B. J. Nucl. Mater. 2005, 342, 164.
during the welding process. The welded component must be resolution annealed and aged to recover some of the initial strength in the joint. Recrystallization in the melt layer degrades the mechanical property of the weldment. DS copper alloys cannot be welded by conventional welding processes because of the loss of oxide particles and recrystallization in the weld zone. Brazing is the most common method for joining copper alloys. All conventional brazing techniques can be used to join copper and copper alloys, including furnace brazing, torch brazing, induction brazing, resistance brazing, and dip brazing. A wide range of filler metals are available, and the most common brazing filler metals are Cu–Zn, Cu–P, Cu–Ag–P, and Ag- and Au-based alloys.118 Ag- and Au-based filler metals are unacceptable in fusion reactor environments because of concerns of high radioactivity from neutron-induced transmutation.119 Copper alloys are typically brazed at temperatures between 600 and 950 C with hold times at the brazing temperature ranging from 10 s (torch, resistance, or induction brazing) to 10 min (furnace
brazing).2 The brazing process can significantly soften PH copper alloys as a result of the adverse precipitation process. To reduce the softening effect, a fast induction brazing technique has been developed to minimize the holding time at high temperature to retain sufficient mechanical properties.120 Alternatively, the brazed component can be aged following furnace brazing to restore part of its initial strength. Complete recovery of high strength after furnace brazing by heat treatment in PH alloys is rather difficult in practice as the component must be heated to a temperature greater than typical brazing temperatures and rapidly quenched to create a supersaturation of solute prior to aging. Oxide DS copper has been successfully joined using torch, furnace, resistance, and induction brazing.2 Softening is not a serious concern for the base metal of DS copper alloys because of their high recrystallization temperature. The brazed copper joints show good fatigue properties and relatively low ductility.2 Diffusion bonding is a viable technique to produce joints with high mechanical strength for DS copper alloys, but cannot be used to produce high-strength
Physical and Mechanical Properties of Copper and Copper Alloys
joints in PH alloys because of significant softening of the base metal during high-temperature exposure. The DS CuAl15 and CuAl25 alloys can be joined by diffusion bonding with acceptable bond strengths under the diffusion bonding conditions similar to the normal HIPing conditions.121 Techniques for joining copper alloys to beryllium or austenitic stainless steels have been developed for the ITER plasma-facing components. A review of the joining technology was given by Odegard and Kalin.119 Recent work has focused on small- and medium-scale mock-ups and full-scale prototypes of the ITER first wall panels.122 The first wall panels of the ITER blanket are composed of a composite Cu alloy/316L(N) SS water-cooled heat sink structure with Be tile clad. A number of joining techniques have been explored for joining copper alloys to austenitic stainless steel, 316L(N), including diffusion bonding, brazing, roll bonding, explosive bonding, friction welding, and HIP.123 HIP joining is by far the most desirable technique. For the PH CuCrZr alloy, the heat treatment must be integrated with the bonding cycle, and a high cooling rate (>50 C min1) is required to obtain good mechanical properties of CuCrZr after subsequent aging treatments. Two alternative processes are recommended124: the HIP cycle (1040 C and 140 MPa for 2 h) followed by quenching in the HIP vessel, or a normal HIP cycle with a subsequent heat treatment in a furnace with fast cooling. Gervash et al.125 studied alternative SS/Cu alloy joining methods, for example, casting, fast brazing, and explosion bonding. Cast SS/CuCrZr joint may be suitable for some ITER applications. Brazing and diffusion bonding have been considered for joining the beryllium armor to a copper alloy heat sink. The Be/DS copper alloy joints can be made by high-temperature HIPing and furnace brazing.126 Results from shear tests on small-scale specimens and from high heat flux tests of the first wall mock-ups showed good performance of joints brazed with STEMET 1108 alloy at 780 C for less than 5 min.122 The Be/Cu-Al25 solid HIPing (e.g., 730 C and 140 MPa for 1 h) showed good performance from shear tests, high heat flux tests, and neutron irradiation.122 The development of joining techniques for PH CuCrCr alloy must consider the loss of mechanical strength because of overaging at high temperatures. The HIPing temperature must be reduced to be as close as possible to the aging temperature. The best results obtained so far is for HIPing at 580 C and 140 MPa for 2 h.126 A fast induction brazing technique has also been developed to minimize the
687
holding time at high temperature. Diffusion bonding of Be/CuCrZr joints gives much better high heat flux performance than brazing, and has been selected as the reference method for the European Union ITER components.120 A low-temperature Be/Cu alloy bonding process has also been developed that is compatible with both DS and PH copper alloys.124,127 In the United States, several different joint assemblies for diffusion bonding a beryllium armor tile to a copper alloy heat sink have been evaluated.128 To prevent formation of intermetallic compounds and promoting a good diffusion bond between the two substrates, aluminum or an aluminum–beryllium composite (AlBeMet-150) has been used as the interfacial material. Explosive bonding was used to bond a layer of Al or AlBeMet-150 to the copper substrate that was subsequently HIP diffusion bonded to an Al-coated beryllium tile. A thin Ti diffusion barrier (0.25 mm) was used as a diffusion barrier between the copper and aluminum to prevent the formation of Cu–Al intermetallic phases. The Be/Cu alloy joints showed good strength and failure resistance.
4.20.7 Summary High heat flux applications for fusion energy systems require high-strength, high-conductivity materials. Selection of materials for high heat flux applications must consider thermal conductivity, strength and tensile ductility, fracture toughness, fatigue and creep–fatigue, and radiation resistance. Pure copper has excellent conductivity but poor strength. PH and DS copper alloys have superior strength and sufficient conductivity, and are prime candidates for high heat flux applications in fusion reactors. These two classes of alloys have their own advantages and disadvantages with regard to fabrication, joining, and inservice performance. PH copper alloys, such as CuCrZr, are heattreatable alloys. Their properties are strongly dependent on the thermomechanical treatments. They possess high strength and high conductivity in the prime-aged condition, and good fracture toughness and fatigue properties in both nonirradiated and irradiated conditions. However, this class of alloys is susceptible to softening at high temperatures because of precipitate overaging and recrystallization. Their properties can be significantly degraded during large component fabrication because of their inability to achieve rapid quenching rates. DS copper alloys such as GlidCop Al25 have excellent thermal stability, and
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Physical and Mechanical Properties of Copper and Copper Alloys 10.
retain high strength up to temperatures near the melting point. The main disadvantages of this class of alloys are their relatively low fracture toughness and difficulty to join. The effect of neutron irradiation in copper alloys depends largely on the irradiation temperature. At irradiation temperatures below 300 C, radiation hardening occurs along with loss of strain hardening capability and complete loss of uniform elongation. Radiation hardening saturates at about 0.1 dpa in this temperature regime. At higher temperatures, radiation-induced softening can occur. Void swelling takes place between 180 and 500 C, and the peak swelling temperature is 300–325 C for neutron irradiation at damage rates near 10–7 dpa s1. PH and DS copper alloys are more resistant to void swelling than pure copper. Irradiation slightly reduces the fracture toughness of copper alloys, and the effect is stronger in CuAl25 than in CuCrZr. Irradiation has no significant effect on fatigue and creep–fatigue performance. Transmutation products can significantly change the physical properties and swelling behavior in copper alloys. Significant R&D efforts have been made to select and characterize copper alloys for high heat flux applications. The ITER Material Property Handbook provides a comprehensive database for pure copper, CuCrZr, and CuAl25. For the ITER first wall and divertor applications, CuCrZr has been selected as the prime candidate. Current focus is on fabrication, joining, and testing of large-scale components.
19. 20.
References
29. 30.
1. 2. 3. 4. 5.
6. 7. 8. 9.
Butterworth, G. J.; Forty, C. B. A. J. Nucl. Mater. 1992, 189, 237. Zinkle, S. J.; Fabritsiev, S. A. Atomic and PlasmaMaterials Interaction Data for Fusion (Supplement to the Journal Nuclear Fusion 1994, 5, 163. Davis, J. R., Ed. ASM Specialty Handbook: Copper and Copper Alloys; ASM International: Materials Park, OH, 2001; p 276. Atrens, A.; Nairn, J.; Fernee, H.; FitzGerald, K.; Skennerton, G.; Olofinjana, A. Mater. Forum 1997, 21, 57. Taubenblat, P. W.; Smith, W. E.; Graviano, A. R. In High Conductivity Copper and Aluminum Alloys; Ling, E., Taubenblat, P. W., Eds.; Metallurgical Soc of AIM: Warrendale, PA, 1984; p 19. Lu, L.; Shen, Y.; Chen, X.; Qian, L.; Lu, K. Science 2004, 304, 422. Hertzberg, R. W. Deformation and Fracture Mechanics of Engineering Materials, 4th ed.; Wiley: New York, 1995; p 142. Kanno, M. Zeitschrift fuer Metallkunde 1988, 79, 684–688. Suzuki, H.; Itoh, G.; J. Japan Inst. Metals 1984, 48, 1016.
11. 12. 13. 14. 15. 16. 17. 18.
21. 22. 23. 24. 25. 26. 27. 28.
31. 32. 33. 34. 35. 36. 37. 38.
Tang, N. Y.; Taplin, D. M. R.; Dunlop, G. L. Mater. Sci. Technol. 1985, 1, 270. Kalinin, G.; Barabash, V.; Cardella, A.; et al. J. Nucl. Mater. 2000, 283–287, 10–19. Edwards, D. J.; Singh, B. N.; Toft, P.; Eldrup, M. J. Nucl. Mater. 1998, 258–263, 978. Edwards, D. J.; Singh, B. N.; Xu, Q.; Toft, P. J. Nucl. Mater. 2002, 307–311, 439. Li, M.; Sokolov, M. A.; Zinkle, S. J. J. Nucl. Mater. 2009, 393, 36. Singh, B. N.; Edwards, D. J.; Eldrup, M.; Toft, P. Risø-R-971(EN); Risø National Laboratory: Roskilde, Denmark, 1997. Singh, B. N.; Edwards, D. J.; Eldrup, M.; Toft, P. Risø-R-937(EN) (Cu13); Risø National Laboratory: Roskilde, Denmark, 1997. Singh, B. N.; Edwards, D. J.; Eldrup, M.; Toft, P. J. Nucl. Mater. 1997, 249, 1. Singh, B. N.; Stubbins, J. F.; Toft, P. Risø-R-991(EN); Risø National Laboratory: Roskilde, Denmark, 1997. Rioja, R. J.; Laughlin, D. E. Acta Metall. 1980, 28, 1301. Guha, A. In High Conductivity Copper and Aluminum Alloys; Ling, E., Taubenblat, P. W., Eds.; The Metallurgical Society of AIME: Warrendale, PA, 1984; p 133. Ageladarakis, P. A.; O’Dowd, N. P.; Webster, G. A. Tensile and fracture toughness tests of CuNiSi at room and cryogenic temperatures; JET-R (99)01, 1999. Fabritsiev, S. A.; Pokrosky, A. S.; Zinkle, S. J.; Ostrovsky, S. E. J. Nucl. Mater. 2002, 306, 218. Singh, B. N.; Edwards, D. J.; Toft, P. J. Nucl. Mater. 1996, 238, 244. Singh, B. N.; Stubbins, J. F.; Toft, P. Risoe-R-1128(EN); Risoe National Laboratory: Roskilde, Denmark, 2000. Zinkle, S. J.; Nesterova, E. V.; Barabash, V. R.; Rybin, V. V.; Naberenkov, A. V. J. Nucl. Mater. 1994, 208, 119. Zinkle, S. J.; Horsewell, A.; Singh, B. N.; Sommer, W. F. J. Nucl. Mater. 1994, 212–215, 132. ASM Handbook: Vol. 2, Properties and Selection: Nonferrous Alloys and Special-Purpose Materials; ASM International: Materials Park, OH, 1990; p 392. Lide, D. R.; Kehiaian, H. V.; CRC Handbook of Thermophysical and Thermochemical Data; CRC Press: Boca Raton, FL, 1994; p 28. Piatti, G.; Boerman, D. J. Nucl. Mater. 1991, 185, 29. Zinkle, S. J.; Eatherly, W. S. Fusion Materials Semi-annual Progress Report; DOE/ER-0313/22; Oak Ridge National Laboratory, 1997; p 143. Li, G.; Thomas, B. G.; Stubbins, J. F. Metall. Mater. Trans. A 2000, 31A, 2491. Singh, B. N.; Edwards, D. J.; Horsewell, A.; Toft, P. Risø-R-839(EN) (Cu07); Risø National Laboratory: Roskilde, Denmark, 1995. Singh, B. N.; Toft, P. Risø-R-1008(EN); Risø National Laboratory: Roskilde, Denmark, 1998. Zinkle, S. J.; Eatherly, W. S. Fusion Materials Semiannual Progress Report; DOE/ER-0313/20; Oak Ridge National Laboratory, 1996; p 207. Fabritsiev, S. A.; Rybin, V. V.; Kazakov, V. A.; Pokrovsky, A. S.; Barabash, V. R. J. Nucl. Mater. 1992, 195, 173. Fabritsiev, S. A.; Pokrovsky, A. S. J. Nucl. Mater. 1997, 249, 250. Fabritsiev, S. A.; Pokrovsky, A. S. J. Nucl. Mater. 1997, 249, 239. Zinkle, S. J. Fusion Materials Semiannual Progress Report for Period; DOE/ER-0313/28; Oak Ridge National Laboratory, 2000; p 171.
Physical and Mechanical Properties of Copper and Copper Alloys 39. 40. 41. 42. 43. 44. 45. 46.
47. 48. 49. 50. 51.
52. 53. 54. 55. 56. 57. 58.
59. 60. 61. 62. 63. 64. 65. 66. 67. 68.
Singh, B. N.; Edwards, D. J.; Tahtinen, S. Risø-R-1436 (EN); Risø National Laboratory: Roskilde, Denmark, 2004. Edwards, D. J.; Singh, B. N.; Tahtinen, S. J. Nucl. Mater. 2007, 367–370, 904. Holzwarth, U.; Pisoni, M.; Scholz, R.; Stamm, H.; Volcan, A. J. Nucl. Mater. 2000, 279, 19. Ivanov, A. D.; Nikolaev, A. K.; Kalinin, G. M.; Rodin, M. E. J. Nucl. Mater. 2002, 307–311, 673. Suzuki, H.; Kanno, M. J. Japan Inst. Metals 1971, 35, 434. Toda, T. Trans. Japan Inst. Metals 1970, 11, 24. Toda, T. Trans. Japan Inst. Metals 1970, 11, 30. Edwards, D. J. Temperature and strain rate effects in high strength high conductivity copper alloys tested in air; Fusion Reactor Materials Semiannual Progress Report; DOE/ER-0313/23; Oak Ridge National Laboratory, 1997; p 213. Stephens, J. J.; Bourier, R. J.; Vigil, F. J.; Schmale, D. T. Sandia National Lab Report, SAND88-1351, 1988. Alexander, D. J.; Zinkle, S. J.; Rowcliffe, A. F. J. Nucl. Mater. 1999, 271–272, 429. Suzuki, R.; Saito, M.; Hatano, T. Fusion Sci. Technol. 2003, 44, 242. Tahtinen, S.; Pyykkonen, M.; Karjalainen-Roikonen, P.; Singh, B. N.; Toft, P. J. Nucl. Mater. 1998, 258–263, 1010. Nadkarni, A. V. In High Conductivity Copper and Aluminum Alloys; Ling, E., Taubenblat, P. W., Eds.; The Metallurgical Society of AIME: Warrendale, PA, 1984; p 77. Broyles, S. E.; Anderson, K. R.; Groza, J. R.; Gibeling, J. G. Metall. Mater. Trans. 1996, 27A, 1217. Leedy, K. D. Ph.D. Thesis, University of Illinois at Urbana-Champaign, Urbana, IL, 1997. Singh, B. N.; Li, M.; Stubbins, J. F.; Johansen, B. S. Risø-R-1528 (EN); Risø National Laboratory: Roskilde, Denmark, 2005. Li, M.; Stubbins, J. F. J. ASTM Int. 2005, 2, 251. Li, M.; Singh, B. N.; Stubbins, J. F. J. Nucl. Mater. 2004, 329–333, 865. Wu, X.; Pan, X.; Singh, B. N.; Li, M.; Stubbins, J. F. J. Nucl. Mater. 2007, 367–370, 984. Tahtinen, S. Crack growth properties of unirradiated and irradiated CuAl25-IG0 and CuCrZr Alloys and HIP Joints with 316L(N)-IG0 Steel; ITER Doc. G73 RE 1298–05–26 F 1, 1997. Nomura, Y.; Suzuki, R.; Saito, M. J. Nucl. Mater. 2002, 307–311, 681. Zinkle, S. J. In Effects of Radiation on Materials, 15th International Symposium, ASTM STF 1125; Stoller, R. E., et al., Eds.; ASTM: Philadelphia, PA, 1992; p 813. Fabritsiev, S. A.; Zinkle, S. J.; Singh, B. N. J. Nucl. Mater. 1996, 233–237, 127. Fabritsiev, S. A.; et al. Fusion Materials Semi-annual Progress Report for Period Ending December 31, 1995; DOE/ER-0313/19 (April 1996); 1996; pp 177–188. Zinkle, S. J. J. Phys. F Metals Phys. 1988, 18, 377. Edwards, D. J.; Singh, B. N.; Bilde-Sørensen, J. B. J. Nucl. Mater. 2005, 342, 164. Zinkle, S. J.; Gibson, L. T. Fusion Materials Semi-annual Progress; Report DOE/ER-0313/27; Oak Ridge National Laboratory, 1999; p 163. Zinkle, S. J.; Victoria, M.; Abe, K. J. Nucl. Mater. 2002, 307–311, 31. Fabritsiev, S. A.; Pokrovsky, A. S. J. Nucl. Mater. 2009, 386–388, 268. Singh, B. N.; Edwards, D. J.; Toft, P. J. Nucl. Mater. 2001, 299, 205.
69. 70. 71. 72.
73. 74. 75. 76. 77. 78.
79. 80.
81. 82. 83. 84. 85. 86. 87. 88. 89. 90. 91. 92. 93.
94. 95. 96. 97.
689
Fabritsiev, S. A.; Pokrovsky, A. S. J. Nucl. Mater. 2002, 307–311, 431. Pan, X.; Wu, X.; Li, M.; Stubbins, J. F. J. Nucl. Mater. 2004, 329–333, 1088. Fabritsiev, S. A.; Pokrovsky, A. S.; Zinkle, S. J.; et al. J. Nucl. Mater. 1996, 233–237, 526. Tahtinen, S.; Pyykkonen, M.; Singh, B. N.; Toft, P. In 19th International Symposium on Effects of Radiation on Materials, ASTM STP 1366; Hamilton, M. L., et al. Eds.; American Society for Testing and Materials: West Conshohocken, PA, 2000; p 1243. Li, M.; Stubbins, J. F. Fusion Sci. Technol. 2003, 44, 186. Leedy, K. D.; Stubbins, J. F.; Singh, B. N.; Garner, F. A. J. Nucl. Mater. 1996, 233–237, 547. Singh, B. N.; Stubbins, J. F.; Toft, P. J. Nucl. Mater. 1999, 275, 125. Srivatsan, T. R.; Anand, S. Eng. Fract. Mech. 1993, 46, 183. Singh, B. N.; Ta¨htinen, S.; Moilanen, P.; et al. Risø-R-1571(EN); Risø National Laboratory: Roskilde, Denmark, 2007. Ibragimov, S. S.; Aitkhozhin, E. S.; Pyatiletov, Y. S. In 13th International Symposium (Part II), ASTM STP 956; Garner, F. A., Henager, C. H., Igata, N., Eds.; ASTM: Philadelphia, PA, 1987; p 5. Jung, P. J. Nucl. Mater. 1993, 200, 138. Pokrovsky, A. S.; Barabash, V. R.; Fabritsiev, S. A.; et al. Fusion Materials Semi-annual Progress Report; DOE/ER-0313/22; Oak Ridge National Laboratory, 1996; p 194. Singh, B. N.; Horsewell, A.; Gelles, D. S.; Garner, F. A. J. Nucl. Mater. 1992, 191–194, 1172. Witzig, W. F. J. Appl. Phys. 1952, 23, 1263. Zinkle, S. J.; Farrell, K. J. Nucl. Mater. 1989, 168, 262. Zinkle, S. J.; Farrell, K.; Kanazawa, H. J. Nucl. Mater. 1991, 179–181, 994. Gamer, F. A.; Hamilton, M. L.; Shikama, T.; Edwards, D. J.; Newkirk, J. W. J. Nucl. Mater. 1992, 191–194, 386. Zinkle, S. J.; Lee, E. H. Metall. Mater. Trans. 1990, 21A, 1037. Zinkle, S. J.; Wolfer, W. G.; Kulcinski, G. L.; Seitzman, L. E. Phil. Mag. 1987, A55, 127. Mukouda, I.; Shimomura, Y.; Iiyama, T.; et al. J. Nucl. Mater. 2000, 283–287, 302. Zinkle, S. J. Fusion Reactor Materials Semiannual Progress Report; DOE-ER-0313/3; Oak Ridge National Laboratory, 1987; p 86. Xu, Q.; Yoshiie, T.; Sato, K. J. Nucl. Mater. 2009, 386–388, 363. Edwards, D. J.; Anderson, K. R.; Garner, F. A.; Hamilton, M. L.; Stubbins, J. F.; Kumar, A. S. J. Nucl. Mater. 1992, 191–194, 416. Gamer, F. A.; Brager, H. R.; Anderson, K. R. J. Nucl. Mater. 1991, 179–181, 250. Edwards, D. J.; Newkirk, J. W.; Garner, F. A.; Hamilton, M. L.; Nadkarni, A. Proceedings of the 16th ASTM International Symposium on the Effect of Radiation on Materials, ASTM STP 1175, June 1992; Kumar, A. S., Gelles, D. S., Nanstad, R. K., Little, E. A., Eds.; 1994; p 1041. Spitznagel, J. A.; Doyle, N. J.; Choke, W. J.; et al. Nucl. Inst. Meth. 1986, B16, 279. Gamer, F. A.; Brager, H. R.; Anderson, K. R. J. Nucl. Mater. 1991, 179–181, 250. Singh, B. N.; Zinkle, S. J. J. Nucl. Mater. 1993, 206, 212. Zinkle, S. J.; Knoll, R. W. A Literature Review of Radiation Damage Data for Copper and Copper Alloys; Report UWFDM-578; University of Wisconsin Fusion Technology Institute, June 1984.
690 98. 99. 100. 101. 102. 103. 104. 105. 106. 107. 108. 109. 110. 111. 112. 113. 114. 115. 116.
Physical and Mechanical Properties of Copper and Copper Alloys Zinkle, S. J.; Matsukawa, Y. J. Nucl. Mater. 2004, 329–333, 88. Ruhle, M.; Crump, J. C., III Phys. Status Solidi A 1970, 2, 257. Zinkle, S. J.; Kulcinski, G. L.; Knoll, R. W. J. Nucl. Mater. 1986, 138, 46. Eyre, B. L. J. Phys. F Metal Phys. 1973, 3, 422. Luft, A. Prog. Mater. Sci. 1991, 35, 97. Wechsler, M. S. In The Inhomogeneity of Plastic Deformation; Reed-Hill, R. E., Ed.; American Society for Metals: Metals Park, OH, 1972; p 19. Greenfield, L. G.; Wilsdorf, H. G. F. J. Appl. Phys. 1961, 32, 827. Diehl, J.; Hinzner, F. Phys. Status Solidi 1964, 7, 121. Essmann, U.; Mader, S.; Seeger, A. Z. Metallk. 1961, 52, 443. Essmann, U.; Seeger, A. Phys. Status Solidi 1964, 4, 177. Sharp, J. V. Phil. Mag. 1967, 19, 77. Sharp, J. V. In Proceedings of the 4th European Regional Conference on Electron Microscopy; Bocciarelli, D. D., Ed.; Rom, 1968; p 417. Sharp, J. V. Radiat. Eff. 1972, 23, 181. Howe, L. M. Radiat. Eff. 1974, 23, 181. Adamson, R. B. Phil. Mag. 1968, 17, 681. Adamson, R. B. Acta Metall. 1969, 17, 1169. Sharp, J. V. Acta Metall. 1974, 22, 449. Edwards, D. J.; Singh, B. N. J. Nucl. Mater. 2004, 329–333, 1072. Johnson, E.; Hirsch, P. B. Philos. Mag. A 1981, 43, 157.
117. 118. 119. 120. 121. 122. 123. 124. 125. 126. 127. 128.
Robach, J. S.; Robertson, I. M.; Wirth, B. D.; Arsenlis, A. Phil. Mag. 2003, 83, 955. Davis, J. R. Ed. ASM Specialty Handbook: Copper and Copper Alloys; ASM International: Materials Park, OH, 2001; p 276. Odegard, B. C.; Kalin, B. A. J. Nucl. Mater. 1996, 233–237, 44. Lorenzetto, P.; Boireau, B.; Boudot, C.; et al. Fusion Eng. Design 2008, 83, 1015. Samal, P. K. In The Metal Science of Joining; Cieslak, M. J., Ed.; The Minerals, Metals and Materials Society: Warrendale, PA, 1992; p 295. Lorenzetto, P.; Boireau, B.; Boudot, C.; et al. Fusion Eng. Design 2005, 75–79, 291. Tsuchiya, K.; Kawamura, H. J. Nucl. Mater. 1996, 233–237, 913. Sherlock, P.; Peacock, A. T.; McGallum, A. D. Fusion Eng. Design 2005, 75–79, 377. Gervash, A.; Mazul, I.; Yablokov, N. Fusion Eng. Design 2001, 56–57, 381. Lorenzetto, P.; Cardella, A.; Daenner, W.; et al. Fusion Eng. Design 2002, 61–62, 643. Sherlock, P.; Peacock, A. T.; Rodig, M. Fusion Eng. Design 2007, 82, 1806. Pukskar, J. D.; Goods, S. H.; Cadden, C. H. Trends in Welding Research, Proceedings of the 8th International Conference; David, S. A., DebRoy, T., DuPont, J. N., Koseki, T., Smartt, H. B., Eds.; ASM International, Materials Park, OH, 2009; p 604.
4.21 Ceramic Coatings as Electrical Insulators in Fusion Blankets T. Muroga National Institute for Fusion Science, Oroshi, Toki, Gifu, Japan
ß 2012 Elsevier Ltd. All rights reserved.
4.21.1
Introduction
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4.21.2 4.21.3 4.21.3.1 4.21.3.2 4.21.3.3 4.21.3.3.1 4.21.3.3.2 4.21.3.3.3 4.21.3.3.4 4.21.3.4 4.21.3.5 4.21.3.6 4.21.4 References
Magnetohydrodynamic Issues and the Requirement for Insulator Coatings Development of Insulator Coating for Liquid Li Blanket In Situ Formation and Healing with CaO In Situ AlN Coating Er2O3 and Y2O3 as New Candidates Scoping by bulk immersion tests In situ coating with Er2O3 Physical coating processes Other coating technologies Two-Layer Coatings Radiation Effects FCI Concept as an Alternative to Insulator Coating Summary and Remaining Issues
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Abbreviations EB-PVD FCI IFMIF ITER-TBM
MHD MOCVD MOD RF RIC TBR TEM
Electron beam-physical vapor deposition Flow channel insert International Fusion Materials Irradiation Facility International Thermonuclear Experimental Reactor-Test Blanket Module Magnetohydrodynamic Metal-organic chemical vapor deposition Metal-organic deposition Radiofrequency Radiation-induced conductivity Tritium breeding ratio Transmission electron microscope
4.21.1 Introduction The use of a ceramic coating for electrical insulation is a key technology for fusion blanket systems using liquid metals as breeding and coolant material and solid metals as the structural material. Particularly for the blanket system using liquid lithium and
vanadium alloys (Li–V blankets), coating development is a major feasibility issue (see also Chapter 4.12, Vanadium for Nuclear Systems for vanadium alloy and liquid Li blankets). Overviews of the coating development for liquid lithium blankets are available in recent publications.1–3 It should, however, be noted that, with the development of a more fundamental understanding of coating behavior and blanket design, there has been a paradigm shift in coating development. This chapter describes the present status of insulator coating R&D, in addition to a historical overview of its development.
4.21.2 Magnetohydrodynamic Issues and the Requirement for Insulator Coatings Breeding blankets for fusion reactors are categorized into solid breeder and liquid breeder concepts. The liquid breeder blankets have certain advantages over the solid breeder blankets such as continuous chemical control of the breeding material including isotopic control of Li, impurity control and tritium recovery, and immunity to irradiation effects. However, some issues such as compatibility of the breeder with structural materials are more serious for liquid breeder blankets4. In addition, blanket structure 691
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can be simplified significantly if the liquid breeder functions as the coolant as well (self-cooled liquid breeder blanket). At present, the major candidate liquid breeder materials are Li, Li–Pb, and molten salt Flibe (LiF–BeF2). In the cases of self-cooled liquid Li and Li–Pb blankets in magnetic confinement fusion systems, the high-speed flow of these materials perpendicular to the strong magnetic field causes an electric current, which then produces an electromagnetic force as a result of interaction with the magnetic field. This force changes the velocity profile in the cooling ducts and acts to retard the coolant flow, leading to what is called a magnetohydrodynamic (MHD) pressure drop. This process is schematically shown in Figure 1(a). The MHD pressure drop may result in loss of flow control and mechanical stresses exceeding the allowable limits of the structural materials. The problems arising from the MHD pressure drop are critical feasibility issues for self-cooled liquid metal breeder blanket concepts with metallic structures. The quantification of the MHD pressure drop requires a rather complex numerical analysis. However, in simple cases such as straight and constant area cross-section flow in conductive ducts with a uniform magnetic field in a traverse direction, the pressure gradient along the flowing direction, dp/dx, is given as follows5: dp=dx ¼ ksUB 2 where s, U, and B are the electrical conductivity, flow velocity of the liquid metal, and magnetic flux density, respectively, and k is a positive function of electrical conductivity of the wall. The equation implies that the MHD pressure drop is an issue in the case of high magnetic field and high velocity flow of conductive liquid metals. In the case of a low flow rate such as would occur in a helium-cooled Li–Pb blanket, the MHD pressure drop will not be an issue.
To reduce the MHD pressure drop, optimization of the coolant flow path by enhancing the flow fraction parallel to the magnetic field may have some effect. However, a more effective way to reduce the MHD pressure drop would be to electrically insulate the coolant flow from the surrounding walls.6 The reduction of MHD pressure drop by an insulator coating is schematically illustrated in Figure 1(b). The requirements for the coating can be summarized as follows: 1. compatibility with liquid breeder under flowing conditions with a temperature gradient, 2. high electrical resistivity under irradiation, 3. robustness and/or an effective self-healing capability, 4. potential for covering large and complex surfaces, and 5. fundamental requirements for in-vessel materials such as radiation resistance, low activation properties, and low tritium inventories in blanket conditions. Quantitative evaluation of the required electrical resistance and an allowable crack fraction are subject to overall blanket design including flow channel structures. A recent model calculation showed that the ratio of electrical resistivity of the insulator to the wall needs to be ≳106 and crack areal fraction to be ≲106 to maintain the pressure drop within tolerable range, assuming Li wets cracks.7,8 For the Li–Pb blankets, the insulator coating should be a critical issue if a self-cooled Li–Pb blanket with metallic structural materials is to be designed. However, current blanket design options with Li–Pb are (1) helium cooled with slow-flowing Li–Pb, (2) dualcoolant Li–Pb with fast-flowing Li–Pb but electrically insulated from the wall by a SiC/SiC flow channel insert (FCI), or (3) self-cooled Li–Pb using SiC/SiC as the structural material. None of these concepts needs the insulator coating. However, development of a ceramic coating, necessary mostly for tritium permeation reduction and possibly for corrosion protection, is still a critical issue.9
Magnetic field Vanadium duct
4.21.3 Development of Insulator Coating for Liquid Li Blanket Lorentz force
Liquid Li flow Insulator coating
Figure 1 Schematic illustration of magnetohydrodynamic pressure drop (left) and the role of insulator coating (right).
Liquid Li is a strong reducing agent, and thus a coating layer of many common oxides is not stable on the wall material. Therefore, intentional coating with insulator ceramics, which are stable in liquid Li, is necessary. Figure 2 shows free energy for oxidation and nitridation for various elements. From the
Ceramic Coatings as Electrical Insulators in Fusion Blankets
V-alloy (O-doped)
ΔG (kJ per mol atomic O)
–450 Al2O3
693
Liquid Li (M-doped)
MgO
–500
BeO
Li2O
Er2O3
CaO
OV
MLi
–550 Y2O3 –600 200
O2–
M¥+ M2O¥
300
400
500 T (⬚C)
600
700
800
ΔG (kJ per mol atomic N)
0 –50
Li3N
Oxide coating
Si3N4
–100 –150
Figure 3 Schematic illustration of the mass transport for in situ oxide coating in Li. Ov: oxygen in vanadium substrate; MLi: metal doped in Li for producing oxide coating.
BN
–200 AlN –250 –300 200
TiN 300
400
500 T (⬚C)
600
700
800
Figure 2 Free energy of oxide and nitride formation for selected ceramics.
thermodynamic viewpoint, CaO, Y2O3 and Er2O3, and AlN are expected to be stable. (Note that the resistivity of TiN is not sufficient for the coating.) The instability, however, is a function of the impurity level of Li.10 Early development efforts focused on CaO and AlN by in situ formation during exposure to chemically controlled lithium. 4.21.3.1 In Situ Formation and Healing with CaO This technology aims at using a corrosion product that forms at the interface between Li and structural materials as an insulating layer. By careful control of the corrosion reaction, the insulation layer can be formed uniformly on the inner surfaces of complex components. The corrosion layer may also be formed on the cracked area of the coating, thereby repairing insulation defects. The in situ coating with CaO and AlN have been studied in the United States11 and the Russian Federation,12 respectively.
A CaO insulator layer forms during the immersion of vanadium alloys in Ca-doped Li. For enhancing this reaction and stabilizing the layer, the O level in the alloy was increased by prior doping. The process is schematically shown in Figure 3. Careful control of the Ca level in Li and the O charging condition of the vanadium alloys made in situ formation and healing of the CaO coating possible at <500 C.13 Because of the successful demonstration, in situ CaO coating was adopted as the means to mitigate the MHD pressure drop in the ARIES-RS power plant study.14 This technology was, however, shown to be inapplicable at temperatures exceeding 600 C because of the unacceptably short lifetime of the coating induced by increased dissolution.2 4.21.3.2
In Situ AlN Coating
AlN is an alternative candidate for the in situ coating because of its high electrical resistivity and high compatibility with Li. In this case, Al and N are doped into Li to enhance AlN formation at the surface of vanadium alloy substrates. The coating has shown sufficient resistivity and stability under thermal cycling.12 In this process, however, Li needs to be saturated with N to prevent the dissolution of the AlN layer. With a high N level in Li, vanadium alloy not covered with the AlN layer will get N from Li and become brittle. This issue was observed in high-temperature
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Ceramic Coatings as Electrical Insulators in Fusion Blankets
capsule tests with Mo and vanadium alloys.15,16 Thus, unless the vanadium alloy channel walls are fully covered with AlN, the use of this coating technique seems problematic. For the same reason, the use of AlN layers, sheets, or plates in Li as the insulator requires full coverage of the vanadium alloy structures with AlN. 4.21.3.3 Er2O3 and Y2O3 as New Candidates 4.21.3.3.1 Scoping by bulk immersion tests
Exploration of new candidates having superior performance compared to CaO and AlN was carried out using static immersion tests. Figure 4 shows the mass loss of various insulator ceramics due to exposure to static Li at high temperatures. As predicted by thermodynamics, Er2O3 and Y2O3 showed stability superior to that of CaO.2,15 For these materials, formation of LiXO2 (X ¼ Er or Y) during exposure to Li was reported,17,18 although the impact of these changes on the coating properties remains to be assessed. In particular, the effects of Li flow on the stability of the corrosion products on the surfaces will be the key issue.
On the other hand, formation of Er2O3 was confirmed.19 Because the solubility of Er in Li is much lower than that of CaO, the stability of the coating, once formed, is much higher compared with the CaO in situ coating. The cross-section of the coating with compositional profile and coating thickness with time and temperature are shown in Figures 5 and 6, respectively. In the effort to optimize the precharging condition of oxygen, the microstructural process for restoring oxygen in vanadium alloy substrate was clarified.20 Figure 7 shows the depth profile of hardness before and after oxygen charging, after heat treatment and
V-4Cr-4Ti
V
Er
4.21.3.3.2 In situ coating with Er2O3
Based on the experience with CaO, in situ coating with Er2O3 and Y2O3 was attempted on V–4Cr–4Ti by doping Er or Y in Li and O into V–4Cr–4Ti. Coating with Y2O3 was shown to be difficult, probably because there was almost no solubility of Y in Li.
30
10 μm Figure 5 Cross-section of in situ Er2O3 coating on V–4Cr–4Ti after exposure to Li(Er) for 300 h at 600 C. Reproduced from Yao, Z.; Suzuki, A.; Muroga, T.; Katahira, K. J. Nucl. Mater. 2004, 329–333, 1414–1418, with permission from Elsevier.
Li 1000 h
V-4Cr-4Ti (NIFS-HEAT-2) Oxidation - 700 ⬚C, 6 h Annealing - 700 ⬚C, 16 h Exposure - Li (Er)
1.5
20 CaO (p-crystal)
15 10
CaO (s-crystal)
5
CaZrO3 Y2O3
0 –5
Thickness of Er2O3 (μm)
Mass loss (mg cm–2)
25
1.0
500 ⬚C 550 ⬚C 600 ⬚C 650 ⬚C 700 ⬚C
0.5
Er2O3 0.0
–10 500
600 700 Li temperature (⬚C)
800
Figure 4 Change of mass after exposure to static Li for 1000 h for a bulk of candidate ceramics. Adapted from Pint, B. A.; Tortorelli, P. F.; Jankowski, A.; et al. J. Nucl. Mater. 2004, 329–333, 119–124, with permission from Elsevier.
0
100 200 300 400 500 600 700 800 Exposure time (h)
Figure 6 Growth of the Er2O3 layer during the in situ coating. Reproduced from Yao, Z.; Suzuki, A.; Muroga, T.; Yeliseyeva, O. I.; Nagasaka, T. Fusion Eng. Des. 2006, 81, 951–956, with permission from Elsevier.
Ceramic Coatings as Electrical Insulators in Fusion Blankets
695
(b)
(c) [200] [020]
Ti-O (d)
Vickers hardness (HV)
1500 As-received (a) Oxidized - 6 h, 700 ⬚C (b) Oxidized - 6 h, annealed - 16 h, 700 ⬚C (c) Oxidized - 6 h, annealed - 16 h, Li(Er) - 100 h, 700 ⬚C (d)
1000
V - 4Cr - 4Ti (NIFS-HEAT-2)
500
(a) 0
0.2 μm
Ti-C-O-N
0
50
100 150 200 250 Depth from surface (μm)
300
Figure 7 Depth distribution of hardness and transmission electron microscope microstructure near the surfaces in the V–4Cr–4Ti substrate. (a) Before oxidation, (b) after oxidation, (c) after annealing, and (d) after in situ coating (coating was removed). Reproduced from Yao, Z.; Suzuki, A.; Muroga, T.; Yeliseyeva, O. I.; Nagasaka, T. Fusion Eng. Des. 2006, 81, 951–956, with permission from Elsevier.
after in situ coating, together with the transmission electron microscope (TEM) microstructures near the surfaces. The hardness is known to follow the approximate level of O in V–4Cr–4Ti.21 After oxidation for 6 h at 700 C, the surface was covered with a complex oxide layer. After the subsequent heat treatment for 16 h at 700 C, the matrix was composed of a high density of needle-shaped Ti–O (mostly TiO2) precipitates oriented in the h100i directions (net structure). This structure was most prominently observed after annealing at 700 C. This is consistent with the results of the precipitation study, which showed that, although Ti interacts with impurity O already at 200 C, Ti–O precipitates start to form at 700 C in V–4Cr–4Ti alloys.22 Figure 7 also shows that the net structure near the surface disappeared after exposure to Li for 100 h at 700 C because of the loss of oxygen. A recent study showed in situ healing capabilities with Er2O3, but further optimization of the process is required to obtain a reliable healing function.23 The mechanism of the net structure formation and supply of oxygen for the coating was elucidated using a kinetic model. The model successfully explained the experimental trends.24
4.21.3.3.3 Physical coating processes
More conventionally deposited coatings of Er2O3 or Y2O3 were developed such as arc source plasma deposition,25 electron beam-physical vapor deposition (EB-PVD),2 and radiofrequency (RF) sputtering.26 Because of the variation in the quality of the coating and the test conditions, reported stability in liquid lithium was different in different tests. Er2O3 produced by EB-PVD was heavily damaged in Li already at 500 C,27 and Er2O3 and Y2O3 coatings produced by RF sputtering were also exfoliated at 500 C.28 However, Er2O3 fabricated by arc source plasma deposition showed promising results, as shown in Figure 8. Deposition on a higher temperature substrate produced a highly crystalline Er2O3 coating, which was shown to be stable in Li for 1000 h at 700 C.1 The stability in Li is known to be enhanced by improving the purity and crystallinity of the coating. An oxide layer at the coating/substrate interface may cause extensive exfoliation because Li introduced through cracks would preferentially attack the oxidized interface. 4.21.3.3.4 Other coating technologies
The efforts in using the physical coating processes explained so far are essential for establishing the
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Ceramic Coatings as Electrical Insulators in Fusion Blankets
coating concepts. However, further efforts are also necessary to enhance the engineering feasibility of coating on large and complex surfaces. For this purpose, the sol–gel method (in other words, metal-organic deposition, MOD) and metal-organic chemical vapor deposition (MOCVD) have been explored. Er2O3 coating was formed on stainless steel by the sol–gel method. Crystallinity of the coating depended on annealing atmosphere and temperature.29 MOCVD was applied to coating Er2O3 on V-alloys and other materials. Successful coating on the inner surface of a tube was demonstrated.30 4.21.3.4
Two-Layer Coatings
Forming an overlayer on the ceramic coating with ductile metals is an alternative approach to in situ healing to compensate for cracks in the insulator layer. Current efforts for fabricating the two-layer
Intensity (a.u.)
XRD (low angle) Substrate temperature
coating use pure V, V-alloys, and pure Fe, as shown in Figure 9. In these cases, only modest compatibility with Li would be required for the insulator ceramics. Instead, the compatibility of V, V-alloys, and pure Fe with Li needs to be verified. It is also necessary to avoid shorting between the overlayer and the substrate. The Er2O3–V and Y2O3–V two-layer coatings on a V-alloy substrate fabricated using the EB-PVD process showed sufficient resistivity in Li.31 Recently, a mono-metallic thermal convection Li loop was constructed using V–4Cr–4Ti alloy. The two-layer coatings were tested in addition to uncoated V-alloy substrates at 700 C for >2000 h in flowing Li. The tentative results showed that the corrosion loss, degradation of the coating, and mechanical property changes to the substrate were small.32 However, a full characterization remains to be completed.
Exfoliated
RT 850 K 15
30
Er2O3
45 2q
60
773 K
873 K 973 K 1000 h in liquid Li
Figure 8 Change of Er2O3 coating on V–4Cr–4Ti by exposure to Li at 773, 873, and 973 K for 1000 h. The coating was made using arc source plasma deposition method with substrate temperature at room temperature and 850 K. Crystalline structure of Er2O3 was observed only in the case of the high substrate temperature. Remarkable change was not observed in the coating at high substrate temperature after exposure to Li. Reproduced from Muroga, T.; Chen, J. M.; Chernov, V. M.; et al. J. Nucl. Mater. 2007, 367–370, 780–787, with permission from Elsevier.
Fe
V outer layer
Er2O3 Er2O3 V-4Cr-4Ti substrate
5 μm
V-4Cr-4Ti substrate
5 μm
Figure 9 Cross-section of Er2O3–V two-layer coating produced by electron beam-physical vapor deposition and Er2O3–Fe two-layer coatings produced by radiofrequency sputtering on V–4Cr–4Ti. Reproduced from Pint, B. A.; Moser, J. L.; Jankowski, A.; Hayes, J. J. Nucl. Mater. 2007, 367–370, 1165–1169, with permission from Elsevier. Suzuki, A.; Pint, B.; Li, M.; Muroga, T.; Terai, T. Compatibility of MHD coating candidate materials with liquid lithium under neutron irradiation. Presented at 13th International Conference of Fusion Reactor Materials, Nice, France, Dec 10–14, 2007, ICFRM2007/488.
Ceramic Coatings as Electrical Insulators in Fusion Blankets
4.21.3.5
Radiation Effects
V–4Cr–4Ti were not considered in this calculation.) With the coating, the dose rate increases because of the contribution from Er but still satisfies the remote recycling limit.38 Er is a neutron-absorbing element and can reduce the TBR especially for an in situ coating where Er is doped into Li. However, because only 0.15% Er is needed in Li for the in situ coating,19 the impact of Er in Li on the TBR is not an issue.39,40
Radiation-induced conductivity (RIC) is the loss of insulation only during irradiation, which is a common issue for insulator ceramics in irradiation environments. Historically, evaluation of RIC has been carried out mostly for Al2O333 (see also Chapter 4.22, Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors). Figure 10 shows RIC as a function of the dose rate for Er2O3, Y2O3, and CaZrO3 bulk specimens for 14 MeV neutron, fission neutron, and g-ray irradiation, in comparison with data on Al2O3.34,35 For Y2O3 and AlN, results for the coating are also shown. The RIC of the candidate materials of Er2O3, Y2O3, and CaZrO3 are comparable with Al2O3. According to these results and the expected dose rate in an Li–V fusion blanket36, the expected induced conductivity in the fusion blanket condition is much lower than the maximum allowable value of 102 S m1.37 The effects of the nuclear properties of Er on radioactivity and tritium breeding ratio (TBR) of V–Li blankets were also investigated. Figure 11 shows the contact dose rate of a V–Li blanket with and without neutron multiplier Be in the cases of (1) no coating, (2) 10 mm, and (3) 1 mm coating with Er2O3. Without the coating, the dose rate was dominated by V–4Cr–4Ti substrates reaching the hands-on recycle limit after several decades of cooling. (Note that impurities in
4.21.3.6 FCI Concept as an Alternative to Insulator Coating In recent years, FCIs made of ceramic materials such as silicon carbide composite (SiC/SiC) have been proposed for both electrical and thermal insulation between the liquid breeder and the channel walls. Although the electrical resistivity of SiC/SiC is lower than that of the candidate insulator coating materials, the use of an insert that is much thicker than the coating allows for sufficient reduction of induced electrical currents. An FCI is attractive for application to dualcoolant Li–Pb blanket concepts in which heat is removed by both a high flow rate of He and Li–Pb.41,42 This concept, however, may not be applied to self-cooled liquid metal blankets because these blankets need to avoid thermal insulation between the coolant and the first wall or the blanket structural components and coolants.
g-ray
DT neutron
697
First wall in V-Li blanket
Fission neutron
Radiation-induced conductivity (S m–1)
10–5 Er2O3
10–6 10–7
Er2O3 CaZrO3
10–8 Y2O3
10–9 10–10 10–11 10–12 10–13
Y2O3
+ CaZrO3
Er2O3
Past data on Al2O3
AIN Y2O3
10–14 10–4
10–3
Bulk samples Coatings
10–2
10–1
Variation range
100 101 102 –1) Dose rate (Gy s
103
104
105
Figure 10 Radiation-induced conductivity as a function of dose rate measured by DT-neutron (FNS), fission neutron (JMTR), and g-ray irradiations for bulk and coating of magnetohydrodynamic insulator coating candidate ceramics. Previous data on Al2O3 are also shown in comparison. Adapted from Tanaka, T.; Shikama, T.; Narui, M.; Tsuchiya, B.; Suzuki, A.; Muroga, T. Fusion Eng. Des. 2005, 75–79, 933–937, with permission from Elsevier.
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Ceramic Coatings as Electrical Insulators in Fusion Blankets
104
2 Neutron wall load:1.5 MW m–2 10 years operation
102
Dose rate (Sv h–1)
100 10–2
Coating thickness 10 μm
Structural materials (V-4Cr-4Ti)
1 μm
Recycling limit
10–4 Hands on limit 10–6 10–8
Without coating (Li-V) Without coating (Li-Be-V) With coating (Li-V) With coating (Li-Be-V)
10–10 10–4
10–3
10–2 10–1 100 101 Time after shutdown (year)
102
103
Figure 11 Contact dose rate of V–Li blanket with and without neutron multiplier Be in the cases of (1) no coating, (2) 10 mm, and (3) 1 mm coating with Er2O3. Reproduced from Muroga, T.; Tanaka, T.; Kondo, M.; Nagasaka, T.; Xu, Q. Fusion Sci. Technol. 2009, 56, 897–901, with permission from Elsevier.
4.21.4 Summary and Remaining Issues Development of an insulator coating is the critical feasibility issue for self-cooled Li blankets. A very limited number of ceramics are stable under long-term exposure to Li at high temperatures. The present candidates are Er2O3, Y2O3, and, under limited conditions, AlN. In addition to concept verification studies using several physical coatings, chemical or reactive coatings have been explored as a potential means to cover large and complex surfaces. Considering the very low tolerable crack fraction (<106), in situ healing and two-layer coating with metallic overlayer are promising candidates. Recent loop tests with high impurity control have demonstrated that a two-layer coating with a vanadium overlayer is stable in flowing lithium. Thus, verification of in situ coatings, including the healing function, and two-layer coatings for large and complex surfaces would be the next necessary step in development. Er2O3 is also of interest as a candidate for the required tritium permeation barrier coating.43 Thus, collaborative effort to develop this material for application in fusion blankets, either for electrical insulation or for tritium permeation reduction, seems to be an efficient development strategy. Studies on the effects of radiation on the coating (resistivity and mechanical properties) are limited and
further research is necessary. In particular, the permanent effects of radiation need to be assessed by controlled irradiation experiments. In addition to the use of fission neutrons and charged particles, the opportunity to use the International Thermonuclear Experimental Reactor-Test Blanket Module (ITER-TBM) and International Fusion Materials Irradiation Facility (IFMIF) is anticipated for integrated coating function tests and high fluence 14 MeV neutron irradiation tests, respectively.
References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10.
Muroga, T.; Chen, J. M.; Chernov, V. M.; et al. J. Nucl. Mater. 2007, 367–370, 780–787. Pint, B. A.; Tortorelli, P. F.; Jankowski, A.; et al. J. Nucl. Mater. 2004, 329–333, 119–124. Smith, D. L.; Konys, J.; Muroga, T.; Evitkhin, V. J. Nucl. Mater. 2002, 307–311, 1314–1322. Malang, S.; Leroy, P.; Casini, G. P.; Mattas, R. F.; Strebkov, Y. Fusion Eng. Des. 1991, 16, 95–109. Kirillov, I. R.; Reed, C. B.; Barleon, L.; Miyazaki, K. Fusion Eng. Des. 1995, 27, 553–569. Liu, Y. Y.; Smith, D. L. J. Nucl. Mater. 1986, 141–143, 38–43. Hashizume, H.; Usui, Y.; Kitajima, S.; Hida, Y.; Sagara, A. Fusion Eng. Des. 2002, 61/62, 251–254. Hashizume, H. Fusion Eng. Des. 2006, 81, 1431–1438. Konys, J.; Aiello, A.; Benamati, G.; Giancarli, L. Fusion Sci. Technol. 2005, 47, 844–850. Hubberstey, P.; Sample, T. J. Nucl. Mater. 1997, 248, 140–146.
Ceramic Coatings as Electrical Insulators in Fusion Blankets 11. Smith, D. L.; Natesan, K.; Park, J. H.; Reed, C. B.; Mattas, R. F. Fusion Eng. Des. 2000, 51–52, 185–192. 12. Vertkov, A. V.; Evtikhin, V. A.; Lyublinski, I. E. Fusion Eng. Des. 2001, 58–59, 731–735. 13. Park, J. H.; Kassner, T. F. J. Nucl. Mater. 1996, 233–237, 476–481. 14. Najmabadi, F. The ARIES Team. Fusion Eng. Des. 1997, 38, 3–25. 15. Pint, B. A.; DeVan, J. H.; DiStefano, J. R. J. Nucl. Mater. 2002, 307–311, 1344–1350. 16. Suzuki, A.; Muroga, T.; Pint, B. A.; Yoneoka, T.; Tanaka, S. Fusion Eng. Des. 2003, 69, 397–401. 17. Nagura, M.; Suzuki, A.; Muroga, T.; Terai, T. Fusion Eng. Des. 2009, 84, 1384–1387. 18. Terai, T.; Mitsuyama, T.; Yoneoka, T.; Tanaka, S. J. Nucl. Mater. 1998, 253, 219–226. 19. Yao, Z.; Suzuki, A.; Muroga, T.; Katahira, K. J. Nucl. Mater. 2004, 329–333, 1414–1418. 20. Yao, Z.; Suzuki, A.; Muroga, T.; Yeliseyeva, O. I.; Nagasaka, T. Fusion Eng. Des. 2006, 81, 951–956. 21. Heo, N. J.; Nagasaka, T.; Muroga, T.; Matsui, H. J. Nucl. Mater. 2002, 307–311, 620–624. 22. Hoelzer, D. T.; West, M. K.; Zinkle, S. J.; Rowcliffe, A. F. J. Nucl. Mater. 2000, 283–287, 616–621. 23. Chikada, T.; Suzuki, A.; Yao, Z.; Sawada, A.; Terai, T.; Muroga, T. Fusion Eng. Des. 2007, 82, 2572–2577. 24. Yeliseyeva, O.; Muroga, T.; Yao, Z.; Tsisar, V. J. Nucl. Mater. 2009, 386–388, 696–699. 25. Koch, F.; Brill, R.; Maier, H.; et al. J. Nucl. Mater. 2004, 329–333, 1403–1406. 26. Sawada, A.; Suzuki, A.; Maier, H.; Koch, F.; Terai, T.; Muroga, T. Fusion Eng. Des. 2005, 75–79, 737–740. 27. Pint, B. A.; Moser, J. L.; Tortorelli, P. F. Fusion Eng. Des. 2006, 81, 901–908. 28. Sawada, A.; Suzuki, A.; Terai, T. Fusion Eng. Des. 2006, 81, 579–582.
29. 30.
31. 32. 33. 34. 35. 36. 37. 38. 39. 40. 41. 42. 43.
699
Yao, Z.; Suzuki, A.; Levchuk, D.; et al. J. Nucl. Mater. 2009, 386–388, 700–702. Hishinuma, Y.; Tanaka, T.; Nagasaka, T.; et al. Development of Er2O3 coating on liquid blanket components synthesized with MOCVD process. In 14th International Conference on Fusion Reactor Materials, Sapporo, Japan, Sept 6–11, 2009. Pint, B. A.; Moser, J. L.; Jankowski, A.; Hayes, J. J. Nucl. Mater. 2007, 367–370, 1165–1169. Pint, B. A.; Pawel, S. J.; Howell, M.; et al. J. Nucl. Mater. 2009, 386–388, 712–715. Shikama, T.; Zinkle, S. J.; Shiiyama, K.; Snead, L. L.; Farnum, E. H. J. Nucl. Mater. 1998, 258–263, 1867–1872. Tanaka, T.; Shikama, T.; Narui, M.; Tsuchiya, B.; Suzuki, A.; Muroga, T. Fusion Eng. Des. 2005, 75–79, 933–937. Tanaka, T.; Nagayasu, R.; Sawada, A.; et al. J. Nucl. Mater. 2007, 367–370, 1155–1159. El-Guebaly, L. A.; The ARIES Team. Fusion Eng. Des. 1997, 38, 139–158. Mattas, R. F.; Smith, D. L.; Reed, C. B.; et al. Fusion Eng. Des. 1998, 39–40, 659–668. Muroga, T.; Tanaka, T.; Kondo, M.; Nagasaka, T.; Xu, Q. Fusion Sci. Technol. 2009, 56, 897–901. El-Guebaly, L. A. Fusion Eng. Des. 2006, 81, 1327–1331. Muroga, T.; Tanaka, T.; Sagara, A. Fusion Eng. Des. 2006, 81, 1203–1209. Noajitra, P.; Buehler, L.; Fischer, U.; Malang, S.; Reimann, G.; Schnauder, H. Fusion Eng. Des. 2002, 61–62, 449–453. Wong, C. P. C.; Malang, S.; Sawan, M.; et al. Fusion Eng. Des. 2006, 81, 461–467. Levchuk, D.; Levchuk, S.; Maier, H.; Bolt, H.; Suzuki, A. J. Nucl. Mater. 2007, 367–370, 1033–1037.
4.22 Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors E. R. Hodgson Euratom/CIEMAT Fusion Association, Madrid, Spain
T. Shikama Tohoku University, Sendai, Japan
ß 2012 Elsevier Ltd. All rights reserved.
4.22.1 4.22.2 4.22.3 4.22.4 4.22.5 4.22.6 4.22.7 4.22.8 References
Introduction Fusion-Relevant Radiation Damage in Insulating Materials Simulation Experiments Degradation of Insulator Electrical Resistance Degradation of Insulator AC/RF Dielectric Properties Degradation of Insulator Thermal Conductivity Degradation of Optical Properties Concluding Remarks
Abbreviations AC/RF BA CDA CIEMAT CVD DC DEMO ECRH EDA EVEDA FIRE H&CD HFIR HFR ICRH IEA IFMIF IMR IR ITER
JET
Alternating current/radio frequency Broader approach Conceptual design activity Centro de Investigaciones Energe´ticas, Medioambientales, y Tecnolo´gicas Chemical vapor deposition Direct current Demonstration Electron cyclotron resonant heating Engineering design activity Engineering Validation and Engineering Design Activities Fusion ignition research experiment Heating and current drive High Flux Isotope Reactor (Oak Ridge, USA) High Flux Reactor (Petten, Holland) Ion cyclotron resonant heating International Energy Agency International Fusion Materials Irradiation Facility Institute for Materials Research Infrared International Thermonuclear Experimental Reactor (Cadarache, France) Joint European Torus (Culham, UK)
KfK KU1, KS-4V LAM LAMPF LH LIDAR MACOR MI NBI ORNL OSIRIS PIE RAFM RIA RIC RIED RIEMF RF RL SCCG TEM UV
702 703 705 706 712 715 717 720 721
Kernforschungszentrum Karlsruhe (Germany) Russian radiation-resistant quartz glasses Low-activation materials Los Alamos Meson Physics Facility (USA) Lower hybrid Light Detection and Ranging Machinable Glass Ceramic (Corning Incorporated) Mineral insulation/insulated Neutral beam injector Oak Ridge National Laboratory (USA) From the Greek for Us-yri ‘the king’ (Reactor at Saclay, France) Postirradiation examination Reduced activation ferritic martensitic Radiation-induced absorption Radiation-induced conductivity Radiation-induced electrical degradation Radiation-induced electromotive force Russian Federation Radioluminescence Subcritical crack growth Transmission electron microscope Ultraviolet
701
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Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors
4.22.1 Introduction It is envisaged that early in the twenty-first century ITER (International Thermonuclear Experimental Reactor) will come into operation, and it is expected that this intermediate ‘technology’ machine will help to bridge the gap between the present-day large ‘physics’ machines and the precommercial DEMO power reactor, thus paving the way for commercial fusion reactors to become available by the end of the century. Although this ‘next-step’ device will undoubtedly help to solve many of the problems, which still remain in the field of plasma confinement, it will also present additional operational and experimental difficulties not found in present-day machines. These problems are related to the expected radiation damage effects as a result of the intense radiation field from the ‘burning’ plasma. This ignited plasma will give rise to high-energy neutron and gamma fluxes, penetrating well beyond the first wall, from which one foresees a serious materials problem that has to be solved. In the initial physics phase of operation of such a machine, it is the radiation flux, which will be of concern, whereas in the later technology phase, both flux and fluence will play important roles as fluence (dose)-dependent radiation damage builds up in the materials. For structural metallic materials, radiation damage in ITER is expected to be severe, although tolerable, only near to the first wall. However, the problem facing the numerous insulating components is far more serious because of the necessity to maintain not only the mechanical, but also the far more sensitive physical properties intact. An additional concern arises from the need to carry out inspection, maintenance, and repair remotely because of the neutron-induced activation of the machine. This ‘remote handling’ activity will employ machinery, which requires the use of numerous standard components ranging from simple wires, connectors, and motors, to optical components such as windows, lenses, and fibers, as well as electronic devices such as cameras and various sophisticated sensors. All these components use insulating materials. It is clear, therefore, that we face a situation in which insulating materials will be required to operate under a radiation field, in a number of key systems from plasma heating and current drive (H&CD), to diagnostics, as well as remote handling maintenance systems. All these systems directly affect not only the operation, but also the safety, control, and long-term reliability of the machine. Even for ITER, the performance of some potential insulating materials appears marginal. In the
long term, beyond ITER, the solution of the materials problem will determine the viability of fusion power. The radiation field will modify to some degree all of the important material physical and mechanical properties. Some of the induced changes will be flux dependent, while others will be modified by the total fluence. Clearly, the former flux-dependent processes will be of concern from the onset of operation of future next-step devices. The fluence-dependent effects on the other hand are the important parameters affecting the component or material lifetime. The properties of concern which need to be considered for the many applications include electrical resistance, dielectric loss, optical absorption, and emission, as well as thermal and mechanical properties. Numerous papers have been published discussing both general, and more recently, specific aspects of radiation damage in insulating materials for fusion applications, and those most relevant to the present chapter are included.1–26 In recent years, because of the acute lack of data for insulators and the recognition of their high sensitivity to radiation, most work has concentrated on the immediate needs for ITER. A comprehensive ceramics irradiation program was established to investigate radiation effects on a wide range of materials for essentially all components foreseen for H&CD and diagnostics in ITER, and to find solutions for the problems which have been identified. A large number of relevant components and candidate materials have been, and are being, studied systematically at gradually increasing radiation dose rates and doses, under increasingly realistic conditions. A considerable volume of the work discussed here was carried out within the ITER framework during the CDA, EDA, and EDA extension (Conceptual and Engineering Design Activities 1992–2002) as specific tasks assigned to the various Home Teams (T26/28 and T246; EU, JA, RF, US; T252/445 and T492; EU, JA, RF).27,28 Since these last ITER tasks, no new coordinated tasks related to insulators have been formulated. However, despite the lack of an official framework in which to develop and assign further common tasks following the end of the ITER-EDA extension, collaborative work has continued between the EU, JA, RF, and US Home Teams on both basic and applied aspects of radiation damage in insulator materials. This has resulted in considerable progress being made on the understanding of the pertinent effects of radiation on in-vessel components and materials in particular for diagnostic applications. Problems which have been addressed and for which irradiation testing
Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors
has been performed include comparison of absorption and luminescence for different optical fibers and window materials, RIEMF (radiation-induced electromotive force) and related effects for MI (mineral insulated) cables and coils, alternative bolometers to the reference JET type gold on mica, hot filament pressure gauges, and electric field effects in aluminas. One must however remember that ITER is only an intermediate ‘technology’ machine on the road to a precommercial power reactor. Such power reactors will face the same radiation flux problems as anticipated in ITER, but the fluence problems will be far more severe. It is also important to note that the radiation flux and fluence levels will be different from one type of device to another depending on the design (e.g., in ITER and the Fusion Ignition Research Experiment (FIRE)26), and also on the specific location within that device. Because of the general uncertainty in defining radiation levels, most radiation effects studies have taken this into account by providing where possible data as a function of dose rate (flux), dose (fluence), and irradiation temperature. Although the task ahead is difficult, important advances are being made not only in the identification of potential problems and operational limitations, but also in the understanding of the relevant radiation effects, as well as materials selection and design accommodation to enable the materials limitations to be tolerated. Following a brief introduction to the problem of radiation damage in both metals and insulators, the important aspect of simulating the operating environment for the component or material under examination will be presented, with reference to present experimental procedures. The chapter will then concentrate on the problems facing the use of insulators, with the radiation effects on the main physical properties being discussed, concentrating in particular on the dielectric properties.
4.22.2 Fusion-Relevant Radiation Damage in Insulating Materials The study of intense radiation effects in metals has been closely associated with the development of nuclear fission reactors, and as a result at the beginning of the 1980s when the urgent need to consider radiation damage aspects of materials to be employed in future fusion reactors was fully realized, a considerable amount of knowledge and expertise already existed for metallic materials.29 This was not the case
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for the insulating materials, mainly because of the fact that the required use of insulators in fissiontype reactors is in general limited to low radiation regions, well protected from the reactor core. However, despite the late start and the reduced number of specialists working in related fields at the time, together with the complexity of the mechanisms involved in radiation damage processes in insulators, considerable progress has been made not only in assessing the possible problem areas, but also in finding viable solutions. Several general reviews give a good introduction to the specific problem of radiation damage in insulators.30–36 The materials employed in the next-step fusion machine will be subjected to fluxes of neutrons and gammas originating in the ignited plasma. The radiation intensity will depend not only on the distance from the plasma, but also in a complex way on the actual position within the machine because of the radiation streaming along the numerous penetrations required for cooling systems, blanket structures, heating systems, and diagnostic and inspection channels, as well as the radiation coming from the water in the outgoing cooling channels due to the 16O(n, p)16N nuclear reaction. However one-, two-, and even three-dimensional models are now available, which enable the neutron and gamma fluxes to be calculated with confidence at most, if not all, machine positions.37–40 Radiation damage is generally divided into two components: displacement damage and ionization effects. In a fusion environment, displacement damage, which affects both metals and insulators, will result from the direct knock-on of atoms/ions from their lattice sites by the neutrons, giving rise to vacancies and interstitials. Those primary knock-on atoms (PKAs) with sufficient energy may go on to produce further displacements, so-called cascades. The numerous point defects thus produced may either recombine, in which case no net damage results, or they may stabilize and even aggregate producing more stable extended defects. These secondary processes which determine the fate of the vacancies and interstitials are governed by their mobilities. These mobilities are highly temperature dependent, and in the case of insulators even depend on the ionizing radiation level (radiation-enhanced diffusion). Displacement damage is measured in ‘dpa’ (displacements per atom) where 1 dpa is equivalent to displacing all the atoms once from their lattice sites. At the first wall of ITER, the primary displacement dose rate will be of the order of 106 dpa s1.
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Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors
In contrast, ionizing radiation although absorbed by both metals and insulators, in general, only produces heating in metals. However, certain aspects of radiation damage in metals, such as radiation-enhanced corrosion and grain boundary modification are related to ionization. The effects of ionization on insulators are in comparison quite marked because of the excitation of electrons from the valence to the conduction band giving rise to charge transfer effects. Ionizing radiation is measured in absorbed dose Gy (Gray) where 1 Gy ¼ 1 J kg1. At the first wall of ITER, the dose rate will be of the order of 104 Gy s1. The response of insulators to both displacement and ionizing radiation is far more complex than in the case of metals. Apart from a few specific cases (diamond for example), insulating materials are polyatomic in nature. This leads to the following: (i) We have in general two or more sublattices which may not tolerate mixing. (ii) This gives rise to more types of defects than can exist in metals. (iii) Because of the electrically insulating nature, the defects may have different charge states, and hence different mobilities. (iv) The displacement rates and thresholds, as well as the mobilities, may be different on each sublattice. (v) We may have interaction between the defects on different sublattices. (vi) Defects can be produced in some cases by purely electronic processes (radiolysis); however, in the insulating materials of interest for fusion, this is generally not the case. As a consequence of these factors, while radiation damage affects all materials, the insulators are far more sensitive to radiation damage than metals. While stainless steel, for example, can withstand several dpa and GGy with no problem, some properties of insulating materials can be noticeably modified by as little as 105 dpa or a few kGy. Because of this, the present ongoing programs of radiation testing for diagnostics are concentrating mainly on the insulating components of the systems. The results of these radiation damage processes are flux- and fluencedependent changes in the physical and mechanical properties of the materials, which may be particularly severe for the insulators. The properties of concern which suffer modification are the electrical and thermal conductivity, dielectric loss and permittivity, optical properties, and to a lesser extent the mechanical strength and volume. The effects of such changes are that the insulators may suffer Joule heating
because of the increased electrical conductivity or lower thermal conductivity, and absorption in windows and fibers can increase from the microwave to the optical region and they emit strong luminescence (radioluminescence, RL); in addition, the materials may become more brittle and may suffer swelling. Clearly, some materials are more radiation resistant than others. The organic insulators, which are widely used in multiple applications in general, degrade under purely ionizing radiation and are not suitable for use at temperatures above about 200 C; as a result their use will be limited to superconducting magnet insulation and remote handling applications during reactor shutdown. Inorganic insulators of the alkali halide class have been widely studied and are used as optical windows; however, they are susceptible to radiolysis (displacement damage induced by electronic excitation) and in general become opaque at low radiation fluences. Of the numerous insulating materials, it is the refractory oxides and nitrides, which in general show the highest radiation resistance, and of these the ones which have received specific attention within the fusion program include MgO, Al2O3, MgAl2O4, BeO, AlN, and Si3N4. In addition, different forms of SiO2 and materials such as diamond and silicon have been examined for various window and optical transmission applications. One other aspect of radiation damage that should be mentioned is nuclear transmutation. The highenergy neutrons will produce nuclear reactions in all the materials giving rise to transmutation products.1 These will build up with time and represent impurities in the materials, which may modify their properties. The physical properties of insulators are particularly sensitive to impurities. Furthermore, some of these transmutation products may be radioactive and give rise to the need for remote handling and hot cell manipulation in the case of component removal, repair, or replacement. For the structural materials, in the present concepts mainly steel alloys, considerable work has been carried out on the development of so-called low or reduced activation materials (LAM, RAFM – reduced activation ferritic/ martensitic) for possible use in DEMO and future commercial fusion reactors.41–45 This work with the aim of reducing the amount of nuclear waste has studied not only the substitution of radiological problem alloying elements such as Mo and Nb in steels, but also the viability of other materials such as vanadium and SiC/SiC composites. In the case of the insulating materials, no equivalent study or
Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors
development has been carried out, in part because of the small fraction of the total material volume represented by the insulators, and also because the important physical properties of these materials are expected to be degraded before the transmutation products become of concern. Certainly, for a next-step machine such as ITER, transmutation products, with the possible exception of hydrogen and helium, are not expected to present a serious problem.
4.22.3 Simulation Experiments Within the fusion community, there is an acute awareness of the necessity to construct a suitable irradiation testing facility for materials, which will enable both testing and development of materials for future fusion reactor devices with a fusion-like neutron spectrum. Within this context, both conceptual and engineering design activities were undertaken during the 1990s within the IEA framework with the view of providing such a facility, the IFMIF (International Fusion Materials Irradiation Facility).46–50 This work has been recently renewed under the EU-Japan Broader Approach (BA) activities with the EVEDA (Engineering Validation and Engineering Design Activities) tasks.51,52 However, at the present time no entirely suitable irradiation testing facility exists, and as a consequence experiments have been performed in nuclear fission reactors and particle accelerators, as well as g- and X-ray sources, in an attempt to simulate the real operating conditions of the insulating materials and components. The experiments required must simulate the neutron and g radiation field, that is, the displacement and ionization damage rates, the radiation environment, that is, vacuum and temperature, and also the operating conditions such as applied voltage, or mechanical stress. As will be seen, for the insulator physical properties, it is furthermore essential that in situ testing is carried out to determine whether or not the required physical properties of the material or component are maintained during irradiation. Examples of this include the electrical conductivity, which can increase many orders of magnitude due to the ionizing radiation, or optical windows, which may emit intense RL. Experimental nuclear fission reactors clearly have the advantage of producing a radiation field consisting of both neutrons and g-rays, although in most cases the actual neutron energy spectrum and the dpa to ionization and He ratios are not those which will
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be experienced in a fusion reactor.50 However, it is worthwhile noting that to date experimental fission reactors have mainly been used for irradiations in the metals programs where the emphasis is on the neutron flux and little consideration is given to the g field. As a result, the irradiation channels have in general been designed and installed with this criterion. However, it should be possible to select positions within the reactors which, together with suitable neutron absorber materials and neutron to g converters, provide acceptable radiation fields. The main difficulties with in-reactor experiments come from the inaccessibility of the radiation volume and are concerned with the problem of carrying out in situ measurements and achieving the correct irradiation environment. While considerable success has been attained in the in situ measurement requirement, with parameters such as electrical conductivity, optical absorption and emission, and even radiofrequency dielectric loss being determined, the problem of irradiating in vacuum still remains, with most experiments being performed in a controlled He environment. Irradiation in a controlled atmosphere such as He causes an immediate problem for in situ electrical and dielectric measurements because of the radiation-enhanced electrical conductivity of the gas,53 and even in the case of irradiation in vacuum at about 103 mbar spurious leakage currents will occur.54 Furthermore, many in-reactor experiments rely on nuclear heating to reach the required temperature, and hence have difficulty maintaining a controlled temperature, in part because of the changes in the reactor power, and also because of the problem of calculating the final sample or component temperature. These aspects will be further discussed later. One additional difficulty comes from the nuclear activation of the sample or component, which generally means that postirradiation examination (PIE) has either to be carried out in a hot cell or postponed until the material can be safely handled. Particle accelerators, on the other hand, are ideal for carrying out in situ experiments in high vacuum and at well-controlled temperatures because of the easy access and the very localized radiation field. High levels of displacement damage and ionization can be achieved with little or no nuclear activation. It is however in the nonnuclear aspect of the radiation field where their disadvantage is evident, and great care has to be taken to ensure that appropriate displacement rates are deduced to enable reliable comparison with the expected fusion damage. A further serious disadvantage is due to the limited irradiation
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Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors
volume and particle penetration depth. This in general means that only small thin material samples or components can be tested. The present-day situation of materials and component radiation testing for fusion applications takes full advantage not only of fission reactors and particle accelerators, but also 60Co g irradiation facilities and even X-ray sources. The use of such widely different radiation sources can be justified as long as the influence of the type of radiation on the physical parameter of interest is known. This, in certain cases, is true for radiation-induced electrical conductivity and RL for example, where for low total fluences it is the ionizing component of the radiation field which is important. In situ measurements can now be made during irradiation of the important electrical, dielectric, and optical properties. In addition other aspects such as mechanical strength and tritium diffusion are being assessed during irradiation. Undoubtedly, successful modeling could be of help to address this diverse use of irradiation sources; however, general modeling for the insulators has hardly got off the ground because of the difficulties associated with describing radiation effects in polyatomic bandstructured materials. As a result, in contrast to the extended activity for metallic structural materials, to date there has been no coordinated activity for the insulators, with only specific models for aspects such as electrical and thermal conductivity being developed.
4.22.4 Degradation of Insulator Electrical Resistance Electrical resistance, more generally discussed in terms of the electrical conductivity (the inverse of the resistance), is an important basic parameter for numerous systems and components including the NBI (neutral beam injector) heating system, ICRH (ion cyclotron resonant heating) windows and supports, magnetic coils, feedthroughs and standoffs, MI cables, and wire insulation. Any reduction in the electrical resistance of the insulator material in these components may give rise to problems such as increased Joule heating, signal loss, or impedance change. The main candidate material for these applications is Al2O3 and is also the one which has been most extensively studied, both in the polycrystalline alumina form and as single crystal sapphire. To a lesser extent, MgO, BeO, MgAl2O4, AlN, and SiO2 have also been studied. At the present time, three types of electrical degradation in a radiation
environment are recognized and have been investigated; these are radiation-induced conductivity (RIC), radiation-induced electrical degradation (RIED), and surface degradation. Of these types of degradation, RIC was the first to be addressed in a fusion context, as this enhancement of the electrical conductivity is flux dependent and hence a possible cause for concern from the onset of operation of any fusion device. Fortunately, RIC had been studied for many years, and a sound theoretical understanding already existed.55–59 The ionizing component of the radiation field causes an increase in the electrical conductivity because of the excitation of electrons from the valence to the conduction band and their subsequent trapping in levels within the band gap near to the conduction band from where they are thermally excited once again into the conduction band. Figure 1 shows schematically RIC as a function of irradiation time and ionizing dose rate (flux). The increase in saturation depends not only on the dose rate as indicated, but also in a complex way on the temperature and sample impurity content, as may be seen in Figure 2 for MgO:Fe.60 Nevertheless, such behavior, including the initial step, is well predicted by theory.57 However, at the dose rates of interest for fusion applications, in the range of approximately 1 Gy s1 to >100 Gy s1, saturation is reached within minutes to seconds, and it is this saturation level which is usually the value of interest. The RIC process can lead to increases in the electrical conductivity of many orders of magnitude. For example, a standard high-purity alumina has a room temperature conductivity of generally less than 1016 S m1, which increases to approximately 1010 S m1 for an ionizing dose rate of only 1 Gy s1.61 The first experiments carried out within a fusion application context, that is, refractory oxide materials, high-dose rates, and temperatures, gave an insight into the effects of dose rate, temperature, and material impurity, and established the well-known relationship at saturation, between the total electrical conductivity measured during irradiation and the ionizing dose rate: stotal ¼ s0 þ KRd where s0 is the conductivity in the absence of radiation, R is the dose rate, and K and d are constants.59,61–63 Although d 1, the detailed studies found temperature, dose, and dose rate dependence in this parameter, with extreme values in certain cases ranging between 0.5 and 1.5, and in addition a temperature dependence was observed for K. At the present time, extensive RIC data are available for materials irradiated with X-rays, g-rays, electrons, protons, positive ions, and fission and
Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors
707
Schematic RIC 2.5
Increasing dose rate
RIC (a.u.)
2
1.5
1
0.5
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Figure 1 Schematic RIC. Saturation is reached more rapidly at higher dose rate. For fusion applications, it is generally the saturation level that is of interest.
MgO:Fe 0.1 Gy s–1 4 14 ⬚C, 180 ppm 172 ⬚C, 180 ppm 14 ⬚C, 650 ppm 136 ⬚C, 650 ppm
RIC
2 1.0
n n(t = ¥)
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Irradiation time (min) Figure 2 RIC for single crystal MgO, doped with 180 and 650 ppm Fe. g irradiation at 0.1 Gy s1 for different temperatures (14, 136, and 172 C).60 Theoretical predictions are shown inset. Reproduced from Huntley, D. J.; Andrews, J. R. Can. J. Phys. 1968, 46, 147.
14 MeV neutrons. Many of the additional results, although in some cases limited to one temperature, and/or one dose rate, add confirmation to the earlier extended studies, but more importantly show that RIC is essentially a function of the ionization, independent of the irradiating particle or source. With very few
exceptions, all the data taken together over a range of dose rates from <1 Gy s1 to about 104 Gy s1 show d 1, as may be seen in Figure 3, and lie within a narrow band with the spread in conductivity values at any given dose rate being about two orders of magnitude13; see also, for example, Noda et al.,66
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Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors
10-2 Superconducting magnet
Electrical conductivity (W-m)-1
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Van Lint et al.64 Klaffky et al.59 Pells et al.65
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4 MgAl2O4
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Pells et al.65
102 104 Dose rate (Gy s–1)
106
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Figure 3 Representative data for RIC as a function of dose rate for different oxide materials. Irradiation with electrons, protons, and neutrons. Reproduced from Shikama, T.; Pells, G. P. J. Nucl. Mater. 1994, 212–215, 80.
1.8 MeV e- 450 ⬚C 700 Gy s–1 10-10 dpa s–1 Al2O3 MgO MgAl2O4 BeO
RIC (S m−1)
10-6
10-7
10-8 100
1000 Impurity content (ppm)
104
Figure 4 RIC for different single and polycrystalline materials measured during 1.8 MeV electron irradiation at 700 Gy s1, 450 C, plotted as a function of the estimated total impurity content. The line is of slope 1. Reproduced from Hodgson, E. R. J. Nucl. Mater. 1998, 258–263, 226.
where 14 MeV neutron results are given together with a small selection of other RIC data. For all the RIC data available, because of the different experimental conditions, it is difficult to draw any conclusions as to the reason for the spread in RIC values at any given dose rate. However, data obtained from electron
irradiations of different aluminas and other materials under identical conditions of dose rate and temperature give an indication that the RIC is inversely proportional to the sample impurity content.19 From these results (Figure 4), two general conclusions/ indications may be drawn:
Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors
RIC ðsingle crystalÞ > RIC ðpolycrystalÞ and RIC ðpureÞ > RIC ðimpureÞ However, the indication on the impurity dependence needs to be qualified, as certain impurities introduce levels near to the conduction band, and increase the RIC.59,60 This would imply therefore that the vast majority of the impurities in the materials act as recombination centers for the electrons and holes, thereby reducing the free charge carrier lifetimes, and do not introduce electron levels near to the conduction band. The reduction of the electron lifetime in the conduction band has important consequences for the RIED effect in different materials, as discussed below. From all the data available, at the present time one can safely say that RIC is sufficiently ‘well understood’ to allow this type of electrical degradation to be accommodated by the design, and that materials exist which give rise to electrical conductivities 106 S m1 for ionizing dose rates of up to >103 Gy s1. One only expects possible problems or influence near the first wall. Unfortunately, this is precisely the region where magnetic coil diagnostics that can tolerate only very low leakage conductivity will be employed. It is important to remember that RIC is a flux-dependent effect and will be present from the onset of operation of the next-step machines. Hence, devices which are sensitive to impedance changes, which will occur for example in MI cables,
must take RIC into account. Furthermore, as RIC is strongly affected by impurity content, the buildup of transmutation products will modify the RIC with irradiation time (fluence), although this is not expected to be of serious concern for ITER. In contrast to RIC, RIED is a more serious problem because it has been observed under certain conditions to permanently increase, that is, degrade, the electrical conductivity with irradiation dose. Figure 5 shows a schematic RIED-type degradation. The initial increase in the conductivity corresponds to the RIC as described above. Following a certain irradiation time, or accumulated dose, the conductivity again begins to increase as s0 degrades. In Al2O3 for which most work has been performed, RIED is observed as a permanent increase or degradation of the electrical conductivity (s0) when a small electric field (100 kV m1) is applied during irradiation at moderate temperatures (450 C). At considerably higher temperatures and voltages, but without an irradiation field,67 or for irradiations performed without an applied electric field,68 no degradation occurs. Even at the present time, this type of degradation is still not fully understood; nor is there general agreement as to whether RIED is a real degradation in the volume. Following the first report of RIED effect in electron-irradiated sapphire (Al2O3) and MgO,8 numerous experiments were carried out to assess its possible relevance to fusion insulator applications. These addressed the effect of the applied electric
2 RIED influence
RIC + RIED (a.u.)
1.5 RIC dominates 1
Permanent degradation
0.5
0
0
20
40
709
60
80
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Irradiation time (a.u.) Figure 5 Schematic RIED. Initially, during irradiation RIC dominates, but with irradiation time (dose) the measured conductivity increases because of permanent degradation.
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Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors
field, DC or AC/RF69 and voltage threshold,70 the irradiation temperature,71,72 and the ionizing dose rate,73 as well as observations that in addition to electrons, RIED occurred with protons (Figure 674), as,75 and neutrons,76–78 and the observation of RIED effects in other materials, for example, MgAl2O4.74 In addition, further experiments were performed in which RIED-like effects were also observed in sapphire that was electron irradiated in air,79 for thin Al2O3 films,80 and MgO insulated cable.81 In contrast, some experiments did not observe any RIED effect, with some reporting enhanced surface conductivity or even cracking of the material.82–88 This led to suggestions that the RIED degradation is not a real volume effect, but is caused by surface contamination.82,86 Because of the potential importance of electrical degradation and the uncertainty, extensive discussions on RIED were held at several IEA Workshops,89,90 including the experimental techniques employed in the irradiations to separate and identify volume degradation from surface effects. It was pointed out at an early stage of the discussions that important factors such as dose rate, and in particular material-type differences, and irradiation temperature, all of which could cause RIED not to be observed were not being taken into account.73 For example, under identical conditions RIED was observed in Vitox alumina but not in Wesgo AL995 alumina,75 strongly suggesting a material (possibly impurity and/or grain size) dependence, and further
observations showed that the low purity, large grain size Wesgo AL995 material was highly susceptible to surface degradation when irradiated in high vacuum.91 The in-reactor RIED experiment in HFIR at ORNL also threw light on the complex RIED problem.92,93 Initial results indicated no significant increase in electrical conductivity for 12 different samples. However, moderate to substantial electrical degradation was later reported for some of the sapphire and alumina samples, so material type is an important parameter.94 One of the major difficulties for in-reactor experiments is the determination of s0, the conductivity in the absence of radiation, and its temperature behavior. The use of nuclear heating and the residual reactor radiation level mean that changes in this parameter with temperature and its corresponding activation energy are not generally measured, although these are the main indicators for the onset of degradation; hence, RIED only becomes measurable when the material conductivity in the absence of radiation is larger than the RIC; that is, s0 KRd. Furthermore, some experiments were performed at temperatures either near room temperature85 or above 600 C,95 considerably outside the expected effective temperature range for RIED of approximately 400–500 C. In an attempt to clarify the situation, work was performed to identify possible basic causes of RIED. These experiments detected specific volume effects in Al2O3 that are observed only for irradiations carried out with an applied electric field. A marked
Log10 displacements per atom -4.0 -3.5 -3.0
-4.5
-2.5
Log10 electrical conductivity (W-1 m-1)
0 Vitox Al2O3
500 ⬚C
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-10 400 ⬚C -15
8.5
9.0
9.5
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Log10 ionization dose (Gy) Figure 6 RIED observed in alumina during 18 MeV proton irradiation, with an applied field of 0.5 MV m1. Reproduced with permission from Pells, G. P. J. Nucl. Mater. 1991, 184, 177.
Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors
enhancement of the well-characterized Fþ-center (oxygen vacancy with one trapped electron) was observed,71 and TEM identified large regions of g-alumina within the bulk of RIED degraded Al2O3.96 The increase in Fþ-center production gave rise to enhanced oxygen vacancy mobility, and led to vacancy aggregation and aluminum colloid formation, as may be seen in Figure 7.97 This clarified the observed close similarity between the RIED effect and colloid production in the alkali halides,68 and helped to explain the formation of g-alumina and associated bulk electrical and mechanical degradation.96 The combined work led to a RIED model being formulated, which successfully explained the role of the electric field (both DC and AC/RF), the ionization, and the anion (oxygen) vacancies.98 The model predicted a threshold electric field for degradation depending on the impurity/defect concentration which, as mentioned above in the discussion of RIC, reduces the free electron lifetime. This helps to explain the negative RIED results for Wesgo AL995 alumina where the applied experimental field was below the predicted value of >0.6 MV m1.75,87 It also highlighted the importance of the ionization, in agreement with earlier conclusions.73,84 Additional support for the model, and RIED as a volume effect, came with the TEM identification of aluminum colloids, as well as previously observed g-alumina, in Al2O3 irradiated with an electric field applied.99 At that time, an alternative model based on charge buildup and breakdown was also developed, but
was not extended to explain many of the important observations.100 During the intense activities related to RIED during the 1990s, two important factors emerged, one concerned with surface electrical degradation, and the other related to the importance of the experimental irradiation environment. For insulating components in future fusion devices, surface electrical degradation may prove to be more serious than the RIC and RIED volume effects. At that time, two types of surface degradation were reported, a contamination caused by poor vacuum, sputtering, or evaporation,83,88 and a real surface degradation related to radiation-enhanced surface vacuum reduction and possibly impurity segregation.101,102 Both forms are affected by the irradiation environment and ionizing radiation. However, the real surface degradation effect is strongly material dependent, and occurs in vacuum but not in air or helium.102 This stresses the extreme importance of a representative irradiation environment for material testing. Most insulating materials required for fusion applications in ITER and beyond must indeed operate in high vacuum, and in consequence accelerator experiments to study electrical conductivity have been performed in vacuum, whereas to date, with few exceptions,76–78,103,104 in-reactor experiments for technical reasons have been performed in helium. Another significant aspect of in-reactor experiments performed in helium is the radiation-induced leakage current in the gas,53 which makes it difficult to
Optical absorption (OD cm-1)
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Energy (eV) Figure 7 Aluminum colloid band in sapphire irradiated with 1.8 MeV electrons at different temperatures with an electric field of 0.2 MV m1 applied. Reproduced from Moron˜o, A.; Hodgson, E. R. J. Nucl. Mater. 1997, 250, 156.
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Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors
determine volume conductivity.81,104 One should also mention that severe electrical surface degradation has recently been observed when oxide insulator materials are bombarded with keV H and He ions.105 The mechanism giving rise to such surface degradation is believed to be the loss of oxygen from the vacuum insulator surface region due to preferential radiolytic sputtering. Similarly, in future fusion devices such as ITER ceramic insulators and windows may also degrade, as they will be bombarded by energetic H isotope and He ions because of ionization of the residual gas by g radiation and acceleration by local electric fields.54 At the present time, the role of the irradiation environment in electrical degradation clearly requires further study. Additional difficulties experienced in performing in-reactor experiments include temperature control and also component testing.104,106–108 It is also important to note that several in-reactor experiments have suffered from electrical breakdowns related to the difficulty of maintaining high voltages in a radiation field, precisely what is required for some H&CD and diagnostics systems in a next-step device. Whether or not these are due to RIED, temperature excursions, He gas breakdown, or problems with the MI cables, terminations, and feedthroughs remains unexplained.
4.22.5 Degradation of Insulator AC/RF Dielectric Properties As with the DC electrical properties, it soon became apparent, even before ITER CDA, that data for radiation effects on the AC/RF dielectric properties (dielectric loss and permittivity) of suitable insulating materials for fusion applications were almost nonexistent. Such materials will be needed for both H&CD and diagnostic applications, where they will be required to maintain their dielectric properties from kHz to GHz under a radiation field in high vacuum. Initial work concentrated on the characterization of candidate materials (Al2O3, MgAl2O4, BeO, AlN, and Si3N4), and also PIE of neutron- and protonirradiated materials.109–114 In general, changes in permittivity were observed to be small (5%) and considered to be acceptable for fusion applications. However, results for dielectric loss (loss tangent measurements) showed orders of magnitude variation for similar materials (105–102 for different forms of alumina at 100 MHz) even before irradiation. To address this problem, a standard material (MACOR) was distributed and measured by the main
laboratories involved (EU, JA, US) to check the different measuring systems used. However, the results showed good agreement,115 and the large variation in reported loss tangent values was later shown to be real, in part because of the effect of the different impurity contents of the materials.116,117 This may be clearly seen in Figure 8, where loss tangent data for different aluminas over a wide frequency range are given, showing marked absorption band structures due to polarizable defects (impurities).116 During the early postirradiation loss tangent measurements, there was an indication of recovery, suggesting that loss during irradiation could be significantly higher.65,109–111 This implied that the already difficult measurements should be made in situ during irradiation. In a simple way, dielectric loss can be considered as being due to two contributions: Loss a ðDC conductivityÞ=Frequency þ Polarization term Clearly, both terms can be modified by the radiation. RIC and RIED will increase the DC conductivity and give rise to dose rate (flux) and dose (fluence) effects, although the contribution will decrease with frequency. The polarization term depends on the defects in the material, which exist as, or can form, dipoles through charge transfer processes due to ionization (impurities, vacancies), and produces the absorption band structure observed in the loss as a function of frequency (Figure 8). This term also gives rise to both flux and fluence effects. Furthermore, defects which are modified by radiation-induced charge transfer processes, for example, levels in the band gap occupied by electrons from the conduction band, are unstable and decay after irradiation. This process is responsible for the slow decrease in electrical conductivity observed at the end of RIC experiments, and will similarly cause a slow decrease in the polarization term. Hence, the initial observations of recovery in dielectric loss are to be expected, and the effort required to make measurements during irradiation fully justified. Following the earlier measurements made during X-ray and proton irradiation,65,109,118 work concentrated on the needs for ICRH at about 100 MHz with the first measurements being made during pulsed neutron irradiation (Figure 9).119,120 These pulsed neutron experiments with ionizing dose rates >104 Gy s1 found increases in loss of only about a factor 4. Such a small increase is not compatible with the PIE results, which indicated that the order of
Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors
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Time (s) Figure 9 The first in-reactor loss measurements at 100 MHz during a narrow neutron pulse (14 ms FWHM), showing the slow recovery for AlN. Reproduced from Stoller, R. E.; Goulding, R. H.; Zinkle, S. J. J. Nucl. Mater. 1992, 191–194, 602, with permission from Zinkle.
magnitude increases during irradiation. This discrepancy may be related to the pulsed nature of the irradiation; although the peak dose rate was high, the integrated dose is only about 500 Gy per pulse, far too low for RIC to reach saturation.59–63 However,
recent results indicate that for low dose (fluence), that is, at the beginning of operation, the influence of the DC conductivity term (RIC) is small for frequencies above about 1 MHz even for dose rates >1 kGy s1.121 Furthermore, in these pulsed experiments, the dpa per
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Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors
pulse (107 dpa) is too small to affect either the DC conductivity (RIC) or the polarizable defects, even though this term at these dose rates becomes important even down to 100 kHz. Candidate RF heating systems for ITER (IC, ion cyclotron; LH, lower hybrid; EC, electron cyclotron) operating at about 100 MHz, 5 GHz, and 200 GHz will require insulators (feedthroughs, standoffs, windows) to operate with large electric fields in a radiation field. In general, the in situ experiments employed low-voltage RF, and the question then arises as to whether RIED could possibly affect the dielectric loss.120 At a time of intense RIED activity, two quite different theoretical models were presented in an attempt to explain why the application of a relatively small electric field during irradiation can substantially modify the damage production process and lead to volume electrical degradation.98,100 The earlier model was based on charge buildup and breakdown, that is, a DC mechanism, but failed to explain many of the results observed during RIED experiments.100 The later model however explained the role of the ionization taking into account the production of highly unstable Fþ-centers,122 the electric field threshold, as well as g-alumina and colloid production, but more importantly predicted that RIED could occur for applied fields at frequencies >100 GHz.98 This was in agreement with early observations of RIED from DC to >100 MHz, and indications for RIED at frequencies above 1 GHz.69 Dielectric loss measurements at 15 GHz, made during electron irradiation at 2 kGy s1, and postirradiation from 1 kHz to 15 GHz, for sapphire, alumina, BeO, and MgAl2O4, show very varied results.123,124 Sapphire, the purest alumina grade, and BeO showed no prompt increase in loss, nor with a dose up to 50 MGy. However, the 999 and 997 alumina grades showed significant prompt and dose-dependent increases in loss, consistent with a modification in the polarization term. Furthermore, these in situ measurements show postirradiation recovery similar to the early reports for proton- and neutron-irradiated materials.65,109–111 In addition, sapphire samples, which had been preirradiated to 7 MGy, 106 dpa at 450 C with a DC electric field (210 kV m1) to produce RIED showed a significant increase in the loss (2 increase), and also in the prompt dielectric loss (5 increase). Similar increases have only been observed for sapphire neutron irradiated, without an electric field applied, to >103 dpa.9 In this context, one should also mention recent work concerned with RF ion sources for NBI systems,
where in situ measurements of dielectric loss during and following electron irradiation of alumina (Deranox 999) to 110 MGy with a 1 MHz RF voltage (0.8 MV m1) applied indicate a permanent increase in loss for irradiation at 240 C, but not at 120 C, as expected from previous RIED studies.125 While various alumina and BeO grades were available with adequate initial properties (dielectric loss, thermal conductivity, and mechanical strength) before irradiation for NBI, IC, and even LH applications, and with potential to withstand the expected ITER radiation levels, this was not the case for ECRH windows. Sapphire or high-purity alumina, the initial ECRH window reference materials with low dielectric loss in the MHz to GHz range,116,126–128 exhibit increasing loss with increasing frequency reaching 104 (loss tangent) by 100 GHz. Hence, to transmit the megawatts of RF power that will be required,9 these materials would have to be employed at cryogenic temperatures, and furthermore with a very low neutron tolerance level, 1020 n m2.128 However, in recent years, considerable progress has been made with CVD diamond, a material with the required combination of low dielectric loss, high thermal conductivity, and mechanical strength.19,25,129–134 In this context, initial work began to examine both high-purity silicon and diamond homopolar crystalline materials which as a result of their decreasing loss with increasing frequency offered the possibility for operation at frequencies above 150 GHz with loss tangents 105, at room temperature.129 These two materials required development in completely opposite directions. The initial high-resistivity silicon had very low loss but extreme radiation sensitivity. Because of its perfection, electrons excited into the conduction band by purely ionizing radiation had very long lifetimes (no defect recombination sites) leading to high dielectric loss through the high electrical conductivity. In contrast, the CVD diamond, initially almost black in color, had high loss because of the numerous defects in the material giving rise to polarization losses, but was almost insensitive to ionizing radiation because of the extremely short lifetime of the conduction band electrons. Although the radiation sensitivity of silicon could be notably reduced by electron irradiation and also by Au doping because of the introduction of recombination defects, the main limitation for silicon comes from its small 1.1 eV band gap. This allows electrons to be readily thermally excited into the conduction band at temperatures only slightly above room temperature,
Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors
Window grade CVD diamond 145 GHz
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Temperature (K) Figure 10 Diamond dielectric loss at 145 GHz while being unirradiated, and neutron irradiated at 320 K (pool temperature) to 1021 and 1022 n m2. Reproduced with permission from Thumm, M.; Arnold, A.; Heidinger, R.; Rohde, M.; Schwab, R.; Spoerl, R. Fusion Eng. Des. 2001, 53, 517.
which rapidly increases the dielectric loss.135–138 In the case of CVD diamond, the progress has been remarkable, available samples going from black and irregular in shape to almost transparent 2 mm thick 100 mm diameter disks, with room temperature loss 1 105 at 145 GHz, comparable with sapphire at 77 K, and furthermore increasing only to about 5 105 by 450 C.130,132 Loss measurements during electron and X-ray irradiation at 18 and 40 GHz, respectively of the developed CVD diamond, show almost negligible contributions of conductivity (RIC) and polarizable defects, and successful high-power transmission tests have now been carried out.132,133 As may be seen in Figure 10, PIE loss tangent measurements of neutron-irradiated ‘window grade’ CVD diamond indicate that even by 1022 n m2 (103 dpa), the room temperature loss only increases to 5 105 at 145 GHz (6 105 at 190 GHz).134 During the intense activity to find suitable materials for the high-power IC, LH, and EC heating applications, work was also being carried out on materials for diagnostic systems. In particular, KU1 quartz glass provided by the Russian Federation within the ITER-EDA task sharing agreement was shown to be highly radiation resistant with respect to its optical properties for use in both diagnostic and remote handling applications, and became the main reference material not only for optical windows, but also fibers.26,139,140 In view of this, the material was also examined for possible use in DC and RF applications. Both RIC and RIED, together with dielectric
loss and permittivity, have been determined for as-received, as well as electron and neutron irradiated material. A large number of different experimental setups were employed to obtain the dielectric spectrum of KU1 over a very wide frequency range (10 mHz to 145 GHz), and where possible, values were obtained during electron irradiation. In addition, data have been obtained for samples neutron irradiated to 104 dpa. The results indicate that for low radiation doses the electrical and dielectric properties are only slightly degraded, and in particular the use of KU1 for electron cyclotron emission (ECE) windows and low-loss DC applications is feasible.134,141
4.22.6 Degradation of Insulator Thermal Conductivity Work began at an early stage to assess the thermomechanical properties of candidate insulating materials for fusion applications. In an attempt to determine the best combination of mechanical, thermophysical, and dielectric properties for the demanding H&CD applications, Al2O3 (both alumina and sapphire), AlN, Si3N4, BeO, and MgAl2O4 in numerous different grades were examined ‘as-received’ and following irradiation.142–149 At room temperature, the unirradiated thermal conductivity of a typical alumina is of the order of 30 W m1 K1, and that of BeO about 280 W m1 K1. These values are sufficiently high
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Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors
for IC and LH heating systems to ensure adequate cooling in most cases; however, the thermal conductivity in ceramics is reduced because of increased phonon scattering, by the presence of point defects and to a lesser extent by extended defects or aggregates. Hence, one expects a reduction in thermal conductivity on irradiation, together with a notable influence of the irradiation temperature, that is, irradiation above temperatures at which the radiationinduced defects become mobile and can either recombine or aggregate should lead to a lower degradation of the thermal conductivity, while lowtemperature irradiation should have a marked effect because of the increased point defect stability. The expected general behavior was confirmed by the early data (Figure 11), and indicated that a maximum reduction to about one-third of the room temperature thermal conductivity value could be expected.142–145 This will occur for a neutron fluence value (dpa), which strongly depends on the irradiation temperature. For near room temperature irradiation (300 K), reduction to the lower saturation level was observed by about 1023 n m2 (0.01 dpa), whereas at 600 K this lower saturation level was only reached following a fluence of above 1024 n m2. Within reasonable margins, these values applied for Al2O3, AlN, and MgAl2O4. Similar PIE results were obtained at a later date for reactor irradiations at different temperatures of a wide range of ceramic materials.150 Because of the importance of point defects in the reduction of thermal conductivity, it is reasonable
to expect that postirradiation measurements may underestimate the effect due to possible postirradiation annealing. An attempt to measure thermal conductivity in situ during reactor irradiation, although unable to quantify the degradation, did highlight a very rapid decrease in thermal conductivity by 1022 n m2 (0.001 dpa) at the startup of irradiation, followed by saturation.151 Finally, one should mention the specific case of sapphire and CVD diamond, the original and the present reference materials for ECRH. For sapphire, the need for low-temperature (<100 K) operation to minimize dielectric loss also provided a gain in thermal conductivity (200 W m1 K1 at 100 K, c.f. about 30 W m1 K1 at room temperature). However, in addition to the dielectric loss showing a very low neutron tolerance (1020 n m2) at this low temperature,128 the high thermal conductivity was reduced by over two orders of magnitude also by 1020 n m2 (105 dpa), because of the enhanced point defect stability.147,152 In the case of CVD diamond, the increase in the room temperature dielectric loss was still tolerable up to 1022 n m2 (103 dpa).134 Unfortunately, although the extremely high thermal conductivity at room temperature (1800 W m1 K1) already began to degrade by 1020 n m2 (105 dpa), it was at the tolerance limit by 1021 n m2 (Figure 12).134 Almost identical results were reported after electron irradiation to 3 106 dpa where the thermal conductivity was reduced by about 9%, confirming the importance of point defects.153
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Figure 11 Effect of neutron and a particle irradiations at different temperatures and doses on alumina (AL23: Degussit) thermal conductivity (KfK 700 K 0.001 dpa, LAMPF 600 K 0.5 dpa, Petten 473 K 0.4 dpa, OSIRIS 823 K 5 dpa). Reproduced from Rohde, M.; Schulz, B. J. Nucl. Mater. 1990, 173, 289, with permission from Rohde.
Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors
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Figure 12 Thermal conductivity reduction for different diamond grades as a function of neutron dose. Reproduced from Heidinger, R.; Rohde, M.; Spo¨rl, R. Fusion Eng. Des. 2001, 56–57, 471, with permission from Heidinger.
4.22.7 Degradation of Optical Properties Within the fusion program, another area of concern is related to the effects of radiation on the optical properties of the dielectric materials to be used as transmission components (windows, lenses, and optical fibers) for the UV, visible, and near-IR wavelength diagnostic systems needed for control and safety, as well as maintenance (remote handling).21,26,140,154,155 Radiation-induced optical absorption (RIA) and light emission or RL impose severe limitations on the use of any optical material within an intense radiation field. For remote handling applications, the optical components will be expected to maintain their transmission properties under high levels of ionizing radiation (1 Gy s1) during hundreds of hours. For such applications, RIA imposes the main limitation, but can be tolerated. However, in the case of diagnostic applications, in addition to a higher level of ionizing radiation (tens to hundreds Gy s1), the materials will also be subjected to atomic displacements 109 dpa s1. It soon became clear that both RIA and RL would impose severe limitations on the main candidate materials (sapphire and silica). Of these two materials, sapphire is by far the most resistant to ionizing radiation. Although ionizing radiation can cause an increase in optical absorption because of trace impurities and vacancy defects present in the material, it is in general the displacement damage mechanism which induces absorption at first in the UV region as a result of oxygen vacancy-related
defects.30,33,156–158 This fluence (dose) effect reduces the transmission in the UV region to essentially zero for doses above about 104 dpa, and more slowly in the visible as the tails of the absorption bands begin to overlap into this region. Although sapphire shows more radiation resistance than SiO2 in terms of optical absorption, the material was found to be unsuitable for many diagnostic applications because of its intense RL, as will be seen below. As with RIC, RL is ionizing flux (dose rate) dependent and hence will be a problem from the onset of operation of future fusion devices. Furthermore, to assess RL clearly requires in situ measurements during irradiation. While many studies had been carried out on luminescence phenomena in SiO2 and sapphire, the problem was only addressed in a quantitative way because of fusion application requirements.159–164 Sapphire was quickly excluded from high-dose rate applications when it was shown that the photon emission for a typical diagnostic window dose rate would be comparable with the photon emission from the plasma.159 In contrast, certain grades of silica show virtually no RL in the UV-visible region, the emission being limited almost to the Cherenkov background. Quantitative luminescence data comparing UV grade sapphire and two types of silica, both of which show low RL, are given in Figure 13, indicating that suitable materials do exist in which the RL can be reduced to a minimum, although there are limited data on RL as a function of fluence.162–164 In particular, the KU1 and KS-4V quartz glass materials, provided by the Russian
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Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors
1011 photons/(s.Å.sr.cm3)
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Figure 13 Quantified RL emission for sapphire and two silica grades during 1.8 MeV electron irradiation at 700 Gy s1, 15 C. Reproduced from Moron˜o, A.; Hodgson, E. R. J. Nucl. Mater. 1998, 258–263, 1889.
Federation for the ITER diagnostics radiation testing program, have proved to be highly resistant to RL and RIA because of ionizing radiation and displacement damage, and are now reference materials.26,165–170 For ionizing radiation doses up to at least 100 MGy and for temperatures at or above about 100 C, very little absorption is induced in the KU1 material over the whole visible range; one must keep in mind however that with irradiation displacement dose the optical absorption related to oxygen vacancies in SiO2, as in all oxide materials, eventually renders them opaque in the UV and visible range.171–175 In an analogous way to the ECRH transmission windows, mention should be made of windows required for high-power laser transmission, that is, the LIDAR (light detection and ranging) system. This demanding diagnostic system being considered for ITER will require very high-quality transmission windows for the high-power laser pulses at about 500 and 1000 nm. It is estimated that transmission losses of the order of 5% may cause problems with the window integrity because of laser damage. However, such small decreases in the transmission corresponding to an optical density increase of only 0.02 are extremely difficult to measure by standard PIE of irradiated optical materials. Such measurements have to be performed in situ. In situ measurement is also required in order to determine possible radiation-enhanced absorption which can easily reach such small values. The possibility of radiation-enhanced dielectric breakdown due to the intense laser pulse and the
ionizing radiation has also to be considered. However, such a determination requires an elaborate in situ experiment. Work on laser-induced damage in KU1 and KS-4V has confirmed the limited influence of RIA and RIC on the damage threshold for highpower laser transmission.176 However, metallic deposition due to sputtering or evaporation can seriously reduce the damage threshold even for a few nanometer thickness, as may be seen in Figure 14. The effect is strongly material dependent, and furthermore selfcleaning with subthreshold laser pulses is not effective for all deposited materials.177,178 Although in general RL is considered to be a problem for diagnostic systems in future devices, it may be employed as a detector/converter for X-ray, UV, and particle emission from the plasma. The intense RL from Al2O3:Cr, for example, has been used for many years in ceramic fluorescent screens for accelerator beam alignment,179 and is now being developed with improved radiation resistance and rapid decay times for fusion applications, along with other alternative luminescent materials (Figure 15).180–182 Furthermore, RL is a potentially powerful tool capable of monitoring material modification during irradiation, but has been largely neglected within the fusion materials activities, in part because of the difficulty in interpreting the resulting emission spectra. However, the technique is now being successfully employed to study insulating materials such as aluminas and silicas, as well as breeding ceramics for fusion applications.183,184
Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors
719
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Figure 14 Effect of a thin sputtered steel layer on the laser induced damage for KU1 and KS-4V quartz glasses. Reproduced from Martin, P.; Moron˜o, A.; Hodgson, E.R. J. Nucl. Mater. 2004, 329–333, 1442.
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Figure 15 14 MeV neutron-induced luminescence in doped aluminates. Reproduced from Toh, K.; Shikama, T.; Katsui, H.; et al. J. Nucl. Mater. 2009, 386–388, 1027.
Finally, in connection with optical transmission components, one should note the flexibility and simplification in diagnostic design that the use of optical fibers would allow. However, this is not straightforward; although RIA and RL are problems for optical window and lens components, in the case of optical fibers the situation is far worse because of the length of the optical path. Furthermore, because of the
manufacturing techniques, fibers with characteristics as good as those observed for the KU1 quartz glass for example have not been obtained. This has prompted an extensive collaborative research program to find the most suitable types of radiation-resistant fiber. Several different optical fibers have been examined to assess RIA and light emission, the viability of high-temperature operation and annealing, jacketing
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Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors
material, and the influence of hydrogen loading. In addition, parallel work is being carried out on the possibility of photobleaching using high-intensity lasers to recover transmission, ‘holey’ fibers for improved transmission and radiation resistance, and fibers with extended blue – UV transmission.26,185–190 Irradiations have been carried out to total doses above 10 MGy and 1022 n m2, and temperatures from about 30 to 300 C. The most promising fibers are the hydrogen loaded KU1 and KS-4V, where above 400 nm they show the lowest RIA, as may be seen in Figure 16.139 Although the KU1 is the slightly better material up to about 700 nm, the intrinsic OH band and its harmonics notably affect transmission above 800 nm, so for a fiber required to transmit in the visible and IR regions, the hydrogen loaded KS-4V may be a better choice. For silica materials up to about 10 MGy, the main radiation damage mechanisms involve electron and holetrapping; hence, the wide differences observed in induced absorption of the fibers tested are due to variations in intrinsic trapping centers (defects and impurities). In general, these trapping centers are thermally unstable, hence the effective thermal annealing for irradiation at higher temperature, or postirradiation thermal annealing. For higher doses, displacement damage leading to extensive structural damage begins to dominate, but by this time the fibers are of little use for diagnostic applications. Limited work is underway to examine the possibility
Radiation-induced absorption (dB m-2)
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of in situ photobleaching of the radiation-induced damage using high-intensity UV lasers, the potential of so-called ‘holey’ fibers (fibers containing an array of vacuum, air, or liquid filled holes) to improve radiation resistance, as well as fibers to extend transmission into the blue – UV region.
4.22.8 Concluding Remarks Since the end of the 1970s, the fusion materials community has been providing the necessary insulator research and database for H&CD, and diagnostic systems for a next-step burning plasma device (ITER). As described above, the continuous research has identified and highlighted the limitations and potential problems related to electrical, dielectric, and thermal conductivity degradation, window and fiber absorption, and luminescence, while at the same time providing possible solutions for the ECRH windows with the development of diamond for example, identifying safer operating conditions for the insulators, assessing optical materials for low RL, and characterizing dielectric properties over a wide range of frequencies (IC to EC). In addition to channeling the necessary expertise, numerous unique experiments and installations have been developed to study candidate materials under relevant conditions, in particular during irradiation. All this has produced data of direct relevance to both ITER and future fusion
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Figure 16 In-reactor radiation-induced optical absorption (RIA) for four types of fiber. The hydrogen loaded KU1 (KU-H2G) shows the best performance. Reproduced from Brichard, B.; Fernandez Fernandez, A.; Ooms, H.; et al. J. Nucl. Mater. 2004, 329–333, 1456.
Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors
devices. As a result of the in-reactor experience, research was then able to concentrate on the required irradiation testing and screening of prototype components, urgently required for ITER. These have included bolometers, Hall probes, MI cables, coated mirrors, pressure gauges, as well as the optical fibers discussed above.26,108,139,191–193 The irradiation testing required for ceramic insulator materials is far more complex than that required for structural materials. For almost all of the properties of interest, in situ testing is mandatory. The difficulties associated with in situ testing together with the high cost of reactor irradiations have meant that g sources and particle accelerators have been used to full advantage, where the behavior with irradiation temperature and dose rate can be more easily assessed. The need for irradiation in vacuum is an added difficulty; one must remember that most in-reactor irradiations are performed in He. However, despite the difficulties, several experimental systems have been developed to enable testing in both static and dynamic (pumped) vacuum, thereby fully simulating the expected environment. In addition, the hostile, inaccessible, and noisy environment of experimental fission reactors makes the measurement of often small but important effects difficult. However, over the years considerable in-reactor testing expertise has been gained and much needed experiments performed. One must remember that ITER is only a ‘next step’; the final goal is to provide a safe and reliable fusion reactor within a reasonable time. Undoubtedly, beyond ITER the use of insulators will be severely restricted to those essential to operation and maintenance, but they will be of paramount importance to the success of fusion power. Hence, future ceramic insulator research activity, while keeping in mind the short-term ‘urgent’ ITER needs, must address the expected fluence degradation effects on all the material properties and enable viable solutions to be available in time. The problems to be addressed are related to long-term degradation of the required properties because of aggregation and segregation of defects and impurities. For high fluence, not only H and He, but also other transmutation elements will begin to play a role in the modification of the material physical properties. Although not discussed here, mechanical property degradation has been studied since the beginning of the fusion materials activity. During this early work, considerable attention was paid to the mechanical properties of refractory oxides and nitrides, where PIE indicated that
721
significant degradation of the mechanical strength would only occur for radiation damage levels of the order of 1 dpa or above.148 However, evidence was found for two types of radiation-enhanced degradation of the mechanical strength, ‘enhanced’ implying degradation of damage levels <<1 dpa. One is related to RIED, where the formation of small regions of g-alumina within the a-alumina matrix caused materials to become more fragile.75,96 The other type is related to subcritical crack growth (SCCG), where for certain aluminas the ‘time to fracture’ was markedly shortened for tests performed during g-irradiation.194 Clearly, these aspects of radiation damage related to synergistic effects of radiation plus an electric field and mechanical stress should be reexamined to be able to predict their long-term importance, as also should be RIED itself which still remains an unresolved potential longterm problem. Finally, one should note that to date none of the in-reactor experiments has been performed without encountering some unexpected difficulty, related to, for example, feedthroughs, MI cables, electrical contacts, applied voltages, gas leakage currents, poor vacuum, and temperature control. These are precisely the types of problem that must be avoided in future fusion devices, and a renewed effort will be required to overcome them. (See also Chapter 1.02, Fundamental Point Defect Properties in Ceramics).
References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16.
Rovner, L. H.; Hopkins, G. R. Nucl. Technol. 1976, 29, 274. Phillips, D. C. AERE Harwell Report R-8923, 1978. Clinard, F. W., Jr. J. Nucl. Mater. 1979, 85–86, 393. Clinard, F. W., Jr.; Hurley, G. F. J. Nucl. Mater. 1981, 103–104, 705. Pells, G. P. J. Nucl. Mater. 1984, 122–123, 1338. Clinard, F. W., Jr. Cryst. Latt. Def. Amorph. Mater. 1987, 14, 241. Pells, G. P. J. Nucl. Mater. 1988, 155–157, 67. Hodgson, E. R. Cryst. Latt. Def. Amorph. Mater. 1989, 18, 169. Heidinger, R. J. Nucl. Mater. 1991, 179–181, 64. Zinkle, S. J.; Hodgson, E. R. J. Nucl. Mater. 1992, 191–194, 58. Clinard, F. W., Jr.; Farnum, E. H.; Griscom, D. L.; et al. J. Nucl. Mater. 1992, 191–194, 1399. Kinoshita, C. J. Nucl. Mater. 1992, 191–194, 67. Shikama, T.; Pells, G. P. J. Nucl. Mater. 1994, 212–215, 80. Hobbs, L. W.; Clinard, F. W., Jr.; Zinkle, S. J.; Ewing, R. C. J. Nucl. Mater. 1994, 216, 219. Kinoshita, C.; Zinkle, S. J. J. Nucl. Mater. 1996, 233–237, 100. Zinkle, S. J. Mat. Res. Soc. Symp. Proc. 1997, 439, 667–678.
722 17. 18.
19. 20. 21. 22. 23. 24. 25. 26. 27. 28.
29. 30. 31. 32. 33. 34. 35. 36. 37. 38. 39. 40. 41. 42. 43. 44. 45. 46. 47.
Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors Zinkle, S. J.; Kinoshita, C. J. Nucl. Mater. 1997, 251, 200–217. Hodgson, E. R. In Diagnostics for Experimental Thermonuclear Fusion Reactors; Stott, P. E., Gorini, G., Prandoni, P., Sindoni, E., Eds.; Plenum Press: New York, 1998; Vol. 2. Hodgson, E. R. J. Nucl. Mater. 1998, 258–263, 226. Shikama, T.; Yasuda, K.; Yamamoto, S.; Kinoshita, C.; Zinkle, S. J.; Hodgson, E. R. J. Nucl. Mater. 1999, 271, 560. Yamamoto, S.; Shikama, T.; Belyakov, V.; et al. J. Nucl. Mater. 2000, 283–287, 60. Shikama, T.; Zinkle, S. J. Nucl. Mater. 2002, 307–311, 1073. Hodgson, E. R. Nucl. Instrum. Meth. Phys. Res. B 2002, 191, 744. Decreton, M.; Shikama, T.; Hodgson, E. R. J. Nucl. Mater. 2004, 329–333, 125. Ibarra, A.; Hodgson, E. R. Nucl. Instrum. Meth. Phys. Res. B 2004, 218, 29. Vayakis, G.; Hodgson, E. R.; Voitsenya, V.; Walker, C. I. Fusion Sci. Technol. 2008, 53, 699. Yamamoto, S. ITER Document G 55 DDD 27 97–12–05 W0.1. Yamamoto, S.; et al. In Diagnostics for Experimental Thermonuclear Fusion Reactors; Stott, P. E., Gorini, G., Prandoni, P., Sindoni, E., Eds.; Plenum Press: New York, 1998; Vol. 2. Schiller, P.; Ehrlich, K.; Nihoul, J. J. Nucl. Mater. 1991, 179–181, 13. Sonder, E.; Sibley, W. A. In Point Defects in Solids; Crawford, J. H., Slifkin, L. M., Eds.; Plenum Press: New York, 1972; Vol. 1. Clinard, F. W.; Hobbs, L. W. In Physics of Radiation Effects in Crystals; Johnson, R. A.; Orlov, A. N., Eds.; North Holland: Amsterdam, 1986; p 387. Hughes, A. E. Rad. Eff. 1986, 97, 1. Agullo-Lopez, F.; Catlow, C. R. A.; Townsend, P. D. Point Defects in Materials; Academic Press: London, 1988. Itoh, N. Ed. Defect Processes Induced by Electronic Excitation in Insulators; Directions in Condensed Matter Physics; World Scientific: Singapore, 1989; Vol. 5. Stoneham, A. M. Nucl. Instrum. Meth. Phys. Res. 1994, A91, 1. Itoh, N.; Stoneham, A. M. Materials Modification by Electronic Excitation; Cambridge University Press: Cambridge, 2001. Santoro, R. T.; Iida, H.; Khripunov, V. ITER Report IDoMS No. NAG-47-8-1-97, 1997. Khripunov, V.; Santoro, R. T. ITER Report IDoMS No. NAG-50, 1997. Nightingale, M.; Taylor, N. ITER NBI Design Task N53 TD 11 FE D322, 1996. Costley, A. E.; Campbell, D. J.; Kasai, S.; Young, K. E.; Zaveriaev, V. Fusion Eng. Des. 2001, 55, 331. Harries, D. R.; Butterworth, G. J.; Hishinuma, A.; Wiffen, F. W. J. Nucl. Mater. 1992, 191–194, 92. Kohyama, A.; Hishinuma, A.; Gelles, D. S.; Klueh, R. L.; Dietz, W.; Ehrlich, K. J. Nucl. Mater. 1996, 233–237, 138. van der Schaaf, B.; Gelles, D. S.; Jitsukawa, S.; et al. J. Nucl. Mater. 2000, 283–287, 52. Baluc, N.; Gelles, D. S.; Jitsukawa, S.; et al. J. Nucl. Mater. 2007, 367–370, 33. Nishitani, T.; Tanigawa, H.; Jitsukawa, S.; et al. J. Nucl. Mater. 2009, 386–388, 405. Kondo, T.; Doran, D. G.; Ehrlich, K.; Wiffen, F. W. J. Nucl. Mater. 1992, 191–194, 100. Kondo, T.; Shannon, T. E.; Ehrlich, K. J. Nucl. Mater. 1996, 233–237, 82.
48. 49. 50. 51. 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. 77. 78. 79. 80. 81. 82. 83. 84. 85.
Noda, K.; Ehrlich, K.; Jitsukawa, S.; Mo¨slang, A.; Zinkle, S. J. J. Nucl. Mater. 1998, 258–263, 97. Mo¨slang, A.; Antonnucci, C.; Daum, E.; et al. J. Nucl. Mater. 1998, 258–263, 427. Ehrlich, K.; Bloom, E. E.; Kondo, T. J. Nucl. Mater. 2000, 283–287, 79. Garin, P. J. Nucl. Mater. 2009, 386–388, 944. Garin, P.; Sugimoto, M. Fusion Eng. Des. 2009, 84, 259. Hodgson, E. R.; Moron˜o, A. J. Nucl. Mater. 2002, 307–311, 1660. Hodgson, E. R.; Moron˜o, A.; Gonzalez, S. M. J. Nucl. Mater. 2009, 386–388, 999. Rose, A. Phys. Rev. 1955, 97, 322. Marshall, J. F.; Pomerantz, M. A.; Shatas, R. A. Phys. Rev. 1957, 106, 432. Huntley, D. J.; Andrews, J. R. Can. J. Phys. 1968, 46, 147. Hughes, R. C. Phys. Rev. 1979, B19, 5318. Klaffky, R. W.; Rose, B. H.; Goland, A. N.; Dienes, G. J. Phys. Rev. 1980, B21, 3610. Hodgson, E. R.; Clement, S. Rad. Eff. 1986, 97, 251. Pells, G. P. Rad. Eff. 1986, 97, 39. Pells, G. P.; Hill, G. J. J. Nucl. Mater. 1986, 141–143, 375. Hodgson, E. R.; Clement, S. J. Nucl. Mater. 1988, 155–157, 357. van Lint, V. A. J.; Harrity, J. W.; Flanagan, T. M. IEEE Trans. Nucl. Sci. 1968, NS-15(6), 194. Pells, G. P.; Buckley, S. N.; Agnew, P.; Foreman, A. J. E.; Murphy, M. J.; Staunton-Lambert, S. A. B. Radiation Effects in Electrically Insulating Ceramics; Harwell AERE Report 13222, 1988. Noda, K.; Nakazawa, T.; Oyama, Y.; Yamaki, D.; Ikedo, Y. J. Nucl. Mater. 1996, 233–237, 1289. Weeks, R. A.; Sonder, E.; Narayan, J. Am. Ceram. Soc. Bull. 1978, 57, 316. Hodgson, E. R. J. Nucl. Mater. 1991, 179–181, 383. Hodgson, E. R. J. Nucl. Mater. 1992, 191–194, 552. Hodgson, E. R. Nucl. Instrum. Meth. 1992, B65, 298. Moron˜o, A.; Hodgson, E. R. J. Nucl. Mater. 1994, 212–215, 1119. Pells, G. P.; Sowden, B. C. J. Nucl. Mater. 1995, 223, 174. Hodgson, E. R. J. Nucl. Mater. 1994, 212–215, 1123. Pells, G. P. J. Nucl. Mater. 1991, 184, 177. Mo¨slang, A.; Daum, E.; Lindau, R. In Proceedings of the 18th Symposium on Fusion Technology, Karlsruhe, August 1994; Fusion Technol. 1994; p 1313. Shikama, T.; Narui, M.; Endo, Y.; Sagawa, T.; Kayano, H. J. Nucl. Mater. 1992, 191–194, 575. Shikama, T.; Narui, M.; Endo, Y.; Ochiai, A.; Kayano, H. J. Nucl. Mater. 1992, 191–194, 544. Shikama, T.; Narui, M.; Kayano, H.; Sagawa, T. J. Nucl. Mater. 1994, 212–215, 1133. Zong, X. F.; Shen, C. F.; Liu, S.; et al. Phys. Rev. B 1994, 49, 15514. Hunn, J. D.; Stoller, R. E.; Zinkle, S. J. J. Nucl. Mater. 1995, 219, 169. Farnum, E. H.; Shikama, T.; Narui, M.; Sagawa, T.; Scarborough, K. J. Nucl. Mater. 1996, 228, 117. Kesternich, W.; Scheuermann, F.; Zinkle, S. J. J. Nucl. Mater. 1993, 206, 68. Jung, P.; Zhu, Z.; Klein, H. J. Nucl. Mater. 1993, 206, 72. Farnum, E. H.; Clinard, F. W., Jr.; Sommer, W. F.; Kennedy, J. C.; Shikama, T. J. Nucl. Mater. 1994, 212–215, 1128. Snead, L. L.; White, D. P.; Zinkle, S. J. J. Nucl. Mater. 1994, 212–215, 1107.
Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors 86. 87. 88. 89. 90. 91. 92. 93. 94. 95. 96. 97. 98. 99. 100. 101. 102. 103. 104. 105. 106. 107. 108. 109. 110. 111. 112. 113. 114. 115. 116. 117. 118. 119. 120.
Kesternich, W.; Scheuermann, F.; Zinkle, S. J. Nucl. Mater. 1995, 219, 190. Snead, L. L.; White, D. P.; Zinkle, S. J. J. Nucl. Mater. 1995, 226, 58. Kesternich, W. J. Nucl. Mater. 1998, 253, 167. Chen, Y.; Clinard, F. W.; Evans, B. D.; et al. J. Nucl. Mater. 1994, 217, 32. Zinkle, S. J.; Hodgson, E. R.; Shikama, T. IEA Workshop on Radiation Effects in Ceramic Insulators, Cincinnati, May 1997 (ORNL/M-6068, DOE/ER-0313/22). Moron˜o, A.; Hodgson, E. R. J. Nucl. Mater. 1996, 233–237, 1299. Shikama, T.; Zinkle, S. J.; Shiiyama, K.; Snead, L. L.; Farnum, E. H. J. Nucl. Mater. 1998, 258–263, 1867. Shiiyama, K.; Howlader, M. M. R.; Zinkle, S. J.; et al. J. Nucl. Mater. 1998, 258–263, 1848. Shikama, T.; Zinkle, S. J. J. Nucl. Mater. 1998, 258–263, 1861. Chernov, V. M.; Khorasanov, G. L.; Plaksin, O. A.; Stepanov, V. A.; Stepanov, P. A.; Belyakov, V. A. J. Nucl. Mater. 1998, 253, 175. Pells, G. P.; Hodgson, E. R. J. Nucl. Mater. 1995, 226, 286. Moron˜o, A.; Hodgson, E. R. J. Nucl. Mater. 1997, 250, 156. Hodgson, E. R.; Moron˜o, A. J. Nucl. Mater. 2000, 283–287, 880. Yasuda, K.; Tanaka, K.; Shimada, M.; Yamamoto, T.; Matsumura, S.; Kinoshita, C. J. Nucl. Mater. 2004, 329–333, 1451. Zong, X. F.; Shen, C. F.; Liu, S.; et al. Phys. Rev. B 1996, 54, 139. Moron˜o, A.; Hodgson, E. R. J. Nucl. Mater. 1996, 233–237, 1299. Moron˜o, A.; Hodgson, E. R. J. Nucl. Mater. 1998, 258–263, 1798. Narui, M.; Shikama, T.; Endo, Y.; Sagawa, T.; Kayano, H. J. Nucl. Mater. 1992, 191–194, 592. Ooms, H.; Hodgson, E. R.; Decre´ton, M.; et al. Fusion Eng. Des. 2007, 82, 2531. Gonza´lez, S. M.; Moron˜o, A.; Hodgson, E. R. Fusion Eng. Des. 2007, 82, 2567. Narui, M.; Sagawa, T.; Endo, Y.; et al. J. Nucl. Mater. 1994, 212–215, 1645. Narui, M.; Sagawa, T.; Shikama, T. J. Nucl. Mater. 1998, 258–263, 372. Gusarov, A.; Vermeeren, L.; Brichard, B.; et al. Fusion Eng. Des. 2005, 75–79, 819. Pells, G. P.; Hill, G. J. J. Nucl. Mater. 1986, 141–143, 375. Frost, H. M.; Clinard, F. W., Jr., J. Nucl. Mater. 1988, 155–157, 315. Buckley, S. N.; Agnew, P. J. Nucl. Mater. 1988, 155–157, 361. Heidinger, R.; Ko¨niger, F. J. Nucl. Mater. 1988, 155–157, 344. Heidinger, R. J. Nucl. Mater. 1990, 173, 243. Dienst, W.; Fett, T.; Heidinger, R.; Ro¨hrig, H. D.; Schulz, B. J. Nucl. Mater. 1990, 174, 102. Pells, G. P.; Heidinger, R.; Ibarra-Sanchez, A.; Ohno, H.; Goulding, R. H. J. Nucl. Mater. 1992, 191–194, 535. Molla´, J.; Heidinger, R.; Ibarra, A. J. Nucl. Mater. 1994, 212–215, 1029. Vila, R.; Gonzalez, M.; Molla´, J.; Ibarra, A. J. Nucl. Mater. 1998, 253, 141. Farnum, E. H.; Kennedy, J. C.; Clinard, F. W.; Frost, H. M. J. Nucl. Mater. 1992, 191–194, 548. Stoller, R. E.; Goulding, R. H.; Zinkle, S. J. J. Nucl. Mater. 1992, 191–194, 602. Goulding, R. H.; Zinkle, S. J.; Rasmussen, D. A.; Stoller, R. E. J. Appl. Phys. 1996, 79, 2920.
121. 122. 123. 124. 125. 126. 127. 128. 129. 130. 131. 132. 133. 134. 135. 136. 137. 138. 139. 140. 141. 142. 143. 144. 145. 146. 147. 148. 149. 150. 151. 152.
153.
154. 155.
723
Vila, R.; Hodgson, E. R. J. Nucl. Mater. 2000, 283–287, 903. Moron˜o, A.; Hodgson, E. R. J. Nucl. Mater. 1997, 249, 128. Molla´, J.; Ibarra, A.; Hodgson, E. R. J. Nucl. Mater. 1994, 212–215, 1113. Molla´, J.; Ibarra, A.; Hodgson, E. R. J. Nucl. Mater. 1995, 219, 182. Vila, R.; Hodgson, E. R. J. Nucl. Mater. 2011; doi:10.1016/j.jnucmat.2011.01.080. Heidinger, R.; Hofmann, A.; Nickel, H. U.; Norajitra, P. Fusion Eng. Des. 1991, 18, 337. Ibarra, A.; Heidinger, R.; Molla, J. J. Nucl. Mater. 1992, 191–194, 530. Heidinger, R. J. Nucl. Mater. 1994, 212–215, 1101. Heidinger, R. In Proceedings of 19th International Conference on Infrared and Millimeter Waves, Sendai, Japan, 1994. Braz, O.; Kasugai, A.; Sakamoto, K.; et al. Int. J. Infrared Millim. Waves 1997, 18, 1495. Ibarra, A.; Gonzalez, M.; Vila, R.; Molla, J. Diamond Relat. Mater. 1997, 6, 856. Thumm, M.; Arnold, A.; Heidinger, R.; Rohde, M.; Schwab, R.; Spoerl, R. Fusion Eng. Des. 2001, 53, 517. Imai, T.; Kobayashi, N.; Temkin, R.; Thumm, M.; Tran, M. Q.; Alikaev, V. Fusion Eng. Des. 2001, 55, 281. Heidinger, R.; Rohde, M.; Spo¨rl, R. Fusion Eng. Des. 2001, 56–57, 471. Molla´, J.; Ibarra, A.; Heidinger, R.; Hodgson, E. R. J. Nucl. Mater. 1995, 218, 108. Vila, R.; Ibarra, A.; Hodgson, E. R. J. Nucl. Mater. 1996, 233–237, 1340. Molla, J.; Vila, R.; Heidinger, R.; Ibarra, A. J. Nucl. Mater. 1998, 258–263, 1884. Heidinger, R.; Ibarra, A.; Molla, J. J. Nucl. Mater. 1998, 258–263, 1822. Brichard, B.; Fernandez Fernandez, A.; Ooms, H.; et al. J. Nucl. Mater. 2004, 329–333, 1456. Donne, A. J. H.; et al. Nucl. Fusion 2007, 47, S337. Vila, R.; Molla´, J.; Heidinger, R.; Moron˜o, A.; Hodgson, E. R. J. Nucl. Mater. 2002, 307–311, 1273. Schulz, B. J. Nucl. Mater. 1988, 155–157, 348. Dienst, W. J. Nucl. Mater. 1990, 174, 102. Rohde, M.; Schulz, B. J. Nucl. Mater. 1990, 173, 289. Dienst, W.; Fett, T.; Heidinger, R.; Ro¨hrig, H. D.; Schulz, B. J. Nucl. Mater. 1990, 174, 102. Dienst, W. J. Nucl. Mater. 1992, 191–194, 555. Burghartz, St.; Schulz, B. J. Nucl. Mater. 1994, 212–215, 1065. Dienst, W. J. Nucl. Mater. 1994, 211, 186. Hazelton, C.; Rice, J.; Snead, L. L.; Zinkle, S. J. J. Nucl. Mater. 1998, 253, 190. Snead, L. L.; Zinkle, S. J.; White, D. P. J. Nucl. Mater. 2005, 340, 187. Snead, L. L.; Yamada, R.; Noda, K.; et al. J. Nucl. Mater. 2000, 283–287, 545. Schulz, B.; Haase, G. T246 Final Report Ceramics for Heating and Current Drive, and Diagnostic Systems; Hodgson, E. R., Ed.; EUR-CIEMAT 92; European Fusion Technology Programme, 1998. Heidinger, R.; Danilov, I.; Meier, A.; Rohde, M. Final Report D13: Irradiation Effects in Ceramics for Heating and Current Drive, and Diagnostic Systems; EFDA Task TW4-TPDC-IRRCER; European Fusion Technology Programme, 2006. Costley, A. E. Rev. Sci. Instrum. 1995, 66, 296. Ramsey, A. T. Rev. Sci. Instrum. 1995, 66, 871.
724
Radiation Effects on the Physical Properties of Dielectric Insulators for Fusion Reactors
156. Arnold, G. W.; Compton, W. D. Phys. Rev. Lett. 1960, 4, 66. 157. Lee, K. H.; Crawford, J. H. Phys. Rev. 1977, B15, 4065. 158. Evans, B. D.; Stapelbroek, M. Solid State Commun. 1980, 33, 765. 159. Moron˜o, A.; Hodgson, E. R. J. Nucl. Mater. 1995, 224, 216. 160. Cooke, D. W.; Bennett, B. L.; Farnum, E. H.; Thomas, D. E.; Portis, A. M. J. Nucl. Mater. 1998, 255, 180. 161. Sato, F.; Oyama, Y.; Iida, T.; Maekawa, F.; Ikeda, Y. J. Nucl. Mater. 1998, 258–263, 1897. 162. Moron˜o, A.; Hodgson, E. R. J. Nucl. Mater. 1998, 258–263, 1889. 163. Gorshkov, A.; Orlinski, D.; Sannikov, V.; et al. J. Nucl. Mater. 1999, 273, 271. 164. Moron˜o, A.; Hodgson, E. R. J. Nucl. Mater. 2007, 367–370, 1048. 165. Orlinski, D. V.; Vukolov, K. Y. Plasma Device Oper. 1999, 17, 195. 166. Garcia-Matos, M.; Moron˜o, A.; Hodgson, E. R. J. Nucl. Mater. 2000, 283–287, 890. 167. Sugie, T.; Nishitani, T.; Kasai, S.; Kaneko, J.; Yamamoto, S. J. Nucl. Mater. 2002, 307–311, 1264. 168. Vukolov, K. Yu.; Levin, B. A. Fusion Eng. Des. 2003, 66–68, 861. 169. Vukolov, K. Yu. Fusion Eng. Des. 2009, 84, 1961. 170. Leo´n, M.; Martı´n, P.; Vila, R.; Molla, J.; Ibarra, A. Fusion Eng. Des. 2009, 84, 1174. 171. Griscom, D. L. J. Non-Cryst. Solids 1985, 73, 51. 172. Friebele, E. J.; Griscom, D. L.; Marrone, M. J. J. Non-Cryst. Solids 1985, 71, 133. 173. Griscom, D. L.; Ceram, J. Soc. Jpn 1991, 99, 923. 174. Orlinski, D. V.; Altovsky, I. V.; Bazilevskaya, T. A.; et al. J. Nucl. Mater. 1994, 212–215, 1059. 175. Gritayna, V. T.; Bazilevskaya, T. A.; Voitsenya, V. S.; Orlinski, D. V.; Tarabrin, Yu. A. J. Nucl. Mater. 1996, 233–237, 1310. 176. Martin, P.; Moron˜o, A.; Hodgson, E. R. J. Nucl. Mater. 2000, 283–287, 894. 177. Martin, P.; Moron˜o, A.; Hodgson, E. R. J. Nucl. Mater. 2004, 329–333, 1442.
178. 179. 180. 181. 182. 183. 184. 185. 186.
187. 188. 189. 190. 191. 192. 193. 194.
Gorbunov, A. V.; Klassen, N. V.; Orlinski, D. V.; Vukolov, K. Yu. Fusion Eng. Des. 2005, 74, 815. Johnson, C. D. CERN/PS/90–42 (AR), 1990. MacCarthy, K. J.; Baciero, A.; Zurro, B.; et al. J. Appl. Phys. 2002, 92, 6541. Toh, K.; Shikama, T.; Nagata, S.; Tsuchiya, B.; Yamauchi, M.; Nishitani, T. J. Nucl. Mater. 2007, 367–370, 1128. Toh, K.; Shikama, T.; Katsui, H.; et al. J. Nucl. Mater. 2009, 386–388, 1027. Moron˜o, A.; Martı´n, P.; Gusarov, A.; Hodgson, E. R. J. Nucl. Mater. 2009, 386, 1030. Nagata, S.; Katsui, H.; Tsuchiya, B.; et al. J. Nucl. Mater. 2009, 386, 1045. Cooke, D. W.; Bennett, B. L.; Farnum, E. H. J. Nucl. Mater. 1996, 232, 214. Deparis, O.; Me´gret, P.; Decre´ton, M.; Blondel, M.; Golant, K. M.; Tomashuk, A. L. In Diagnostics for Experimental Thermonuclear Fusion Reactors; Stott, P. E., Gorini, G., Prandoni, P., Sindoni, E., Eds.; Plenum Press: New York, 1998; Vol. 2, p 291. Kakuta, T.; Sakasai, K.; Shikama, T.; Narui, M.; Sagawa, T. J. Nucl. Mater. 1998, 258–263, 1893. Kakuta, T.; Shikama, T.; Nishitani, T.; et al. J. Nucl. Mater. 2002, 307–311, 1277. Toh, K.; Shikama, T.; Nagata, S.; et al. J. Nucl. Mater. 2004, 329–333, 1495. Sporea, D.; Sporea, A.; Constantinescu, B. Fusion Eng. Des. 2005, 74, 763. Nishitani, T.; Shikama, T.; Reichle, R.; et al. Fusion Eng. Des. 2002, 63–64, 437. Shikama, T.; Nishitani, T.; Kakuta, T.; et al. Nucl. Fusion 2003, 43, 517. Orlovskiy, I. I.; Vukolov, K. Yu. Fusion. Eng. Des. 2005, 74, 865. Pells, G. P.; Boothby, R. M. J. Nucl. Mater. 1998, 256, 25.