Understanding and Mitigating Ageing in Nuclear Power Plants: Materials and Operational Aspects of Plant Life Management (PLiM) (Woodhead Publishing Series in Energy)

i

Understanding and mitigating ageing in nuclear power plants

© Woodhead Publishing Limited, 2010

ii

Related titles: Geological repository systems for safe disposal of spent nuclear fuels and radioactive waste (ISBN 978-1-84569-542-2) The long-term fate of spent nuclear fuel and radioactive waste materials is of critical importance to the nuclear industry. Long-term safety must be assured without active human oversight, based on the requirement that we do not pass the burden of nuclear waste onto future generations. Geological disposal systems and technology, utilising both natural geological barriers and engineered barrier systems, have therefore been developed to isolate spent nuclear fuel and radioactive materials from the human environment. This book critically reviews state-of-the-art technologies, scientific methods and engineering practices directly related to the design, operation and safety of geological repositories. Advanced separation techniques for nuclear fuel reprocessing and radioactive waste treatment (ISBN 978-1-84569-501-9) Advanced separation technology is key to ensuring the environmental protection, criticality-safety and non-proliferation (security) of radioactive materials. Separation techniques are applied in reprocessing nuclear fuels for further power generation as well as in radioactive waste treatment for radiological-content reduction of materials due for storage and disposal. Initial sections review materials management issues and the characterisation and behaviour of the main radioactive materials produced by nuclear power generation. Further sections cover irradiated fuel treatment, such as the PUREX process, and radioactive waste treatment, such as supercritical fluid extraction. Final sections cover next generation processes, including partitioning and transmutation, and applications. Materials for energy conversion devices (ISBN 978-1-85573-932-1) The term electroceramic is used to describe ceramic materials that have been specially formulated with specific electrical, magnetic or optical properties. Electroceramics are of increasing importance in many key technologies including: microelectronics, communications and energy conversion. This innovative book is the first comprehensive survey on major new developments in electroceramics for energy conversion devices. It presents current research from leading innovators in the field. Details of these and other Woodhead Publishing books can be obtained by: ∑ visiting our web site at www.woodheadpublishing.com ∑ contacting Customer Services (e-mail: [email protected]; fax: +44 (0) 1223 893694; tel.: +44 (0) 1223 891358 ext. 130; address: Woodhead Publishing Limited, Abington Hall, Granta Park, Great Abington, Cambridge CB21 6AH, UK) If you would like to receive information on forthcoming titles, please send your address details to: Francis Dodds (address, tel. and fax as above; e-mail: francis. [email protected]). Please confirm which subject areas you are interested in.

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Woodhead Publishing Series in Energy: Number 4

Understanding and mitigating ageing in nuclear power plants Materials and operational aspects of plant life management (PLiM) Edited by Philip G. Tipping

Oxford

Cambridge

Philadelphia

New Delhi

© Woodhead Publishing Limited, 2010

iv Published by Woodhead Publishing Limited, Abington Hall, Granta Park, Great Abington, Cambridge CB21 6AH, UK www.woodheadpublishing.com Woodhead Publishing, 525 South 4th Street #241, Philadelphia, PA 19147, USA Woodhead Publishing India Private Limited, G-2, Vardaan House, 7/28 Ansari Road, Daryaganj, New Delhi – 110002, India www.woodheadpublishingindia.com First published 2010, Woodhead Publishing Limited © Woodhead Publishing Limited, 2010 The authors have asserted their moral rights. This book contains information obtained from authentic and highly regarded sources. Reprinted material is quoted with permission, and sources are indicated. Reasonable efforts have been made to publish reliable data and information, but the authors and the publisher cannot assume responsibility for the validity of all materials. Neither the authors nor the publisher, nor anyone else associated with this publication, shall be liable for any loss, damage or liability directly or indirectly caused or alleged to be caused by this book. Neither this book nor any part may be reproduced or transmitted in any form or by any means, electronic or mechanical, including photocopying, microfilming and recording, or by any information storage or retrieval system, without permission in writing from Woodhead Publishing Limited. The consent of Woodhead Publishing Limited does not extend to copying for general distribution, for promotion, for creating new works, or for resale. Specific permission must be obtained in writing from Woodhead Publishing Limited for such copying. Trademark notice: Product or corporate names may be trademarks or registered trademarks, and are used only for identification and explanation, without intent to infringe. British Library Cataloguing in Publication Data A catalogue record for this book is available from the British Library. ISBN 978-1-84569-511-8 (print) ISBN 978-1-84569-995-6 (online) ISSN 2044-9364 Woodhead Publishing Series in Energy (print) ISSN 2044-9372 Woodhead Publishing Series in Energy (online) The publisher’s policy is to use permanent paper from mills that operate a sustainable forestry policy, and which has been manufactured from pulp which is processed using acid-free and elemental chlorine-free practices. Furthermore, the publisher ensures that the text paper and cover board used have met acceptable environmental accreditation standards.  Typeset by Replika Press Pvt Ltd, India Printed by TJI Digital, Padstow, Cornwall, UK

© Woodhead Publishing Limited, 2010

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Contents

Contributor contact details

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Woodhead Publishing Series in Energy

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Foreword

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Y. Dou, Shanghai Nuclear Energy Research and Development Institute (SNERDI), P. R. China

Executive summary

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Ph. G. Tipping, Nuclear Energy and Materials Consultant, Switzerland

Part I Introduction to plant life management (PLiM), safety regulation and economics of nuclear power plants 1

Introduction to nuclear energy, and materials and operational aspects of nuclear power plants



Ph. G. Tipping, Nuclear Energy and Materials Consultant, Switzerland

1.1 1.2 1.3 1.4

Introduction Age as a relative term The importance of nuclear energy Learning from experience to continually improve safety in nuclear power plants (NPPs) Global situation of the status of installed nuclear power in 2010 The importance of keeping nuclear power plants (NPPs) operating safely and reliably Political and climate change issues and disposal of radioactive waste Energy resources: a comparison

1.5 1.6 1.7 1.8

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3 4 5 7 8 10 10 11

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Contents

1.9 1.10

Further ageing aspects in nuclear power plants (NPPs) Historical evolution of nuclear power and some materials aspects Overview of two important materials issues in older design nuclear power plants (NPPs) Conclusions Sources of further information References

1.11 1.12 1.13 1.14

12 13 14 16 17 17

2

Key elements and principles of nuclear power plant life management (PLiM) for current and long-term operation



Ph. G. Tipping, Nuclear Energy and Materials Consultant, Switzerland

2.1 2.2

Introduction Nuclear power plant ageing terminology and associated definitions Overview of ageing and its effects in nuclear power plants Overview of systems, structures and components (SSC) safety classes Setting up and scoping ageing degradation and surveillance programmes in nuclear power plants (NPPs) Safety culture and human factors and knowledge management Trends and issues in nuclear power plant (NPP) life management Past, current and future nuclear power plant (NPP) concepts and designs Conclusions Sources of further information Acknowledgements References

47 52 53 54 54

3

Safety regulations for nuclear power plant life management and licence renewal

56



A. Alonso, Universidad Politécnica de Madrid, Spain

3.1 3.2 3.3 3.4 3.5

Introduction Safety review/licence renewal Surveillance, operation and maintenance programmes Integration of plant life management Ageing degradation mechanisms, and time-limited structures, systems and components

2.3 2.4 2.5 2.6 2.7 2.8 2.9 2.10 2.11 2.12

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19 22 24 31 34 39 41

56 57 65 70 72

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Main areas of concern for plant designers, operators and regulators Future trends References

76 80 85

4

Probabilistic and deterministic safety assessment methods for nuclear power plant life management

88



P. Contri and A. Rodionov, European Commission DG-JRC Institute for Energy, The Netherlands

4.1

Introduction – plant safety assessment in a plant life management (PLiM) framework The plant life management (PLiM) problem – definitions and selected experience cases A unified proposal for a plant life management (PLiM) model integrating maintenance optimization Probabilistic safety assessment of components and systems Impact of ageing effects at system and plant level Conclusions References

103 108 109 113 114

5

Assessing the socio-economic impacts of ageing and plant life management (AM-PLiM) programmes for long-term operation (LTO) of nuclear power plants (NPPs)

117



Ph. G. Tipping, Nuclear Energy and Materials Consultant, Switzerland

5.1

Nuclear power as part of the global energy mix: energy demand, environmental issues and manpower Aspects of current and future nuclear fuel supply and its impact on the viability of nuclear power Economic overview of the nuclear power plant (NPP) lifecycle Cost drivers of nuclear power plant (NPP) operation Basic economic requirements for sustainable operation of nuclear power plants (NPPs) Assessing the costs and economics of nuclear power plant (NPP) operation and the impact of ageing and plant life management (AM-PLiM) programmes for long-term operation (LTO) Conclusions Sources of further information and advice Acknowledgements References

3.7 3.8

4.2 4.3 4.4 4.5 4.6 4.7

5.2 5.3 5.4 5.5 5.6

5.7 5.8 5.9 5.10

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117 119 120 122 123

124 126 127 127 127

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Part II Ageing degradation of irradiated materials in nuclear power plant systems, structures and components (SSC): mechanisms, effects and mitigation techniques 6

Failure prevention and analysis in nuclear power plant systems, structures and components (SSC): a holistic approach



Ph. G. Tipping, Nuclear Energy and Materials Consultant, Switzerland

6.1 6.2

Introduction Reducing failure probablility and consequences thereof in nuclear power plant (NPP) systems, structures and components (SSCs) Latent failure conditions (LFCs) and failure terminology Holistic approach to analysing nuclear power plant systems, structures and components (NPP SSC) failure events Discussion Conclusions Sources of further information References

138 141 142 143 144

7

Impact of operational loads and creep, fatigue and corrosion interactions on nuclear power plant systems, structures and components (SSC)

146



M. Bakirov, Center of Material Science and Lifetime Management Ltd, Russia

7.1 7.2 7.3 7.4 7.5

Introduction Nuclear power plant (NPP) equipment materials Medium and corrosion Stress-corrosion cracking Evaluation of impact of thermo-mechanical loading on strength of equipment materials Equipment condition monitoring, prediction and testing Conclusions and future trends Acknowledgements References and further reading

164 176 184 185 185

8

Microstructure evolution of irradiated structural materials in nuclear power plants

189



M. Hernández-Mayoral, CIEMAT, Spain and M. J. Caturla, University of Alicante, Spain

8.1

Introduction

6.3 6.4 6.5 6.6 6.7 6.8

7.6 7.7 7.8 7.9

131

131 133 136

146 149 151 158

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8.2 8.3 8.4 8.5 8.6 8.7 8.8 8.9 8.10

Structures and materials affected Environmental and other stressors Changes in the microstructure and degradation mechanisms Mitigation paths Application of research and operational experience to the practical solution of problems Acknowledgements Definitions Sources of further information and advice References

9

Stress corrosion cracking (SCC) of austenitic stainless steels in high temperature light water reactor (LWR) environments



P. L. Andresen, GE Global Research Center, USA

9.1 9.2 9.3

Introduction Historical problems and structures affected Stress corrosion cracking (SCC) dependencies – introduction Stress corrosion cracking (SCC) dependencies – materials and water chemistry Stress corrosion cracking (SCC) dependencies – cold work, stress intensity factor and irradiation Stress corrosion cracking (SCC) dependencies – miscellaneous Mechanism of stress corrosion cracking (SCC) Stress corrosion cracking (SCC) mitigation Prediction of stress corrosion cracking (SCC) and irradiation assisted stress corrosion cracking (IASCC) Future trends Sources of further information and advice References

9.4 9.5 9.6 9.7 9.8 9.9 9.10 9.11 9.12 10

Void swelling and irradiation creep in light water reactor (LWR) environments



F. A. Garner, Radiation Effects Consulting, USA

10.1 10.2

Introduction to void swelling and irradiation creep Potential for swelling and irradiation creep in light water cooled reactors (LWRs) Predictions of void swelling and associated uncertainties Potential swelling/creep consequences Second-order effects associated with or concurrent with void swelling

10.3 10.4 10.5

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236 236 238 243 257 270 281 291 294 296 299 302 302 308 308 318 332 338 340

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10.6 10.7

Conclusion References

349 349

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Irradiation hardening and materials embrittlement in light water reactor (LWR) environments

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M. Brumovsky, Nuclear Research Institute Rez plc, Czech Republic

11.1 11.2 11.3 11.4 11.5 11.6 11.7

Introduction Irradiation conditions Nature of radiation damage Irradiation hardening and embrittlement Main factors Predictive formulae Detection and measurement of irradiation hardening and embrittlement 11.8 Conclusions 11.9 Sources of further information and advice 11.10 References 12

Reactor pressure vessel (RPV) annealing and mitigation in nuclear power plants



M. Brumovsky, Nuclear Research Institute Rez plc, Czech Republic

12.1 12.2 12.3 12.4 12.5

Introduction Structures and materials affected Main mitigation measures Mitigation mechanisms including microstructure changes Application of research and operational experience to the practical solution of problems Conclusions Sources of further information References

12.6 12.7 12.8

Part III Analysis of nuclear power plant materials, and application of advanced systems, structures and components (SSC) 13 Characterization techniques for assessing irradiated and ageing materials in nuclear power plant systems, structures and components (SSC)

357 358 359 362 365 367 368 371 371 372 374 374 375 375 379 381 385 385 385

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S. Lozano-Perez, University of Oxford, UK

13.1 13.2 13.3 13.4

Introduction Non-destructive techniques Destructive techniques Recent advances, future trends and new techniques

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13.5

References

412

14

On-line and real-time corrosion monitoring techniques of metals and alloys in nuclear power plants and laboratories

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L. Yang and K. T. Chiang, Southwest Research Institute, USA

14.1 14.2 14.3 14.4 14.5 14.6 14.7

Introduction General corrosion monitoring Localized corrosion monitoring Electrochemical potential (ECP) monitoring Conclusion Acknowledgements References

417 418 431 443 448 450 451

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Multi-scale modelling of irradiation effects in nuclear power plant materials

456



L. Malerba, SCK.CEN, Belgium

15.1 15.2 15.3 15.4 15.5 15.6 15.7 15.8 15.9 15.10 15.11 15.12

Introduction An overview of radiation effects Multi-scale modelling Nuclear- and atomic-level interactions Atomic-level modelling Microstructure evolution modelling Mechanical property modelling Example of application: the PERFECT example Discussion Conclusion Acknowledgements References

456 459 474 478 483 495 503 512 519 523 524 524

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Development and application of instrumentation and control (I&C) components in nuclear power plants (NPPs)

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H. M. Hashemian, Analysis and Measurement Services Corporation, USA

16.1 16.2

Introduction Instrumentation and control (I&C) components in nuclear power plants (NPPs) Key instrumentation and control (I&C) components Ageing and instrumentation and control (I&C) Mitigating ageing in instrumentation and control (I&C) components Online monitoring (OLM)

16.3 16.4 16.5 16.6

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16.7 16.8 16.9 16.10

Online monitoring (OLM) methods and ageing management 573 Future trends 575 Sources of further information and advice 578 Bibliography 579

17

Development and application of nano-structured materials in nuclear power plants



W. Hoffelner, Paul Scherrer Institut, Switzerland

17.1 17.2 17.3

Introduction Ferritic-martensitic 9–12% Cr steels Dispersion strengthened ferritic and ferritic-martensitic steels Other routes for nano-particle strengthening Mechanical properties Components Application of research and operational experience to the practical solution of problems (relation to plant life management, PLiM) Conclusions References

17.4 17.5 17.6 17.7 17.8 17.9

581 581 584 586 587 590 594 596 601 602

Part IV Plant life management (PLiM) practices in nuclear power plants 18

Plant life management (PLiM) practices for pressurized light water nuclear reactors (PWR)



Ph. G. Tipping, Nuclear Energy and Materials Consultant, Switzerland

18.1 18.2

Introduction Ageing-related terminology and descriptions of major pressurized water reactor (PWR) components Overview of fuel and control of core power in pressurized water reactors (PWR) Discussion Conclusions Sources of further information References

18.3 18.4 18.5 18.6 18.7 19

Plant life management (PLiM) practices for water-cooled water-moderated nuclear reactors (WWER)



T. J. Katona, Paks Nuclear Power Plant Ltd, Hungary

19.1

Introduction

609

609 612 623 624 627 628 630

633 633

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19.2 19.3 19.4 19.5 19.6 19.7 19.8 19.9 19.10 19.11 19.12 19.13 19.14

Description of water-cooled water-moderated nuclear reactors (WWERs) Plant life management (PLiM) policy of water-cooled water-moderated nuclear reactor (WWER) operators Mechanical components relevant for safe long-term operation Structures and structural components relevant for safe long-term operation Electrical, instrumentation and control equipment relevant for safe long-term operation Regulatory requirements for continued operation Integration of plant life management (PLiM) programmes for water-cooled water-moderated nuclear reactor (WWERs) Feedback of operational experience Research needs in area of ageing of water-cooled watermoderated nuclear reactor (WWER) components Role of international organizations and programmes Future trends Sources of further information References

20

Plant life management (PLiM) practices for boiling water nuclear reactors (BWR): Japanese experience



N. Sekimura, University of Tokyo, Japan and N. Yamashita, Tokyo Electric Power Company, Japan

20.1 20.2

Introduction Features and types of boiling water reactors – boiling water reactor (BWR) and advanced boiling water reactor (ABWR) 20.3 Major ageing mechanisms significant for boiling water reactor (BWR) systems, structures and components (SSCs) 20.4 Ageing management practices against major significant ageing mechanisms 20.5 Major component replacement/refurbishment programmes 20.6 Technical subjects to be facilitated for ageing management 20.7 Current direction for more effective and systematic ageing management programmes 20.8 Knowledge management and research and development (R&D) 20.9 References 20.10 Abbreviations

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635 639 647 657 666 673 675 693 695 696 701 702 702 706

706 708 709 714 720 723 723 726 730 730

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Plant life management (PLiM) practices for pressurised heavy water nuclear reactors (PHWR)

732

R. K. Sinha and S. K. Sinha, Bhabha Atomic Research Centre, India and K. B. Dixit, A. K. Chakrabarty and D. K. Jain, Nuclear Power Corporation of India Ltd., India

21.1 21.2

Introduction Pressurised heavy water reactor (PHWR)/Canadian Deuterium Uranium (CANDU) 21.3 Critical components of Indian pressurised heavy water reactor (PHWR) 21.4 Reactor ageing issues: pressure tube, end shields and calandria tube 21.5 Reactor ageing issues: reactivity mechanisms and fuel handling systems 21.6 Reactor ageing issues: feeders, secondary side piping, steam generators and heat exchangers 21.7 Reactor ageing issues: civil structures, cables and sea water systems 21.8 Regulatory issues associated with plant life management (PLiM) 21.9 Application of research and operational experience to find the practical solution to problems 21.10 Future trends 21.11 Acknowledgements 21.12 References 22

Plant life management (PLiM) practices for sodium cooled fast neutron spectrum nuclear reactors (SFRs)



B. Raj, P. Chellapandi, T. Jayakumar, B. P. C. Rao and K. Bhanu Sankara Rao, Indira Gandhi Centre for Atomic Research, India

22.1 22.2 22.3 22.4 22.5

Introduction Sodium cooled fast neutron spectrum reactors (SFRs) Design approach Safety and regulatory perspective In-service inspection (ISI) and robotics in life assessment based on research and development (R&D) and applications 22.6 Life extension aspects of international sodium cooled fast neutron spectrum reactors (SFRs) 22.7 Future trends 22.8 Conclusion 22.9 Acknowledgements 22.10 References

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732 739 748 755 763 765 770 773 780 784 791 791 795

795 796 805 812 813 820 826 834 834 835

Contents

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Plant life management (PLiM) practices for gas-cooled, graphite-moderated nuclear reactors: UK experience



G. B. Neighbour, University of Hull, UK

23.1 23.2

Introduction UK gas-cooled reactor types (Magnox and advanced gas-cooled reactor (AGR)) 23.3 Nuclear graphite 23.4 Effects of reactor environment on the graphite moderator 23.5 The UK nuclear regulatory regime 23.6 Maintaining the safety of graphite moderator cores 23.7 Regulatory requirements for continued operation 23.8 Future trends 23.9 Sources of further information 23.10 Useful websites 23.11 References

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838 838 840 848 852 860 862 868 870 871 872 872

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Outlook for nuclear power plant life management (PLiM) practices – summary, conclusions, recommendations



Ph. G. Tipping, Nuclear Energy and Materials Consultant, Switzerland

24.1 24.2

876

24.6 24.7

Introduction Further elements to consider for nuclear power plant ageing and plant life management (PLiM-AM) Discussion Current and projected requirements of the nuclear power industry List of topical issues of current and future relevance to nuclear power plants (NPPs) Conclusions References



Index

889

24.3 24.4 24.5

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877 879 880 881 885 888

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Contributor contact details

(* = main contact)

Chapters 1, 2, 5, 6, 18 and 24 Ph. G. Tipping Nuclear Energy and Materials Consultant – NE&MC CH 5200 Brugg Switzerland E-mail: [email protected]

Chapter 3

Chapter 7 M. Bakirov Center of Material Science and Lifetime Management Ltd Office 5 Kirova str. 7 Lubertsy City Moscow Region 140002 Russia E-mail: [email protected]

A. Alonso Universidad Politécnica de Madrid Departamento de Ingenieria Nuclear Rafael Calvo 3 2F 28010 Madrid Spain E-mail: [email protected]

Chapter 4 P. Contri* and A. Rodionov European Commission, DG-JRC Institute for Energy Safety of Current Reactors Unit PO Box 2 1755 ZG Petten, The Netherlands Westerduinweg 3 1755 LE Petten The Netherlands

Chapter 8 M. Hernández-Mayoral Division of Structural Materials Department of Technology CIEMAT 28040 Madrid Spain E-mail: [email protected]

M. J. Caturla* Department Fisica Aplicada Facultad de Ciencias, Fase II Universidad de Alicante 03690 Alicante Spain E-mail: [email protected]

E-mail: [email protected] © Woodhead Publishing Limited, 2010

xviii

Contributor contact details

Chapter 9

Chapter 14

P. L. Andresen GE Global Research Center One Research Circle CE2513 Schenectady, NY 12309 USA

L. Yang and K. T. Chiang Department of Earth, Material, and Planetary Sciences Southwest Research Institute® (SwRI®) 6220 Culebra Rd San Antonio, TX 78228 USA

E-mail: [email protected]

Chapter 10

E-mail: [email protected]

F. A. Garner Radiation Effects Consulting Richland, WA 99354 USA

Chapter 15

E-mail: [email protected]

Chapters 11 and 12 M. Brumovsky Nuclear Research Institute Rez plc, Division of Integrity and Technical Engineering 250 68 Rez Czech Republic E-mail: [email protected]

L. Malerba Structural Materials Group Institute of Nuclear Materials Science Studiecentrum voor Kernenergie ∑ Centre d’Etude de l’Energie Nucléaire (SCK ∑ CEN) Boeretang 200 B-2400 Mol Belgium E-mail: [email protected]

Chapter 16

Chapter 13 S. Lozano-Perez Department of Materials University of Oxford Parks Road Oxford OX1 3PH UK

H. M. Hashemian Analysis and Measurement Services Corporation AMS Technology Center 9111 Cross Park Drive Knoxville, TN 37923 USA E-mail: [email protected]

E-mail: sergio.lozano-perez@materials. ox.ac.uk

© Woodhead Publishing Limited, 2010

Contributor contact details

xix

Chapter 17

Chapter 21

W. Hoffelner Paul Scherrer Institut CH-5232 Villigen PSI Switzerland

R. K. Sinha* and S.K. Sinha Reactor Design and Development Group Bhabha Atomic Research Centre Mumbai India

E-mail: [email protected]

Chapter 19

E-mail: [email protected]

T. J. Katona Paks Nuclear Power Plant Ltd P.O. Box 71 Paks 7031 Hungary

K. B. Dixit, A. K. Chakrabarty and D. K. Jain Engineering and Procurement Nuclear Power Corporation of India Limited India

E-mail: [email protected]

Chapter 22

Chapter 20 N. Sekimura* Department of Nuclear Engineering and Management University of Tokyo 7-3-1 Hongo Bunkyo-ku 113-8656 Tokyo Japan E-mail: [email protected]

N. Yamashita Reactor Mechanical Maintenance Group Maintenance Department (Unit 3&4) Fukushima Dai-ichi Nuclear Power Station Tokyo Electric Power Company 22 Kitahara Ottozawa Ohkumamachi Futaba-gun, 979-1301 Fukushima Prefecture Japan

B. Raj, P. Chellapandi, T. Jayakumar, B. P. C. Rao and K. Bhanu Sankara Rao Indira Gandhi Centre for Atomic Research Kalpakkam TN – 603 102 India E-mail: [email protected]

Chapter 23 G. B. Neighbour Materials and Process Performance Department of Engineering University of Hull Hull HU6 7RX UK E-mail: [email protected]

E-mail: [email protected]

© Woodhead Publishing Limited, 2010

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xxi

Woodhead Publishing Series in Energy

1 Generating power at high efficiency: Combined cycle technology for sustainable energy production Eric Jeffs 2 Advanced separation techniques for nuclear fuel reprocessing and radioactive waste treatment Edited by Kenneth L. Nash and Gregg J. Lumetta 3 Bioalcohol production: Biochemical conversion of lignocellulosic biomass Edited by K. W. Waldron 4 Understanding and mitigating ageing in nuclear power plants: Materials and operational aspects of plant life management (PLiM) Edited by Philip G. Tipping 5 Advanced power plant materials, design and technology Edited by Dermot Roddy 6 Stand-alone and hybrid wind energy systems: Technology, energy storage and applications Edited by J. K. Kaldellis 7 Biodiesel science and technology: From soil to oil Jan C. J. Bart, Natale Palmeri and Stefano Cavallaro 8 Developments and innovation in carbon dioxide (CO2) capture and storage technology Volume 1: Carbon dioxide (CO2) capture, transport and industrial applications Edited by M. Mercedes Maroto-Valer 9 Geological repository systems for safe disposal of spent nuclear fuels and radioactive waste Edited by Joonhong Ahn and Michael J. Apted

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Woodhead Publishing Series in Energy

10 Wind energy systems: Optimising design and construction for safe and reliable operation Edited by John D. Sørensen and Jens N. Sørensen 11 Solid oxide fuel cell technology: Principles, performance and operations Kevin Huang and John Bannister Goodenough 12 Handbook of advanced radioactive waste conditioning technologies Edited by Michael I. Ojovan 13 Nuclear reactor safety systems Edited by Dan Gabriel Cacuci 14 Materials for energy efficiency and thermal comfort in buildings Edited by Matthew R. Hall 15 Handbook of biofuels production: Processes and technology Edited by Rafael Luque, Juan Campelo and James Clark 16 Developments and innovation in carbon dioxide (CO2) capture and storage technology Volume 2: Carbon dioxide (CO2) storage and utilisation Edited by M. Mercedes Maroto-Valer 17 Oxy-fuel combustion for power generation and carbon dioxide (CO2) capture Edited by Ligang Zheng 18 Small and micro combined heat and power (CHP) systems: Advanced design, performance, materials and applications Edited by Robert Beith 19 Hydrocarbon fuel conversion technology: Advanced processes for clean fuel production Edited by M. Rashid Khan 20 Modern gas turbine systems: High efficiency, low emission, fuel flexible power generation Edited by Peter Jansohn 21 Concentrating solar power (CSP) technology: Developments and applications Edited by Keith Lovegrove and Wes Stein 22 Nuclear corrosion science and engineering Edited by Damien Féron

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Woodhead Publishing Series in Energy

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23 Power plant life management and performance improvement Edited by John Oakey 24 Direct-drive wind and marine energy systems Edited by Markus Mueller 25 Advanced membrane science and technology for sustainable energy and environmental applications Edited by Angelo Basile and Suzana Nunes 26 Irradiation embrittlement of reactor pressure vessels (RPVs) Edited by Naoki Soneda 27 High temperature superconductors (HTS) for energy applications Edited by Ziad Melhem 28 Infrastructure and methodologies for the justification of nuclear power programmes Edited by Agustín Alonso Santos

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© Woodhead Publishing Limited, 2010

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Foreword Y. D ou, Shanghai Nuclear Energy Research and Development Institute (SNERDI), P. R. China

The development and commercial use of nuclear energy in the form of costcompetitive electricity and district heating was one of the most significant industrial achievements of the 20th century. The 21st century is now approaching its second decade, and global energy consumption continues to grow rapidly, and projections show a doubling of world electricity demand by 2030, creating the need for considerable amounts of newly installed generating capacity (of all origins) in the next 25 years. The current global nuclear energy capacity is about 367 GWe, and between 524 GWe and 740 GWe is expected to be needed before 2050. This will necessitate building between 200 and 400 new nuclear power plants (NPPs) worldwide to just replace the lost capacity of electrical power from decommissioned NPPs and to provide the new capacity that will be necessary by 2050, not only to satisfy the needs of growing energy consumption, but also to ease environmental pressure by reducing dependence on fossil-based energy. For example, in China, where the economy has continuously increased by 9% per year, on average, over the past 30 years, the current portion of nuclear-generated energy is only around 2%, which is a relatively small percentage in comparison with the world’s current nuclear-generated energy portion of about 16%. To satisfy the demand for energy in an environmentally friendly way, free from greenhouse gas emissions, China is thus launching an ambitious nuclear power plan to raise the portion of nuclear energy by 4–5% by the year 2020. That means that 40–60 new NPPs of 1000 MWe should be built in a decade or more. Such ambitions for new-build requirements for NPPs create considerable logistics and planning tasks, having due regard for the industry’s overall current and projected ability to supply the heavy equipment required for NPPs (e.g. pressure vessels, quality alloys for piping, steam generators, core shrouds, pumps, high quality cement for containment buildings and cooling towers, etc.). Furthermore, action must be taken now to ensure that sufficiently trained, educated and experienced personnel are available to operate the plants and to regulate them. Also, the need to have fossil-fuel free low ‘greenhouse gas’ energy sources puts the focus squarely on the nuclear option, as long as it can remain safe.

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Foreword

Two crucial and decisive factors to sustain long-term operation of nuclear power plants are their safety and profitability, which can be achieved through a combination of applying optimum management strategies with an understanding of the ways in which the safety-related systems, structures and components (SSCs) perform and interact in their respective operating environments. The SSCs are made of various materials (e.g. metals, alloys, concrete, plastics) and it is the behaviour of materials due to their operational conditions (e.g. temperature, pressure, loading, irradiation, coolant chemistry) that can lead to ageing degradation with attendant impact on NPP availability, SSC reliability, or a lessening of safety margins and possible attendant operational constraints. It is the duty of NPP operators to ensure that their NPPs are safe as well as profitable. Operators can be assisted in achieving these goals by using ageing management and plant-life management approaches, based on science and technology and global experience. Ageing management and plant-life management covers various aspects of knowledge on SSCs, such as information on requirements of design basis, manufacture, installation, commissioning and operation, understanding of material degradation mechanisms, inspection programmes, evaluation and robust implementation of associated methodologies for assessing SSC fitness-for-service, use of database techniques, etc. For those countries having various designs and types of NPP, the available experts and plant personnel may need information or have to deal with tasks or problems from the different units synchronously. It is therefore extremely useful to have a single book which provides a wide range of the most recent information covering all types of NPPs on ageing and plant-life management techniques and also gives insights and guidance on how to keep NPPs operating safely and reliably. This is facilitated by understanding how ageing degradation occurs in SSCs and then by using this knowledge to develop scientifically based methods to mitigate or eliminate it. This book has been produced through a multinational team of globally recognized experts in their respective fields. The scope of this modern book is therefore immense and records the knowledge and experience gained with materials in NPPs over the last 50 or so years of nuclear power development. The suggestions for further reading and the references provided give the reader access to even more information. Accordingly, this book is a practical and technical manual for engineers, technologists and specialists involved in all aspects of NPP operation. Students and younger technologists studying nuclear technology, or those just embarking on their careers in nuclear power, will find here a source of inspiration and current information to help them achieve their academic and career goals. Furthermore, this book is a record of the current knowledge and experience that needs to be kept for future generations. This is even more so now, due to the gradual but steady loss of the “pioneer” generation of nuclear technologists and researchers, as

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they go into retirement. External organizations, including design institutes, technical supporters, sub-contractors of nuclear steam supply systems, materials suppliers and regulators will also find valuable information in this book to enable them to carry out their respective tasks. In a word, this book is essential reading for anyone associated with nuclear power. In China we have a proverb, namely, ‘Experience is the best teacher’. I believe that this book will act as one of the best teachers currently available, since it is based on the cumulative experience of more than 750 personyears when the career years of all the contributing authors are taken into consideration. It remains to be seen what the next generation of NPPs will demand in terms of specialist knowledge and operational practices. One thing is certain: as new knowledge is obtained, it must also be recorded in a book such as this.

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Executive summary P h. G. T I P P I N G, Nuclear Energy and Materials Consultant, Switzerland

Overview of the book This reference book is a comprehensive state-of-the-art, science and technology record of the current knowledge base concerning materials ageing degradation (AD), and its mitigation and elimination in systems, structures and components (SSCs) used in commercially operated nuclear power plants (NPPs). Accordingly, it covers a wide range of subjects relating to NPP SSCAD, and so the phrase ‘from atoms to zirconium’ could be a fitting one when attempting to capture the very essence of this book. It traces the historical development of commercial nuclear power, while illustrating the way our understanding and mitigation of SSC-AD has continually increased through basic and applied research approaches. The effectiveness and integration of validated SSC-AD mitigation methods (which have been largely furnished by basic research efforts) into the daily operation of NPPs is exemplified throughout. The necessity for always keeping SSCs ‘fit-for-service’ is a vital safetyrelated aspect, but it is also an important commercial requirement as well, since NPPs may thereby retain the technological basis, and thus the regulatory possibility (in satisfying licence requirements), to continue safe and reliable operation, even in excess of their original design life. The chance that this ‘long-term operation’ (LTO) can be realized is significantly enhanced when NPP plant life management (PLiM), ageing management (AM), ageing surveillance programmes (ASPs) and standard operational practices (OPs) are optimized, and when plant-specific and worldwide lessons learned are continually and robustly integrated into the operational management of NPPs. The effectiveness of PLiM, AM, ASPs and OPs is also shown to depend significantly on the level of safety culture prevailing in the NPP’s workforce, as well as on efficient plant knowledge management (KM) and associated job succession and training strategies for personnel. These essential themes are discussed throughout the book. Each chapter may be regarded as a ‘stand-alone’ contribution, providing in-depth information concerning the specific subject matter dealt with. The

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authors responsible have provided concise abstracts, summaries, conclusions, references and further information sources regarding SSC-AD, PLiM, AM and ASPs from their perspectives. The book is structured to allow the reader to select specific subject areas on the topic of interest in order to provide detailed information on SSC-AD, and also to give a perspective on the commercial nuclear power industry as a whole.

Overview of Part I Part I reviews the role of nuclear power in the global energy mix, and the importance and relevance of plant life management (PLiM) for the safety regulation and economics of nuclear power plants. In Chapter 1, by Tipping, a strong argument for the use of commercial nuclear power is provided with respect to its role as a non-fossil based energy source. Accordingly, aspects concerning the world’s climate, and the part nuclear power has in providing safe, cost-effective, and low environmental impact energy, are presented. Owing to climate and carbon dioxide emission issues, current and future use of nuclear power is considered to be a vital factor in the global effort to improve on the Earth’s environmental ‘balance of health’. (Viewed relatively, the carbon dioxide content of the atmosphere is currently about 70% more than it was 40 years ago, and this trend is continuing.) Comparison is made between nuclear and other forms of energy, not only in terms of carbon footprint and greenhouse gas aspects, but also in terms of costs and availability of fuels. Attention is brought to the finite nature of fossil-based fuel resources compared to nuclear fuel cycles that can ‘breed’ further fuel. Cost-competitive fissionable materials suitable for fuel in current and future fission-based nuclear technology are conservatively estimated to be sufficient for at least the next 500 years at projected usage rates/requirements estimates. Features and characteristics of next generation NPPs have also been provided to introduce the reader to this evolving aspect of commercial NPP development, and fusion-based nuclear power, although beyond the scope of this book, is also briefly addressed in terms of a future fossil-fuel free energy source. The key elements, principles and approaches to NPP-PLiM for plant designs most commonly in current use are explained in Chapter 2, by Tipping. Goals and essential features of PLiM, AM, ASPs and standard OPs are shown to be the result of logical and safety-based approaches to ensure that resources are primarily invested into the most important SSCs to maintain adequate safety margins and to continuously increase safety levels. Common NPP SSC-AD terminology and definitions are listed. Many SSCs can be routinely maintained or replaced, but it is the condition of the large, passive SSCs, which are deemed irreplaceable (due to practical and cost issues), that will ultimately decide the operational life of a NPP. Failures in

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NPP components, systems, machinery and structures can occur. Similarly, administrative and operational oversights can also take place and potentially contribute to a failure. However, most NPP-SSC failures cannot cause a direct threat to safety or operation, since extensive defence-in-depth (DID) design principles, redundant/back-up systems and tested accident/emergency management strategies are in place. The DID and accident management approaches are explained to give the reader an appreciation of the extent designers, operators and regulators interact to ensure, and assure, the highest levels of safety, even under accident/emergency conditions. Safety awareness, appropriate safety training courses and questioning attitudes of all NPP personnel are vital to the overall operational success of NPPs. It is shown that PLiM, AM, ASPs and OPs, based on understanding of SSC-AD mechanisms, and their mitigation, are a way to cost effectively achieve all these goals for both current and long-term operation. Safety in the operation of NPPs always has first priority. Specific safety regulations, regulatory requirements and licensing aspects of NPPs obviously vary with the country in which the plants are operating. However, as discussed in Chapter 3, by Alonso, while national structures of regulation and the way legally binding requirements are implemented/enforced may differ in detail, they share overall objectives in common, i.e. to ensure that the use of nuclear power does not affect the population or environment at any time. The main areas of concern for plant designers, operators and regulators are identified as neutron irradiation embrittlement, stress corrosion cracking (SCC) and irradiation assisted stress corrosion cracking (IASCC). International activities and research efforts are addressed, as well as aspects and methods for integration of PLiM into operating plants. Plant safety assessment for nuclear power plants, and in particular the development and application of probabilistic and deterministic methods for NPP and SSC safety assessment, are explained in Chapter 4, by Contri and Rodionov. The chapter provides definitions of terms commonly used, and highlights the main issues of SSC-AD. The field of probabilistic safety assessment (PSA) is explained with reference to the evaluation of time-dependent ageing effects. The importance of having a robust NPP organizational structure for supporting PLiM is discussed in depth. The socio-economic impacts of ageing and PLiM programmes for NPPLTO are discussed in Chapter 5, by Tipping. The importance of nuclear power in the world’s energy resource mix is presented and fuel availability and pricing is examined for current and long-term operation. Cost drivers of NPP operation are analysed, since the business case has to be adjusted to accommodate costs associated with possible (unforeseen) additional regulatory burdens/restrictions, procurement of large, expensive SSCs (e.g. steam generator replacements) and the business case necessity to amortize the plants before they are phased out of service. Significant investments in

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replacement SSCs, systems upgrading/modernization and back-fitting costs can be more readily amortized if the NPPs go into LTO. Power uprates and associated issues are explained and identified.

Overview of Part II Part II reviews the mechanisms and effects of ageing degradation of materials in nuclear power plant systems, structures and components (SSC), as well as routes taken to characterize and analyse the degradation of materials and mitigate degradation effects. Chapter 6, by Tipping, presents a holistic approach to failure analysis of NPP-SSCs, and is given as a guide to researchers, designers, operators, investigators and regulators for formulating appropriate questions to ultimately explain why failures occurred and to facilitate a logic-based methodology to determine their root causes and thereby to provide routes to take to avoid such failures in future. The phenomenon of latent failure conditions (LFCs) is outlined to provide a further concept in understanding the root causes of SSC failures. Understanding how failures occur in SSCs, and how they impact NPP operation and safety, is shown to be a many-faceted challenge, since the interaction of man-machine and human factors may contribute in subtle ways to SSC-AD types, rates and levels. Chapter 7, by Bakirov, explores the impact of operational loads and AD interactions on SSCs. Creep, fatigue and corrosion interactions are examined in depth, as well as equipment condition monitoring, prediction and testing routes used to confirm whether SSCs are in the optimum physical and chemical condition to fulfil their technical specifications, and to check that they remain resistant to the continuing operational stressors they are subjected to. In particular, the most efficient in situ (specimen-free) nondestructive monitoring and testing methods to detect levels of SSC-AD are identified Chapter 8, by Hernández-Mayoral and Caturla, reviews the evolution of microstructures in reactor pressure vessel (RPV) steels and other reactor internal structural materials, describing this evolution in terms of the physical and chemical processes taking place due to the effects of operational stressors and stresses, time and temperature. A comparison is made between ferritic/ martensitic and austenitic steels with regard to the damage processes that occur. Tools available to study irradiation-induced damage are described. Practical measures that can be taken to mitigate the effects of neutron irradiation damage are identified, being the result of robust integration of research results to the solution of problems in NPPs. Chapter 9, by Andresen, reviews stress corrosion cracking (SCC) in light water reactor (LWR) environments, covering in particular austenitic stainless steels and nickel-base alloys. This AD mechanism is examined

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in depth, and the highly complex interaction between the materials, their heat treatment, microstructure, level and type of residual stress and specific operating environment (water chemistry of coolant, neutron fluence) is explained. Furthermore, comparison is made between product forms (e.g. wrought austenitic stainless steel and cast ferritic or martensitic stainless steels) in terms of their SCC propensity in LWR coolant and environments (i.e. in boiling and pressurized water reactors (BWRs and PWRs)). Corrosion fatigue and operating environmental effects on fracture are also discussed, since they are related forms of AD. Chapter 10, by Garner, reviews the mechanisms and effects of void swelling (VS) and irradiation creep (IC) in LWR environments. The phenomena of VS and IC have the potential to cause distortion in internal structural components, and since neutron fluences increase with operational time and swelling and creep depend, among other things, on the actual level of neutron irradiation damage present, it is expected that these AD mechanisms will become significant issues in future, especially when NPPs go into LTO. Owing to lower neutron flux levels in BWRs, the problems of VS and IC appear less of an issue compared to the situation in PWRs. The development of VS is identified as being non-linear with neutron dose, thus an acceleration of this effect is expected with LTO. Predictive equations regarding deformation levels for AISI 304 austenitic stainless steel are discussed. Chapter 11, by Brumovsky, outlines examples taken from practice to illustrate how ferritic low alloy reactor pressure vessel (RPV) steel and welds may embrittle to varying degrees, depending on impurity levels (e.g. copper, phosphorus) or alloying constituents (e.g. nickel, manganese), irradiation temperature and the level of neutron fluence. The effect of tensile yield stress and hardness increase as a result of RPV alloy and weld matrix hardening through neutron irradiation induced point defects and precipitate formation is discussed, as is the concomitant loss in fracture toughness levels (Charpy notch impact or other fracture toughness measures) in susceptible RPV materials. The importance of well-characterized irradiation conditions for creating databases is highlighted. The issue of neutron irradiation induced embrittlement in RPVs is extended in Chapter 12, by Brumovsky, which explains the principles and the approaches to RPV annealing and other procedures used for mitigating and managing embrittlement rates and levels. Methods such as fuel management and RPV shielding are explained. In particular, the practical annealing of RPVs, and selection of the optimum time-at-temperature heat treatment schedule are examined in depth. The rate of re-embrittlement of RPVs has generally been shown to be less after annealing and re-irradiation, thus allowing the so-called ‘lateral shift’ approach to be adopted.

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Overview of Part III Part III reviews analysis, monitoring and modelling techniques applicable to the study of NPP materials, as well as the application of advanced systems, structures and components in NPPs. Chapter 13, by Lozano-Perez, presents science-based approaches used to detect and understand irradiation damage and related ageing in NPP-SSC materials. The distinction is made between non-destructive and destructive methods available to researchers. Aspects of both volumetric and surface techniques are also provided, so that the optimum choice of method and approach can be made to study the various AD mechanisms. The microstructures and chemical compositions of aged/irradiated materials, even at the 100 nm scale, are shown to be a key to better understanding SSC-AD. Advances in specialized tools allow atomic resolutions in two or three dimensions. Chapter 14, by Yang and Chiang, provides information concerning on-line and real-time corrosion monitoring techniques of metals and alloys in nuclear power plants and laboratories. The underlying principles and applications of general or localized corrosion monitoring are explained in detail, and the variety of tools available are described and discussed in depth. Advantages and disadvantages of different monitoring methods are critically reviewed. Specifically, the importance of accurately measuring the electrochemical potential (ECP) of alloys in the actual NPP environment of high pressure and temperature coolant is discussed, and the range of suitable electrodes that may be used is presented. Chapter 15, by Malerba presents multi-scale modelling of irradiation effects in NPP materials. Irradiation effects are treated as a multi-scale problem, and microstructural features are linked to mechanical property changes. Computer-based multi-scale modelling approaches at the atomic, microstructural and mechanical property levels are explained, and an example of practical application of modelling is provided. In particular, irradiation induced hardening and embrittlement effects in steels used in NPPs are discussed. Chapter 16, by Hashemian, discusses the significant developments in, and applications of, advanced instrumentation and control (I&C) in NPPs, as exemplified in the adoption of digital signal-based systems to replace the original analogue ones in older NPPs. The transition from analogue to digital I&C has generally enhanced safety levels and shortened response times for operational correctional procedures or emergency actions, for example. Temperature sensors, pressure transmitters, neutron detectors and associated cables are key I&C components, and may be subject to ageing. Both high and low frequency methods for on-line monitoring (OLM) are examined. However, attention is brought to the possible corruption of data and signals, and wilful interference/tampering aspects that could be an issue

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with digital-based technologies, if robust counter-measures and back-up systems are not in place. The development and application of nanostructured materials in nuclear power plants is traced and discussed in Chapter 17, by Hoffelner. Both fusion and fission-based future NPPs will operate at higher temperatures and pressures than currently operating NPPs, and the development of specially tailored alloys to resist ageing degradation is a key aspect. High temperature alloy properties include the need to resist creep, and the way to manufacture oxide-dispersion strengthened (ODS) materials is explained. The mechanical properties of ODS alloys, created by various routes – ceramic oxide dispersion, carbo-nitrides (with thermo-mechanical treatment) and internal oxidation – are discussed in depth.

Overview of Part IV Part IV reviews the particular ageing degradation issues, plant designs, and application of plant life management (PLiM) practices in a range of commercial nuclear reactor types. All reactor types operate under conditions dictated by their design concepts and engineered systems, and their SSCs are thus subjected to system-specific conditions and stressors. Some materials used in SSCs are found in all types of NPP-SSCs, but they may perform quite differently according to the specific operational conditions and environments they are exposed to. From both safety and technical standpoints, the overwhelming majority of the world’s current fleet of NPPs will be able to continue operation in excess of their original design lifetime (i.e. LTO). This is a direct consequence of effective PLiM, AM, ASPs, and OPs and focussed attention to SSC repairs, replacements or back-fitting, as required. Management of ageing effects in large, irreplaceable SSCs logically lies at the centre of NPP operational and commercial strategies, and to this end, coverage of the application of knowledge management strategies as a vital elemental aspect of overall NPP operations is also provided. The designs, plant-specific PLiM, AM, ASPs needs, and standard OPs for the main commercially operating nuclear reactor types are reviewed, including: pressurized water reactors, with PWRs covered in Chapter 18, by Tipping, and WWERs covered in Chapter 19, by Katona; boiling water reactors (BWRs), as covered in Chapter 20, by Sekimura and Yamashita; pressurized heavy water reactors (PHWRs), as covered in Chapter 21, by Sinha, Sinha, Dixit, Chakrabarty and Jain; sodium-cooled fast neutron spectrum reactors (SFRs), as covered in Chapter 22, by Raj, Chellapandi, Jayakumar, Rao and Rao, and gas-cooled graphite-moderated reactors (Magnox, AGRs), as covered in Chapter 23, by Neighbour. All of these chapters show that, irrespective of NPP type, safety and integrity of NPP-SSCs

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are closely connected to commercial viability, since SSCs are kept within design specifications and regulatory requirements, which are precursors to reliable and continuous operation. The longer a NPP can operate (i.e. free from prolonged maintenance shut-downs or forced outages), the more electrical energy (or district heating) can be sold. A NPP which has well-maintained SSCs (i.e. kept within their design and technical specifications and with sufficient safety margins), and has a good safety culture and vigilance in its workforce will, de facto, have a good safety record and a high availability and reliability.

Concluding remarks: Quo vadis nuclear power? The information in this book shows the immense amount of work done to date to understand, and mitigate, NPP SSC-AD. Indeed, thanks to PLiM, AM, ASPs, OPs and materials research, many older NPPs are now in the position to move into the LTO phase of their lives. This is the true value of such programmes, since both safety and economic aspects are optimized. As outlined in Chapter 24, by Tipping, by building on what is known today, it is expected that next generation NPPs will greatly benefit in terms of lessons learned concerning design, manufacture, choices of materials, inspection/monitoring methods and operational practices, to give just a few examples. Notwithstanding this, PLiM, AM, ASPs and OPs will still be essential for assuring safe, long and reliable operation of next generation NPPs. However, until these new NPPs come into service, it remains essential to closely follow, understand and mitigate the effects of SSC-AD in the current fleet of NPPs, as well as managing other less conventional ageing phenomena (e.g. retirement of experienced personnel, requirement to update documentation). A list of topical current and future NPP SSC-AD aspects and issues for NPPs is provided to focus on the main areas that require further research, monitoring and continued attention by designers, operators and regulators. With optimized PLiM, AM, ASPs and OPs in place, owners and operators will continue to provide their customers with environmentally clean and reliable energy at competitive prices whilst protecting their overall investment and will concomitantly provide licensing authorities with the proof that their plants are always being operated safely, irrespective of their chronological age.

Acknowledgements I would like to acknowledge and thank the international team of experts who have assisted me in creating this comprehensive book on how to understand ageing and degradation issues in nuclear power plants. The authors have

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made hereby a signal contribution to science and technology, exemplified by the high quality of the work they provided. Accordingly, this book will serve as a valuable reference and guide for all those currently involved with nuclear power generation and regulation, as well as those contemplating, or just starting, a career in nuclear power plant technology or nuclear materials research. Without the spirit of openness, the professional approaches taken and the collaboration between all contributing experts from around the world, this book could never have been written. My thanks also go to Mr Ian Borthwick (Commissioning Editor) for his encouragement, advice and support in creating this book, and to Mrs Diana Gill, Ms Ceridwen McCarthy and Ms Nell Holden, all of Woodhead Publishing Limited, for their highly competent assistance in the development and handling of this book project.

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Part I Introduction to plant life management (PLiM), safety regulation and economics of nuclear power plants

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Understanding and mitigating ageing in nuclear power plants

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Introduction to nuclear energy, and materials and operational aspects of nuclear power plants

P h. G. T i pp i n g, Nuclear Energy and Materials Consultant, Switzerland

Abstract: The importance of nuclear power is presented in terms of its significant contribution to the world’s overall generating capacity, and its essentially ‘carbon-free’ nature is highlighted. The age distribution of the current fleet of nuclear power plants (NPPs) is explained and also the need for ensuring safety in operation, irrespective of the NPP’s age or design. Global resources of nuclear fuel and comparison with other energy carriers are discussed and attention is drawn to the requirement to condition and dispose of radioactive waste correctly. Stress corrosion cracking and neutron irradiation embrittlement issues are taken as two examples of ageing degradation that have occurred in NPPs. Key words: nuclear power plants, climate change, safety, nuclear fuel, radioactive waste, stress corrosion cracking, neutron embrittlement.

1.1

Introduction

This book contains more than 750 person-years of experience in all aspects of commercial nuclear power, basic research and applied technology. The contributing authors have various backgrounds, including engineering (civil, electrical, mechanical), physics, chemistry, metallurgy, radiology and safety regulation, for example. Specialist areas within the above general categories of science and technology have evolved over the years to address issues concerned with nuclear power generation, and specifically, this book deals with key aspects of research into how structural materials used in nuclear power plants (NPPs) behave and age in operation. The book also identifies and analyses many other important issues that have continued to dominate and shape the commercial nuclear power industry during its development, evolution and regulation since the mid-1950s. The world’s political and environmental situation has changed drastically over this time, reflecting global concerns about dwindling supplies of fossil-based energy resources (oil, gas and coal) and the possible impact their use has on climate and pollution levels in the atmosphere. The importance of nuclear energy, which currently supplies about 16% of the global electrical power capacity, is further 3 © Woodhead Publishing Limited, 2010

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highlighted in terms of its low carbon footprint (CF) and no greenhouse gas emissions (GGEs) in operation. Recognizing that many NPPs constructed in the 1965–80 nuclear power ‘boom-era’, have already reached, or are approaching, the end of their original design lives, the publication of this book at this point in time appears fortuitous. However, the information in this book is not concerned with the shutting down, decommissioning, dismantling and final disposal of NPPs; on the contrary, it deals with how these older NPPs may still continue to operate safely and reliably, benefitting from lessons learned, good operational and standard maintenance practices plus the robust implementation of system, structure and component (SSC) ageing surveillance programmes (ASPs), ageing management (AM) and plant-life management (PLiM) programmes. This book has an underlying goal, namely to record relevant and updated knowledge concerning how ageing degradation (AD) in NPP-SSCs occurs, and how it may be effectively detected, avoided, eliminated or mitigated to facilitate safe, reliable and profitable current operation, and even long-term operation (LTO) (i.e. operation of NPPs in excess of their original design lives). Consequently, this book will furnish all those involved in any sector of commercial NPP operation and regulation, with a deeper understanding of the many facets and issues concerning nuclear power technology as a whole. The topics and issues addressed in this book are experience and sciencebased sources of modern information concerning the commercial nuclear power industry, and its regulation. Safety, operational reliability and profitability of NPPs will all benefit when such information is robustly implemented as technically perfected (and regulatory approved) methodologies and remedies for SSC-AD. It is expected that all NPPs, irrespective of their design or age, that are operated by a qualified workforce, who conscientiously apply all the tenets of safety culture and diligently ensure that standard operational practices (OPs), ASPs, AM and PLiM programmes are followed, will be safe, reliable and profitable to operate at every stage of their lives, including LTO, and until, at the very end of their chronological lives, decommissioning, dismantling and disposal eventually become necessary.

1.2

Age as a relative term

The design life of a NPP has generally little in common with its real-life operational safety status and fitness-for-service, and thus the amount of years it can operate until the true end of operational life (EOL) is reached. Barring political or public acceptance issues, for example, the NPP’s EOL will be reached when safety requirements can no longer be satisfied or, alternatively, despite having a good safety status, it is simply no longer profitable to continue operation. Specifically, good standard plant OPs, ASPs, AM and

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PLiM programmes make it possible to delay significantly this point in time, since both safety and profitability of NPPs are maintained and optimized. Again, NPP personnel safety culture attitudes will be significant in assuring the degree of effectiveness of the above-mentioned programmes. Depending on various aspects, NPPs may possess several types of ‘lifetime’, for example: ∑ ∑ ∑ ∑ ∑ ∑ ∑

operational lifetime (when power is being produced safely and reliably and can be sold at a profit); political/legal lifetime (operation ends when there is no more public acceptance or when legally binding international agreements to shut down particular NPPs have to be respected); business/economic lifetime (operation ends when the overall NPP running costs can no longer be covered by sales of electricity); technical lifetime (operation ends when SSC refurbishment and repairs are no longer technically feasible, or possible); conceptual lifetime (operation ends when the NPP can no longer be kept at the state-of-the-art, science and technology, since it has far too many outdated and obsolete SSCs); safety lifetime (operation ends when statutory and regulatory requirements for continuing operation (fulfilment of licensing conditions) are too extensive, expensive or not possible to implement practically); the chronological lifetime of a NPP may be taken as the time-span between site selection, constructing and operating the NPP, ceasing operation, removing the fuel and finally decommissioning and dismantling it, and returning the site back to ‘greenfield’ status.

Irrespective of their age, NPPs that have benefitted from continual improvement in management and operational practices (MOPs), maintenance, monitoring technologies, inspections, repairs, refurbishing or replacement of SSCs, and have cost-effectively upgraded plant control and safety systems in full accordance with regulatory requirements, are in good overall condition to continue operation and go eventually into the LTO phase.

1.3

The importance of nuclear energy

As previously noted, nuclear energy is already an important contributor (about 16%) to the total global electricity requirement, and as demand for energy increases, the nuclear-generated share of the electrical power market is expected to rise correspondingly. Potentially, nuclear-generated power could provide up to 1.7 terawatts (TWe) of electricity to the world by 2050, but this will depend, among other things, on the number and capacity rating thereof of new generation NPPs coming on-line [1]. However, if the nuclear power industry is to maintain its current position, and wishes to increase its

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market share in the future, it must continue to be accepted by the public. A crucial aspect of this is that NPPs must remain a safe and reliable source of energy. Acceptance of nuclear-generated electrical power will also depend considerably on its perceived and real impact on the quality of life and the environment. It is therefore appropriate to introduce the concepts of ‘greenhouse gases’ and the ‘carbon footprint’ (CF) already here, since human activities are now frequently classed in terms of a CF, which is a measure of the amount of carbon dioxide (CO2) (a greenhouse gas) produced by the combustion of fossil fuels. The CF is often expressed as tonnes of CO2 or tonnes of carbon emitted, usually on an annual basis. However, it must be stated here that any activity requiring or generating energy (e.g. power plants (hydro, fossil-fuelled, solar, wind, nuclear), mining (extracting coal or uranium, pumping oil), manufacturing (refining metals, enriching uranium for nuclear fuel, production of solar panels or wind turbines, making cement), transport (petrol and diesel combustion), electrically driven motors (batteries charged via coal-fired electricity stations or solar panels)) will leave varying sizes of CF. Thus, although NPPs generate negligible direct carbon emissions, they nevertheless still have a finite CF. The method used to obtain a CF is the ‘life cycle assessment’ (LCA), which is accredited by the International Organization for Standardization (ISO) 14000 standards. Calculating exact values for CFs is a fairly complex task, since basically every process and sub-process involved in the item under consideration must be assessed for its own CF and then added to get the CF total. However, taking all CF aspects (mining, fuel enrichment, manufacturing, decommissioning and disposal, etc.) into consideration, nuclear power generation may still be classed as having a very low CF, and thus nuclear energy remains an extremely low source of environmentally harmful greenhouse gases. For example, even if lower grade uranium-bearing ores (e.g. 0.03% content of fissionable 235U) will have to be used in future, necessitating more energy to extract and refine the fuel, this would only raise the current total CF emissions of the United Kingdom’s NPPs from 5 to 6.8 grams of CO2 equivalent per kilowatt-hour (CO2e/ kWh). This increase, however, would still keep nuclear power technology at a level comparable with other low-carbon power technologies, and will be anyway well below the CFs of fossil-fuelled power technologies [2]. (Note: The CO2e is the carbon dioxide equivalent. Each of the greenhouse gases addressed by the Kyoto Protocol (see later) can be identified in terms of its climate change impact relative to that of carbon dioxide.) The commonly used unit for emission reductions is one tonne of carbon dioxide equivalent. Combined power plants using gas turbines and hydro-power produce 420, solar-photo-voltaic 62 and coal-fired ones about 900 (all in approximate gram-equivalent CO2e/kWh). It has been reported that a pressurized water reactor (PWR) NPP in Switzerland had a value of 3.04 CO2e/kWh, which is on a par with hydro-produced only electricity [3]. This value even included

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mining the uranium fuel and final disposal of the entire plant. Although electricity production in NPPs is esentially CO2 free, over the entire plant lifetimes about 6–8 CO2e/kWh is nevertheless expected to be generated through various processes. This is still very low compared to other energysupply systems, as exemplified above. Not only is nuclear-generated power environmentally benign, it is cheap: at a representative price of 5 US cent/ kWh (3.4 Euro-cent//kWh at the average exchange rate in October 2009), a cost-analysis gives roughly: 1.97 US cent (1.3 Euro-cent) for capital costs, 1.85 US cent (1.26 Euro-cent) for operating costs, 1 US cent (0.7 Eurocent) for decommissioning/disposal and 0.38 US cent (0.26 Euro-cent) for fuel. (It can be noted that if all fossil-fuelled power plants were to be replaced by NPPs (to get the same amount of energy), it would mean that the world’s NPP fleet would then have to number about 4316 by the year 2050 [4]. This is roughly ten times more than the number of NPPs operating in 2010.) A detailed analysis of CO2 issues is beyond the scope of this book, but it is worth mentioning that carbon capture and storage (CCS) and associated costs are currently (2009) running between 30 and 35 Euros per tonne, and this cost is likely to increase. For comparison, Norway currently levies a duty of about 44 Euros per tonne of CO2 released in the process of extracting oil from the North Sea. The European Energy Exchange, located in Leipzig, Germany, currently quotes prices between 20 and 30 Euros per tonne of CO2 [5]. In other words, just like nuclear waste, fossil-fuel-generated CO2 waste also has a price and sequestering (storage) of ‘carbon’ underground requires further environmental impact, site selection and viability assessments. From this, it can be seen that continued and increased use of nuclear power remains justifiable in terms of its low CF and GGE emissions profile, and it is in a favourable position to contribute significantly to both current and future global requirements for energy.

1.4

Learning from experience to continually improve safety in nuclear power plants (NPPs)

Creating and maintaining safety in operation of NPPs arises through many factors including good design, appropriate materials selection for SSCs, highest manufacturing standards and approved methods, quality assurance, operational practices, maintenance, inspection and human resource training and actions. However, continually improving safety at all levels lies in the necessity to learn constantly from our own, or others, experience. As mentioned in the Foreword, ‘Experience is the best teacher’, but when dealing with human and materials-related issues in NPPs, particularly those that have impacted safety, it is evident that ‘Detrimentum magister nos’ (‘Damage teaches us’). As with any other industry, it is necessary to put all lessons learned, including

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practical experience combined with analysis and inspections, robustly into practice to avoid future problems in the operation of NPPs and the SSCs in them. This facilitates continual improvement of safety, as well as reliable operation. Such actions taken in a NPP can not only serve the regulatory and corporate goals of safety and profitability respectively, but also may pave the way for continued operation even after the NPP’s original design life has been reached (LTO). The use of nuclear energy is subject to regulatory control, and licences to operate (or permission granted for continued operation) will be issued only when the safety case is clearly demonstrated. Thus, proof of safety remains a mandatory regulatory requirement for the current and LTO of existing NPPs, as well as for any future new-build NPPs. Lessons learned and experience gained with all types and designs of NPP over the last 55 years of using commercial nuclear power now serve to form a solid base of knowledge concerned with regulation, design, choice of materials, inspection, monitoring techniques and best operational practices, including nurturing a safety-oriented attitude in the workforce, namely the implementation of a good safety culture. Although the main scope and focus of this book is on understanding and mitigating SSC-AD mechanisms or issues which have arisen in the past and that have impacted safety, reliability and profitability of NPPs, a further goal is to provide guidance and information derived from research and operational experience, based on sound science and proven methods, to ensure that NPPs may continue to function in full accordance with their design specifications and to thus produce electricity as safely and as reliably as possible at a competitive price. Implementation of best possible practices and methods also has the potential to facilitate LTO operation of NPPs, since safety margins may be kept at sufficient and effective levels, and SSCs remain reliable. The way in which best practices and methods can be incorporated into the overall operation of NPPs is discussed in this book by drawing on experience gained and introducing the concepts, features and goals of ASPs, AM and PLiM programmes. Accordingly, this book also serves the broader goals of knowledge management, which is necessary for retaining both managerial and technical competence in the nuclear power industry, now and in the future.

1.5

Global situation of the status of installed nuclear power in 2010

The world’s fleet of NPPs currently has an average operational age of around 23 years. In June 2010 there were 438 NPPs operating in 30 countries, with an installed electric net capacity of about 370 GWe, and another 59 NPPs, with a planned capacity of 40 GWe, were in various stages of planning and construction, for example, 6 new NPPs were being built in mainland China

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and India. The currently installed capacity is expected to increase to at least 530 GWe by the year 2025 as new NPPs connect to the grid. The amount of installed nuclear power capacity, relative to other energy sources, varies between countries, as shown, as an example, in the following abbreviated list, which details the current situation (2010), in per cent installed nuclear power capacity relative to other sources (approximate values): ∑ ∑ ∑ ∑ ∑ ∑ ∑ ∑ ∑ ∑ ∑ ∑

Belgium: 7 NPPs, 52%; Canada: 18 NPPs and 2 NPPs in refurbishment, 15%; Peoples Republic of China: 11 NPPs, <3% (>5% by 2020 is planned); France: 58 NPPs, 75%; Hungary: 4 NPPs, 20%; India: 17 NPPs; 3% (with the goal for 25% in 2050); Japan: 55 NPPs, 30%; Russian Federation: 32 NPPs, 16%; Sweden: 10 NPPs, 46%; Switzerland: 5 NPPs, 40%; United Kingdom: 19 NPPs, 19%; United States of America: 104 NPPs, 20%;

In June 2010 there were 195 NPPs in operation in continental Europe (with an installed electric net capacity of about 170 000 megawatt (MWe)) and 19 units were under construction with a projected net capacity of about 13 000 MWe; the new NPP constructions being in Bulgaria (2), Finland (1), France (1), Russian Federation (11), Slovak Republic (2) and Ukraine (2) [6]. Other countries within Europe currently have no commercial NPPs (e.g. Denmark, Greece). Countries outside Europe (e.g. Egypt) are keeping the nuclear power option open, whilst other countries (e.g. Iran) have already embarked on nuclear power schemes. Some countries are increasing their nuclear energy capacity (e.g. China, India, Russia, Finland and France), whilst, in contrast, Germany is planning to gradually phase out nuclear power. In February 2009, Italy and France agreed to the development, construction and initial set-up of four new generation European Pressurized Reactor (EPR) power plants in Italy, with the first one projected to start operation in 2020. Italy’s aim is to eventually produce 25% of the country’s electricity from nuclear energy. (Aspects of new generation NPPs, such as the EPR are discussed in Chapter 2 of this book.) Some countries (e.g. United Kingdom, Switzerland) are planning and assessing options to install new NPPs to replace the capacity of NPPs that could eventually be phased out. Despite the aforementioned international differences, political issues and goals, the overall global trend for nuclear power is seen as one of expansion. In June 2010, 149 new NPPs were in planning.

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1.6

Understanding and mitigating ageing in nuclear power plants

The importance of keeping nuclear power plants (npps) operating safely and reliably

A vital contribution to the overall level of nuclear generated power (irrespective of any new plants coming on-line in future, and the extra capacity presently being generated by NPP power uprates (PUs)), lies in assuring that the existing NPPs can continue to operate. This is, of course, with the requirement that they can do so safely and profitably, despite their chronological age or original design life. As will be shown later in this book, thanks to improved materials, maintenance, MOPs, ASPs, AM and PLiM programmes and ageing degradation mitigation, these presently operating NPPs are still capable of fulfilling safety, technical and operational requirements. The question of whether a NPP may go to LTO is addressed through licensing procedures which require a priori proof of safety. It may be stated here that, despite different approaches in Europe and the USA, for example, the principles and goals are basically the same: first prove the safety case to the relevant regulatory authority and demonstrate that the overall plant practices are effective (e.g. maintenance, inspections, monitoring) and, if the business case is favourable (i.e. total market cost per kilowatt-hour (kWh) of nuclear generated electricity is competitive (or cheaper) compared to other sources, and profit margins remain favourable), and public acceptance remains high, continue operation.

1.7

Political and climate change issues and disposal of radioactive waste

Since the first ‘oil shock’ in the early 1970s it became clear that the world had become heavily dependent on this fossil-based fuel, and shortages in supply would quickly disrupt the world’s economies. Compounding this, oil prices have increased many-fold (a price is deliberately omitted here, due to the extremely volatile nature of the crude oil market) between 1970 and 2009 [7], and wildly fluctuating prices and shortages of supply have caused sharp inflation spikes and even civil protest. At the time of the first oil shock, there was generally a lack of global awareness concerning the possible harmful side-effects of burning fossil-based fuel and the production of GGEs with the attendant potential impact on the environment through pollution and global warming. Things have changed over the last 40 years, and environmental awareness and issues have now attracted increased public, political, energy policy-makers’ and media attention. The Kyoto Protocol [8] made provisions and pledges to reduce the signatory nations’ GGEs. As of February 2009, 183 states had signed and ratified the Protocol. Of these, 36 developed countries (plus the European Union, as a party in its own right) are required to reduce GGEs to specified levels. Some 137 developing

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countries have ratified the protocol, but they currently have no obligation beyond monitoring and reporting GGEs. As an example to focus attention on CO2 issues, it is noted that the concentration of CO2 was at a level of 372 parts per million in April 2005. This was higher than that calculated to have been present in the Earth’s atmosphere during the past 42 000 years. However, in September 2008 it was reported that scientists had thought that the general economic downturn would have slowed energy use, but instead, CO2 output actually jumped 3% from 2006 to 2007. This amount exceeds the most pessimistic outlook for emissions from burning coal and oil and related activities, as projected by scientists in 2007 [9]. It is worth noting that Great Britain has now pledged to cut CO2 emissions by at least 50% by 2050 and the ambitious new generation NPP construction scheme, currently now under plan, is seen as critical to achieving this target [10].

1.8

Energy resources: a comparison

The world’s demand for energy is increasing in pace with the growing human population (from approximately 1.5 billion in 1900, to about 6.7 billion in May 2009, and projected to reach about 9 billion by 2050 [11]), and with the emergence of more export-based economies in developing and industrializing nations. This means that the fossil-fuel based energy carriers (and these are finite resources) will eventually have to be replaced with alternatives that have lower environmental impact, are cheap and safe to use and, ideally, are ‘infinite resources’. Obvious such energy resources lie in harnessing hydro-, wind-, solar- (the sun shines on the earth about 35 000 times more power than the population uses daily, but presently only about 0.2% of the world’s energy comes from solar power) or direct tidal power. However, most of the world’s ideal hydropower sites have been used up, and environmental concerns and impact issues, such as flooding of marshlands and valleys and covering large land and sea areas with solar panels or wind turbines respectively may anyway preclude expansion, or restrict use, of these alternative sources in some cases. Costs of solar, wind and tidal energy are currently not competitive with nuclear energy (e.g. solar electricity is presently (2010) selling at a price six times more expensive than nuclear-generated electricity). So how does nuclear power really measure up? Uranium, a main source of nuclear fuel, is clearly also a finite quantity, but relatively little of the fissionable isotope 235U is actually needed for fuel and the current conservative estimate is that the resources are more than sufficient for at least the next 150 years, even allowing for increased demand through currently operating NPPs, including those in the LTO stage, and new NPPs projected to be operational in the next 5–40 years from now. Currently known uranium reserves, allowing extraction at a nominal benchmark price of US$130 per kilogramme of

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natural uranium metal (i.e. fuel enrichment price not included), or less, are estimated to be about 5.5 million tonnes. Consumption is about 66 500 tonnes per year. In total, about 35 million tonnes of uranium are probably available [12]. With advanced reactor and new fuel cycle technologies (breeder reactor technology), plus the eventual conversion of plutonium to fuel and reprocessing of spent fuel elements, the availability of nuclear fuel is likely to be assured for at least 1000 years [13]. Use of thorium as a breeder for fuel (discussed later) will potentially double the known reserves of fuel for the presently used fission-based reactor technology [12]. In the future, it is expected that other totally new ‘fossil-free’ or ‘virtually infinite’ energy sources will be further developed (e.g. fusion-based reactor technology, hydrogen-based fuels), and improved ways of energy utilization and distribution efficiency will anyway further conserve all other energy resources. Such developments will not take place in the near future, but prototype fast breeder reactors, for example, are currently being developed or under construction. Recognizing this situation, the only logical option left is to use low or zero GGE energy sources for as long as possible, whilst developing alternatives to fossil-based energy carriers in parallel. Although a NPP produces essentially no GGEs in operation, it does, however, create its own special waste products, but this radioactive waste (‘radwaste’) is not jettisoned into the environment; it is collected, sorted, conditioned and stored according to nationally approved practices and legislation. Aspects concerned with radwaste will be examined later, but radwaste treatment technology is already advanced and deemed capable of preventing environmental damage by keeping harmful substances out of the biosphere. It has to be remembered that radwaste must be handled, conditioned and disposed of so as to span very long geological, political and social time frames. Furthermore, by closing the nuclear fuel cycle, with the separation of minor actinides and long-lived radio-isotopes/fission products, the life span for unacceptable radiotoxicity and monitoring can be brought down to less than 1000 years.

1.9

Further ageing aspects in nuclear power plants (NPPs)

Nuclear power plants, just like any other industrial plants or transport systems (e.g. trains, aircraft, ships), are made of SSCs that undergo ‘ageing’, and this term may even be extended to include the personnel who operate them. Ageing effects may be seen as some finite rate of continuous, time-dependent mechanical, physical or chemical degradation of SSCs, level of personnel training, loss of knowledge (due to retirement of experienced personnel and inadequate succession planning), inadequacies in documentation (e.g. omission to follow exactly any modifications that have occurred in the NPP) and other

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aspects associated with mindsets and complacency. The effects of ageing can thus be seen to impact humans, plant procedures, practices and management, materials, computer software, hardware, documentation and machinery alike. Ageing degradation may occur at almost imperceptible or relatively rapid rates, depending on the specific subject area under consideration and the effectiveness of MOPs, ASPs, AM and PLiM programmes in place. A key to safe operation of NPPs is to be aware of all the possible types of ageing that could take place, and to understand their root causes, and their potential impact on safety and reliability. Knowledge of the ways of how to eliminate or mitigate ageing degradation effects, as presented in this book, will then effectively ‘close the loop’, to create optimized NPP operation. While safe plant operation is vital for public acceptance, reliable operation is essential for the supply of competitively priced energy and thus profitability of the plant. These aspects will also be dealt with in this book, but it is important to mention here that the nuclear power industry is quite aware of SSC-AD and other ageing issues, and continues to implement strategies to address them. One important tool used to keep NPPs in a safe condition is to apply standard (routine) OPs, inspections, repairs and replacements, supported further by ASPs, AM and PLiM programmes. Their role and benefits will become increasingly apparent in the following chapters, but it is sufficient to mention here that such programmes not only protect the utilities’ business goals and investments, but also maintain safety at the highest level for a reasonable cost–benefit relationship and also pave the way for continued NPP operation, even after the original design life has been reached. This represents a significant financial incentive for utilities to invest resources in reliable and safe operation now to assure the future (LTO) of their plants.

1.10

Historical evolution of nuclear power and some materials aspects

On 20 December 1951, the Experimental Breeder Reactor EBR-1 in Arco, Idaho, USA, succeeded in supplying enough electricity to illuminate four low-power light bulbs. (The EBR-1’s task was, however, only to demonstrate the concept of a breeder reactor.) The Russian APS-1, 5 MWe reactor was the first to supply electricity to the national grid, on 26 June 1954. Not long after this, the nuclear power era gained considerable momentum when the world’s first commercial-scale NPP, ‘Calder Hall 1’, situated at Sellafield, England, was connected to the national electricity grid on 27 August, 1956. The Sellafield site eventually featured four Magnox-type reactors, each with a rated power of 50 MWe. (Apart from supplying power, the Magnox reactors also served to produce plutonium for military purposes.) The Magnox reactors were natural uranium metal fuelled, graphite moderated, and carbon dioxide gas cooled. They featured magnesium–aluminium

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(non-oxidizing – hence ‘Magnox’) alloy fuel cladding. The fuel cladding alloy had a low neutron capture cross-section, but limited the operating temperature to about 390 °C and thus also the thermal efficiency. The heat exchanger was outside the containment. Early Magnox reactors featured a steel containment, whilst later ones had a reinforced concrete one. The United Kingdom’s Magnox reactors were designed for 20 years of operation but the oldest one at Calder Hall actually achieved nearly 47 years of service before being shut down on 31 March 2003. Decommissioning of the Calder Hall site, to eventually return it to ‘greenfield’ status, is projected to take about 100 years to complete. By comparison, most NPPs operating today usually have a full power operation original design life of 30–40 years (with a relicensing or open-license situation for the option of operation of 60 years, and possibly more), produce between 220 and 1600 MWe, and, depending on the design or type, use either, for example, natural uranium, plutonium or uranium–plutonium mixed oxides or carbides as fuel, have zirconium alloy (e.g.‘Zircaloy-4’ (Zry-4) with 98.23 wt% zirconium, 1.45% tin, 0.21% iron, 0.1% chromium, and 0.01% hafnium (impurity hafnium very low, since it has a high neutron capture cross-section)) fuel cladding materials, and use either heavy water, light water, helium gas, carbon dioxide gas or molten sodium metal for coolant. The pressure vessels in light water cooled reactors (i.e. pressurized and boiling water reactors – PWR and BWR, respectively), for example, are made from tough, low alloy ferritic steel (cladded inside with austenitic stainless steel for corrosion protection). It can be seen that nuclear power systems technology, fuel and materials of SSCs have all evolved considerably over the last 55 years.

1.11

Overview of two important materials issues in older design nuclear power plants (npps)

Each type and design of NPP has its own special characteristics, and as operational experience was gained over the years, solutions to specific problems associated with SSC ageing degradation mechanisms had to be found. Design concepts and materials which had functioned well for other industries (e.g. in fossil-fuelled power plants, oil refining and chemical industries) were natural choices for the first commercial NPPs, but with time, it became clear that the nuclear environment presented other, and special, challenges to SSCs and the materials involved. For example, operational conditions and parameters such as temperature, pressure, neutron irradiation, coolant water chemistry (including radiolysis thereof) and other factors often combined to cause degradation in materials previously thought to be resistant to ageing degradation. Two examples of unexpected material degradation are given below.

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1.11.1 Stress corrosion cracking Despite a good service record in other industries, stress corrosion cracking (SCC) occurred in the well-known ‘Alloy 600’ (a nickel-base alloy approximately 72% Ni, 14% Cr and 10% Fe), used extensively for reactor pressure vessel (RPV) closure head penetrations cladding and steam generator (SG) tubing in pressurized water reactors (PWRs). Technical and engineering solutions were found to repair, plug, sleeve or surface-treat affected SG tubes, but when the so-called ‘plugging-rate value’ was reached (typically 10–20% of the total number of tubes in the SGs) it was no longer possible to operate at full power, having regard for emergency cooling requirements. Thus the balance between safety (emergency cooling capacity) and economic operation (power had to be reduced) eventually lead to the replacement of affected SGs. The action of ‘overplugging’ SGs is an alternative temporary solution to SG replacement, but necessitates a significant further reduction of plant power and this translates to the loss of even more money for each operational cycle: in other words, a natural limit is reached where it becomes economically attractive (i.e. cheaper in the long term) to replace the SGs. (Note: SCC degradation is dealt with extensively in Chapter 9 of this book.) New SGs usually feature tubing made from Alloy 690 TT (nickel-base alloy approximately 60% Ni, 30% Cr and 10% Fe), which is thermally treated (TT) to create a favourable microstructure and optimum mechanical properties, which has a much improved resistance to SCC in PWR coolant. It is noted here that, in general, nickel–chromium–iron alloys, such as Alloy 600, are quite resistant to corrosion, but if they contain high internal tensile stress (fabrication aspects), are not heat treated in an optimum way (microstructural effects) and are subjected to pressurized borated water (i.e. the coolant used in PWRs) at around 300 °C, primary water stress corrosion cracking (PWSCC) may occur, leading to costly repairs or eventual replacement of SGs and RPV closure heads. In some PWRs, the leakage of borated coolant water through PWSCC-affected head penetrations (for control rods and instrumentation) onto the external RPV closure head caused problems. In the most severe case discovered, boric acid crystals left on the external surfaces of the ferritic steel RPV head had caused extensive corrosion. The large area of corrosion (also known as ‘wastage’) had reached the stainless steel cladding of the RPV inner surface [14, 15]. It is interesting to note here that in the case mentioned, there were also managerial (human factor) inadequacies and ambitious power production goals that had also combined to aggravate the situation [16]. The NPP involved was the 873 MWe PWR ‘Davis-Besse’, situated at Oak Harbour, OH, USA. The forced outage lasted about two years and cost around US$600 million. The NPP underwent a successful recovery process, including purchase of a new RPV head, and

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also implemented comprehensive operational, monitoring and management strategies to avoid the problem in future. Replacement of RPV closure heads, usually featuring Alloy 690 TT penetration claddings, is an on-going task in many PWR-NPPs. Appropriate monitoring systems are also being implemented to detect if leakage should, nevertheless, occur.

1.11.2 Neutron irradiation embrittlement enhancement caused by impurities introduced during welding A further example of evolution of materials technology in NPPs can be found in the development of improved ferritic RPV alloys and welds. In the 1960s, NPPs were constructed using the current state-of-the-art welding technologies. However, there was no awareness about the effect copper (introduced into the weld material as an impurity via copper-coated welding rods) would have in increasing the RPV weld metal’s propensity to fast neutron embrittlement damage. The main embrittling mechanism, basically ‘age-hardening’ as matrix-coherent minute copper-rich particles come out of solid solution and block dislocation movement, coupled with atomic-scale point-defect damage, was researched extensively and now specifications for new RPVs limit the amount of impurities (e.g. copper restricted to <0.08 wt%) in materials in the beltline area of RPVs (opposite the core and thus highly irradiated with neutrons). (Various works are cited here to provide information on how research results may be used to understand neutron irradiation embrittlement and to provide practical solutions to mitigate it [17–22].) For some NPPs it became necessary to restore fracture toughness by thermal annealing the RPV beltline area (see Chapters 11 and 12 of this book).

1.12

Conclusions

It can be stated that developed countries depend on the reliable supply of safe and cheap energy to maintain vital industrial and socially based infrastructures, sustain economic growth and to maintain, or increase, living standards. Developing countries are striving to improve their standard of living and to become more commercially competitive, and they will also require energy if they are to be successful in these endeavours. Nuclear-generated power thus has a crucial role to play to provide cost-competitive energy, at a low CF and GGE penalty, to both developed and developing countries. Strategies to maintain, and even increase, the role of nuclear power and to assure sufficient energy supplies in the future must be instigated today, since new NPP construction times, depending on design, can take up to five or more years. It is also absolutely necessary to promote the training and

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qualification of personnel in nuclear power technology, and its regulation, so that sufficient human resources will be available to the industry in the future. With a cumulative operating experience exceeding 14  000 reactoryears (in 2010) and having supplied more than 64 billion kWh of electricity over this integrated time, the nuclear power industry is clearly a mature one and capable of supplying an important part of the current and future world’s clean energy requirements. Ongoing research and operational experiences continue to contribute to all aspects of maintaining safety and reliability in NPPs. Additionally, profitable current and LTO operation is expected to be a characteristic feature of NPPs having optimized MOPs, ASPs, AM and PLiM programmes in place.

1.13

Sources of further information

Nuclear Power Plant Life Management and Longer Term Operation. OECD/ NEA, NEA No. 6105, 2006. PLEX: Information Around Plant Lifetime Extension, Austrian Ecological Institute (Oesterreichisches Ökologie Institut, Vienna, 2006–2007), see: http://www.ecology.at/files/berichte/E22.547-1.pdf.

1.14

References

1. World Energy Technology Outlook – 2050, European Commission: EUR 22038, 2006. 2. ‘Carbon Footprint of Electricity Generation’, Parliamentary Office of Science and Technology Postnote, Number 268, October 2006. 3. KKW Beznau: Kernkraftwerk mit ‘klimafreundlichstem Strom’, Aargauer Zeitung, Switzerland, Friday, 31 October 2008. 4. Nielsen, R., The Little Green Handbook, Picador, New York, 2006. 5. European Energy Exchange (EEX), ‘Carbon dioxide prices’, Leipzig, Germany (http:// www.eex.com/en/EEX/Cooperation%20EEX%20and%20Eurex/Background). 6. European Nuclear Society – ENS, ‘Nuclear Power Plants in Europe: Status in January 2009’ (http://www.euronuclear.org/info/maps.htm). 7. Energy Information Administration, Official Energy Statistics from the US Government (http://www.eia.doe.gov/neic/aboutEIA/aboutus.html). 8. Kyoto Protocol to the International Framework Convention on Climate Change. Adopted 11 December 1997 by 3rd Conference of the Parties, Kyoto, Japan, in force, 16 February 2005. 9. ‘Carbon dioxide output jumps to record level in 2007’, report from CNN-News AP, 25 September 2008. 10. Winnett, R., ‘Eight new nuclear power stations planned for England’, Telegraph, 13 July 2008. 11. International Atomic Energy Agency (IAEA) Bulletin 49–2, March 2008. 12. Schiermeir, Q., Tollefson, J., Scully, T., Wizte, A. and Morton, O., ‘Energy alternatives: electricity without carbon’, Nature, Vol. 454, Issue 7206, 14 August 2008. 13. Uranium 2007: Resources, Production and Demand. A Joint Report by the OECD

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14. 15. 16. 17.

18.

19.

20.

21.

22.

Understanding and mitigating ageing in nuclear power plants Nuclear Energy Agency and the International Atomic Energy Agency, OECD, Paris, 2008. US Nuclear Safety Commission, Bulletin 2002-01 ‘Reactor pressure vessel head degradation and coolant pressure boundary integrity’, January, 2002. NRC Information Notice 2002-11, ‘Recent experience with degradation of reactor pressure vessel head, 12 March 2002. ‘Material degradation and related managerial issues at nuclear power plants’, IAEAProceedings Series, STI/PUB/1260, 15–18 February, 2005, Vienna, Austria. Nanstad, R.K., Tipping, Ph., Kalkhof, R.D. and Sokolov, M., ‘Irradiation and postannealing re-irradiation effects on fracture toughness of RPV steel heat JRQ’, in Grossbeck, M.L. (ed.), ASTM Special Technical Publication on Effects of Radiation on Materials, 21st. International Symposium, ASTM STP 1447, 2003. Ghazi-Wakili, K., Zimmermann, U., Brunner, J., Tipping, Ph., Waeber, W. and Heinrich, F., ‘Positron annhilation studies on neutron irradiated pressure vessel steels’, Phys. Stat. Sol. (a), Vol. 102, pp. 153–163, 1987. Ghazi-Wakili, K., Tipping, Ph., Zimmerman, U. and Waeber, W., ‘Investigation of neutron irradiated Fe-0.8wt% Cu alloys by means of positron annhilation and microhardness measurements’, Zeitschrift für Physik B Condensed Matter, Vol. 79, pp. 39–45 (1990). Tipping, Ph., Waeber, W. and Mercier, O., ‘A study of the mechanical property changes of irradiation embrittled pressure vessel steels and their response to annealing treatments’, Int. Journal of Pressure Vessels and Piping, Vol. 46, pp. 133–148, 1991. Tipping, Ph. and Cripps, R., ‘Annealing for plant-life management: hardness, tensile and Charpy toughness properties of irradiated, annealed and re-irradiated mock-up low alloy pressure vessel steel’, Int. Journal of Pressure Vessels and Piping, Vol. 60, pp. 217–222, 1994. Tipping, Ph., ‘Isothermal annealing behaviour and Vickers microhardness of irradiated mock-up pressure vessel weld material’, PROMETEY/IAEA Third International Conference: Material Science Problems in NPP Equipment Production and Operation, Vol. 1, pp. 71–82, 17–22 June 1994, St. Petersberg and Moscow, Russia.

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Key elements and principles of nuclear power plant life management (PLiM) for current and long-term operation

P h. G. T i pp i n g, Nuclear Energy and Materials Consultant, Switzerland

Abstract: Key elements and principles of nuclear power plant (NPP) life management for current and long-term operation (LTO) are presented. Concepts and methodologies of NPP ageing surveillance programmes (ASPs), ageing management (AM) and plant life management (PLiM) are discussed. Ageing degradation (AD) terminology concerning NPP systems, structures and components (SSCs) is provided. Safety classification of SSCs, precursors for scoping and applying optimum OPs, ASPs, AM and PLiM programmes, design requirements of NPPs, knowledge management, human factors, safety culture, radiological protection, SSC-AD mechanisms and AD mitigation, obsolescence in equipment, regulatory aspects, radioactive waste disposal and radiological protection are discussed. Concepts for future NPP designs and PLiM are introduced. It is shown that by implementing the stateof-the-art, science and technology into the existing standard plant operational practices (OPs), which include ASPs, and AM and PLiM programmes, SSCAD can be managed to keep the option open for NPPs to reach, and then exceed, their original design lives in a safe and profitable way. Key words: nuclear power plants, ageing and plant life management, long-term operation, design, structures, systems and components, reverse engineering, obsolescence, safety class, safety culture, knowledge management, regulatory aspects, radiological protection, power uprates, radwaste, future generation reactors.

2.1

Introduction

The following deals with NPPs using nuclear fission technology. The basic principles and features of standard OPs, ASPs, AM and PLiM programmes used to achieve NPP design life safely and economically, and then to potentially facilitate LTO, are generally valid no matter what design of NPP is under consideration. However, plant-specific aspects and features, as well as operational experiences reported from similarly designed and operated NPPs elsewhere in the world, will always necessitate objective assessments, adjustments and updates of plant life ageing and management programmes to ensure that the scope, focus, applicability and depth thereof remains at the state-of-the-art, science and technology. 19 © Woodhead Publishing Limited, 2010

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Nuclear power plants, irrespective of design, function in a conceptually simple way: they extract heat arising from the fission processes occurring in the fuel via a suitable medium (water, gas, liquid metal) and transfer this to heat exchangers (steam generators (SGs)), or directly as steam or very high temperature gas (e.g. helium), to drive turbines connected to generators to make electricity. Depending on the type and design, a NPP may be characterized by primary and secondary circuits and also a reinforced concrete and steel containment to ensure that any leakages arising from the primary circuit cannot cause radioactive contamination to the environment. A NPP is further characterized by having components that can be classed as passive and non-replaceable (e.g. reactor pressure vessel (RPV), containment) and thus are deemed ‘life determining’ since, conceptually and practically, it is virtually impossible to replace them or it is prohibitively expensive to do so. Nuclear power plants also have active but replaceable components (e.g. SGs, RPV closure heads, piping, pressurizers, core shrouds and pumps), which are conceptually and practically replaceable in nature, albeit at a cost. Other categories of components are those that are relatively easy to replace on a routine basis, or as the necessity arises (e.g. control rods/drives, pump impellers, seals, bolts, switches, cables, valves and fuses). Irrespective of the classification ‘plant life-determining’ or ‘replaceable’, safety relevant components obviously must have first priority in terms of inspection, monitoring, repair, replacement and assessment as far as ‘fitness-for-service’ is concerned. Furthermore, it may be very expensive to replace a major component, and the need to replace it will either be dictated directly by regulatory requirements or by the remaining projected plant operating life in order to assess the chance of amortizing the investment of the replacement SSC in question. Thus the business case analysis will determine the point in time where profitable operation of the NPP is no longer possible. The real cost of nuclear-generated electricity compared to other sources is a nominal value, since environmental costs must also be brought into the total equation if an objective comparison is to be made. Electricity price analysts usually use conservative assumptions about NPP construction costs, the length of construction time, delay penalties and then make projections on how efficiently (e.g. availability and operational and maintenance costs, etc.) the NPP would operate over its design life. Prolonged shutdowns, regulatory restrictions and unexpected operational factors will obviously affect such estimates. However, experience to date (mid-2010) shows that most NPPs can usually generate electricity at prices between 3 and 4 Euro cents (about 4–6 US cents) per kilowatt hour (kWh), but even if fossil power were marginally cheaper than nuclear-generated power, it still carries with it a far greater potential penalty concerned with greenhouse gas emissions (GGEs) and the associated environmental issues thereof. It is a generally recognized fact that humanity must drastically ‘decarbonize’ its electrical

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energy production if the currently perceived rate of global warming is to be reduced. It is therefore necessary to invest resources in NPPs, which are essentially GGE-free, and to keep safety-relevant SSCs, as well as passive, non-replaceable SSCs, in good functional condition. For example, items such as SGs in pressurized water reactors (PWRs) and core shrouds in boiling water reactors (BWRs) have been replaced or repaired due to stress corrosion cracking (SCC) issues, but it was expensive and difficult to do so. Core shrouds in BWRs have also been fitted with tie-rods to maintain mechanical integrity, despite the presence of SCC in the austenitic stainless steel material used. Furthermore, changes to water chemistry (e.g. hydrogen injection) have further been implemented to improve the BWR working environment conditions. However, it was worth the cost and effort for such plants affected, since they could then continue reliable operation at full power and with good prospects to amortize the investment. Furthermore, apart from the problems associated with SCC, which were not expected, the earlier designs of most NPPs did not take into account the scenario to allow for a future possible need to remove and replace massive components such as SGs, core shrouds and RPV closure heads. Replacement of SGs, for example, has generally necessitated major work tasks and planning to be done (i.e. containment access hole drilling, work scheduling and logistics, radiological protection planning, procurement of replacement SGs, storage of ex-service SGs and re-sealing and pressure-testing the containment). It is noted here that ease of replacement, or for that matter, inspection and repair, of SSCs will also facilitate that the radiation dose-penalties incurred by those personnel who carry out the work are ‘as low as reasonably acceptable/ achievable’ (ALARA-principle). It is a good practice to perform mock-up runs using dummy equipment to ensure that doses in the real situation remain ALARA. Doses can be efficiently reduced or practically eliminated by applying radiological protection principles (i.e. a person’s distance to the source should be as far as is practically possible, the time spent on the task should be as short as possible and lead shielding and remote tools should be used wherever practicable). The annual dose a worker in a NPP may acquire (person annual dose – PAD) is subject to national legal limits and regulation. Permitted occupational exposures, for nuclear workers and others (e.g. medical profession and researchers), are considerably higher than those for the general public. However, the ALARA principle still applies. For example, a maximum PAD of 20 millisievert (mSv) (1 mSv = 10–3 Sv) is allowed in Switzerland. It can be noted that even PADs of up to 100 mSv are not expected to cause health hazards, and so the above limit is conservative in nature. (Note: background doses for non-professionally exposed persons vary around the world due to altitude and presence or absence of natural uranium ores and the depth in the ground in which they are present, for

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example, but a PAD of 2 mSv is a commonly accepted average value.) A consequence of PAD limits is that when NPP personnel have reached their PAD, they will not be allowed to work in ionizing radiation fields for a year. Numbers of sufficiently trained/qualified reserve personnel in NPPs may be inadequate to cover for such eventualities, so ALARA planning is essential to ensure that PAD limits are not exceeded and that the available personnel will be sufficient to finish the task in hand in good time. Over the years, various ways and approaches have evolved to take into account AD effects in NPPs and how to implement experience gained into good practices through appropriate mitigation strategies. It is fortuitous that even in very early design NPPs, basic engineering principles had foreseen that conservative safety margins used for NPP-SSCs (dimensions and properties of materials), or the stipulated operational limits and conditions (e.g. temperatures, pressures, loading rates, coolant chemistry purity) would ensure inherently safe operation, even if significant levels of SSC-AD were present.

2.2

Nuclear power plant ageing terminology and associated definitions

The expressions ‘ageing management, ageing surveillance, plant life management and long-term operation’, which are defined below, are frequently used in nuclear power terminology [1]. Definitions ∑

Ageing management (AM): engineering, operational and maintenance actions to control, within acceptable limits, SSC-AD. Since AM is focused on keeping a high level of integrity in safety-related SSCs, it could be regarded as a quality assurance (QA) activity, because QA deals with quality systems focused on the relevant aspects. Thus, AM is a part of a total quality system. ∑ Ageing management programmes (AMP): any programme or activity that adequately manages the effects of ageing (e.g. maintenance programme, inspection or surveillance activities, all ageing mitigation strategies, repair, refreshment and replacement of SSCs). ∑ Ageing surveillance programmes (ASPs): these identify and document all known and possible SSC-AD and address the inspection, maintenance and mitigation aspects, as part of standard NPP-OPs. ∑ Plant life management (PLiM): a NPP programme integrating safety and non-safety issues, aimed at effective management of the NPP assets and resources. It is therefore the integration of AM and economic planning to:

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– maintain a high level of safety; – optimize the operation, maintenance and service life of SSCs; – maintain an acceptable plant performance (availability and reliability) level; – maximize return on investment over the entire service life of the NPP; – provide operating organizations and owners with the optimum preconditions for achieving long-term operation (LTO). (Note: within the context of this chapter, LTO means operation of the NPP in excess of its original design-life.)

Plant life management is a thus a methodology whereby all expenses are optimized to favour commercial profitability and competitiveness, while safe and reliable supplies of electric power are being produced. It can be noted here that throughout its development, the nuclear power industry has continuously implemented various programmes and strategies to control and mitigate SSC-AD. The traditional approach has been to adjust standard in-service inspection (ISI) and OPs and to focus monitoring, testing, repair, replacement and maintenance programmes to keep up with the state-ofthe-art, science and technology, as new knowledge and operational experience became available. The underlying reason for specific AM programmes is to have a better integration of existing programmes, driving the feedback from ISI to maintenance and vice versa, which may not be optimized in some cases. The ultimate objective of this feedback loop is to control and maintain the level of design and safety margins in SSCs to an acceptable level as operational time is accrued. A further facet and characteristic of AM is to address the impact of AD mechanisms on safety and reliability that were not addressed (or inadequately so) in the original design concept of the NPP, and to accordingly adapt ISI and maintenance activities to accommodate for them. For example, concrete creep and settling (structures), RPV and closure head bolt fatigue usage were addressed in original designs, but issues such as possible concrete carbonation, electrical cable ageing (embrittlement and loss of insulating properties) and certain SCC issues (e.g. intergranular and irradiation-assisted stress corrosion cracking in austenitic stainless steels (IGSCC and IASCC)) were not taken enough into account, since they had not yet manifested themselves in practice. The possibility of enhanced local piping erosion (wall thinning) was also not taken into full account in some cases. Other SSC-AD mechanisms may appear in future, so it is important to have in place a way to address them accordingly. Additionally, the AM programmes address the control of those environmental factors that have been identified to act as stressors assisting the SSC-AD. Thus AM programmes are tools to counter any deficiencies that have become evident in original NPP design assumptions and concepts.

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The goal of PLiM programmes is to facilitate an overall evaluation of the NPP’s assets. It integrates both safety and economic aspects. In particular, the economics aspects in the PLiM programme address such items as the management and procurement of spare parts and inventories thereof, and other fields within the PLiM deal with when to replace SSCs, fuel management, the possibility (and level) of power uprates (PUs), training of personnel and optimization of outage planning, for example. If they are not properly focused, scoped or managed, it is expected that AM and PLiM programmes, and even extended maintenance programmes, will bring with them extra cost burdens to the NPP’s operator. Therefore, a robust cost analysis of all AM and PLiM programmes and operation and maintenance costs, including ASPs, will be necessary to identify the real benefits to safety, reliability and overall profitability of the plant. To date, such analyses have shown all-round improvements to the safety, operational aspects and profitability, and have correspondingly encouraged operators to accept and implement such programmes.

2.3

Overview of ageing and its effects in nuclear power plants

All NPP types and their SSCs have undergone AD to varying degrees, since this is a natural phenomenon found in any engineered system, whether it is an aircraft, boat, chemical plant, oil refinery, bridge or building, for example. The nuclear environment, however, is characterized by special features whereby exposure of materials, especially to high energy neutrons (>1 MeV energy) arising from the fission processes in the fuel, can potentially cause some embrittlement in the ferritic RPV or IASCC in austenitic stainless steel core shrouds, for example. The coolant in PWRs and BWRs may also cause corrosion if specific material-coolant chemistry (electrochemical aspects, including water radiolysis), stress and flow-rate combinations and conditions are present. The reduction or mitigation of SSC-AD depends on understanding the often complex underlying mechanisms and synergies that act. This clearly depends on the results obtained from fundamental scientific research, whereby the interaction of the SSC materials with their respective stressors and operating environments may be understood in terms of the way they strive to achieve thermodynamic equilibrium. This underlines the importance of on-going research using realistic experimental parameters (i.e. equivalent to NPP conditions found in practice). This latter point is important since, apart from the convenience of simulated experiments, the goal of such research is to provide transferable data to the practical case and also to obtain regulatory acceptance of such results and the AD mitigation methods derived. Once the AD mechanism is understood, the appropriate actions can be robustly implemented to mitigate or even eliminate its effects.

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2.3.1 Concepts of defence-in-depth (DID) and common-cause failure (CCF) avoidance A requirement of the use of nuclear power is that it is of paramount importance to prevent harmful radioactive products from being released into the environment. It is for this reason that NPP designs have been conceived to function using the defence-in-depth (DID) principle. This multi-barrier approach can be likened to the cross-section of an onion, with each ring going out from the centre representing an engineered barrier. To achieve optimum safety, NPPs also feature multiple safety systems that supplement the natural features of the reactor core (e.g. negative void coefficient). Safety systems actually account for about one quarter of the capital cost of modern NPPs. The success and effectiveness of DID necessitate the highest quality in design principles and materials used and optimum construction principles. It also requires that equipment and layout thereof is installed to prevent operational disturbances or events from developing into more serious problems (i.e. avoidance of CCF). (Note: further aspects of DID, CCFs and latent failure conditions (LFCs) are discussed in Chapter 6.) Further features of DID are redundant and diverse systems to quickly detect problems, control damage to the fuel and prevent significant radioactive releases. Overall safety provisions thus feature physical barriers between the radioactive reactor core (or active materials storage facilities, such as the spent fuel pool) and the environment; they provide multiple safety systems, each with backup, and also have a concept to allow for human error. A robust severe accident management strategy will further support the DID principles and thus lessen the consequences of an accident should it occur. The engineered features and layered levels of protection behind DID designs obviously must take the safety classification and functions of the NPP-SSCs into account. From a logical and safety standpoint, the DID principle is a robust one, and OPs, ASPs, AM and PLiM programmes should also contribute to maintaining the integrity of DID barriers at all levels of operation, including design base accidents. It is fitting to mention here the final physical DID barrier, namely the reinforced concrete containment, which is designed to prevent the uncontrolled release of radioactive materials to the environment in the case of a severe accident. The containment is a life-limiting structure, as it is not practical to replace it and it is therefore also essential to ensure the overall economic and operational life of the NPP. As with any other material, reinforced concrete will also degrade with time when exposed to specific stressors. The effects of weather and pollution (acid rain) can cause degradation of concrete via chemical attack (e.g. carbonate formation, swelling, cracking and weakening). Cracks in the containment’s concrete can then allow moisture and corrosive chemicals to reach the steel reinforcement, which can then corrode, weakening

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the structure further or causing loss of leak-tightness. When water freezes it expands about 9% in volume and this can cause further damage due to spalling. Containment AD may be seen as discoloured patches (brown rust) or areas of white powder (calcium carbonate and other chemical species) on the surface. Severely weakened containments may not be able to withstand even design base seismic or impact events. Chemically weakened and severely physically degraded containments (cracks) may also lose some of their radiological shielding or retention capacity (leak-tightness). Repairs and refurbishments of containments, including attention to anchor bolts, tendons and auxiliaries, must be carried out during NPP outages for safety reasons.

2.3.2 Requirements on repair, replacement, inspection and monitoring of SSCs Ideally, whenever modifications to the NPP’s OPs are introduced, or when SSCs are repaired or replaced, the overall levels of reliability and safety should increase or, at least, be maintained. Accepting that some finite amount of SSC-AD will inevitably take place as a consequence of operating conditions, it remains to examine which SSCs are impacted for safety reasons and overall economic viability of the NPP. Also, the extent and rate of SSC-AD may be slower or faster in practice compared to that allowed for in the design provisions, including the safety and operational margins used. It is therefore the task of monitoring and testing systems to follow and quantify degradation rates, using appropriately sensitive and robust methods, and to thus determine the ‘margin of safety’ (e.g. degree of fatigue usage and thickness of pipewalls) or operational flexibility (e.g. RPV pressure-temperature limits for start-up/shut-down) still available in the SSCs. When this is known, it may be compared to design or regulatory requirements and realistic assessments made concerning ‘residual safe life’ of the SSC. Changes in operational parameters (e.g. coolant water chemistry, flow rate and temperature of coolant (the specific issues of PUs are dealt with later) and replacing SSCs with new designs and materials have a potential to influence degradation rates, and AM and PLiM programmes, including monitoring and inspection time schedules, have to be adjusted accordingly. Too frequent or in-depth inspections may not be necessary, but they will nevertheless cost money. In contrast, cursory checks or too lengthy time intervals between inspections may lead to spontaneous failure of a SSC. This will usually lead to even higher costs, since such forced outages, particularly those involving radiological contamination, will necessitate implementation of decontamination and associated radiation protection tasks, which can be time-consuming (thus costly) procedures. Thus it can be seen that inspection intervals, and the depths they go to (depending on safety class, importance to reliability, etc.) must be optimized, using the feedback from operational experience and results from the appropriate monitoring tools.

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2.3.3 Basic design principles, considerations and strategies in NPPs Before dealing with principles, strategies and goals of OPs, ASPs, AM and PLiM in detail, it is necessary to first identify some basic design requirements of NPPs, since design impacts not only the way the plant functions, but also how the SSCs are accessible for monitoring, testing, inspection and, eventually, replacement. Nuclear power plant site selection is determined by many factors, including local and political acceptance thereof, hydrological (e.g. river flooding or consequences of any possible dam-break upstream of the NPP and tsunami threat) and environmental impact (e.g. heat pollution of rivers or availability of sufficient cooling water) issues. Seismic activity is also a prime consideration, and NPP designs must take earthquakes into account, even if the probability of earthquakes is statistically very low as derived from the seismic chronological record of the region under consideration. Design and, for that matter, operational concepts, should identify and robustly eliminate the possible creation of CCF paths (i.e. where a series of human errors/weaknesses in procedures or the failure of one SSC leads to a ‘domino’ or ‘knock-on effect’ of progressive malfunction or failures in other SSCs vital for safety or control of the plant). It serves no purpose to have procedures or SSCs that individually fulfil design and regulatory requirements, but then are prevented from doing so due to CCF paths. The same approach is valid for thwarting terrorist attacks or lessening the possibility and consequences of aircraft crashes onto NPPs. Experience has shown that some SSCs are difficult to inspect, monitor, repair or replace (e.g. vessel internals), so (future) ideal NPP designs should allow for ease of access in general. Ease of inspection and monitoring will also increase the quality of inspection data and results, thus facilitating objective decisions to be made regarding residual life and when to repair or replace SSCs. Dedicated tools, such as robotic crawlers with integrated video systems, have been developed for currently operating NPPs to facilitate inspection of otherwise inaccessible areas. Realistic operational lives of some NPPs currently operating are now being regarded as 60 or more years. Regarding operation, increased online monitoring, condition-based and risk-based approaches are gradually supplementing the traditional method of fixed interval periodic inspections, since they are cheaper to implement and excessive conservatism may be reduced, to benefit operational costs and enhance safety through better awareness of the SSC’s real-time state regarding fitness-for-service. Much money can be saved by avoiding replacement of components that are still sufficiently good (within design limits), just as a consequence of old established habits or practices [2].

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2.3.4 Precursors for successful implementation of ASPs, AM, PLiM programmes and achieving LTO The basic concepts, methodologies and goals of ASPs, AM and PLiM programmes, and the utilities’ wish to achieve safe and profitable current and LTO should be well established already at the design stage of a NPP, and before it goes into operation. Basic programmes, including standard OPs, must be regarded as ‘living documents’, being amenable to adjustment and flexible to take into account experiences gained and the response of the plant’s SSCs as operational time is accrued. The nuclear steam supply system (NSSS) and all associated SSCs should feature the best materials known at the time of design, manufacture and construction and design concepts should also foresee the possibility for easy repair or replacement of SSCs, if necessary. Whilst a utility/operator’s goal is to run a NPP safely, reliably and profitably, and thus protect the overall plant investment, a regulator’s goal is to protect the public and environment from any possible negative consequences arising from the use of nuclear energy. These goals are, per se, complementary to each other, since operational safety is an inseparable facet of reliability and fitness-for-service of NPP-SSCs, and reliable operation facilitates high NPP load and availability factors, meaning more electricity sold. As in any other large and relatively expensive industrial project, a NPP might only begin to make a real profit towards the latter part of its original design life. The point in time where this takes place obviously depends, for example, on the level of current and outstanding capital costs/debts, number and type of SSC replacements carried out to date, degree of exceptional costs caused by mandatory back-fitting requirements and modifications and upgrading of safety systems, as well as the overall level of NPP amortization. (Note: further aspects of the economics of nuclear power are dealt with in Chapter 5 of this book.) It is logical that all operational actions and practices should be focused to create conditions to maintain compliance with the NPP’s licence conditions, optimize OPs, ASPs, AM and PLiM programmes to reach the full original design life and then make possible the option of further operation (LTO) under the applicable national licensing legislation. However, independent of other factors such as political and public acceptance, a NPP which has had freedom from forced outages due to spontaneous failure of SSCs (which is a reflection of the level of success of monitoring, inspection, maintenance, standard plant OPs, including ASPs, and AM and PLiM programmes), and demonstrates a proactive approach to safety (safety culture and training of personnel aspects), will be well placed to obtain regulatory permission to continue operation and go to the LTO phase of its life. Programmes to manage any form of AD in NPPs will have an expected level of success, but in practice this can either exceed or fall short of the expectations, depending on the suitability of the routine and standard OPs,

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maintenance, replacements, repairs and the effectiveness of ASPs, AM and PLiM programmes used. The success of such programmes will thus depend to a great extent on the competence of the workforce involved in their creation and effective implementation. Actions taken within standard OPs, ASPs and AM and PLiM programmes should lead to reliable and quantitative detection and monitoring of SSC-AD. When SSC-AD is discovered, the actions of repair or replacement will be guided by safety and business-case requirements. Practical implementation of AD mitigation or elimination remedies will depend on the availability of well-proven technologies, usually coming from basic research. Despite some diversity with respect to their separate goals and concepts, the OPs, ASPs and AM and PLiM programmes create mutually supporting conditions to ensure that NPPs operate safely, reliably and cost effectively, whilst they attain their original design lives. Since the SSCs, including non-replaceable items (life-determining), are managed and kept sufficiently within specification (engineering, reliability and safety margin aspects), even for times potentially well in excess of their original design lives, LTO will become a real option, at least from the safety, reliability and engineering standpoints. The business case, however, will decide the economic viability of further operation, whilst the NPP’s safety record will greatly influence the degree of public and political acceptance thereof. Ageing degradation arises due to the presence of stressors such as stresses and strains (constant or varying), temperature (level, constant or rapidly fluctuating), irradiation (primarily high-energy neutrons, gamma rays) and coolant/operational environment characteristics (impurities, level of pH, flow velocities and turbulence) causing ageing mechanisms. Typical ageing mechanisms are creep (through mechanical loading at elevated temperature), fatigue (high or low cycle, mechanically or thermally induced), thermal ageing (change in microstructure with time at temperature as new metallic phases are formed or when diffusion of elements to, or from, grain boundaries in alloys occurs), relaxation of pre-tension levels, corrosion, irradiation damage, wear and flow-induced erosion-corrosion in piping, for example. Loss of electrical properties in insulators or conductors may also be traced to ageing mechanisms whereby plastics degrade through heat or exposure to ionizing radiation (e.g. polymers will decompose into smaller molecular fragments and carbon dioxide gas). Connectivity or signal quality may be affected through oxide formation on electrical contact points of switches. Degradation usually leads to loss of material, cracks, deformation/distortion, changes in physical, mechanical, electrical and chemical properties and a corresponding reduction in SSC reliability and, eventually, safety margins. Those NPPs that have been designed, operated and maintained with the tenets of safety-first and effective implementation of OPs, ASPs, AM and PLiM in all areas will have the best chance of achieving safety, design life and economic goals and also LTO. Nuclear power plants that possess

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all original and comprehensive design documentation are in good basic condition, have no (or few) unresolved problems, have kept updated and relevant documentation on design features or SSC changes that have been implemented since start-up, have a true record of all design-base stressors that have occurred (e.g. transients, fast shutdowns (scrams) and fatigue usage on relevant SSCs), are ideally suited for focused AM and PLiM programmes. When original documentation is missing or inadequate, reverse engineering has to be performed to better serve the scoping, assessment and goals of OPs, ASPs, AM, PLiM programmes and, eventually, the possibility of LTO. Reverse engineering is the logic-based process of rediscovering and then documenting the technological principles of a device, object or system through analysis of its structure, function and operation. New RPV surveillance capsule specimens have, for example, been made from equivalent material and re-introduced into RPVs for irradiation damage studies in order to better quantify the extent and rate of degradation, since some original surveillance specimen sets were found to be inadequately documented or placed in an unfavourable irradiation position. Since NPP-AM and PLiM programmes, together with standard OPs and their ASPs may be regarded as management tools to achieve design life, optimized economic safe current and eventual LTO, and their goals are to deal with and manage all facets of SSC-AD in an efficient and cost-effective way, the extent and depth of such programmes must also take into consideration such further LTO-related aspects as the following: ∑

Any projected safety issues must be identified and analysed with respect to severity and the cost of implementing appropriate mitigation or upgrading methodologies. In LTO, the safety relevant and other SSCs will operate longer and perhaps under changed or more demanding conditions than were foreseen in the original design (e.g. when PU has been implemented). Nevertheless, safety margins must still be kept at a sufficient level. As a result of basic engineering and safety principles, SSC designs already feature conservative safety margins, thus allowing for the effects of design-life SSC-AD, as well as that expected, or projected, to arise in the LTO phase. ∑ The economic risk of a stranded investment when large, expensive SSCs are replaced or extensive back-fitting is carried out. As a result of the deregulation of electricity markets, a cost-benefit analysis will be an even more critical task to perform. However, safety must still have priority over economic aspects. The operator thus needs assurance that large investments (e.g. new SGs) will be amortized at least by the end of the NPP’s original design life, and that the regulatory climate will be favourable to allow LTO with relicensing or continued operation on the existing licence, as appropriate.

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∑

The availability of enough trained and qualified personnel to carry out the specialized tasks within the OPs, ASPs, AM and PLiM programmes for both current and LTO. It is no good having a state-of-the-art, science and technology NPP if there are insufficient or inadequately trained personnel to run it. Succession planning must be thought of well in advance to facilitate a smooth and efficient transfer of expertise. This is an aspect of knowledge management (KM). ∑ Identifying any additional SSCs for AM and PLiM that may be crucial for safety and economic viability of the NPP in the LTO phase. The focus of LTO-AM and PLiM programmes will not only be on managing and mitigating SSC-AD in large, passive ‘life-determining’ SSCs, but will also continue to address those SSCs vital for safe and reliable operation, which could be expensive to monitor, test, repair, maintain or replace. Furthermore, the safety principle of maintaining the engineered DID barriers must also be upheld in the NPP’s LTO phase. Future environmental and operational stressors must be identified (or anticipated, as far as possible), and their control or mitigation addressed in the standard OPs, ASPs, AM and PLiM programmes, respectively. ∑ Projected needs and capacity (both human and financial) for spent fuel (SF) storage management tasks, since increased quantities of SF are a natural consequence of LTO, final NPP decommissioning, dismantling, radioactive waste (radwaste) conditioning and disposal, and associated plant actions required (e.g. additional SF pool construction or modifications to increase the capacity of existing SF pools, and commissioning of plasma ovens for radwaste volume reduction) and to identify areas where additional safety barriers and controls (e.g. filters and radiological monitoring) and adjustments to OPs, ASPs, AM and PLiM programmes may be applicable.

2.4

Overview of systems, structures and components (ssc) safety classes

2.4.1 Safety classification of items and SSCs Safety classification is a vast, sometimes complex and always a time-consuming subject and requires comprehensive systems and engineering knowledge as well as a holistic approach to exactly capture the levels of safety importance and inter-dependencies of SSCs in NPPs. The following provides an overview of the subject, and gives some examples of safety classification of items and SSCs. The main source of information used here has been drawn from STUK [3]. Other supporting information may be obtained from appropriate references and sources of further information (e.g. Männisto, 2005) provided at the end of this chapter).

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Although certain commonalities and generic approaches exist, the scopes and depths of OPs, ASPs, AM and PLiM programmes are NPPspecific, reflecting the actual operating experiences and needs of the NPP, and design, under consideration. Correspondingly, safety classes of SSCs, although broadly similar for varying NPP designs, will reflect the diversity in function, operational modes and reactor control systems in NPPs. Whilst AM programmes are critical for safety, and detailed ageing management analysis is restricted to a few critical plant life-determining SSCs, the PLiM programmes address almost all SSCs in a NPP. The standard OPs, ASPs, maintenance, inspection, testing, repair and replacement programmes in NPPs are supplemented accordingly by AM and PLiM programmes. Operational conditions of NPP-SSCs necessitate that they possess certain levels of inherent good quality (e.g. works certification, sufficient design and safety margin features and regulatory approvals) and that they are suitable for use in the specified environment. Mechanical components can be conveniently placed into three basic groups: 1. those where integrity has to be guaranteed (e.g. RPV and shells and nozzles of SGs); 2. those where preventative maintenance is used to preserve the initial levels of quality and function (e.g. removal of sludge deposits in SGs to lessen stress corrosion cracking, sleeving and plugging of defect SG tubes); 3. those that may be allowed to fail (failure-based maintenance/run-tofailure), and then are repaired or replaced even on a routine basis to re-establish the initial quality (e.g. replacement of small diameter piping, switches/fuses, pump seals and bearings). Safety-relevant components and systems are in the first two groups, and these are also usually the subject of AM programmes. Safety classifications in NPPs have traditionally been focused at the system and component level of equipment. However, the evolving approach for currently operating NPPs is to focus more on areas that are important to their safe, reliable and profitable operation and also the effectiveness of maintenance used to keep the items and SSCs fit-for-service. The NPP current licensing basis (CLB) can be taken as a starting point to evaluate the NPP’s SSCs in terms of their importance to safety. Safety class 1 A key purpose of safety classification of items and SSCs in NPPs is to identify those which are potentially hazardous in nature or those that must maintain their integrity under design base accident scenarios and to ultimately focus OPs, ASPs, inspections, monitoring, testing, maintenance, repairs and replacements

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on these to ensure they are handled and maintained according to their design and regulatory requirements. Safety class 1 items would thus include the reactor fuel, since it is fissile, radioactive and toxic, and if misused has the potential to cause damage to the population, environment and to the NPP itself (contamination). Primary circuit components (PCCs) retain large and fast flowing volumes of gas, water or steam under high pressures and temperatures, and major breaks in such PCCs that lead to a loss of coolant, the extent of which being so great that it is impossible to compensate for it using the available supply of emergency make-up water, for example, must therefore be avoided. A RPV is a safety class 1 component, since its catastrophic (sudden) fracture could lead to conditions that exceed loss-of-coolant accident (LOCA) design base allowances. However, not all PCCs are assignable to safety class 1, since their failure may be readily controlled. Such items are small diameter pipes (<20 mm inside diameter) and components integrally connected to the reactor coolant system by a passive flow-limiting device and which, even if ruptured, do not cause a leak any larger than that caused by the rupture of a 20 mm inside diameter pipe. Other PCCs, which in the event of failure can be rapidly and automatically isolated (multi-valve closure) from the reactor to facilitate normal shutdown and cooling, are not included in Safety class 1. Safety class 2 Safety class 2 items and components are those not assigned to Safety class 1. They cover items and SSCs needed for reactor trip, the systems to deal with emergency core cooling in case of a LOCA, boron supply to control core reactivity/criticality and decay heat removal, for example. Further examples are: steam and feed-water systems, reactor containment and systems needed to protect containment integrity in a postulated accident (e.g. containment spray, pressure relief and isolation valves), primary circuit supporting structures that protect Safety class 1 components (e.g. restraints, missile barriers) and RPV internals needed to support the fuel core or to facilitate the cooling thereof. Other Safety class 2 items include storage racks for new or spent fuel and electrical components and essential electrical power supply equipment to Safety class 1 and 2 items or SSCs. Safety class 3 Some typical systems in this safety class are: boron supply, reactor volume control (PWRs), parts of the emergency feed-water supply not in Safety class 2, primary circuit cooling and pressure relief, reactor and spent fuel decay heat removal and also those systems needed to start up or operate Safety class 2 and 3 items and SSCs. Other systems in this safety class are the reactor cooling water clean-up, treatment and storage for liquid wastes and

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radioactive gas treatment systems, nuclear fuel and handling systems that may cause fuel damage should they malfunction, for example. Further items are hoisting and transport systems for nuclear fuel, pools and tanks, reactor power limitation systems, monitors for reactor criticality when refuelling, primary circuit chemistry and leak monitors, and systems for controlling containment releases, as well as those to prevent formation of explosive oxygen–hydrogen gas mixtures. Safety class 4 This safety class contains the following items and SSCs: fire protection, turbine generator components (bearings, rotor trip valves, and oil systems), vibration monitors, hydrogen cooling, generator circuit and field breaker, turbine and generator protection systems. Various instrumentation and control (I&C) and computer systems for plant control and operation, safety significant information systems dealing with operation and maintenance. (Note: a comprehensive listing of typical safety classes is given in STUK [3]).

2.5 Setting up and scoping ageing degradation and surveillence programmmes in nuclear power plants (NPPs) The creation of SSC-AD surveillance and management methodologies in NPPs necessitates a comprehensive appreciation of the varied AD mechanisms that could occur. The approaches have to thus take into account the many types of SSCs present, the different materials used and the operational environments, as well as their significance to safety (safety classification), their impact on reliability, profitability and their relevance to the NPP for assuring attainment of the original design life and, eventually, LTO. Standard OPs, which include routine maintenance, monitoring, controls, repairs and replacements, as well as ASPs, are normally quite sufficient to address most issues that arise in the plant on a day-to-day basis. The standard OPs are generally based on written procedures and recommendations of the NPP manufacturer, having due regard for any additional tests, monitoring or other supplementary actions that are prescribed by the regulator. One approach to identify, characterize and document AD is to create dossiers which describe each specific component in a structure or system. Component-specific dossiers are a methodical way to document all phases of the component’s life. Such dossiers should contain details and information such as: ∑ Basic description of component, its function and location in the NPP. ∑ Manufacturer, supplier or sub-contractor.

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Materials used (e.g. alloys, ceramics, plastics, rubbers and cements) and, where relevant, applied heat treatments, surface conditioning (e.g. peening and protective coating). Works reports, including final quality assurance certificates, will be an integral part of the SSC ageing surveillence dossier. Particular details, such as approved lubricants to use must also be included. Where applicable, notes on the avoidance to exposure of the SSC material to chloride or sulphide ions (e.g. danger of stress corrosion cracking occurring in some austenitic stainless steels), or to other environmental conditions (e.g. excessive temperature, moisture, radiation), must be included and highlighted. Approved substances for SSC decontamination tasks must be listed. ∑ Assigned safety class of the item or SSC with regard to its function (e.g. no RPV fracture allowed, piping leak-before-break to be demonstrated, assurance of functionality in emergency situations, redundant systems and importance to DID) and supporting documentation. ∑ Specific operational environment (e.g. coolant (water, heavy water, borated water, hot gas and quality/purity thereof), steam, neutron irradiation, raw/service water). ∑ All known, or possible, AD mechanisms (drawn from plant’s own and global operational experiences, failure reports (root causes) and published research literature) relevant for the component and its materials in the specific operating environment. (The AM programme will address SSCAD issues that were unforeseen or inadequately addressed in the design concept.) ∑ Validated and approved AD mitigation actions relevant to the SSC being documented (some examples of which are given below): – neutron embrittlement: reduction of RPV neutron flux, hence fluence with time, either by ‘inside-out’ fuel management, whereby the older fuel elements are placed on the outside of the fuel core and the fresh elements are placed in the centre of the core, or by the use of dummy fuel elements situated on the periphery of the core; RPV beltline annealing, with parameters of time and temperature for recovery of properties, derived from tests on the plant’s own RPV surveillance programme materials, and any necessary supplementary test specimen programme to quantify fracture toughness levels (level of success of annealing), including provision to follow re-embrittlement trends as additional service time is accumulated; – water chemistry improvements (e.g. hydrogen injection into boiling water reactor coolant to lower the electrochemical potential of stainless steel in contact with it and thus lessen the propensity toward stress corrosion cracking in some austenitic stainless steels); – shot-peening (mechanical improvement process, MIP) (e.g. to introduce compressive stress in the surface to improve stress corrosion cracking resistance in Alloy 600); © Woodhead Publishing Limited, 2010

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– use of tie-rods (to enhance the structural stability in cracked austenitic stainless steel core shrouds in boiling water reactors, even if through-wall cracks are present); the necessity to use approved methodologies to check the integrity of such tie-rods and inspection intervals should be highlighted. ∑ The optimum inspection, measurement and monitoring methods for detecting the most probable AD mechanisms should be documented. Ultrasonic testing, dye-penetration, magnetic particle, X-ray and eddy-current (as non-destructive tools) and destructive testing of RPV surveillance capsule specimens (Charpy, tensile and fracture mechanical specimens) will have application in specific fields of SSC-AD detection, monitoring and testing. ∑ The frequency and depth of inspections required, based on the SSC safety class, the expected rate of AD and operational experience to date must be recorded. A physical graphical record could be included, showing, for example, pipe wall thickness change, obtained from periodic measurements, as a function of time. Computerized systems, using continuous wireless signals and direct data transfer to monitors, are also finding increasing application. Where necessary, references to Codes and obligatory (regulatory) procedures to be followed must also be provided in the dossier. ∑ Where applicable, any special requirements for personnel charged with inspection and testing of safety-related SSCs (e.g. licensing of personnel who perform work on such SSCs) must be highlighted in the dossiers and accompanied by supplementary documentation dealing with training, qualification and licensing of personnel. ∑ A chronologic record of the SSC operational experience (e.g. sporadic malfunction or spontaneous failure and root causes), and provision for keeping updated records of plant transients and ‘scrams’ (essential for RPV and bolt fatigue usage monitoring, according to design specification allowances) is an integral part of the SSC ageing surveillance dossier. The effect of PUs in NPPs requires a comprehensive systems analysis, and NPP-SSC modifications or changes that have become necessary on both primary and secondary sides of the NPP (e.g. adjustments to valve set-point values) must be duly documented. Any changes to standard plant OPs (e.g. frequency, depth and scope of inspections, monitoring, maintenance and periodic testing) have to be noted. Furthermore, the increased core-power density and higher coolant core-exit temperatures, as well as faster steam flow rates will influence the time schedules between SSC inspections in NPPs with PU. Documentation of adjusted operational prescriptions or emergency procedures relevant to the plant’s PU status is also necessary. The creation of comprehensive SSC-AD dossiers, in addition to the standard plant documentation on OPs, enables the plant personnel to recognize © Woodhead Publishing Limited, 2010

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possible SSC-AD issues objectively and to instigate appropriate measures to address them in a timely manner. Such dossiers also show a regulator that the NPP personnel are aware of SSC-AD and that a safety-oriented, pragmatic, science-based and quantitative approach is being used in the general operation and maintenance of the NPP. Depending on the country involved, there are different approaches to the NPP licence status and thus the permission to operate. In Europe, licences are mostly not time-limited, and are granted for as long as the NPP can operate safely, and the periodic safety review (PSR) [4], usually performed every 10 years by the NPP owners and controlled by the regulator, is the main tool used to provide proof of safety. In the United States, the licence is time-limited, necessitating a re-licensing process, after the NPP brings proof that all maintenance has been effective and all regulatory requirements have been fulfilled.

2.5.1 Living document nature of ASP, AM and PLiM programmes Despite rigorously applied standard OPs, ASPs, AM and PLiM programmes, SSCs may still fail unexpectedly, or become prematurely unfit for service, but it is then not only essential to find the reason why a given SSC fails, it is also necessary to determine how the circumstances arose that allowed it to do so. Failures in NPP-SSCs can be caused by incorrect human actions or they can be a result of direct physical and chemical causes, for example, but it is important, sometimes crucial, to identify if there are any underlying weaknesses in the original design and in the inspection, testing and monitoring programmes used. Only a full in-depth analysis, leading to root-causes, will lead to robust and fact-based actions to be taken in order to avoid a repetition of the failure in future operation. The standard OPs, ASPs, and AM and PLiM programmes, must then be adjusted accordingly to take into account the (usually) plant-specific new knowledge, in accordance with a ‘living document’ approach. Problems can also be caused through inadequately performed or neglected inspections (e.g. insufficient AD detection, low sensitivity of inspection tools used, or time interval between inspections being too long), inaccuracies in SSC monitoring equipment (e.g. calibration/signal drift of on-line monitors sends incorrect information to the plant control systems, thus misleading the operator concerning the real situation or threat to the plant’s safety) or incorrect maintenance procedures (e.g. wrong lubrication or decontamination procedure used, or excessive (or insufficient) torque applied to bolts). From the above examples, it can be seen that updated information exchange with the NPP’s manufacturer/SSC supplier and other operators with similar plants, as well as with regulators, is an essential part of an operator’s duties. Following published literature on

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all aspects of NPP technology, reportable events that have occurred and basic research into SSC-AD are further related tasks to perform to ensure OPs, ASPs, AM and PLiM programmes continually benefit from new knowledge, experience and lessons learned.

2.5.2 Important systems, structures and components in nuclear power plants Different designs of NPPs will feature their own specific SSCs that, despite accruing various levels and types of service-induced AD, still have to be kept functional within, and according to, their design specifications and regulatory requirements. Conservative operational design safety margins of SSCs are originally in place to counter AD, so the preservation of sufficient safety margins is a way to keep NPP-SSCs safely and reliably in operation. When OPs, ASPs, AM and PLiM programmes are optimized, and all SSCAD is effectively managed, mitigated or eliminated, and safety margins kept within specification, the NPP should reach its original design life in a safe and profitable manner. Thereafter, when the SSCs, including the containment and RPV (irreplaceable items), continue to fulfil their design and regulatory requirements, there are prima facie no technological or safety issues to prevent the NPP from going into LTO. A typical list of important SSCs that can undergo AD in Western-designed PWRs would feature the RPV and closure head, primary system large diameter pipes, steam generators, primary pump casings, pressurizer, control rod drive mechanisms, RPV internals (e.g. baffle plates), containment, turbine, generator, I&C, electrical cables and the cooling tower. For a BWR, the list would contain items such as the RPV, the vessel internals (core shroud, jet pumps and core spray and supports), reactivity control blades, containment, coolant recirculation piping, steam dryer, tanks and valves, as well as electrical equipment, I&C and cables. For the pressurized water Russian designed WWER-type of reactor, a typical list of SSCs would include the RPV and head closure, control and safety systems, piping, pumps, pressurizer, electrical circuits important for safety, I&C, safety valves, RPV internals and the emergency core cooling system. The Canadian designed CANDU reactor (heavy water cooled and moderated) would include fuel channels, zircaloy pressure tubes, steam generators, calandria vessel, reactor headers, vacuum building, containment, reactor building, calandria end shield and cooling system, cables and spent fuel bay.

2.5.3 Obsolescence and its consequences in SSCs Obsolescence, when applied to SSCs, means that the original manufacturer no longer makes the product and no other source of supply is available.

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The problem can be exacerbated due to limited industry capacity or a manufacturer’s reluctance to re-tool to produce a limited number of ‘outdated’ components. It is therefore good practice to identify critical (also potentially expensive) components in the frame of the PLiM programme, and to build up a strategic inventory of spare parts. The economics of such inventories must be evaluated against the risk of long outages of the NPP if no spare part is readily available. However, non-service ageing may become an issue if the storage conditions are not ideal, since elevated temperatures, dampness, ingress of chemical vapours, corrosion and incidental vibrations may still cause SSC-AD (e.g. set-point drift, change in resistance, loosening of clamps/ connections or oxidation of contacts). Nuclear power plants constructed in the 1960–90 period had mostly analogue I&C systems. These performed adequately, but as technology evolved and the digital era made itself felt in the nuclear industry, it seemed a natural development to implement digital I&C into NPPs, whilst classifying analogue technology as old and ‘obsolete’. However, the general area of ‘electronics’ is characterized by frequent modifications and modernization, creating challenges in assuring that functional equivalence is present. Both proof of quality assurance (QA) and functional equivalence are necessary to obtain regulatory approval for replacement (or new design) safety-relevant components. Use of new technologies and procedures used may also require regulatory approval. Obtaining approval may take time, thus creating reluctance in some operators to ‘modernize’, despite the advantages of digitally-based I&C. There are some special caveats involved with replacing old technology with new. Some components, for example in the electronics sector, may fulfil the specifications to provide a given actuation voltage, but reaction time allowances/windows to actuate timers, relays or switches involved may be reduced compared to that in the replaced component and this new characteristic may be intolerant to the system it is a part of, causing erratic behaviour. Variations such as this arise due to small differences in (sub) components coming from various manufacturers. It is therefore necessary to quality control and validate all component characteristics, including tolerance bands for generated signals and their possible effect on other components and systems.

2.6 Safety culture and human factors and knowledge management 2.6.1 Importance of safety culture and human factors in NPP operation The term ‘safety culture’ can be used to convey the way in which a NPP (or in fact any other plant) is operated with respect to safety. It may be defined as

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the product of individual and group values, attitudes, perceptions, competencies and patterns of behaviour that determine the commitment to, and the style and proficiency of, an organization’s health and safety management. Nuclear power plants with a good safety culture are characterized by communications founded on mutual trust, by shared perceptions of the importance of safety and by confidence in the efficacy of preventive measures. It is thus of paramount importance in NPP operation to have in place an appropriate safety culture, whereby openness to, and for, communication, questioning attitudes, adherence to written procedures, the necessity to analyse and understand the immediate and later consequences of actions performed and a pride in the work done, all have high standing within the workforce. For example, the pressure of electricity production targets must never be allowed to create conditions under which short-cuts in established OPs or mandatory procedures may be contemplated. Apart from legal aspects, such personnel actions may unknowingly put the NPP in an immediate or future dangerous condition (i.e. creation of latent failure conditions, LFCs). The tools of OPs, ASPs, AM and PLiM programmes, combined with a trained workforce instilled with the core values of safety culture, create the best possible conditions for safe current and future NPP operation. It is therefore an important task for NPP management staff to create a strong enabling environment for safety culture development and efficient implementation thereof.

2.6.2 Importance of knowledge management The plant’s management strategy to identify where (specialist) pools of knowledge lie within the NPP’s workforce, to acquire and document it, to efficiently disseminate it and to thus preserve this valuable asset, even when the original specialists have left, is called knowledge management (KM). In fact, KM can be associated with elements of safety culture, since the corporate asset of a NPP’s experienced and technologically specialized workforce is also a key to safe operation. Explicit or generic knowledge from manuals or reports is usually adequately available, but ‘hands-on’ teaching of new personnel, who are accompanied by a senior specialist for the requisite amount of time, is often a more efficient way to disseminate know-how, especially when the work processes are then documented and rationalized in ways which are transparent, logical and easy to understand. This implicit knowledge is based on experience, which has usually been accrued by an individual or team, perhaps over many years, and it is mostly plant-specific and thus extremely valuable to the NPP concerned. An important tool in KM is the exact documentation of all NPP-SSCs (especially those relevant to safety and the NPP’s operational life) reflecting the operational history (e.g. transients, fatigue usage, nature and type of repairs and for what reason). Such sources of information may be readily obtained from the

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component-specific dossiers featured in the ASPs, or given deeper focus in AM and PLiM programme documentation. It is opportune to mention here the necessity for KM within the NPP, even when outsourcing is required, i.e. when external specialist companies are called in to perform specific tasks. The NPP’s own personnel should be involved as much as possible, and be ready to learn new skills. Such ‘home-grown’ and acquired knowledge creates a competent workforce, and those involved have the potential to develop their careers further as specialists. This increases motivation and safety culture in the NPP. Documentation must be kept on all safety-relevant issues and events that arise. This is usually done under the provisions of regulatory requirements concerning reportable events. The details leading up to, during and after the event should provide a clear analysis of the main and contributory factors. The analysis will show, for example, whether the OPs, ASPs, AM and PLiM programmes have any inherent weaknesses to address, or if SSCs have been inadequately inspected, repaired or replaced, or if the supplier of the SSCs changed some detail, originally thought to be irrelevant, or whether human factors were involved. All this information must be used to ensure a similar event does not occur again. Lessons learned (root causes of problems) are a part of KM and provide valuable feedback for enhancing the QA and effectiveness of OPs, ASPs, AM and PLiM programmes.

2.7

Trends and issues in nuclear power plant (NPP) life management

In order to address future issues and trends in NPP life management, it is necessary to first examine past political situations and technological events that have shaped the industry until now. Both the national and international safety record of nuclear power will continue to be a yardstick used by energy policy-makers to assess the justification, or not, for using nuclear power. The international dimension is introduced here to emphasize the collective responsibilities of all involved to operate their NPPs safely, and to highlight the fact that releases of radioactive source terms (contamination) simply do not respect international boundaries. The accidents in Chernobyl (26 April 1986, Ukraine) [5] and Three Mile Island (March 28 1979, Pennsylvania, US) [6], caused widespread rejection of nuclear power and a general distrust in nuclear technology. Since then, much confidence has been restored in nuclear power, thanks to improvements to designs, safety culture and continued freedom from major accidents. With uncompromising vigilance, questioning attitude (safety culture aspects), implementation of effective OPs, ASPs, AM and PLiM programmes and following the current state of knowledge, sharing experiences (open discussions with other operators and regulators), and by using validated good practices and applying lessons learned, safe current and

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LTO operation of NPPs should be facilitated. However, utilities and operators have to prepare, plan and invest for LTO at least five years in advance of reaching the NPP’s original design life, even if political and public support could change in the future, due to other unforeseeable circumstances. This is a challenge to the business case, even if the regulatory climate should remain favourable. In other words, regulators can implement operational and safety rules and can verify that a NPP is in compliance with these, but they are not involved with national energy policies or private sector business goals. Furthermore, many utilities are aiming to keep their NPPs operating for 60 or more years, and this will need increased regulatory capacity (re-licensing aspects) and oversight. A potential future problem may arise due to regulatory capacity being stretched, especially when new NPP designs and associated new operational practices have to be approved and construction of new NPPs starts. Thus, the nuclear power industry and its regulators must ensure that sufficient numbers of suitably trained personnel will be available in both currently operating and future NPPs. It is necessary to implement relevant courses in technical colleges and universities in a timely manner, provide young technicians and engineers with the prospect of a life-career and to realize that basic training and qualification may require two to three years, with some specialist areas (doctorate level at university) requiring up to five or more years. Expertise in decommissioning NPPs, radwaste conditioning and radiological protection will also be needed in future. Training of future radiation protection experts and technicians in ways to plan jobs using the ALARA principle should be given due attention, bearing in mind the time needed to train and qualify such personnel. The prospect of a career for life in an expanding and vital industry, plus the pride and responsibility in working for the benefit of the environment and humanity should be seen as good incentives for the coming generation of nuclear technologists.

2.7.1 Role of NPPs currently in operation Past history shows that despite economic recessions and associated reduction in industrial output and changing political situations, the world will, overall, need more energy as it progresses further into the twenty-first century. This will be against a background of increased environmental awareness concerning the use of fossil-based fuels and possible shortages thereof. In contrast, nuclear energy represents an economically sensible, reliable and environmentally responsible path to follow. Over 300 NPPs (from a current world total of 438 in June 2010) are now approaching, have reached or have already exceeded half of their original design lives, and almost all NPPs in the US have plans to apply for licence renewal, when the time comes. This clearly indicates the renaissance of nuclear power in the US. In Europe, most NPPs will have the possibility to continue operation on their current

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licensing basis through successful completion of the PSR procedure. Even allowing for new NPPs being constructed now and in the future, a significant quantity of the generated capacity in the next 20 years, or so, will still have to come from NPPs that have operated for 30 or more years. These older NPPs generally possess improved safety and economic status compared to when they first went into service, due, in part, to replacements of SSCs, back-fitting and refurbishments and implemented mitigation methods against SSC-AD. Furthermore, OPs, ASPs, AM and PLiM programmes, and monitoring the effectiveness of maintenance would have all contributed significantly to the relicensing and LTO of these NPPs [7].

2.7.2 Contribution of power uprates Power uprates (PUs) are an on-going activity in many NPPs, and they are already a valuable contribution to the total of global nuclear energy produced [8]. In the US, for example, up to 5 GW of electrical power could be added to the total nuclear capacity via PU between 2005 and 2010 [9]. Uprating the power output of nuclear reactors is a highly economically attractive source of additional generating capacity. By modifying the existing turbo-generator (e.g. increasing steam intake capacity) or installing a new turbo-generator, and by taking advantage of existing design margins and installing digital I&C, plant output can be readily increased by up to 15–20%. Power uprates are done for both BWR and PWR. The necessary increase in the fuel core power is achieved in BWRs by increasing the core feed water flows and steam flows. Either the rate of coolant recirculation is kept constant, thus creating larger core steam volumes, or steam volumes are held the same via an increased recirculation rate. A combination of these measures is also possible. In PWRs it is necessary to increase the core coolant flow rate or to raise the mean coolant temperature rise across the core, or both. The secondary steam volume is increased, which is then converted to more electrical power via the turbo-generators. It is usually quicker, easier and cheaper to make the adjustments necessary for PU than to build a new NPP. Using a simple example, if a country operates five similarly rated NPPs and each one was given a 20% PU, the equivalent power generated after PU completion would then correspond to having 6 NPPs in operation. This could be an attractive alternative to building a new NPP, or as an effective and cost-efficient measure to ‘bridge-over’ energy requirements whilst waiting for a new NPP to be built and to connect to the supply grid. There are three categories of PU: measurement uncertainty recapture (MUR) uprates; stretch power uprates (SPU); and extended power uprates (EPU). The MUR uprates are typically no more than 2% and are obtained by improved calculation of the reactor power. The SPU leads generally to

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a maximum of 7% increase and is characterized by altering set-points in the instrumentation. The EPU can be used to obtain up to a 20% increase in power. This, in contrast to the previous PU methods, requires significant changes to the OPs and equipment of the NPP, including a higher power density core, generator renewal or modification, systems to deal with the steam or coolant flow rate increase and focused monitoring and inspection on piping and steam dryers, etc). The NPP balance of plant usually addresses the secondary side situation (i.e. turbines and all else not in the primary side). Although PUs are done within the design tolerances, they can pose new and special challenges to the operation of the NPP, since higher volumes of coolant and steam arise and flow rates increase. Furthermore, the nuclear core is usually modified to create the extra energy via more enriched fuel. Higher core power, increased flow rates, temperatures and pressures not only have the potential to affect AD rates in SSCs (e.g. vibration-induced cracks in steam dryer cover plates in BWRs [10]), but also new operational guides and rules have to be implemented to allow the plant’s personnel time to act accordingly in cases of emergency (i.e. a PU also necessitates an analysis of human factors and operator’s actions, especially under anomalous or transient conditions in the newly configured NPP. Times available for operator action/ intervention may be shorter relative to before the PU was implemented). The possible effects of PU on reliability and safety margins of SSCs must be analysed and proof brought that they remain within regulatory prescriptions [11].

2.7.3 Safety regulation and anticipating future changes Regulatory and safety requirements expected for new types of NPP will be, in principle, no different from those for current NPPs, namely that all plants must be operated from ‘greenfield-back-to-greenfield’ (total life cycle) such that no harm occurs to the public and environment. The traditional approach of NPP regulation has focused on verifying that the safety and operation of NPPs always conforms to their current licence requirements. This pragmatic approach implies that the best possible NPP designs are used from the start of operations, that safety systems upgrades and SSC replacements are implemented as necessary, the DID barriers are kept functional, the OPs, ASPs, AM and PLiM programmes are effective in their respective goals, and the workforce is correspondingly qualified and safety oriented. This is a ‘state-of-plant’ approach. In future, this approach will also be important, but a growing trend is for regulators to verify that the organizational arrangements in place at the NPP can ensure safe operation at all times. The organizational arrangements and OPs obviously cover a multitude of activities within the NPP, and these, in turn, are based on established codes and guidelines and the implementation of lessons learned and good practices. The OPs, ASPs,

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AM and PLiM programmes for new NPPs should clearly reflect a sciencebased approach, backed up with sound engineering principles with a primary goal of safety. Assuring the integrity of primary pressure boundaries (e.g. RPV, SGs, pressurizer and piping) remains a priority task, but DID barriers and robust severe accident management procedures must nevertheless be available to limit any consequences of SSC failure [12]. Risk-based maintenance and inspection methods are also taking on more significance, and the use of probabilistic risk assessment (PRA) is an increasingly valuable tool used for plant-specific assessment of risks to safety. Core damage frequency (CDF) is a term used in PRA. It is a way of expressing how probable it is that an accident would cause damage to the nuclear fuel core. The CDF number is a useful tool for managing the level of risk of core accidents. The CDF is obtained by estimating the frequency of design base accidents that may cause damage to a nuclear reactor fuel core (PRA Level 1). In Western European NPPs, the theoretical CDF is calculated to be around 5 ¥ 10–5 per reactor year (which would be once in 20 000 reactor years). In the US, reactor designs must fulfil a theoretical CDF criterion of 1 event in 10 000 years, but current designs are superior in this respect, and are usually projected to have a CDF of 1 in 100 000 years. However, NPPs commonly have a CDF of about 1 in 1 million years. New generation NPPs (Generation III, III+ and IV) feature many passive safety systems and the theoretical CDF is estimated to be about 1 in 10 million years. It should be noted that nuclear core damage accidents do not necessarily mean the release of radiotoxic substances to the environment, as was shown to be the case for the Three Mile Accident in the US [6]. The value of PRA approaches will increase as methods for PRA Level 2 (addressing plant source-term release) and PRA Level 3 (assessment of injury and economic loss due to release of radioactivity) are further refined. The impacts of external events on the safe operation of a NPP (e.g. seismic, flood, aviation accident, terror attack and hurricane) require continued analysis. Safety regulation in NPPs is applied in a national way, but it has its roots also in internationally derived and approved approaches. Proven design principles and engineering codes and practices also provide input to assist the regulation and assessment of safety in NPPs. Safety regulation can have basically either a prescriptive or a ‘goal-seeking’ character. As time evolves, it is likely that the latter approach will find increased use as the principle of ‘The operator is responsible for safety’, becomes even more firmly established. It is opportune to mention here that electricity markets are becoming increasingly deregulated, thus creating market opportunities but also potential impact on safety levels. Changes in NPP operational procedures (e.g. in order to save time/money/increase output) must first be analysed, and officially approved if safety may be affected.

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A basic principle used in the regulatory field is to ensure that NPPs are kept at the current state-of-the-art, science and technology, and sufficient safety margins are present in safety-relevant SSCs irrespective of the NPP’s chronological age. This principle will be valid also for future (new generation) NPPs, thus assuring that new knowledge and technology will continually benefit the goals of safety-first.

2.7.4 Issues concerning waste arising from the operation of nuclear power plants The operation of a NPP produces radioactive waste (‘radwaste’) and radwaste management has to take into account the low, medium and high level types arising. It is evident that the longer a NPP operates, the more waste and spent fuel it will produce. Thus, even well before the NPP’s LTO phase is entered into, sufficient capacity to store, condition and dispose of the projected amounts of radwaste must be planned for, or already be available. For example, construction of large items such as new spent fuel storage pools (or modifications to existing ones) must be undertaken, and corresponding strategies evolved to enable storage of the fuel elements whilst avoiding criticality accidents. Spent fuel storage pools are a very important part of NPPs since they must ensure that the spent fuel’s decay heat is removed efficiently, they shield personnel from radiation and provide a barrier against radioactive releases. Correspondingly, they must also be continually monitored and maintained chemically (i.e. sufficient boration must be present in the water to absorb neutrons emitted by decays in the fuel) to prevent spontaneous criticality of the fuel [13]. Spent fuel can be re-processed to extract remaining quantities of fissionable materials (e.g. uranium and plutonium) and to separate non-fissionable, but radioactive materials, which can be processed and then encased in suitable containers for storage or final disposal. The overall strategy used for radwaste management is to concentrate or vitrify it as appropriate, reduce its volume and to confine and contain it for as long as it may present environmental or health hazards. High-level waste (HLW) may contain several billion Becquerel (Bq) per cubic centimetre compared to a few tens of Bq per cubic centimetre in low-level waste (LLW). (One Bq is equivalent to 1 radioactive decay per second.) It is important to realize that the volume of HLW (mostly from spent fuel) is small and it can be stored relatively easily, until final disposal is possible. The HLW tends to be short-lived due to its greater radioactivity, whereas LLW could present much longer-term hazards. It is the current generation’s moral obligation to ensure that future generations are not burdened with radwaste issues. The future will demand national political solutions to address the final disposal of radwaste, since technical issues regarding managing, conditioning and packaging have been largely

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resolved. Radwaste and other materials arising from decommissioning and dismantling NPPs may be disposed of finally in deep (400 m) repositories or could be stored in such a manner that the option remains to monitor or even retrieve it even after several hundred years. This may be a viable option when new methods are developed to transmute long-lived isotopes into shorterlived ones or even non-radioactive ones [14]. Although an expensive option, nuclear transmutation of transuranium elements (actinides) such as isotopes of plutonium, neptunium and curium can potentially reduce the inventory of these elements and thus facilitate easier radwaste management. It can be mentioned here that earlier NPP designs often used components made from hard cobalt-containing alloys for applications such as rollers and pins in control rod systems that require very high resistance to wear and freedom from sticking through friction. However, when irradiated with neutrons, 59Co (the stable form) may become activated to form the relatively long-lived isotope 60Co, which is a gamma-radiation emitter with a half-life of 5.27 years. Thus, due to ALARA principles and to lessen activity inventories, Co-containing alloys have now been gradually replaced with other materials not so susceptible to activation. This is an example where experience gained with radiological aspects (and not SSC-AD) has led to the phasing-out of Co-containing alloys used for some components in NPPs.

2.8

Past, current and future nuclear power plant (NPP) concepts and designs

Since Generation I NPPs have been mostly phased out, and Generation II NPPs are discussed in detail elsewhere in this book, the following will deal only with salient features, characteristics and special aspects of some Generation III NPPs and those projected to come into commercial service further along the timescale until about 2030. Since the Generation IV NPPs are a group of theoretical NPP designs for next-generation technologies projected to be available after about 2030, they are beyond the scope of this book. However, those currently under study are the gas cooled fast reactor (GFR), the very high temperature reactor (VHTR), the supercritical water cooled reactor (SCWR), the sodium cooled fast reactor (SFR), the lead cooled fast reactor (LFR) and the molten salt reactor (MSR). Nuclear power plants may be categorized as follows, reflecting their historical evolution and technical development and it should be noted that quite different reactor types exist within each reactor generation: ∑

Generation I (1950s–1970s) NPPs. These were essentially the early prototypes of todays operating NPPs. With the current exception of the United Kingdom (with their carbon dioxide cooled, graphite moderated Magnox-type reactors, currently scheduled to be shut down by 2010), none are in operation elsewhere in the world. © Woodhead Publishing Limited, 2010

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∑

Generation II (1970s–1990s) NPPs. Reactors from this era comprise the majority of commercial NPPs operating today, and some of these NPPs were even constructed using modified nuclear submarine propulsion systems and associated technologies, for example, the PWR. Other designs in this Generation II category are the BWR, CANDU (CANada Deuterium Uranium), WWER (sometimes called VVER) Russian designed PWR, RBMK (Russian designed channel-type graphite moderated reactor) and the advanced gas-cooled reactor (AGR). ∑ Generation III (1990s–2010) NPPs. These reactors have benefitted widely from lessons learned from older NPP types, designs and technologies. They thus may be classed as being at the current state-of-the-art, science and technology and advanced in concept. The EPR, System 80+, AP 600 and ABWR NPPs are examples of the light water-based Generation III NPPs (see below for some further details). ∑ Generation III+ (2010–2030) NPPs. These reactors will have shorter construction times and lower capital costs. They have evolutionary designs and will operate even more efficiently, thus more economically than currently operating NPPs. Examples of this generation are the Westinghouse, AP 1000, APWR and ESBWR NPPs (see below for some examples). ∑ Generation IV (from around 2030 onward) NPPs. Most of the designs are still in the conceptual stage. The NPPs will be characterized by lower construction and operating costs (generically-based and standardized approach), they will have enhanced safety systems, many of them passive, and due to the innovative fuel cycles they will be resistant to proliferation issues and will generate less volumes of radwaste.

2.8.1 Generation III and III+ NPPs The 1600 MWe European pressurized water reactor (EPR) is a Generation III type of NPP, meaning it is derived from the present designs of NPPs, but features evolutionary improvements in design (standardized), including many passive safety systems, better fuel efficiency and materials having high resistance to AD. The EPR features a 1.5 m thick concrete outer wall and a substantial inner containment. Another Generation III design is the advanced boiling water reactor (ABWR), four of which have actually operated in Japan since 1996 and two more are under construction there. The ABWR design is licensed in the US, Japan and Taiwan (with two currently under construction in Taiwan). Since the ABWR is based on many years of BWR experience, it has benefits concerning the construction time (e.g. 39 months from first concrete to first fuelling, reduced operation and maintenance costs and improved safety and operability features, for example.

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The economic simplified boiling water reactor (ESBWR) is a Generation III+ NPP and is a further development on the ABWR, featuring natural circulation with no recirculation pumps or their associated piping. As its name implies, the ESBWR is simplified (e.g. 25% fewer pumps) and presumably more economical to construct and operate due to a projected short (36-month) construction time. These designs can remain safe in accident conditions without the use of any active control components. In principle, it is better to have passive safety features, not depending on other systems, because gravity feeds, evaporation and natural circulation, for example, cannot be affected by inadequate human actions or procedures during an emergency situation. Designs also even feature improved capability to deal with the unlikely rare occurrence of severe accidents, such as when the nuclear fuel core melts through the RPV and gains entry into the containment. The so-called ‘corium’ or fuel-containing material, a mixture of fuel, debris and molten metal, can fall onto a ‘core catcher’ and be thus spread out efficiently and cooled so that its progression is stopped and the containment remains tight, thus maintaining its DID design requirement to prevent the uncontrolled release of radioactive substances to the environment. Depending on the type, the new generation of NPPs are designed to have higher fuel enrichment (5% uranium oxide or mixture with uranium–plutonium oxide) and thus have a higher electrical power output (e.g. the 1600 MWe EPR) compared to earlier design NPPs (e.g. typically 220–1000 MWe WWER, PWR and BWR, with the exception of the Generation II 1500 MWe RBMK NPPs which were operated at Ignalina, Lithuania until December 2009).

2.8.2 Advanced reactor technologies: the pebble bed reactor A possible future NPP will be the ‘helium cooled pebble bed reactor (PBR)’, which has been studied and developed further in South Africa since 1993. A pioneer PBR operated for 21 years in Germany [15]. Future PBRs will use the equivalent of about 450 000 tennis ball-sized mini-reactor cores contained in a vessel. The nuclear fuel is enclosed in pyrolytic graphite, which acts as a moderator, and the fuel (e.g. uranium, plutonium as oxides or carbides) is made of micro-particles, coated with silicon carbide for structural integrity. Other coolants can be nitrogen or carbon dioxide. The high temperature PBR will have increased thermal efficiency (about 50%), compared to that usual in water-moderated NPPs (about 35%). Helium flows through the ‘pebble-bed’, extracting heat, and then the heated gas is directed onto turbines to produce electricity. The helium cools down and is recycled back to the reactor. Compared to the 1600 MWe EPR, these PBRs are of low power, typically 110 MWe. Operating risks are deemed to be very small, due to the low power density and inherent passive safety

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features. Since it works at high temperature (approximately 900 °C), the PBR can cool itself via natural circulation and withstand accident scenarios up to at least 1600 °C. The PBR is classed as a type Generation IV very high temperature reactor (VHTR).

2.8.3 Future reactors: an example from the United Kingdom The UK’s current fleet (2010) of 19 NPPs of various ages, designs and capacities supply about 16% of the country’s total electricity. However, the majority of Magnox and AGRs are now nearing the end of their original design lives, and about half of the NPP fleet are already in, or planning, decommissioning. Up to 35% of the total supplied electrical power may have to be replaced by 2020; by this date the share of nuclear-supplied power will have dropped to about 7% unless the construction of Generation III and III+ reactors is implemented before this time. This is a medium-term possibility, since in contrast, Generation IV reactors are a longer-term prospect (probably available from 2030 onwards). The Generation III+ reactors will feature modular construction, simplified design, passive safety (use of gravity, natural circulation and evaporation, instead of pumps, motors and valves) and the overall construction costs should be low. The Generation III+ reactors will be more efficient and produce much less radwaste compared to earlier NPPs. By July 2007, four designs had been approved for further assessment, namely, the European AREVA-EPR (pressurized water), the US Westinghouse AP 1000 (pressurized water – a feature of the AP 1000 reactor is that the core coolant is forced into the core by gravity and natural convection recirculates it), the Canadian CANDU ACR 1000 (heavy water reactor) and the US General Electric ESBWR (boiling water) reactor. However, on 4 April 2008, the Atomic Energy of Canada Limited (AECL) Company, responsible for the ACR 1000 design, announced that it was withdrawing their CANDU ACR 1000 from the prelicensing process (called ‘Generic Design Assessment, GDA’). The design had successfully completed Step 2 of the GDA (initial assessment carried out by the Health and Safety Executive (HSE) and the Environment Agency) in the UK, but the announcement stated that the decision to withdraw would allow the AECL to focus its marketing and licensing resources on the Canadian market [16]. On 12 March 2008, the UK’s HSE and the US Nuclear Regulatory Commission (NRC) signed an agreement to renew co-operation between the two countries. The five-year agreement covers the exchange of technical information and the development of safety standards. It also allows the two organizations to exchange personnel, information and training. It is noted that two of the new reactor designs under consideration, and being taken through

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GDA in the UK, are from the US, being the General Electric ESBWR and the Westinghouse AP 1000. On 27 March 2008 the French and UK governments announced that they would be working together to improve the efficiency and effectiveness of nuclear development projects, including safety and pre-licensing (GDA) [17]. In particular, the collaboration would also address matters regarding the regulatory assessment of EPRs which are currently (2010) being constructed in France and Finland and undergoing GDA in the UK. It can thus be seen that the UK is collaborating closely with both European and US designers of the future NPPs.

2.8.4 Future AM and PLiM programmes applied to new generation NPPs As new generation NPPs go into operation, it can be expected that even their SSCs will also undergo some level of AD. It is, however, not expected that well-known AD mechanisms such as neutron embrittlement of the RPV beltline material or stress corrosion cracking in steam generators or vessel head penetrations will pose any significant problems, since appropriate design, manufacturing processes, material and operational aspects and lessons learned will have been implemented. The new generation NPPs will generally feature higher working temperatures, greater power rating, higher flow rates of coolant but less reliance on engineered safety features due to the passive design concepts present. For example, as temperatures will be much higher in the helium-cooled PBR, ageing mechanisms such as creep and evolution of microstructures, carburization, decarburization, oxidation and the effect of particulate-laden gases in causing corrosion erosion must all be items to monitor and control in ASPs, AM and PLiM programmes. Research work should continue to investigate the mechanisms acting and how best they can be mitigated or eliminated. Special surface treatments could conceivably reduce ingress of impurities and associated AD. Concerning microstructure effects, it is known that some grades of duplex cast austenitic stainless steel (CASS) used for pump casings and piping may suffer thermal embrittlement due to the development of metallic phases over time within a given temperature range [18]. It is therefore important to allow for such a possibility in new NPPs using appropriate monitoring and surveillance programmes. New tools may need to be developed or existing ones improved upon, to detect microstructure evolution with time and to correlate this with the propensity to corrode or embrittle. Other aspects concerning new generation reactors will be to determine how far risk-based inspection procedures may be used and what level of regulatory acceptance they will find. Risk-based inspection is still being developed for the current fleet of NPPs, so it remains to be seen how far this can be applied to the new generation of NPPs.

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Increased reliance on computerized plant control will necessitate the highest QA of computer programmes and sensors. Sensors must supply the correct information to computer-controlled functions to allow correct operational decisions to be made, and AD monitoring must be extended to cover these technologies. Ageing and obsolescence of sensors and issues of equivalent QA replacements must be included in the OPs, ASPs, AM and PLiM programmes as appropriate. Plant personnel will have to be trained not only to respect electronically supplied plant information, but also not to have total reliance on it, and to question seemingly anomalous signals and plant response and to be able to instigate cross-checks for verification of the true situation [6]. Trust in instrument readings is good, but intelligent human control is better. There has been a gradual decline in the numbers of qualified workforce to regulate, construct and operate NPPs since the mid-1980s, and many pioneer nuclear power workers are already in, or going into, retirement. However, it remains problematic to find and recruit young qualified people to replace them, since corresponding courses have become less available in many universities, thus compounding the problem. The young generation of nuclear power workers are now being faced with the task of keeping current NPPs operating safely, whilst confronting the challenge to build and operate new generation NPPs. The standard OPs, ASPs, AM and PLiM programmes for the new generation NPPs may well have their basic principles based on established and well-proven approaches and methodologies, but this should not nurture complacency. Each new NPP type will require analysis and specific OPs, ASPs, AM and PLiM programmes tailored to suit. However, the new NPPs are going to be constructed on a modular and generic basis, thus facilitating, in principle at least, the direct transfer of basic knowledge and operational principles.

2.9 Conclusions 1. Ageing management and associated programmes assist integration between existing NPP standard operational practices (e.g. maintenance, repairs, replacements and ISI) and also address potentially insufficient conservative design assumptions that have become evident over a NPP’s operational time; AM programmes are basically more safety-oriented in nature. 2. Plant life management (PLiM) programmes facilitate a comprehensive evaluation of the NPP assets and contain elements of both safety and economics; PLiM programmes are basically more economically oriented in nature. 3. It is vital to assign the various NPP-SSCs to their respective safety classes to obtain a focused and objective approach to the OPs, ASPs,

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AM and PLiM programmes, and that resources are used where they are most needed. 4. System, structure and component ageing surveillance dossiers are living documents, to be updated whenever new knowledge becomes available regarding SSC-AD. 5. Diverse AD effects have been observed in all types of NPP SSCs. Understanding the mechanisms and causes of degradation allows implementation of robust mitigation methods. The information from ISI, testing, normal and exceptional maintenance and research remain key to achieve this. 6. Changes to the NPP, such as power uprating, must be assessed in terms of their holistic impact on OPs, ASPs, AM and PLiM programmes, including provision for adjusted emergency procedures. 7. Obsolescence in SSCs and replacements and strategic reserves thereof must be addressed in the PLiM programme in a timely manner. Depending on safety class or reliability issues, new types of SSCs will require qualification for functional equivalence. Owners, operators and regulators must be aware of non-service (storage) AD, and optimum management, elimination or mitigation thereof. 8. New generation NPPs feature many passive safety features and improved designs and materials, but knowledge from current OPs, ASPs, AM and PLiM programmes, modified as necessary, will still provide a basis to achieve these NPP’s design lives safely and profitably and lay the foundations for their LTO. 9. Safety culture and knowledge management enhance the effectiveness of standard OPs and ASPs, and AM and PLiM programmes. 10. Risk-based inspection methods and probabilistic approaches will continue to supplement deterministic methods and their further development is necessary to achieve the overall goal of operating NPPs safely and reliably.

2.10 Sources of further information NEA/CNRA/R, ‘Regulatory aspects of ageing reactors’, OECD Nuclear Energy Agency, CNRA Special Issue Meeting, June 1998. NEA/CNRA/R (99) 1, March 1999. IAEA, ‘Safety aspects of nuclear power plant ageing’, IAEA-TECDOC-540, Vienna, 1990. IAEA, ‘Cost drivers for the assessment of nuclear power plant life extension’, IAEA-TECDOC-1309, Vienna, September 2002. Penny, R.K. (ed.), Ageing of Materials and Methods for the Assessment of Lifetimes of Engineering Plant, CAPE ’97, A.A. Balkema, Rotterdam, Brookfield, 1997.

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Penny, R.K. (ed.), Risk, Economy and Safety, Failure Minimization and Analysis, Failures ’98, A.A. Balkema, Rotterdam, Brookfield, 1998. USNRC, ‘generic ageing lessons learned (GALL) report’, USNRC-NUREG1801, Volume 1, July 2001. IAEA, Site selection and evaluation for nuclear power plants with respect to population distribution: a safety guide’, IAEA Safety Series No. 50SG-54, STI/PUB/569, 1980, IAEA, Vienna, 1980. Major Projects Association, ‘A new generation of uk nuclear power plants – are we ready?’ Seminar No. 123, held at the Institution of Civil Engineers, London, February 2006. IAEA, ‘Periodic safety review of nuclear power plants – safety guide’, Safety Standards Series No. NS-g-2.10, IAEA, Vienna, 2003. Tjernlund, R.M. and Manacsa, G.C., ‘Guidelines for the safety classification of systems, components and parts used in nuclear power plant applications (NGCIG-17)’, Electric Power Research Institute, Palo Alto, CA, EPRINP-6895, OSTI ID: 6372758, 1 February 1991. Männisto, I., ‘Risk-informed classification of systems, structures and components in nuclear power plants’, Master’s Thesis, Helsinki University of Technology, Department of Engineering, Physics and Mathematics, Espoo, 2005. (available at: http://www.sal.hut.fi/Publications/pdf-files/ TMAN05.pdf). Radiation Dose Management in the Nuclear Industry, proceedings of the conference organized by the British Nuclear Energy Society, Windemere, Cumbria, 9–11 October 1995, Thomas Telford, London, 1995. IAEA, ‘Severe accident management programmes for nuclear power plants – safety guide No. NS-G-2.15’, IAEA, Vienna, July 2009.

2.11

Acknowledgements

The author wishes to thank B. Raj of the Indira Gandhi Centre for Atomic Research, Kalpakkam, Tamil Nadu, India, for his help and suggestions, and for reviewing this chapter. Thanks are due also to P. Contri of JRC-Petten/ NL for his valuable comments regarding safety classification and the roles of OPs, ASPs, AM and PLiM programmes.

2.12

References

1. IAEA, Plant life management for long term operation of light water reactors. Principles and guidelines, TRS 448 IAEA, Vienna, Austria, October 2006. 2. Bond, L., ‘Economics of plant life management’, Second International Symposium on Nuclear Power Plant Life Management, Shanghai, China, IAEA, 15–18 October 2007. 3. STUK – Radiation and Nuclear Safety Authority of Finland, Nuclear power plant systems, structures and components and their safety classification, 26 June 2000

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(source extracted from STUKLEX, http://www.edilex.fi/stuklex/en/lainsaadanto/ saannosto/YVL2-1). 4. IAEA Safety Standards Series: Periodic Safety Review of Nuclear Power Plants, Safety Guide No. NS-G-2.10, STI/PUB/1157, IAEA, Vienna, Austria, August 2003. 5. Medvedev, G. (1989). The Truth about Chernobyl. VAAP. First American edition published by Basic Books, New York, in 1991. 6. US Nuclear Regulatory Commission (NRC) Fact Sheet on the Three Mile Island accident. Annual report – 1979 NUREG–0690. 7. 10 CFR 50.65, Requirements for Monitoring the Effectiveness of Maintenance at Nuclear Power Plants. 8. Tipping, Ph., ‘Some power up-rate issues in nuclear power plants’, Nuclear Engineering and Technology, Vol. 40, No. 4, 2008, pp. 251–4. 9. The New Economics of Nuclear Power, World Nuclear Association (WNA) Report (http://www.uic.com.au/neweconomics.pdf). 10. US Nuclear Regulatory Commission, ‘Failure of a steam dryer cover plate after a recent power up-rate’, Information Notice IN 2002-26, September 2002. 11. IAEA, ‘Implications of power up-rates on safety margins of nuclear power plants’, Report of a IAEA and OECD/NEA Technical Meeting, Vienna, Austria, 13–15 October 2003, IAEA-TECDOC-1418, IAEA, Vienna, 2004. 12. IAEA, ‘Overview of training methodology for accident management at nuclear power plants’, IAEA-TECDOC-1440, IAEA, Vienna, 2005. 13. Title 10, ENERGY, Chapter I: Nuclear Regulatory Commission (NRC), Part 70 – Domestic Licensing of Special Nuclear material, subpart d – License applications, 70:24 ‘Criticality accident requirements’, paragraph 50.68. 14. Takibayev, A., Saito, M., Artisyuk, V. and Sagara H., ‘Fusion-driven transmutation of selected long-lived fission products’, Progress in Nuclear Energy, Vol. 47, 2005, pp. 354–60. 15. Association of German Engineers (VDI), AVR – Experimental High-Temperature Reactor, 21 Years of Successful Operation for a Future Energy Technology, The Society for Energy Technologies, Dusseldorf, VDI Verlag, 1990, pp. 9–23. 16. AECL, ‘CANDU-ACR-1000 withdrawn from the generic design assessment’, reported at http://www.aecl.ca/NewsRoom/News/Press 2008/080404/.htm?ebul=newreactor/1may-2008&cr=2. 17. UK Health and Safety Executive, ‘UK–French collaboration on generic design assessment for EPR’, reported at: http://www.hse.gov.uk/newreactors/communique. htm?ebul=newreactor/1-may-2008&cr=3). 18. Pohl, M., Storz, O. and Glogowski, T., ‘Effect of sigma phase morphology on the properties of duplex stainless steels’, Microscopy and Microanalysis Volume 11 (Supplement 2), 2005, pp. 230–1.

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Safety regulations for nuclear power plant life management and licence renewal

A. A l o n s o, Universidad Politécnica de Madrid, Spain

Abstract: This chapter presents current regulations to ensure safety and reliability of ageing nuclear power plants. Management of ageing in structures, systems and components requires an understanding of ageing mechanisms and how to detect and control them through surveillance, feedback of operating experience and maintenance programmes. Nuclear power plants’ longer-term operation needs to be addressed by designers, plant operators and regulators within the umbrella of international organizations to implement, complete and augment research and enact new regulation. Key words: nuclear safety regulations, ageing mechanisms, life management, longer-term operation, role of international organizations.

3.1

Introduction

Nuclear power plant (NPP) longer-term operation (LTO) has to demonstrate compliance with the current licence basis and be founded on a completely robust and satisfactory set of regulations. Compliance with such regulations has to be proven by the applicant and verified by the regulatory organization through well-established licensing procedures. Those activities will be first described for some representative cases. Such regulations require the establishment and conducting of a specific plant life management programme (PLiM), which is based in the understanding, monitoring and management of the ageing processes in structures, systems and components (SSCs). Furthermore, PLiM programmes are continually being improved upon. Although ageing processes may have a generic nature and affect similar designs, the ageing status of each single nuclear power plant, and the corresponding remaining safe life, should be known through plant surveillance, operating experience feedback, maintenance and inspection programmes considering the specific design, operational history and the functional experience of each individual plant. The ageing processes are considered on an introductory basis, as they will be described in detail in other chapters of this book. Different ageing mechanisms shorten the life of the SSCs. Some of those will be part of the safety envelop, while others will be installed in the non-nuclear side. The first will receive major attention without neglecting the others, mainly when they may affect the safety and reliability of operations. PLiM mainly 56 © Woodhead Publishing Limited, 2010

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concerns not only operators, but also designers and regulators. Old designs did not properly consider all possible ageing mechanisms; in fact, on average, a new significant ageing effect has appeared every seven years or so. The ageing of such plant has to be managed while in service to keep the required safety level at all times and also during LTO, while new designs can prevent some of the known ageing mechanisms and facilitate improved managerial procedures. The major areas of concern are related to neutron fluence and stress corrosion cracking, which will be presented in brief, as they will be considered in depth in later chapters. PLiM and LTO are still evolving, so more research is still needed to understand and prevent ageing mechanisms. Information networks, to be discussed later on, have been created globally to evaluate and apply worldwide operating experience and research results. Periodic examination will gradually be substituted by the more effective on-line surveillance; improved digital technology will help in that endeavour. This and other future trends will be presented. A list of references and identification of sources for further information are offered to the reader.

3.2

Safety review/licence renewal

Original NPP designs were created using the deterministic approach based on the defence-in-depth principle, as defined by the International Nuclear Safety Group (INSAG, 1996). Although economic considerations, rather than technical reasoning, were the basis for determining the lifetime of nuclear power plants, pessimistic early approaches limited the operation of light and heavy water reactors to 30–40 years. The regulatory authorities of supplier countries formally established such limits in their operation licences. In the USA 40 years was the standard, while in the USSR 30 years was the limit. Such restrictions where generally copied by the importing countries of such technologies. In other supplier countries such restrictions were not formally applied. In Germany, for example, the plants could operate as long as they could prove their required safety level. The concept of periodic safety review (PSR), to be conducted every ten years on average, started to take root in Germany as well as in many other countries, notably in Spain. That concept did not put any formal life limit to the power plants, provided they could show their safety at any time and after the PSR. The analysis of operating experience clearly showed that the early fatigue considerations in metallic components were not a matter of concern even after 40 years of operation. At the same time, new ageing mechanisms started to appear. Some of them, mainly new forms of corrosion and embritlement by neutron irradiation, clearly indicated that the operating life of some components would have to be limited unless proper corrective actions were taken. These experiences initiated new activities related to surveillance, maintenance and

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operational experience feedback, which later on were integrated in PLiM. This was not only for safety reasons, but also to improve the performance and economics of the plants. Thus the economical advantages of keeping such plants in operation beyond the formally established lifetime soon became evident. Under the pioneering effort of the US Nuclear Regulatory Commission (NRC), those countries with old designs started PLiM activities and the regulatory organizations commenced developing regulations for LTO. International institutions, mainly the International atomic Energy Agency (IAEA) and the Organisation for economic Development/Nuclear Energy Agency (OECD/ NEA) established working groups to consider PLiM and LTO.

3.2.1 IAEA guidance on longer-term operation The IAEA has performed valuable work on PLiM and it is providing guidance on ageing-related regulatory aspects through its Safety Standards Series and Technical Report Series (IAEA, 1992). Although the IAEA has no regulatory authority, it has the statutory function of establishing standards for the protection of individuals, society and the environment in the use of nuclear power for peaceful purposes; such standards could form the basis and guide regulatory development in Member States. So far, LTO has been briefly included in the on-going revision of the IAEA requirements for safety in operation (IAEA, 2000). A current revision of such requirements (IAEA, 2008) includes a new chapter on longer-term operation, equipment qualification and ageing management. The new document defines LTO as: ‘The operation beyond an established time frame set forth in design standards licence and/or regulations that is justified by safety assessment considering processes and features that may limit the life of SSCs’. It adds that: ‘… it shall use the results of periodic safety review (PSR) and approved by the regulatory body on the basis of the analysis of the ageing management program that will assure plant safety during its extended life time’. The reference to PSR is of particular importance because many countries have established such regulatory practice as a means to license current and longer-term operation. The document further requires the establishment and implementation of a comprehensive LTO programme to be presented and approved or agreed to by the regulatory body. The LTO programme shall include such aspects as the current licensing basis (CLB); safety upgrading and verification and operational programmes, scoping and screening of SSCs with respect to degradation and ageing; re-validation of safety analysis that use time-limited assumptions; review of the management programmes, and the implementation of a new ageing management programme for the expected LTO period. As the IAEA document belongs to the requirements category, it should be expected that its contents will be developed into more specific safety guides.

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In the absence of more detailed guidance, the IAEA understands that PSRs provide a comprehensive methodology for reviewing the safety of the plant, including ageing effects and regulatory changes, technical progress and technological obsolescence. The IAEA has addressed such reviews (IAEA, 2003a): The objective of a PSR is to determine by means of a comprehensive assessment of an existing nuclear power plant: the extent to which the plant conforms to current international safety standards and practices; the extent to which the licensing basis remains valid; the adequacy of the arrangements that are in place to maintain plant safety until the next PSR or the end of plant lifetime; and the safety improvements to be implemented to resolve the safety issues that have been identified. The guide recognizes that the PSR can be the basis for acceptability of longer-term operation. It also recognizes that compliance with the current safety standards may not always be possible, but suggests that: ‘practicable improvements should be made as steps towards meeting them’ and ‘that the risk associated with the shortcomings be assessed and that a justification for continued plant operation be provided’. Many countries have adopted this formal regulatory path, although they are also considering the more detailed practices established by the US NRC.

3.2.2 Regulations for longer-term operation in the United States Regulations for longer-term operation have been fully developed and practised in the United States and they have set the examples for other countries. The US NRC decided to limit to 40 years the operating licence of NPPs, based on economic and antitrust considerations. The technical design was based on the assumed lifetime of the plant using very conservative assumptions to cover the then fragmentary knowledge on the expected behaviour of SSCs in service. Further safety analyses based on research results and operating experience feedback have clearly demonstrated that NPPs could operate safely and reliably well beyond the established formal licence limit. After recognizing this fact, the US NRC decided to enact 10 CFR Part 54 (NRC, 2006) to cover up 20 additional years of operation. Plants with 20 or more years of operation are candidates for licence renewal up to a total of 60 years. As of May 2009, there were 104 NPPs licensed to operate in the Unitted States, 52 units had received new licences, 14 units were under review and letters of intention to apply for a new licence were submitted to the US NRC for 52 units. It is expected that a very limited number of operating plants will not apply for a new licence due to individual unfavourable cost–benefit analyses or other causes.

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10 CFR Part 54 requires the licensee to submit some new documents: ∑ ∑

an integrated plant assessment (IPA); changes in the current licence bases (CLB), taken place during the NRC review of the application; ∑ an evaluation of time-limited aging analyses (TLAA); ∑ a supplement to the final safety analysis report (FSAR).

The interest of the regulator is based in the applicant knowing, managing and maintaining ageing within the plant safety envelope. The IPA must identify a complete list of those structures and components requiring ageing management justification for LTO. They are structures and components that perform their intended safety function without moving parts. The regulations include, but are not limited to, a lengthy list of long-lived structures and metallic, electric and electronic components, such as the reactor pressure vessel, the steam generators, pipes, supports and seismic Category 1 structures and cables, among others. It will be necessary to prove that those structures and components could perform their intended safety functions during LTO in the way established within the CLB. The applicant is obliged to submit changes to the CLB that may have taken place during the NRC review of the application. Submissions of such changes must take place each year after the application for a new licence and three months before the scheduled completion of the NRC review. All changes in the CLB of the facility affecting the contents of the new licence must be included in such amendments. TLAA includes the observations, calculations and analyses that the licensee has conducted to predict the effects of ageing on SSCs considered relevant and included in the CLB. It has to be proven that the analyses remain valid, have been projected for the period of additional operation or the effects of ageing can be managed and will not limit the safety functions. Among others, typical TLAA include metal fatigue, environmental qualification and neutron embrittlement. The FSAR supplement must describe the programmes and activities for managing the effects of ageing and the evaluation of TLAA for the new period of operation included in the IPA and in the ageing analyses. The applicant must also include any change or addition to technical specifications of the plant, with the appropriate justification, which may be necessary to manage the effects of ageing. Enacting any new basic regulation, as in this case, requires from the regulator the development of detailed procedures for better compliance. To that effect, the US NRC published Regulatory Guide 1.188 (NRC, 2005a) describing in more detail an acceptable format and content of the documents that justify the application. The Guide endorses a previously developed industry guidance document prepared by the Nuclear Energy Institute (NEI).

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The Regulator also edited a Standard Review Plan to License Renewal (NRC, 2005b) explaining how the US NRC staff should perform the evaluation to ensure that all the items have been properly considered and that the affected SSCs will perform their safety functions as expected during the period of longer-term operation. This document also serves to assure quality and uniformity in the review process. The Regulator has also published an additional document on the generic lessons learned, the so-called GALL report (NRC, 2005c). It includes the technical basis for the Standard Review Plan based on a systematic compilation of ageing information in nuclear power plants. The Gall report constitutes a helpful document to the applicants, since it contains recommendations on those existing ageing programmes that should be augmented for licence renewal. It also serves as a reference for the US NRC review process, and such a review is not necessary when the applicant’s ageing programmes follow those given in the GALL report.

3.2.3 Longer-term operation in other countries Most countries with nuclear power plants have a common approach to regulatory requirements related to PLiM and LTO. Although only a few have established time limits to the operation licence, others have not formally established such limits. The countries who have not established a limited lifetime have enforced PSR, about every ten years, which in many cases, as in Spain, are linked to a formal operation licence for the next ten-year period with no formal limit on the number of renewals. There are also countries with formal operation time limits that also practise PSR; in such cases, the review is reinforced by ageing management programmes (AMPs) and TLAAs. Despite the different models, in all countries the responsible licensee also maintains a continuous evaluation of the safety level of each plant under close supervision from the national regulator. Table 3.1 indicates how LTO is managed in some representative countries. Countries using PSR as the formal way to justify longer-term operation may differ in some specific aspects, but within a core of fixed principles such as: compliance with the current national safety requirements and guides and international recommendations; identification of improvements required to achieve or justify compliance; revision of the PLiM programme to ensure safe and reliable operation during the next ten-year period and definition of the necessary surveillance, maintenance and ageing mitigation measures. The Western European Nuclear Regulators’ Association (WENRA) has studied the level of harmonization of safety regulations followed in 17 countries (WENRA, 2006); by referring to the IAEA reference levels as guidance, it has concluded that only half of them have established formal requirements for PSR and ageing management, although the majority of them implement

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Foreseen life extension (y)

Main regulations for licence renewal

Use of periodic safety review

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Argentina Plant dependent Undetermined

Canadian regulations applied to Embalse NPP

Canada 40 10–20

CNSC Draft, 6-360, ‘Life Extension of Nuclear Power Plants’ IAEA TECDOC 1503, PliM, Guidelines for HWRs

Practised, but not directly linked to LTO, based on: IAEA ‘Periodic Safety Reviews of Nuclear Power Plants’ NS-G2.10 (2003)

Czech Republic As given in the Undetermined Technical Certificate

Proposed development of regulatory requirements based on US NRC practices and IAEA recommendations

PSR applied, based on IAEA NS-G-2.10

France No legal limit 10+

Industry procedures based on IAEA and NEA documents presented to regulatory authority for groups of standard plants

Practised by groups of standard plants and linked to licence operation. Part of the TSN Act

Germany No legal limit Not foreseen

Utilities have initiated AM programmes in accordance with RSK recommendation (2004) KTA 2301, Ageing management of NPP

Based on periodic safety reviews Requires deterministic and probabilistic analyses

Hungary 30 20 Regulations based on 10 CFR Part 54

Practised but not part of licence renewal. PSR based on IAEA NS-G-2.10

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Country Legal design life (y)

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Table 3.1 Regulatory requirements for LTO in different countries

Coordinating Committee Practised, linked to licence on Ageing Management renewal (Industry, Academy, Regulator)

Mexico 30 Under consideration Adoption of US NRC regulations. Participation in EPRI BWR-VIP and in IAEA co-ordinated research projects © Woodhead Publishing Limited, 2010

Netherlands 30 60 No specific policy on LTO (Borssele) IAEA AMAT Review

Practised every 10 years, weakly linked to licence renewal

Russia 30 5+10

NP-017-2000, ‘The main requirements for extension of the operating life of a NPP unit’ NP-24-2000, ‘Requirements for Not considered substantiation of possibility of extension of the operating life of nuclear facilities and installations’

Slovak Republic 60

Safety guide BNS 1.9.2/2001, ‘Ageing Management of NPPs’ Requirements’

Practised

Practised and linked to licence renewal. Based on IAEA 50-SG-012

South Korea No legal limit. 10 + (unlimited) Enhanced PSR developed in 30 y considered MOST Bulletins 2005-31 acceptable In depth review of 11 safety indicators (NUREG-1801)

Practised (since 2000) and linked to licence renewal, based on IAEA 50-SG-012

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Slovenia 40 10+ Regulations based on 10 CFR Part 54

Safety regulations for nuclear power plant life management

Japan 30 10+

Foreseen life extension (y)



Main regulations for license renewal

Use of periodic safety review

KINS/GE-N8 Safety review guidelines for PWR

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Spain 40 10+ Regulations based on 10 CFR Part 54

Prastised and linked to licence renewal. Instruction IS-22

Sweden 40 Not foreseen

SKIFS 2004:1, ‘Regulations concerning safety in nuclear facilities’ ‘SKIFS 2005:2, ‘Regulations concerning mechanical components in nuclear facilities’

USA 40 20

10 CFR Part 54, ‘Requirements Not practised for Renewal of Operating License for Nuclear Power Plants’, amended 2006

Note: This table has been constructed mainly from information presented by licensees and regulators in the IAEA 2nd International Symposium on Nuclear Power Life Management, Shanghai, 15–18 October 2007.

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Table 3.1 Continued

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such practices. In 2007, as a counterpart to WENRA, the European utilities holding nuclear power licenses decided to create the European Nuclear Installations Safety Standards (ENISS) to contribute to the harmonization efforts by providing comments and suggestions to WENRA’s activities. In Spain, for example, the nuclear safety regulatory authority, Consejo de Seguridad Nuclear (CSN) has enacted Instruction IS-22 (CSN, 2009) as the guiding document for PSR. The Guide explicitly recognizes: ∑

the importance of the feedback from national and international operating experience and how has it been applied to the plant to be analysed; ∑ the need to contemplate regulatory changes in the country of origin of the project or new requirements suggested by international organizations and their applicability to the plant in question; ∑ the full understanding of the ageing mechanisms and how to manage their effects on the SSCs to ensure the proper accomplishment of their safety functions; ∑ the value of any design or operational change, either already performed or in planning, to improve safety; and ∑ the potential use of the probabilistic methodology, which is also recommended in the IAEA PSR guidance.

3.3

Surveillance, operation and maintenance programmes

The early observations on SSC ageing determined the premature establishment of specific surveillance, operation and maintenance programmes that have been later integrated into PLiM. Surveillance starts with equipment verification and includes monitoring the behaviour of SSCs as well as regulatory inspection and peer reviews. Operation includes training and the feedback of operating experience; while maintenance, mainly risk informed maintenance, constitutes a substantial part of the ageing management programme, AMP. The most significant aspects of these activities are described in this section.

3.3.1 Surveillance of structures, systems and components Equipment verification is a key factor in the TLAA required by US NRC 10 CFR Part 54 and similar standards in other countries. Its importance was soon recognized by INSAG. In an early report (INSAG, 1988) the Group states: ‘Safety components and systems are chosen which are qualified for the environmental conditions that would prevail if they were required to function. The effects of ageing on normal and abnormal functioning are considered in design and qualification’.

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In an IAEA report (IAEA, 1998), recommendations are given for equipment qualification programmes. Safety-related equipment must be qualified not only for normal operating conditions but also for the so-called ‘harsh environments’ such as the ones created by postulated initiating events (PIEs). Such events may create conditions very different from those occurring during normal operation and affect an individual component or different components with the potential of causing common-cause failures in equipment experiencing ageing effects during normal operation. Equipment qualification is measured by using the concept of ‘qualified life’ or ‘qualified condition’ established by equipment qualification. Qualified life is the ‘period of time during normal operation when ageing does not prevent satisfactory performance during a subsequent PIE’. Before the end of its qualified life, the equipment has to be replaced, life-limiting components renewed or a new longer qualified life established. The qualified condition of equipment ‘is expressed in terms of a measurable condition indicator(s) for which it has been demonstrated that the equipment will meet its performance requirements’. The reactor pressure vessel (RPV) is the major component needing a surveillance programme to monitor changes in the fracture toughness properties of ferritic materials in the vessel beltline region which result from its exposure to neutron irradiation and possibly, but to a lesser extent, from the thermal and chemical environment. The US NRC promptly established a comprehensive surveillance programme (NRC, 2008a) based on standards published by the American Society for Testing and Materials (ASTM), which is now common to other light water reactors (LWRs) and heavy water reactors (HWRs). Fracture toughness test data are obtained from material specimens exposed in surveillance capsules, which are withdrawn periodically from the reactor vessel and tested to examine the evolution of the mechanical properties of the vessel material. For each individual vessel a databank has been created on fracture toughness that permits the expected life of the said vessel to be forecast with a good level of certainty. In 1996 the IAEA created a worldwide database to store the results from national RPV surveillance programmes. To date, results from surveillance programmes from 10 countries are included in this database. Access to the database is controlled by agreements with the IAEA.

3.3.2 Operation and operating experience feedback The importance of operation experience (OE) and its feedback (OEF) has been recognized since the early 1970s within the NEA/OECD Committee on the Safety of Nuclear Installations (CSNI). Such a Committee started to share between the participants of NEA Member States information on the construction and operation experience of nuclear power plants. The value of such an exercise was promptly recognized and offered to the IAEA to

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cover a wider audience. In 1980, the IAEA, in collaboration with the NEA, initiated the formal creation of a databank on operating experience, the so-called Incident Reporting System (IRS) (IAEA/NEA, 1980) with the participation of all IAEA Member States. Every country acquires and shares operating experiences in accordance with well-established procedures. The cited international organizations are responsible for the analysis of such experiences to determine tendencies or to understand root causes behind relevant incidents. Periodically the NEA publishes a report analysing the major experiences and any foreseeable tendencies. The last published report (IAEA/NEA, 2006) covers the large number of recent incidents involving corrosion on secondary side piping and recommends that more attention should be given to these ageing-related events. On its side, the industry also has its own programmes for OEF. The World Association of Nuclear Operators (WANO) and the US Institute on Nuclear Power Operations (INPO) have also established similar banks on OE and OEF, including a larger spectrum of incidents. Moreover, there are also national programmes for OE and OEF mainly in those countries operating a significant nuclear power fleet. More recently, INSAG has expressed its high consideration for OEF in a recently published document (INSAG, 2008) where it has suggested that improvements should be established in the international programmes and the need to improve the IAEA/NEA IRS databank. In the report, INSAG states: ‘The focus of OEF should therefore be shifted from mere reporting of safety significant events to a more comprehensive OEF system aimed at capturing even early signs of deterioration in safety. The key criterion should not be whether the safety event is significant, but rather whether the safety lesson is significant.’ It is also recognized that OEF is of paramount importance on drafting reactor operation standards and their revisions. Any new knowledge coming through OEF is used, on one side, to reduce uncertainties and in cases to simplify standards; on the other side, new phenomena may be discovered requiring new standards; this has been mainly the case with the ageing of SSCs. Experience shows that detailed technical standards have to be revised frequently to cover such new discoveries and observations. To help the Member States in establishing a national system for OEF, the IAEA has published a safety guide (IAEA, 2006a) which also emphasizes the advantages in participating in the IRS system. On its side, the US 10 CFR Part 54 was based on the analysis of more than 500 research and operating experience documents related to the behaviour of SSCs in operation. One of the most effective systems to monitor continuously nuclear plant operation has been recently introduced by the US NRC and followed in other countries. The reactor oversight process (ROP) (NRC, 2000b) covers reactor safety, radiation safety and safeguards. For each area some cornerstones have

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been determined against which the performance of the plant is compared. A satisfactory comparison gives a reasonable assurance of plant safe operation. From the safety point, the three major cornerstones are the initiating events, the mitigating systems and the integrity of the defence-in-depth barriers. Any departure from normal conditions is closely analysed to determine root causes. A colour code has been established to give an intuitive pictorial representation of the safety level of the plant. The Spanish Regulatory Authority, CSN, has established a similar process for plant operation supervision (CSN, 2006). Experience has proven that up to 80% of industrial incidents and accidents are due to human error and 20% to equipment failure. This is the accepted 80:20 rule. Moreover, almost 50% of the human errors are linked to poor or non-existing appropriate training on the related matter. The analysis of the Three Mile Island 1979 accident revealed that there were deficiencies in the way that personnel were trained to perform their duties. The Kemeny report recommended the creation of an Institute on Nuclear Power Operations (INPO) to deal with such matters, among other subjects. First INPO, and later on the IAEA and other national and international organizations, created and developed the so-called systematic approach to training (SAT), and recommended that such an approach should be considered as an international standard for nuclear training. The application of such methodology has improved human performance considerably and so increased the safety and reliability of operating nuclear power plants. The IAEA has issued a safety guide on recruitment, qualification and training of nuclear power plant personnel (IAEA, 2002a). In such a document, SAT is considered as the most effective policy for personnel training and gives clear indication on how such a policy should be included within the national training programmes. More recently, the IAEA has recognized that SAT should also be included in all the phases in the life of a nuclear power plant, including PLiM and LTO. These later considerations are in the course of being developed.

3.3.3 Maintenance programmes It was soon recognized that an efficient and complete NPP maintenance programme is necessary to ensure that the current accepted design assumptions and safety margins are maintained and not unacceptably degraded through ageing mechanisms. To support LTO, it was believed that new maintenance programmes should be established to include: the verification of the performance goals, the root cause analysis of failures and the feedback from maintenance experience. To that effect, in 1991, the US NRC enacted the so-called maintenance rule (MR) (NRC, 1991). The NRC has developed guidance for monitoring the effectiveness of the application of the rule (NRC, 1997) in accordance with programmes previously developed by the

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Nuclear Management and Resources Council (NUMARC) (1993). Later on, the US NRC has applied risk informed regulation to the MR (NRC, 2000a) to facilitate its application and to make it more effective, requiring that risk evaluation should be performed in maintenance activities. The MR became effective in Spain in April 1999 when it was included as one of the requirements in the NPP operation licences. Finland, Slovenia and Hungary have also implemented at least part of the MR at their NPPs, while others are thinking about future implementation of the MR or have already developed equivalent methodologies. In any case, the engineering rationale behind the MR makes it attractive to other countries. The IAEA has developed requirements for maintenance, testing, surveillance and inspection of SSCs important to safety. Such requirements have been further developed into a safety guide (IAEA, 2002b), which contains specific references to ageing. The licensee is made responsible for determining which additional maintenance, surveillance and inspections will be necessary as the plant ages. The maintenance, surveillance and inspection programmes should be capable of identifying and monitoring ageing mechanisms of SSCs. The safety reviews will determine which additions to the maintenance, surveillance and inspection programmes should be necessary, particularly the preventive maintenance programme, or when any specific ageing SSC should be replaced or backfitted. In 2003, the European Joint Research Centre, Institute for Energy (JRC-IE) launched a network on the Safety of European Nuclear facilities (SENUF); a working group within SENUF was dedicated to assist Eastern European countries on NPP maintenance with the objectives of reviewing maintenance issues, promoting well-designed maintenance plans for SSCs, supporting the implementation of advanced maintenance approaches and assisting in its implementation. A benchmark study conducted in 2004, collected and evaluated the available and applied maintenance programmes in the European countries, as well as the Russian Federation and Ukraine. It served the basis for a SENUF workshop held in Madrid on maintenance rules and maintenance effectiveness, reported by Contri (2006). The development of the probabilistic methodology and the value of its results have also been used to optimize plant maintenance. One of such approaches is the reliability centred maintenance (RCM) subject which is receiving increasing technical interest. However, it is recommended to take an integrated approach based on the complementary use of deterministic and probabilistic considerations. The probabilistic methodology could serve to improve plant performance and enhance safety levels by ensuring higher reliability and availability of plant equipment and to optimize the maintenance cost by focusing maintenance on SSC in a manner commensurate with their safety significance. The SENUF network, in collaboration with the IAEA, organized a second workshop, held in Petten, on advanced probabilistic

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methods to optimize maintenance programmes. A summary report was produced by Ranguelova and Contri (2006). It is widely recognized that the application of the probabilistic methodology to maintenance requires a probabilistic safety analysis report (PSA) of great quality. The IAEA has given a framework to assure the quality of a PSA for different applications (IAEA, 1999). The American Society of Mechanical Engineers has also published standards for PSA applications in NPPs (ASME, 2003).

3.4

Integration of plant life management

Ageing management has been defined by the IAEA (2007a) as: Engineering, operations and maintenance actions to control within acceptable limits the ageing degradation of structures, systems and components. Examples of engineering actions include design, qualification and failure analysis. Examples of operations actions include surveillance, carrying out operating procedures within specified limits and performing environmental measurements. Therefore, PLiM integrates a series of policies and activities aiming for safe and reliable plant operation. The IAEA has a new standard on ageing management (IAEA, 2009). The IAEA standard addresses both physical ageing of SSCs and technical obsolescence in comparison with present knowledge, current standards and up-to-date regulation. It also covers all stages in the life of a power plant, i.e. design, construction, commissioning, operation, including LTO and decommissioning. The guide recognizes that effective ageing management is in practice accomplished by co-ordinating maintenance, in-service inspection and surveillance, operations, periodic safety evaluation, feedback of operating experience, research and development. The guide includes a systematic approach to the ageing management process based on the known plan-do-check-act (PDCA) wheel, as given in the mentioned guide and summarily represented in Fig. 3.1. The Guide promotes a proactive attitude towards ageing management during every stage in the life of the plant and recommends that regulatory requirements and guides for ageing management should be established and updated, and that compliance should be verified as well as enforced, when necessary. The Guide clearly indicates that the operating organization should be responsible for demonstrating that the relevant issues of ageing that are specific to the plant are clearly identified and documented in the safety analysis report. New designs will have ample opportunities to prevent and mitigate ageing when the plant comes into operation. The Guide gives suggestions on how to manage ageing during design, construction and fabrication and commissioning. For instance, the designers should consider relevant experience

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Plan Ageing Management Programme

Improve AMP

Act Corrective maintenance

Correct unacceptable degradation

ageing science and technology

Check Survey and assess ageing

71

Minimize expected degradation

DO Preventive measures

Check degradation

3.1 The plan-do-check-act wheel in ageing management.

and research results, the use of advanced materials, the need to incorporate on-line monitoring to assist in the forewarning of degradation, new plant layouts and SSCs design that facilitate inspectability, maintainability and easy access. The operating organization should ensure that manufacturers are well informed on the factors affecting ageing management, fabrication of SSCs takes into account current knowledge about relevant ageing degradations and possible mitigation measures, baseline data are collected and documented, and surveillance specimens for specific ageing monitoring programmes are available and installed. Commissioning should be used to know the real plant conditions and to identify and measure all parameters that can influence ageing degradation during operation. Data collection and record keeping should be established early in the life of the plant in accordance with well-established procedures (IAEA, 1991). To obtain the desired quality and quantity of ageing-related data, maintenance and engineering personnel responsible for operation, systems health, maintenance and training should be involved in the record keeping system. The screening approach should focus resources on those SSCs that can have a negative impact on the safety of the plant and are susceptible to ageing degradation. Also those SSCs that are important to safety systems and influence reliability are to be considered.

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PLiM should also be integrated with economic planning. It is said that a plant must be safe but its operation needs to be economical, therefore PLiM has the potential of integrating safety and economical operation. In the last ten years the safety and operational indexes of many nuclear power plants worldwide have experienced and are maintaining relevant safety and operation improvements. It is now well accepted that the safest nuclear power plants are also the top performing plants economically. A typical measure of the safety and reliability of nuclear power plants is related to the annual rate of significant events or reportable events that the licensee must communicate to the regulator. A significant event is broadly defined as any occurrence that challenges a plant’s safety system. These procedures are country dependent, but they can give an idea of the safety in operation and on the reliability of the plant, as they are related to equipment failures or to violations of operating, maintenance or testing and calibration procedures that challenge the plant protection system. The IAEA Principle 3, on Leadership and management for safety in its Fundamental Safety Principles (IAEA, 2006b) requires that: ‘Effective leadership and management for safety must be established and sustained in organizations concerned with, and facilities and activities that give rise to, radiation risks’. Compliance with such principle requires that due consideration should be given to human performance, including systematic training, safety culture and operating experience feedback. Not following this principle could lead to significant and reportable events. The US NRC publishes a yearly Information Digest giving details of the performance of the operating plants. The average number of significant events per reactor and year has declined from 0.90 in 1989 to 0.05 in 2005 (NRC, 2009). In fact, the average number of significant events has remained at or below 0.10 since 1996. The average number of significant events is now below 0.07 per reactor and year. The IAEA also publishes details on nuclear power plant incidents through the Incident Reporting System and its Incident and Emergency Centre. Regulatory authorities also publish data on the safety performance of the national nuclear fleet.

3.5 Ageing degradation mechanisms, and timelimited structures, systems and components There is a long list of ageing degradation mechanisms, the root causes of which can be classified into two major groups: environmental and operational. Environmental agents include radiation, chemicals in coolant and coolant/ moderator fluids and in the environment and physical parameters such as temperature, pressure and humidity. The operational factors include thermal and pressure gradients produced by transients, flow-induced vibrations, chemical stability of materials, seismic displacements and extreme meteorological or

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hydrological natural events. Most of those agents can also be found in other industries, with radiation being particular to the nuclear industry. Licence renewal regulations required that analyses of the different ageing processes should be applied to a long list of structures and components in what has been called time-limited ageing analysis (TLAA). It is therefore necessary to identify each structure and component relevant to safety, subjected to ageing mechanisms, which were designed for a specific lifetime, and to perform analyses to prove that the item in question will perform the assigned safety function as foreseen in the CLB during the additional operation time. Therefore, TLAA has been defined as: ‘those licensee calculations and analyses that involve structures, systems and components within the scope of the licence renewal, consider the effects of ageing and involve time-limited assumptions defined by the current operating terms’. Basic regulations, such as the US 10 CFR Part 54 (NRC, 2006), gives selection criteria for SSCs that should be the object of TLAA; it also includes a long list of potential candidates for analyses. Such a list can be subdivided into groups as summarized in Table 3.2, where the corresponding major ageing processes are also incorporated. The first group includes components of the pressure boundary. The most relevant is the reactor pressure vessel (RPV), where the major ageing stressor is neutron irradiation producing embrittlement in the ferritic steel, mainly in PWRs. Embrittlement also occurs in boiling water and heavy water reactor RPVs, but generally it is of lesser importance due to the existing larger water gap attenuating the neutrons. Stress corrosion cracking has been found in PWR reactor vessel head Alloy 600 clad penetrations and in BWR reactor vessel bottom around control rod penetrations, as well as in pipe welds of dissimilar metals. Irradiation assisted stress corrosion cracking (IASCC) has been found mainly in some BWR stainless steel internals. Changes in dimensions due to void swelling are possible in stainless steels and nickel-rich alloys. Significant deformations have been found in coldworked Zr-2.5Nb pressure tubes channels in the Candu PHWRs due to thermal creep, irradiation growth and irradiation creep forcing the substitution of such components in some of those reactors. Significant growth is also typical of graphite moderated reactors, mainly when the graphite irradiation takes place at low temperatures. Thermal stresses in graphite blocks have produced longitudinal cracks in UK gas-cooled graphite moderated reactors. Wear and corrosion have been found in the pipes, pump casings, valve bodies and steam generator tubes of many reactors. Mechanical vibration fatigue in small pipes and in socket welds may occur due to pressure pulses produced by pumps. Circumferential cracking in the steam generator tubes in Westinghouse Model 3D steam generators, due to wear produced by fluid induced tube vibrations, was a major rapid ageing mechanism discovered in the 1980s, first in Ringhals and later in the Almaraz plants. This design

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Table 3.2 Major components and structures subjected to time-limited ageing analyses and their ageing mechanisms Major candidates for TLAA

Typology of analyses

Reactor vessel, internals and reactor coolant system

Reactor vessel neutron embrittlement Changes in dimensions due to void swelling Metal corrosion allowance Granular and intergranular stress corrosion cracking Radiation enhanced stress corrosion cracking Cumulative fatigue damage

Engineered safety features

Metal corrosion allowance Cumulative fatigue damage Granular and intergranular stress corrosion cracking

Auxiliary systems

Metal corrosion allowance Reduction of neutron-absorbing capacity and loss of material due to general corrosion in spent fuel pools Cracking due to stress corrosion cracking and cyclic loading

Steam and power conversion system

Containment, structures and component supports

Electrical and instrumentation and control

Cumulative fatigue damage Loss of material due to pitting, crevice, and microbiologically influenced corrosion, and fouling Wall thinning due to flow-accelerated corrosion Reduction of strength of concrete due to elevated temperature Loss of metal containment material due to general, pitting and crevice corrosion Loss of tendon prestress due to relaxation, shrinkage, creep and elevated temperature Loss of leak tightness due to mechanical wear of locks, hinges and closures Reduction in foundation strength, cracking, differential settlement due to erosion of porous concrete corrosion Reduced insulation resistance and electrical failure due to thermal, radiolysis, photolysis and chemical mechanisms Degradation of insulation quality due to presence of salt deposits and surface contamination Conductor fatigue due to ohmic heating, thermal cycling, electrical transients, vibration, chemical contamination, corrosion and oxidation

defect has forced the substitution of the affected steam generators. Thermal fatigue may occur in pipes due to fluid thermal stratification caused by defective closure of isolation valves or the seepage of a fluid at different

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temperatures. A well-studied case occurred in one of the units of the Oconee NPP in the United States in 1997 in the junction of a small water make-up pipe and the main recirculation pipe unprotected by a thermal sleeve. A large leaking crack was produced but pipe fracture did not occur, demonstrating the validity of the leak before fracture approach. The second group includes the engineered safety features, mainly the emergency core cooling systems, the containment spray and isolation components and the standby gas treatment in BWRs, which are exposed to cumulative fatigue damage. The emergency core cooling system and the containment spray systems in PWRs include vessels, heat exchangers and pipes containing corrosive chemicals, mainly boric acid and sodium hydroxide, so that the losses of material due to pitting and crevice corrosion are to be expected with the possibility of reducing the heat transfer capacity in heat exchangers due to fouling. The many auxiliary systems in the nuclear power plant belong to the third group. The most salient systems include new and spent fuel storage and handling facilities and related cooling and cleanup systems, the compressed air system, the residual heat removal and coolant cleanup systems, the closed and open cooling water systems and the ultimate heat sink, the ventilation system and emergency diesel generators. As in the previous case, the different systems remain in contact with different chemicals and environments that may produce loss of material by corrosion, including stress corrosion cracking, and material fatigue due to cycling loading. A case of particular importance is the reduction of neutron-absorbing capacity in the spent fuel storage pool due to the loss of the absorber (boron) produced by general corrosion of the supporting structures. Another case is the external corrosion of pipes in the open circuit cooling system recently discovered in the Spanish Vandellós 2 plant (CSN, 2004, p. 16). The steam and power conversion system also includes the main steam system, steam extraction, steam turbine, blowdown system, condenser, feedwater and the auxiliary feedwater systems. Components in the steam and power conversion system are exposed to steam and condensate water and include a variety of pressures, temperatures, fluid regimens in different metals; therefore, metal fatigue and corrosion are the main ageing agents. Containments, structures and component supports include a large variety of equipment. The most relevant part is the reactor containment, its main safety function being to prevent the release of radioactivity to the environment in the case of accidents. This function requires that the integrity and leak tightness of the containment must be ensured. Containments generally include a large metal vessel and a thick concrete cover providing shielding and protection against external man-made and natural events. The major concerns are linked to corrosion of the metallic parts, mainly in inaccessible areas, the long-term behaviour of concrete at high temperatures and the

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loss of leak tightness due to mechanical wear of locks, hinges and closures. Loss of tendon prestress due to relaxation, shrinkage, creep and elevated temperature are of particular importance for the PWR containments using that technology. NUREG-1800 (NRC, 2005b) defines nine groups of structures and seven groups of component supports. Those structures and component supports may be affected by loss of bond and loss of material due to corrosion of embedded steel, increase in porosity and permeability, cracking due to aggressive chemical attack, expansion and reaction with aggregates, cracks and distortion due to increased stress levels from settlement, reduction in foundation strength and cracking and differential settlement due to erosion of porous concrete sub-foundation, among other causes. The most typical electrical, instrumentation and control components include cables, metal enclosed buses, fuse holders, high voltage insulators, transmission conductors and connections and switchyard buses and connections. The major concerns are to be found in degradation of insulators due to salt deposits, surface contamination and loss of material due to mechanical wear, as well as fatigue due to ohmic heating, thermal cycling, electrical transients, frequent manipulation, vibration, chemical contamination and oxidation.

3.6

Main areas of concern for plant designers, operators and regulators

From the discussion in Section 3.5, it becomes clear that plant designers, operators and regulators should be concerned with ageing mechanisms for both already operating plants and new designs. There are many aspects to be considered. Neutron irradiation and the different types of corrosion and wear are considered of major concern to plant designers, operators and regulators. The science and technology of such phenomena are amply considered in other chapters in this book.

3.6.1 Neutron irradiation The interactions between neutrons and the atomic nucleus of crystalline metallic structures have the potential to displace the atom from its normal position in the lattice creating the so-called Frenkel pairs – interstitials and vacancies – which mainly determine the changes in the physical properties of the irradiated material. When stress is applied to such materials substantial mechanical effects appear. Most present and future reactors will include crystalline metallic materials in their core and its proximities, where irradiation by high energy neutrons (>1 MeV) may be substantial. Internals and the pressure vessel belt line are typical examples of components subjected to high energy neutron fluences. The mechanical properties of relevance include:

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∑

hardening, an increase in the yield stress and tensile strength and a reduction in ductility and fracture toughness, more pronounced in bodycentered cubic crystal ferritic steel; ∑ embrittlement, a reduction in the amount of plastic or creep deformation that occurs before rupture – a ductile fracture requires appreciable plastic deformation prior to and during the propagation of a crack, while a brittle fracture implies a rapid rate of crack propagation; and ∑ radiation enhanced stress corrosion cracking, an increase in the corrosion potential of the material.

The importance of embrittlement of RPV materials was soon recognized and regulated (NRC, 1998). Irradiation hardening and embritlement need to be under control to prevent the RPV from ever suffering a brittle fracture when subjected to over-pressurization and to avoid pressurized thermal shock (PTS), i.e. rupture when at pressure by the addition of feed or emergency cold cooling water. Pressurized thermal shock should also be avoided during start-up and shut-down of the reactor. The Charpy V-notch impact test has been the most common way to measure the transition from ductile to brittle fracture. Another chapter in this book is dedicated to assess the embrittlement in ferritic components. The device measures the energy required to fracture a normalized specimen and determines the type of fracture by observing the resulting surfaces. For a given material, tests are performed at different temperatures and a plot is obtained relating the energy absorbed in producing the fracture, generally measured in joules, versus the temperature of the specimen. Figure 3.2 illustrates such relation for a typical RPV material. The plot delimits the brittle and ductile regions of the material. The temperature at which the inflexion of the curve takes place is called the nil ductility temperature (NDT). At high temperatures, the energy absorbed to produce fracture tends to an asymptotic value, which is called the upper shelf energy (USE). When the material is irradiated by fast neutrons there is a change in the above-mentioned parameters. The corresponding Charpy plot shifts towards higher NDTs, there is also a reduction in the USE and the slope of the curve is reduced; moreover, these effects increase with increasing neutron fluence and are also a function of the type of steel and its impurities, mainly sulphur, phosphorous and copper, which have to be limited as much as possible. The temperature shifts could surpass 100 °C for neutron fluences above a few times 1019 n/cm2 . The Charpy tests do not provide a sharp temperature transition from ductile to brittle behaviour; a reference NDT temperature, RTNDT , should be defined for regulatory purposes; it is defined as that temperature at which fracture initiates with essentially no prior plastic deformation. The temperature at which a Charpy V-notch specimen breaks with a fixed amount of energy, generally 41 J, is often considered as the reference temperature.

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Energy absorbed, E(J)

Unirradiated 80 Duse 60 Irradiated 40 DT 20

0 –100

0 100 Temperature, T(°C)

200

3.2 Charpy V-notch tests showing the transition temperature and upper shelf energy shifts in neutron irradiated ferritic steels.

Regulations define the acceptable decrement in the USE and the allowable increment in the NDT. The first is a general measure of the fracture toughness of the material, while the second is vital in the prevention of PTS in PWRs. Regarding the USE, Appendix G to 10 CFR Part 50 (NRC, 2008b) states that: Reactor vessel beltline materials must have Charpy upper-shelf energy in the transverse direction for base material and along the weld for weld material according to the ASME Code, of no less than 102 J initially and must maintain Charpy upper-shelf energy throughout the life of the vessel of no less than 68 J, unless it is demonstrated … that lower values of Charpy upper-shelf energy will provide margins of safety against fracture equivalent to those required by Appendix G of Section XI of the ASME Code. Appendix G also gives screening criteria to prevent PTS. It is based on a specific reference temperature for PTS, RTPTS, which is equal to the RTDNT evaluated at the location where the material receives the highest neutron fluence. Such temperature is limited to 132 °C for plates, forgings and axial weld materials, and 149 °C for circumferential weld metals. A revision in the PTS rule is being considered to reduce the conservatism included. Should the vessel reach the limit, it can be annealed to recover the mechanical properties of the material. It is globally recognized that the Russian developments on pressure vessel science and technology and the improvements they have introduced in design, material selection, fabrication and surveillance of such vessels are based on

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a deep understanding of the physical effects of irradiation and other agents on the vessel materials. They have also developed a valid methodology for annealing to restore the fracture toughness of vessels affected beyond the accepted limits for service. Such information can be found in a Russian monograph published by the American Nuclear Society (Alekseenko et al., 1997).

3.6.2 Stress corrosion cracking Stress corrosion cracking (SCC) is the premature cracking of an alloy in the presence of a tensile stress and a corrosive environment. It was soon recognized that alloys used in nuclear technology were susceptible to SCC and that the intensity of the corrosion depended on the reactivity of the environment and on the presence of tensile stress. SCC reduces the strain to failure as well as the maximum stress, it can be intergranular or transgranular in nature, and susceptibility to SCC is generally high when overall corrosion rates are low. SCC requires an incubation period for crack initiation; it is followed by steady state crack propagation and ends in a failure. SCC was early observed in nickel-rich Alloy 600 used in steam generator tubes, control rod drive mechanism nozzles, pressurized instrument penetrations and heater sleeves and hot leg penetrations. In many steam generators it was observed that the cumulative fractional number of tubes failed started to increase after some ten effective full power years. The case was analyzed and although there were other causes of failure, SCC was predominant after the incubation period. The tube failure rates increased rapidly and many steam generators had to be replaced by new ones made from alloys with less SCC susceptibility (e.g. Alloy 690TT). Was (2007) includes a good analysis of thermodynamics, kinetics and mechanism of SCC; references to examples of SCC in nuclear power plant are found in the IAEA Tecdoc-1361 (IAEA, 2003b). SCC in austenitic stainless steels is considered in depth in other chapters in this book.

3.6.3 Irradiation assisted stress corrosion cracking Irradiation assisted stress corrosion cracking (IASCC) was observed in the early 1960s. It mainly affected stainless steel fuel element cladding, instrument tubes, control rod followers and other core devices or internals. The most important experience appeared in the early 1990s, mainly in the BWR recirculation loops and core shrouds, the last with weld residual stresses, and PWR baffle former bolts, receiving high neutron fluence. The IASCC mechanism is not yet fully known; Was (2007) indicates that existing theories fall into five categories:

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∑ radiation induced grain boundary chromium depletion, ∑ radiation hardening, ∑ localized deformation, ∑ selective internal oxidation and ∑ irradiation creep. A recent study conducted at the Argonne National Laboratory by Chunk and Shack (2005) on 27 commercial and model laboratory austenitic stainless steel heats irradiated in the Halden reactor under BWR conditions has demonstrated that, in 304 or 316 stainless steel, sulphur atoms play a deleterious effect in IASCC, which could be compensated by an increase in carbon content. For instance, when the material has experienced some three displacements per atom, the sulphur content should be <0.002 wt% S; for such materials to be IASCC resistant under BWR conditions, the carbon content should be >0.03 wt% C. A two-dimensional map has been created in which the susceptibility to IASCC in said materials is shown as a function of S and C content. The authors have developed an IASCC model based on Ni and S segregation and preferential oxidation of chromium and iron atoms over nickel atoms, among other aspects. That research effort will define the materials to be used in future water reactors.

3.7

Future trends

It is foreseeable that future trends will aim at improving the knowledge on ageing through additional research and feedback from operating experience. Such knowledge will serve to develop more effective and scientifically-based regulation for present and future designs. In a recent OECD/NEA study (NEA, 2006) it is found that ‘there are no significant technical challenges which would prevent nuclear plant lifetimes being extended to 50 or 60 years’. Nevertheless, the document adds that: ‘it must always be recognized that there is the possibility of unknown ageing mechanisms, as well as the expected development of a known ageing mechanism, emerging during longer term operation’. Solid fundamental research and effective operating experience feedback are the two cornerstones for gaining new knowledge. Nevertheless, operating experience extended beyond 40 years is not yet available and substantial research is not yet finished. Therefore strong international co-operation is needed in sharing operating experience and in conducting research. The end products of such efforts have to be consolidated in harmonized standards and regulations jointly developed by designers, plant owner operators and regulatory authorities. All this effort has to be based on high level trained human resources. Related international activities and some relevant research efforts are described in brief.

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3.7.1 Activities in the International Atomic Energy Agency The IAEA has created a safety knowledge base for ageing and long-term operation of nuclear power plants (SKALTO). It is a framework for sharing information on PLiM and LTO. It provides documents and information related to these thematic areas created by the IAEA and other national or international organizations. It includes safety standards published by the IAEA and regulatory authorities in other countries, as well as INSAG documents. It also includes IAEA review services; international conferences, meetings and national activities related to the safety aspects of ageing and LTO; IAEA co-ordinated and national research programmes, as well as education and training programmes on the matter. All these activities constitute the backbone for future developments. In 2003 the IAEA created an Extrabudgetary Programme on Safety Aspects of Long Term Operation of Water Moderated Reactors (SALTO), with the objectives of: ‘Reviewing existing national approaches, practices and experience that need to be considered during LTO decision-making; developing guidance for regulators on the identification of the applicable safety criteria and on the establishment of guidelines for plant operators’ LTO submittals; and providing guidance for plant operators on the process and practices related to support safe LTO’. The final report of the programme was published in 2007 (IAEA, 2007b); the programme will be incorporated, together with SKALTO, into a comprehensive knowledge base on LTO and PLiM. SALTO is also providing peer review services on LTO tailored to the specific requests of Member States; one of those services is Ageing Management Assessment Team (AMAT).

3.7.2 Activities of the OECD Nuclear Energy Agency Two Committees in the Nuclear Energy Agency are involved in nuclear power plant ageing activities. The Committee on the Safety of Nuclear Installations (CSNI) constitutes a forum for the exchange of technical information between organizations which contribute in research, development, engineering and regulation. It also reviews the state of knowledge on selected topics of nuclear safety technology and safety assessment, including operating experience and research on ageing. The Working Group on Integrity of Components and Structures (IAGE), and its subgroups, conduct studies and analyse research results on ageing in NPPs. The Committee on Nuclear Regulatory Activities (CNRA), made up primarily of senior nuclear regulators, is a forum for the exchange of information and experience among regulatory organizations and for the review of developments which could affect regulatory requirements, including PLiM and LTO. Both Committees are composed of Working Groups

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that have made comparative studies, celebrated symposia and workshops and have made suggestions on needed research efforts and regulatory activities. The present Strategic Plan for both Committees, covering 2005–9, can be found in NEA (2005). The CNRA held a Special Issue Meeting in June 2000 on the topic of ‘life extension and upgrading’. The published report (NEA/CNRA, 2001a) includes the synthesis of the responses received from Member States and the results and conclusions of the CNRA discussions. At the time of the report most members with operating reactors had already developed, or were developing programmes on longer-term operation (NEA/CNRA, 2001b) Based on the many technical and regulatory issues treated by the CSNI and the CNRA, the Committee for Technical and Economic Studies on Nuclear Energy Development and the Fuel Cycle, known as the Nuclear Development Comittee, NDC, decided to take a holistic view on the matter which resulted in the creation of an ad hoc Expert Group on Nuclear Power Plant Life Management. The Expert Group expressed its considerations and recommendations in a report (NEA, 2006). The Group suggests that research and development capacities have to be kept alive for studying material degradation and unknown phenomena. The activities of the NEA are equally needed to ensure a proper scientific and technical development on ageing and related issues to serve as a basis for regulatory requirements.

3.7.3 The research effort Fundamental research is performed nationally and within international organizations. The European Union is sponsoring relevant research, the IAEA is also maintaining international co-ordinated research activities, the NEA is serving valid efforts on common research efforts, and some national institutions, such as the US Electric Power Research Institute (EPRI), have offered their research efforts to foreign participants. EURATOM research on ageing science and technology started early. In 1993 the Ageing Materials Evaluation and Studies programme (AMES) was established within the Framework Programme 4, FP-4, to bring together the European expertise on nuclear materials ageing. The strategies within AMES were based on understanding the embrittlement causes and thus improve the prediction of irradiated material fracture toughness. In 2001 the European Commission (EC) defined the ageing mechanisms, their potential effects and the available identification and mitigation methods that needed to be explored (EC, 2001). The Concerted Action VERSAFE, as part of the Euratom 5th Framework Programme, FP-5, aimed at an overview on the comprehensive approaches to PLiM in general terms and plant-specific issues of VVER-440 plants. In the Euratom 6th Framework Programme, FP-6, PERFECT was the most significant project. The aim was to build two ‘virtual reactors’ to

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simulate the effect of irradiation on reactor pressure vessels and on internal structures, thus reducing the need for experimental data. The resulting four numerical tools were integrated in a Software Integration Platform. In 2006 a new European Network of Excellence was started under the title Nuclear Plant Life Prediction, NULIFE, to integrate safety-oriented research on materials, structures and systems and to exploit results of this integration through the production of harmonized NPP lifetime assessment methods. An account of this network has been presented by Rintamaa et al. (2007). In December 2006 the European Council approved the 7th Framework Programme, FP-7, for the period 2007–11 (EC, 2006). The document decided to commend: ‘Research to underpin the continued safe operation of all relevant types of existing reactor systems (including fuel cycle facilities), taking into account new challenges such as life-time extension and development of new advanced safety assessment methodologies’. Safety of Ageing Components in Nuclear Power Plants (SAFELIFE) has been the project of reference. SAFELIFE aims at establishing best practices based on deterministic and risk-informed methods for assessing the structural safety of key components in both Western and Russian nuclear power plant designs. The most salient goals are: Provide a scientific and technical basis for harmonization of European codes and standards on key primary components of light water reactors through developing and disseminating best practices. Support long-term EU policy needs on PLiM and advanced reactor concept through enhancing JRC R&D competence and capabilities in nuclear safety technology. Integration of R&D efforts in line with European Research Area (ERA) principles by linking our R&D to utilities, manufacturers, R&D organizations and regulators through continuing exploitation of networks and collaborating with EC and international organizations. Implementation of an effective plan for training, mobility, dissemination and knowledge management and development of competitive activities. Since the early 1990s, the IAEA has conducted a series of co-ordinated research projects, (CRPs), covering the managing of ageing of RPV primary nozzles, motor operated isolation valves, in-containment I&C cables and concrete containment buildings. The IAEA has also conducted up to nine CRPs on irradiation embrittlement of RPV steels; the most significant ones are related to the application of the master curve approach (IAEA, 2005a), the RPV surveillance programmes and their analysis and application (IAEA, 2005b) and the nickel effects in radiation embrittlement of RPV materials (IAEA, 2005c). The Nuclear Energy Agency administers a series of international research programmes suggested by the Member States. The most significant one, going on at present, is the Project on Stress Corrosion Cracking and Cable Ageing

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(SCAP), which is supported by 14 NEA member countries. The project began in 2006 and the current mandate ends in 2010. It has been described by Yamamoto et al. (2007). The project’s main objectives are to: ∑

establish two complete databases with regard to major ageing phenomena for stress corrosion cracking (SCC) and degradation of cable insulation respectively, through collective efforts by OECD/NEA member countries; ∑ establish a knowledge base by compiling and evaluating collected data and information systematically; and ∑ perform an assessment of the data and identify the basis for commendable practices which would help regulators and operators to enhance ageing management. The present near-term (2008–9) EPRI research portfolio includes five major projects related to materials ageing: Boiling Water Reactor Vessel and Internals (BWRVIP); Pressurized Water Reactor Materials Reliability Programme; Primary System Corrosion Research; Steam Generator Management Programme (SGMP) and Water Chemistry Programme. Of particular interest is the corrosion programme in the primary system aimed at analysing IASCC in PWRs and BWRs under a co-operative international research programme. BWRVIP will try to develop effective countermeasures for mitigating stress corrosion cracking of reactor internal components in BWRs, such as using hydrogen water chemistry and noble metal chemical application on fuel, and other advanced mitigation technologies. The programme will also provide acceptable design criteria and unique solutions to repair or replace reactor internals and piping. It will also provide understanding of materials performance exposed to high neutron fluence, weldability of irradiated materials and crack growth rates on IASCC and ISCC. Items of special interest deal with X-750 high strength materials which have been used to repair jet pump beams and core shrouds in BWRs and the susceptibility of the pressure vessel bottom head drain line to flow assisted corrosion; a specific User Group under BWRVIP is developing inspection and repair tools for such components. The Pressurized Water Reactor Materials Reliability Programme aims at optimizing the addition of hydrogen and zinc to mitigate primary water stress corrosion cracking. The programme also tries to better understand the crack initiation and propagation processes and environmental corrosion in the reactor coolant system components and develop better predictive and mitigation technologies. It also addresses IASCC using irradiation samples in PWR conditions. The Primary System Corrosion Research programme includes an internationally sponsored co-operative programme. It aims at a better

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understanding of the crack initiation and early propagation processes involved in SCC and IASCC in nickel base alloys and stainless steels used in PWRs and BWRs. Extensive international collaboration ensures that research findings reflect a wide range of nuclear technologies, operating conditions, and service environments. In SGMP, research is conducted to ensure the safe and economic operation of steam generators in PWRs. Research activities target identification and mitigation of various forms of steam generator degradation, replacement steam generator specifications, water chemistry guidelines, in-service inspections and tube integrity. The Water Chemistry Programme develops and updates water chemistry guidelines for nuclear reactors based on industry research and plant experience. The programme also develops water chemistry optimization tools to mitigate corrosion, achieve and maintain design fuel performance standards, and minimize plant radiation fields.

3.8

References

Alekseenko N S, Amaev A, Gorynin I, Nikolaev V A (1997), Radiation damage of nuclear power plant pressure vessel steels, Russian Materials Monograph Series, La Grange Park, American Nuclear Society. ASME (2003), Standard for probabilistic risk assessment for nuclear power plant applications, RA-S-2002, New York, ASME. Chunk H M and Shack W J (2005), Irradiation-assisted stress corrosion cracking in austenitic stainless steels applicable to LWR core internals, NUREG/CR-6892, ANL04/10, Washington DC, US NRC. Contri P (2006), Workshop on ‘Maintenance rules: Improving maintenance effectiveness’. Summary report, EUR 22603 EN, Luxembourg, EC Publications. CSN (2009), Revisiones periódicas de seguridad de las centrales nucleares, Instruction IS-22, Madrid, CSN. CSN (2004), Informe del Consejo de Seguridad Nuclear al Congreso de los Diputados y al Senado, INF-01.04, Madrid, CSN. CSN (2006), Sistema Integrado de Supervisión de Centrales Nucleares-SISC, Madrid, CSN. EC (2001), Safe management of NPP ageing in the European Union, CE/DG X, EUR 19843 EN, Luxembourg, EC Publications. EC (2006), Council decision concerning the seventh framework programme of the European Atomic Energy Community (Euratom) for nuclear research and training activities (2007 to 2011), Official Journal of the European Union L 460 of 30 December 2006, Luxembourg, EC Publications. IAEA (1991), Data collection and record keeping for the management of nuclear power plant ageing, Safety Series 50-P-3, Vienna, IAEA. IAEA (1992), Methodology for ageing management of nuclear power plant component important to safety, Technical Report Series No. 338, Vienna, IAEA. IAEA (1998), Equipment qualification in operational nuclear power plants: Upgrading, preserving and reviewing, Safety Report Series No. 3, Vienna, IAEA.

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IAEA (1999), Framework for a quality assurance programme for PSA, TECDOC-1101, Vienna, IAEA. IAEA (2000), Safety of nuclear power plants, Safety Standards Series NS-R-2, Vienna, IAEA. IAEA (2002a), Recruitment, qualification and training of personnel for nuclear power plants, Safety Standards Series ND-G-2.8, Vienna, IAEA. IAEA (2002b), Maintenance, surveillance and in-service inspection in nuclear power plants, Safety Standards Series NS-G-2.6, Vienna, IAEA. IAEA (2003a), Periodic safety review of nuclear power plants, Safety Standards Series NS-G-2.10, Vienna, IAEA. IAEA (2003b), Assessment and management of ageing of major nuclear power plant components important to safety, TECDOC-1361, Vienna, IAEA. IAEA (2005a), Guidelines for application of the MasterCurve approach to reactor pressure vessel integrity in nuclear power plants, Technical Report Series No. 429, Vienna, IAEA. IAEA (2005b), Application of surveillance programme results to reactor pressure vessel integrity assessment, TECDOC-1435, Vienna, IAEA. IAEA (2005c), Effects of nickel on irradiation embrittlement of light water reactor pressure vessel steels, TECDOC-1441, Vienna, IAEA. IAEA (2006a), A system for the feedback of experience from events in nuclear installations, Safety Standards Series NS-G-2.11, Vienna, IAEA. IAEA (2006b), Fundamental safety principles, Safety Standards Series SF-1, Vienna, IAEA. IAEA (2007a), IAEA Safety glossary: Terminology used in nuclear safety and radiation protection, 2007 edition, Vienna, IAEA. IAEA (2007b), Safety aspects of long term operation of water moderated reactors, IAEA-EBP-SALTO, Vienna, IAEA. IAEA (2008), Safety of nuclear power plants: Operation, Safety Requirements Draft, DS413, Vienna, IAEA. IAEA (2009), Ageing management for nuclear power plants, Safety Standards Series Na NS-6-214, Vienna, IAEA. IAEA/NEA (1980), Incident Reporting System (IRS), Using operational experience to improve safety, Vienna, IAEA. IAEA/NEA (2006), Nuclear power plant operating experience 2002–2005, No. 6150, Paris, OECD Publications. INSAG (1988), Basic safety principles for nuclear power plants, INSAG-3, Vienna, IAEA. INSAG (1996), Defence in depth in nuclear safety, INSAG-10, Vienna, IAEA. INSAG (2008), Improving the international system for operating experience feedback, INSAG-23, Vienna, IAEA. NEA (2005), Joint CSNI/CNRA strategic plan and mandates, 2005–2009, NEA 6034, Paris, OECD Publications. NEA (2006), Nuclear power plant life management and longer term operation, NEA 6105, Paris, OECD Publications. NEA/CNRA (2001a), Regulatory aspects of life extension and upgrading of NPPs. NEA/ CNRA/R (2001)1, Paris, OECD Publications. NEA/CNRA (2001b), CNRA Special Issue’s Meeting 2000, Member Countries responses to the questionnaire, NEA/CNRA/R (2001)2, Paris, OECD Publications. NRC (1991), Requirements for monitoring the effectiveness of maintenance at nuclear power plants, 10 CFR Part 50.65, Washington DC, US NRC. © Woodhead Publishing Limited, 2010

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NRC (1997), Monitoring the effectiveness of maintenance at nuclear power plants. Regulatory Guide 1.160, rev. 2, Washington DC, US NRC. NRC (1998), Radiation embrittlement of reactor vessel materials, Regulatory Guide 1.99 rev. 2, Washington DC, US NRC. NRC (2000a), Assessing and managing risk before maintenance activities at nuclear power plants, NRC Regulatory Guide 1.182, Washington DC, US NRC. NRC (2000b), NRC Reactor Oversight Process, ROP, NUREG-1649 rev. 3, Washington DC, US NRC. NRC (2005a), Standard format and content for applications to renew nuclear power plant operating licenses, Regulatory Guide 1.188 rev. 1, Washington DC, US NRC. NRC (2005b), Standard review plan for review of license renewal application for nuclear power plant, NUREG-1800 rev. 1, Washington DC, US NRC. NRC (2005c), Generic Aging Lessons Learned (GALL) Report, NUREG-1801, Vol. 1 & Vol. 2, rev 1, Washington DC, US NRC. NRC (2006), Requirements for renewal of operating license for nuclear power plants, 10 CFR Part 54, amended 2006, Washington DC, US NRC. NRC (2009), Information Digest 2009–2010, NUREG-1350 vol. 20, Washington DC, US NRC. NRC (2008a), Reactor vessel material surveillance program requirements, Appendix H (as amended) to 10 CFR Part 5, Washington DC, US NRC. NRC (2008b), Fracture toughness requirements, Appendix G to 10 CFR Part 50, Washington DC, US NRC. NUMARC (1993), Industry guideline for monitoring the effectiveness of maintenance at nuclear power plants, NUMARC 93-01, Washington DC, NEI. Ranguelova V and Contri P (2006), Workshop on ‘Advanced methods for safety assessment and optimization of NPP Maintenance’, Summary report, EUR 22604 EN, Luxembourg, EC Publications. Rintamaa R, Aho-Mantila I, Heikinheimo L and Taylor N (2007), European research network aiming at harmonized plant life prediction procedures, IAEA-CN-155-001, in Shanghai International Symposium on PLiM. Was G S (2007), Fundamentals of radiation materials science, New York, Springer. WENRA (2006), Harmonization of reactor safety in WENRA countries, Reactor Harmonization Working Group. Yamamoto A, Huerta A, Gott K and Koshy T (2007), The NEA Project on Stress Corrosion Cracking and Cable Ageing (SCAP), NEA News, Vol. 25, No. 1, p. 18.

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Probabilistic and deterministic safety assessment methods for nuclear power plant life management

P. C o n t r i and A. R o d i o n o v, European Commission DG-JRC Institute for Energy, The Netherlands

Abstract: This chapter addresses safety assessment methods in relation to ageing effects and develops an integrated safety and economic plant life management system able to effectively manage ageing effects at nuclear power plants. After a short summary with definitions and framework, a reference Plant Life Management (PLiM) model is discussed, as background for a detailed analysis of the main issues related to component and system degradation. A Probabilistic Safety Assessment (PSA) approach where timedependent ageing effects are considered is also discussed, with emphasis on the evaluation of time-dependent component reliabilities. Some case studies on component reliability calculation are described where ageing considerations have a strong impact, providing suggestions on how reliability and data analysis for active components can be carried out in an effective way and incorporation of age-dependent reliability and data into PSA models can be implemented. Key words: ageing, plant life management, probabilistic safety assessment, component reliability, operation safety.

4.1

Introduction – plant safety assessment in a plant life management (PLiM) framework

According to IAEA (2002a), the aim of a nuclear power plant (NPP) safety assessment should be to establish and confirm the design basis for the items important to safety and to ensure that the overall plant design is capable of meeting the prescribed and acceptable limits for radiation doses and releases for each plant condition category. In the case of an operating plant, the safety analysis developed at the design stage is replicated/updated at regular time intervals or every time a major change is made to the plant, in order to provide for the safety justification of a proposed design modification. In this framework, it is necessary that the plant design models and data (which are essential foundations for the safety analysis) are kept up-to-date during the design phase and throughout the lifetime of the plant, including decommissioning. This should be the responsibility of the designer during the design phase and then of the operating organization over the life of the plant. 88 © Woodhead Publishing Limited, 2010

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The safety analysis should formally assess the performance of the plant under various operational and accident conditions (according to definitions in IAEA, 2002a), against goals or criteria for safety and radiological releases as may have been established by the operating organization, the regulatory body, or other national or international authorities, as applicable to the plant. Moreover, the safety analysis should support safe operation of the plant by serving as an important tool in developing and confirming plant protection and control system set points and control parameters. It should also be used to establish and validate the plant’s operating specifications and limits under normal and off-normal operating conditions, procedures used, maintenance and inspection requirements, and verify normal and emergency procedures. In particular, the safety analysis should assess whether: ∑

sufficient defence-in-depth has been provided and the levels of defence are preserved in that potential accident sequences are arrested as early as possible ∑ the plant can withstand the physical and environmental conditions it would experience, including extremes of environmental and other conditions ∑ human factors and human performance issues have been adequately addressed ∑ long-term ageing mechanisms that could affect the plant’s reliability over the plant life are identified, monitored and managed (i.e. by upgrade, refurbishment or replacement) so that safety is not affected and risk does not increase.

4.1.1 Deterministic and probabilistic approaches According to IAEA (2002a), the achievement of a high level of safety should be demonstrated primarily in a deterministic way. However, the safety analysis should incorporate both deterministic and probabilistic approaches. These approaches have been shown to complement each other and both should be used in the decision-making process on the safety and ability of the plant to be licensed. The probabilistic approach provides insights into plant performance, defence-in-depth and risk that are not provided by a deterministic approach. The aim of the deterministic approach should be to address plant behaviour under specific predetermined operational states and accident conditions and to apply a specific set of rules for judging design adequacy. The probabilistic safety assessment (PSA) should set out to determine all significant contributors to risk from the plant and should evaluate the extent to which the design of the overall system configuration is well balanced, there are no risk outliers and the design meets basic probabilistic targets. The PSA should preferably use a best estimate approach.

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4.1.2 Probabilistic plant safety assessment considering ageing effects According to the above-mentioned generic statement, the periodic review of the plant safety assessment should include the assessment of potential time-dependent effects. Concerning this aspect, INSAG-12 (1999) states the safety goal of a probabilistic safety assessment of a nuclear installation and provides references for consideration of time-dependent effects: The target for existing nuclear power plants consistent with the technical safety objective is a frequency of occurrence of severe core damage that is below about 10–4 events per plant operating year. Severe accident management and mitigation measures could reduce by a factor of at least ten the probability of large off-site releases requiring short-term off-site response. In addition to that generic statement, the worldwide tendency to apply riskinformed regulations and procedures to optimize, among others, inspections and maintenance tasks, assigns to the PSA a key role, even in the decisionmaking phase, as shown in Fig. 4.1 (US-NRC, 2002). The possible impact of ageing phenomena on SSC reliability and on overall plant safety is illustrated in the risk–barrier–target diagram in Fig. 4.2. Each of the ‘barriers’ used to decrease or to avoid the impact of ageing on safety is covered in some way by the risk-informed regulation approach. Deterministic analysis

1. Define change

PSA

2. Perform engineering analysis

3. Define implementation/ monitoring

4. Submit a proposal

4.1 Principal elements of risk-informed, plant-specific decision making (US-NRC, 2002). Barrier

Risk

Operating loads + Environmental stressors

SSC degradation due to ageing

• • • • •

Design Qualification Test/Monitoring Maintenance Operating feedback

4.2 Ageing effect on unit/SSC reliability and safety.

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Presently, ageing evaluation-related activities have been or are being carried out as part of the following programmes: ∑ periodic safety review ∑ ageing management ∑ maintenance optimization ∑ operation after the original design life (also called ‘long-term operation’, LTO). There are many national and international standards and guidelines available, but all of them are based on a deterministic approach and describe very limited PSA application. The PSA could be incorporated more into these programmes as a safety evaluation tool to help identify and prioritize ageing issues and optimize ageing management activities. In general, to apply PSA in a risk-informed approach, PSA should be as realistic as practicable and appropriate support data should be available for the review. This would also apply to identifying potential risks associated with ageing effects. In conclusion, the following methodological issues have to be addressed in the attempt to develop time-dependent PSA analysis able to cover ageing effects: ∑ Could PSA be applied to ageing assessments? ∑ How realistically do PSA models reflect important ageing issues? ∑ Are PSA methods and models sensitive enough? ∑ Are any modifications or revisions of PSA assumptions needed to apply a PSA approach to risk-informed decision making with regard to ageing evaluation? ∑ What data are available and how representative are they with regard to important ageing issues? This chapter makes reference to the previous chapters concerning the deterministic assessment of the integrity of the main components, while it addresses methods and procedures for PSA models where ageing effects are explicitly included, making the periodical safety assessment a process fully integrated with the plant ageing evaluation. In particular, an overview is presented of the different PSA phases where ageing considerations have a significant impact, such as: selection of systems, structures and components (SSCs) to be considered, reliability and data analysis for active components and incorporation of age-dependent reliability parameters and data into PSA models. For consistency, a reference PLiM model is described first, as a background for the proposed probabilistic safety assessment approach.

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The plant life management (PLiM) problem – definitions and selected experience cases

4.2.1 Setting the problem The plant life management (PLiM) problem was raised some years ago when it became clear that technological, safety, regulatory, human and economic issues had to be addressed simultaneously for the overall management of the plant assets (IAEA, 2006; EU, 2007). It is a fact that new global approaches have been triggered in recent years by a combination of factors such as: ∑

the generic trend towards plant life operation beyond the original design life, in order to exploit the plant design to the maximum level ∑ the market economy, which is pushing for a more stringent management of the economic assets ∑ the detection of significant ageing phenomena, which may be challenging the original design assumptions ∑ the need for timely preservation of the human knowledge, particularly in countries with growing opposition to nuclear power expansion ∑ the more stringent regulatory requirements in terms of safety assessment and monitoring. However, the PLiM models developed in recent years differ from each other because of national frameworks and regulations and therefore a generalization sometimes appears difficult. Interesting attempts were carried out by the International Atomic Energy Agency in a series of technical documents and papers (see list in EU, 2007), to identify common drivers among the different national programmes, but the discipline was never indeed regulated by binding documents on its Member States, by presenting commonly accepted principles, recognized by all the interested parties. Nevertheless, a large number of IAEA documents are available on basic safety concepts that could be relevant to life management programmes (IAEA, 1992, 1997a,b, 2000a,b,c, 2001, 2002a,b, 2004a,b). In particular, a generic misunderstanding still exists in the engineering community among objectives and content of the different programmes put in place in the different countries which developed experience in the PLiM field. Programmes such as Licence Renewal (LR), Long Term Operation (LTO), Plant Life Extension (PLEX), Periodic Safety Review (PSR), Ageing Management Programme (AMP), for example, were shown to share many technical tasks, but also to meet different objectives and to follow different regulatory frameworks. The European Commission, through its Joint Research Centre at the Institute for Energy (JRC-IE) has spent some research effort in recent years in the clarification of the many issues addressed by the European countries’ programmes and developed some unified models, which received

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very high consensus in many engineering communities and particularly in the research network of European countries interested in this discipline, SENUF (Safety of European-type Nuclear Facilities) (JRC, 2008). A number of scientific papers were also published in order to foster feedback from the engineering community (JRC, 2006; Contri, 2007a,b; Contri et al., 2007a,b; Vaisnys et al., 2007). As an outcome of this effort, a list of generic considerations can be drawn as support to the development of a more unified approach to the common issue of managing the plant assets in time, while meeting the highest safety standards: 1. The PLiM programme appears the type of programme most suitable to address long-lasting safety and economical issues and to present the most comprehensive approach to the plant asset management. 2. The PLiM programme is neither necessarily related to operational plant life after the original design life, nor to license extension of any plant. It is a logical framework on which strategic thinking may find the appropriate answers in relation to safety, economy and human asset management. 3. Related programmes such as LR, AMP, PSR, each with its own objective, may find in the PLiM framework the answers and the background information that they need to meet their specific objectives; however, they definitely represent separate programmes, different from PLiM itself. 4. The PLiM programme is crucially based upon a strong integration of many existing programmes at the plants, such as asset management, plant life beyond design, ageing management, configuration control, predictive maintenance, etc., that share common assumptions and contribute to the same overall objectives. 5. Some special features are required in standard programmes and also some specific programmes are needed to be in place at NPPs in order to supply necessary input to a PLiM programme adequately. These features/programmes create the pre-conditions for a PLiM programme to be successfully applied, that is, the maintenance programme should be mostly reliability based, while the ISI programme should be risk informed, in nature, a fuel management programme should be in place, an outage optimization programme should make available all data in relation to the economic implications of the outage duration, a knowledge management programme should be in place, public acceptance analysis should be available, etc. 6. In order to manage the very complex structure of a PLiM programme, specialized software tools and databases are highly recommended, also for the management of daily work, due to the huge amount of data to be processed and stored. © Woodhead Publishing Limited, 2010

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One example of an approach to PLiM is shown in Fig. 4.3, taken from Finnish practice (JRC, 2008). In this example, the PLiM programme aims at demonstrating that during the design and possibly the continued plant operational life (LTO) (JRC, 2006): 1. the safety and ageing analyses remain valid and could be projected to the end of intended operational lifetime 2. the effects of ageing on the intended safety function(s) are adequately managed for the entire envisaged lifetime 3. there is a mechanism to deal with unexpected ageing mechanisms that may appear 4. there is a pro-active process for decision making, also involving nonsafety equipment significant to plant availability 5. there is a programme to manage human resources and knowledge 6. plant economic assets are properly managed. In this framework certain programmes play a crucial role, namely: ∑ the ageing management programme (AMP) ∑ the maintenance, surveillance and inspection (MS&I) programme ∑ the knowledge management programme (KM) ∑ the asset management programme ∑ major plant upgrading programmes (if in place, such as power uprating, modernization, etc.). In particular, the AMP is a transversal programme (JRC, 2006) cross-cutting maintenance, surveillance, and in-service inspection programmes and other operation-related programmes. It addresses ageing mechanisms prevention, control and consequence mitigation. The operating experience shows that

Safety and licensing

Continuous safety upgrading

Life management of critical SSCs

Production and economy

Strategic key issues: prerequisites for success

50 years operation

Long-term personnel plan Human resources

Long-term investment plan

4.3 Examples of approach to component life management (JRC, 2008).

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active and short-lived SSCs are in general addressed by existing maintenance programmes. Conversely, the performance and safety margins of the passive long-lived SSCs are assumed to be guaranteed by design. However, the analysis of the operating experience worldwide has shown that unforeseen ageing phenomena may occur either because of shortcomings in design, manufacturing or by operating errors, calling for a refined, self-improving programme. The maintenance programme for a nuclear power plant covers all preventive and remedial measures that are necessary to detect and mitigate degradation of a functioning SSC or to restore to an acceptable level the performance of design functions of a failed SSC (IAEA, 1997b). In this sense, the integration with surveillance and in-service inspection is crucial, as the most advanced types of maintenance do integrate the three programmes, which have a common objective: to ensure that the plant is operated in accordance with the design assumptions and within the operational limits and conditions. Therefore in the following, MS&I will address all the three programmes in an integrated form. It is clear that the MS&I programme is a crucial part of PLiM, being by far the main contributor to both operating costs (after operation) and operation planning. However, in order to support a PLiM framework, MS&I should have a specific list of attributes, making both safety assessment and cost optimization possible. In conclusion, the implementation of an AMP and a predictive MS&I (maintenance, surveillance and inspection) programme is definitely a condition for the operation within the limits of design or licensed lifetime and is a condition for a PLiM as well. Knowledge management and asset management are traditionally isolated programmes from MS&I and AMP. PLiM recognizes the need for their integration and sets an overall optimization framework.

4.2.2 Countries’ generic experience with PLiM Thanks to the large survey on countries’ practice carried out at the JRC (JRC, 2008), also through the organization of many international events, it was possible to summarize the most relevant aspects of some countries’ practice in the field of PLiM, with special emphasis on the relationship with other programmes being used in European Union (EU) countries and outside. ∑

The United Sates, Canada, Spain and some other countries accumulated valuable experience in recent years in PLiM issues and related programmes. The interest of the international community of plant operators on reliabilitybased approaches to PLiM and maintenance optimization in particular, is increasing. The US approach is codified in the INPO (2004), which is also closely followed by some European countries (such as Spain and Hungary).

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∑

Other European countries, such as Finland, are more in favour of integrated approaches to PLiM, with a more explicit control of the component degradation and a clear day-to-day basis for the decision makers on replacement, maintenance and operation. ∑ In many European countries, PLiM is accompanied by a periodic safety review (PSR) programme. The combination is not surprising, as PLiM is typically a utility-driven programme, while PSR is driven by the safety authorities. Many technical tasks (those safety related) are similar, but objectives, time frame and regulatory implications are definitely different, even if in some countries the PSR is the time when LTO tasks are carried out. ∑ Some pre-conditions for PLiM in many countries include maintenance optimization, RI-ISI, fuel management, outage optimization, knowledge management, public acceptance, seismic upgrading, etc., sometimes making the programme very complex. In all cases they are assisted by complex software tools and databases, also for the management of daily work. ∑ The relationship among PLiM and the other programmes running at the NPPs is now quite clear in the EU countries: well-known programmes such as component integrity, ageing management (AMP), life extension (PLEX), periodic safety review (PSR) and plant life management (PLiM) are in fact well connected, but definitely not interchangeable. Despite the different names, mostly derived from the national regulatory and engineering frameworks, there is a clear hierarchy among them. In particular, component integrity is a basic science dealing with the failure modes of the different components, their detection and their control. The AMP is an operational programme in place at any NPP, which integrates maintenance, ISI and organizational issues aiming at controlling the component degradation. PLiM addresses safety as well as economics, knowledge management as well as decision making, and provides an overall framework to keep the whole plant in a safe and economically sustainable condition.

4.2.3 PLiM at the design stage for new reactors The comparison of the approach to PLiM among other technology areas suggests some interesting considerations. For example, in the aerospace industry the maintenance and ageing management programme (considered among the most crucial components in PLiM) are optimized at the design phase, in part because of the large number of similar aircraft; in nuclear power practice, time is needed to accumulate statistics and to develop confidence in the optimization procedures. Many types of NPP are operating and generic approaches have limited application, except, maybe in France, where many

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NPPs of similar design are in service. However, it is common judgement that PLiM should be applied already at the design phase of the NPPs, possibly based on the lessons learned from operating fleets. In this sense, practice could be assimilated in a similar way to that in other industrial technologies, as mentioned above. Up to today, the development of standards and design rules for the new generation reactors is lagging behind. This delay also makes the certification of the new reactors problematic. At the same time, also the safety assessment methods and the quality assurance (QA) rules for construction and operation need to be revised. The role of both licensee and regulators is still to be defined in many countries. This generic statement is applicable to all PLiM relevant aspects that deserve an early understanding at the design phase. In particular, ageing considerations should be addressed already at the design stage, for example to provide inspectability, replaceability and access to the most sensitive components and a solid basis for the control of their degradation. Ageing should also be addressed at the beginning of operation, in order to make available a broad range of data for trending and optimization goals. In particular, the following PLiM/ageing relevant issues should be addressed in the pre-design or pre-licensing phase of new reactors: ∑ choice of materials ∑ major drawings ∑ operating conditions ∑ collection of relevant data ∑ monitoring, surveillance ∑ ISI: inspectability/access/ease of replacement of SSCs ∑ radiation protection of workers (as low as reasonably acceptable (ALARA) principles). For example, in the Areva/EPR, the following design actions have been taken in order to improve the PLiM performance (Areva, 2007): ∑

∑ ∑ ∑ ∑ ∑

accessibility of the reactor building during normal operation to perform maintenance tasks and inspections, but also to start refuelling seven days before reactor shutdown and to continue demobilization three days after reactor restart improved main coolant loop cool down, depressurization and vessel head opening after shutdown bringing the standard outage time to 16 days very low radiation level to workers some modifications in steam generators (SGs) or pressurizer or reactor vessel internals (RVIs) improvements to nozzles and tees for thermal fatigue reduction in general FU factors (unavailability factor) have to be less than 0.5 for limited ISI in operation.

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In the Westinghouse AP1000 (Westinghouse, 2007) the following design actions intend to address the PLiM issues: ∑ ∑ ∑ ∑ ∑ ∑ ∑ ∑

large use of passive features, also to reduce MS&I tasks variable speed in the main reactor coolant pump, to shorten start-up and shutdown times special design of the digital instrumentation and control (I&C) which reduces the I&C surveillance testing large use of component standardization to reduce parts inventory and training scope of personnel built-in testing capabilities is provided for many critical components easy access for MS&I tasks and lifting devices few nuclear grade equipment very low radiation level to workers.

4.2.4 The maintenance programme in the LTO perspective: why PLiM needs an optimized maintenance programme In 2003 the JRC-IE carried out a preliminary investigation of the priorities in European countries in relation to the PLiM programmes. The conclusion was that the nuclear power community is generally convinced that the maintenance programme should have specific attributes in order to support a long-term operation (LTO) programme for the plant. In this sense, the international standards (e.g. the IAEA), but also the national experience of the United States, Spain and Hungary, etc., provided confirmation of this. More specifically, the maintenance programmes based on standard preventive maintenance (time-based), not oriented to the monitoring of its effectiveness and to the prediction of damage, are not considered suitable to support LTO programmes. Crucial attributes for maintenance programmes in order to support LTO are considered to be: the verification of the performance goals, the root cause analysis of failures, the feedback from maintenance to the ISI programme, and the feedback on the operational limits and conditions (OLC). All countries implementing a LTO programme applied extensive modifications to their requirements on maintenance as a first step, setting up mechanisms to monitor the effectiveness of the maintenance activities, which is seen as a pre-condition for entering LTO programmes (see, for example, the new requirements for maintenance in the United States contained in the 10 CFR 50-54 document). In particular, the following features are believed to be indispensable for a maintenance programme in a PLiM framework: ∑

monitoring the performance of the SSCs which may have an impact on safety and reliability during all operational statuses of the plants;

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assessing and managing the risk that may result from the proposed maintenance activities in terms of planning, prioritization and scheduling.

In order to implement these requirements, some issues have to be addressed, namely: ∑

The identification of the scope of the condition-based maintenance rules: typically the countries choose the safety-related SSCs, SSCs that mitigate accidents or transients, SSCs interacting with safety-related SSCs, and SSCs that could cause scram or actuation of safety-related systems. Therefore, many non-safety-related SSCs may see the application of such maintenance rules, with augmented efforts in monitoring their performance and planning their reparation. ∑ The setting of performance goals for every component in the scope of the maintenance rules, ranking them according to their risk significance to plant safety. This task may end up very challenging as, when industry experience is not available, either dedicated PSA tasks have to be developed (with special requirements on PSA quality) or special qualification programmes for the evaluation of the component reliability. ∑ The performance monitoring techniques for the very broad categories of SSCs in the scope of the rules. ∑ The assessment of safety during implementation of maintenance actions. ∑ The feedback from the result of the monitoring of the component reliability back into the inspection, surveillance and maintenance procedures. Root cause analysis, equipment performance trend analysis and corrective actions have to be developed on a case-by-case basis. In this sense, for example, the experiences of the United States and Spain (where a LTO programme is well established), Hungary, and Finland (where a PLiM model is in place at the Loviisa NPP) are a confirmation of this generic statement: all these countries modified their regulatory requirements or practice on maintenance, in the direction mentioned above, as one of the preconditions for the operation of their plants after reaching the original design life. As summary of the country practice in the field of maintenance optimization, a quick questionnaire was run by the JRC-IE at one international event organized in 2007 on maintenance optimization issues (JRC, 2008). The results are summarized in Table 4.1. Furthermore, Ukrainian, Slovenian, Czech and Russian representatives expressed on many occasions (Contri and Bieth, 2007) their interest in adopting a maintenance rule (MR) type approach in their countries, even starting on a voluntary basis, most probably closer to the ‘equipment reliability’ model (INPO, 2004). Many of them have already created some training centres,

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Scope of optimized MS&I (no. of systems)

M Cost Reduction Reduction Optimization issues in M cost in CDF process included? after M in place optimization

Reduction in outage duration (%)

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SFW used for M planning/ optimization

Network of spare parts available

No. of Risk indicators monitor on M available

Utility level

10

no

10

Russian Federation – RBMK1000

plans

Desna/ Primavera

Hungary

CBM

Passport/?

CBM

5%

yes

5%

no

40–28

Slovakia CBM <10% RCM - yes 10–15% n.a. 45–26 Arsoz/ EMO CBM Primavera

With 30 Bohunice with big parts

Ukraine

no

Bulgaria – diagnostic VVER1000

plans

plans

Primavera

Circ. plans no 51–45 Primavera With Pumps, Temelin containment, etc.

Czech Rep. CBM Adaptive yes expected - Dukovany maintenance model, RCM, cost benefit, use of PSA, etc.

Overall decrease 48% in outages duration From 2 to 3 types

Yes on living PSA

no Not no systematically

Passport/ yes Primavera/ MNT Graph, use of safety monitor system

Yes on living PSA

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Country Type of non-time- based M

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Table 4.1 Summary of the experience in selected European countries on specific PLiM-related issues

50% (all safety related)

CBM analysis

No, plans



32–22, Passport plans for 20 (every two years)

Few cases 10 with Candu owners

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Lithuania - CBM <10% CBM no 10% no 10% Fobos no ~15 Ignalina analysis (IFS)/ Primavera Note: M: Maintenance CBM: Condition-Based Monitoring RCM: Reliability-Centred Maintenance CDF: Core Damage Frequency (per reactor/year)

Yes based on living PSA Yes, only for pipes (300 mm), main circulation circuits

Probabilistic and deterministic safety assessment methods

Romania - CBM Cernavoda

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which are developing procedures in this direction. In conclusion, in relation to the operating cost reduction as a consequence of the kind of maintenance optimisation programme, the following reductions (Contri and Bieth, 2007) in maintenance costs/tasks were recorded: ∑ In Sweden, 10–20% of the effort, especially for I&C calibration intervals. ∑ In Spain, 20% in work, 30% in number of tasks. ∑ In Hungary, expected, not quantified. ∑ In Czech Republic, 30% on a restricted number of systems selected for a benchmark (according to a pilot project implemented at the Dukovany NPP). ∑ In Slovakia, expected, not quantified. It was noted that the ‘equipment reliability’ programme is not mandatory in most of the countries (including the United States). However, it is gaining growing interest due to its systematic approach to the management of plant safety. In particular, the correlation among the many existing safety-related programmes and the consistent classification of items (important, critical, run-to-failure) seems to be very attractive and practical. The main goals of the maintenance programme in the framework of PLiM can be summarized as follows: ∑ ∑ ∑ ∑ ∑ ∑ ∑ ∑

assure plant safety maintain optimal plant availability optimize operation and maintenance costs assure and develop industrial safety comply with codes, legislation and regulation decrease failures of safety and availability of critical components find and implement performance improvements increase the reliability and maintainability of the machinery and the performance of maintenance support ∑ increase the economic service life of equipment and plant. A precondition for a MS&I programme to be effectively inserted into a PLiM structure is that the maintenance programme is continuously optimized on the basis of the risk importance of any SSC, controlling the overall reliability of operation and preventing functional failures. One way of achieving this goal is the implementation of a rigorous SSC classification and the selection of the most appropriate maintenance tasks (and periodicity) for each type of equipment. The classification should be based on the importance to safety and operation, requirements of nuclear regulations, replacement costs, environmental risks and maintenance experience. Classification can be changed based on operation and maintenance experience. The maintenance unit should be responsible for executing and

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planning normal maintenance work, carrying out the work in refurbishment and investment projects, carrying out inspections and periodical tests, managing spare parts, organizing personnel training and enforcing an appropriate QA system.

4.3

A unified proposal for a plant life management (PLiM) model integrating maintenance optimization

4.3.1 Introduction Previous sections highlighted the main issues behind the development of a PLiM model, its main features and the experience of a few European and nonEuropean countries in this effort. As a consequence, the JRC-IE researchers developed a new version of a new PLiM model that they believed could significantly improve the performance of the nuclear plants. A first draft of this model is available at JRC (2008). The model was subsequently validated at one European plant that is believed to have one of the most advanced PLiM models in place. It is summarized in the following sections.

4.3.2 PLiM objectives PLiM can be defined as a programme (or even a combination of programmes and procedures) aiming at a safe and cost-effective operation of a nuclear power plant over the longest possible time period, including plant long-term operation (LTO). In this sense, it represents a framework for optimized, day-to-day decision making aiming at a plant LTO with optimum utilization of resources. In other words, the objective of PLiM is the development of a consistent framework programme at the plant that enables the plant to produce electricity in a safe and responsible way by continuously improving the power plant operation and safety. This objective is typically achieved with co-ordination of some key programmes at the plant, such as operation, asset management, maintenance surveillance and inspection (MS&I), ageing management, knowledge management and nuclear safety.

4.3.3 Approach to PLiM In order to achieve the goals set up in the previous section, the PLiM programme has to consider the following main components: 1. Nuclear safety and licensing. 2. Production and economy (including fuel and waste management). 3. Human resources.

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The long-term investment plan is the basic tool for managing the investment portfolio where all the technical programmes provide input. The generic PLiM structure is the result of the integration of selected existing programmes at the plant and the development of suitable links and feedback loops. In particular, the following programmes are directly coordinated by PLiM: ∑

maintenance, surveillance and inspection (MS&I), including control of human factors ∑ ageing management, component obsolescence and plant configuration control ∑ knowledge management ∑ asset management and investment planning. Plant modernization, power uprating and fuel management may also be part of PLiM, but they are not necessarily implemented at all plants. This concept is described in Fig. 4.4, where the four main components of PLiM are highlighted in the central programme, the input and the output are in the vertical lines and other programmes are listed in the lateral boxes. These programmes should also meet specific preconditions on their main features, as discussed above and summarized in Fig. 4.5. Other programmes represent a generic background for PLiM, and exchange data with PLiM, but they are not explicitly part of it, such as operation, nuclear safety, fuel management, waste management, licensing (including the continuous updating of the Safety Analysis Report), engineering, etc. Market trends

Other programmes Operation Nucl. safety Licensing Fuel man. Waste man. Safety analysis Etc.

Plex-LR-PSR

Safety regulations

Internal operating experience

R&D Maintenance, surveillance & inspections Ageing management

PLIM

Component integrity Thermal Hydraulics

Asset management

Safety assessment Human resources

Investment plan

Etc.

Plant upgrading

4.4 Approach to PLiM and interfaces with related programmes.

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Reliability-based Optimized

Long-term trends Environmental variables Obsolescence

Outage optimization Fuel management Spare part management

Maintenance, surveillance & inspections

Ageing management

PLIM

Asset management

Human resources

Knowledge management, in time Human reliability programme Public acceptance

4.5 Preconditions for the key programmes to be part of PLiM.

Finally, important programmes may be based upon PLiM, but they are not part of it, such as plant life maintenance and assurance after design life, licence renewal, periodic safety review, plant upgrading (including power uprating), public acceptance, etc. From the technical standpoint, the approach to plant life management consists of: ∑ ∑ ∑ ∑ ∑ ∑ ∑

identification of critical systems, structures and components (SSCs) from the standpoint of the plant operation and safety classification of the identified SSCs identification of operational loadings, stressors and ageing mechanisms development of method for the lifetime prediction identification and implementation of applicable ageing countermeasures feedback to MS&I programmes and other relevant programmes development of the investment planning and business case aspects.

4.3.4 PLiM scope – component classification A consistent application of PLiM suggests the use of a dedicated component classification, where safety, availability and cost issues contribute to form the

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criteria. All structures, systems and components at the plant, regardless of their safety relevance, should be covered by such classification. According to the classification, a suitable grading of measures may be applied and therefore different levels of MS&I and AMP, economic analysis, etc., may be assigned. A proposal (see also JRC, 2008) may group different classes as in the following: ∑

Class A: critical components and structures directly limiting the plant life with their availability/integrity, and essential non-replaceability. Examples include reactor pressure vessel, steam generator (SG are replaceable … but expensive), pressurizer, main coolant pump, containment structures and main buildings. Those items are in principle replaceable, but the impact on both business interruption and budget would be huge, so it is best to discourage their consideration in the list of is replaceable items. Example of MS&I strategy: full scope monitoring and analysis of the degradation. ∑ Class B: critical components, systems and structures from the standpoint of their importance to safety and their cost of replacement/reparation. Examples include primary circuit, high and low pressure safety injection systems, feedwater system, condensers, turbine, generators, diesels. Example of MS&I strategy: condition-based MS&I. ∑ Class C: sensitive components, from both safety and economics standpoint, systems and structures. Examples include nuclear intermediate cooling, sprinkler, drainage and vents, main steam line, residual heat removal, circulating and service water systems, condenser cooling system. Example of MS&I strategy: preventive (time-based) MS&I. ∑ Class D: other components and structures. Examples include condenser purification system, auxiliary boiler plant, drinking water supply, sewerage. Example of MS&I strategy: run-to-failure. It is noted that such an approach is still quite heterogeneous, as it mixes various components, systems and special equipment. Therefore, the proposed classification may be reviewed to provide a more homogeneous approach that would make the interfaces with the maintenance classification or spare part classification much easier and traceable. For components in classes A, B and C, a type of component ‘health certificate or dossier’ is recommended for continuous review and upgrading by the system engineers. The dossier should make reference to the design basis and should collect the results from the AMP, the operation and the ISI programmes, including the pending issues detected whilst previous tasks (e.g. inspections, monitoring) were being performed.

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4.3.5 The organizational structure that supports the PLiM – the system engineers The PLiM programme requires important changes in the traditional plant organization. In particular, the following preparatory organizational tasks should be implemented: ∑ ∑

development of a PLiM supervision and co-ordination unit nomination of system engineers, in full-time charge of selected systems, especially for Class A ∑ identification of research and engineering specialists in different disciplines either at the plant or at the technical support organization (TSO), ready to co-operate with the system engineers to address methodological issues, interpretation of results, interfaces with the scientific and engineering community, etc. System engineers are the main interpreters of the PLiM programme at the plant; they should be responsible for the life management of a particular system, structure or component. They represent the system ‘owners’. These engineers are typically responsible for the following tasks (JRC, 2008): ∑ ∑ ∑ ∑ ∑ ∑ ∑

preparation and control of inspection, monitoring and maintenance activities related to life management of systems, structures and components critical to safety detection and assessment of ageing mechanisms and effects preparation and implementation of improvements in the field of proactive maintenance. maintain and be proactive in the the life management system update the information in the long range planning system keep updated records of the component/system health status in the reference documentation system guarantee the reporting.

Some interfaces between the system engineers and other groups/departments are particularly important in the PLiM framework, namely: 1. The operators: plant technical specifications (TS) and operational limits and conditions (OLC) may be discussed and changed (with the due authorizations) as a consequence of detailed analysis of the operating experience and of the MS&I outcome. 2. The MS&I technicians: objectives, periodicity, scope and other attributes of the programmes may be agreed and modified. 3. The safety specialists (either on-site or at the TSO): they are responsible for the plant safety analysis and therefore all the acceptance criteria for ageing and degradation should be agreed and reviewed with them. 4. The technical support group: the decision to repair/replace/maintain a

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component or structure is taken jointly and approved by the management group of the PLiM.

4.4

Probabilistic safety assessment of components and systems

4.4.1 Applicability to different SSCs When a living PSA has to include ageing effects, the first task is the selection of the SSCs to be scoped. As a generic process, ageing could affect the performance and reliability of all NPP SSCs. However, differences in design, maintenance, operation stressors and environmental conditions make ageing kinetics and effects very different from one component to another. In addition, the risk associated with and the safety importance of the SSCs is not the same. To identify and prioritize SSCs potentially sensitive to ageing, the following technique is proposed: ∑

prioritize the SSCs which are modelled in PSA using risk importance factors ∑ perform trend analysis of available reliability data ∑ use qualitative assessment or ageing failure modes and effect analysis (AFMEA) for a limited number of components. All three steps are consecutive and complementary. Prioritization by risk importance is an effective way of reducing the list of SSCs under consideration. Application of Fussell–Vesely (or risk decreasing) and risk increasing factors has been demonstrated via case studies (Nitoi and Rodionov, 2008, Poghosyan et al., 2008). Trend analysis helps to identify SSCs where failure intensity (failure rate) increases with time. This could be a direct indication of ageing. Given a large population of SSCs and a consolidated collection of operating experience data, statistical methods could be applied. One basic task of statistical analysis is to investigate whether the SSC failure rate is approximately constant. Various statistical tests could be used to validate or to refute the assumption of constant failure rate. Some of them are discussed, for example, in Atwood et al. (2003). Depending on the statistical technique used, they can be divided into three groups: ∑ graphs (visual evaluation) ∑ non-parametric hypothesis tests ∑ parametric hypothesis tests. Currently, most nuclear utilities have a good reliability data collection system and a large amount of data on safety-important components from probabilistic safety assessments. This data could be used as a basis for ageing analysis.

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Several case studies have been performed by the scientific community to develop and demonstrate the applicability of statistical methods (see, for example, Rodionov et al., 2008, Antonov et al., 2009 and Kelly et al., 2008). Qualitative assessment or AFMEA is aimed at identifying and characterizing the link between the potential ageing mechanism and failure modes, and to demonstrate both the effect of failure on system reliability and the effectiveness of the surveillance strategy. In addition, the results of the analysis help to postulate assumptions concerning the periodicity and degree of component renewal after maintenance. This assessment could cover SSCs initially neglected in PSA models and SSCs for which there is not enough failure data for trend analysis. The main problem of qualitative assessment is that it is very time-consuming and resource-intensive. Some case study results demonstrating the quantitative assessment approach are presented in Nitoi et al. (2008) and Poghosyan et al. (2008).

4.5

Impact of ageing effects at system and plant level

4.5.1 Setting up a sample case A sample case can be developed to demonstrate how to integrate time-dependent reliability parameters into a PSA study and to calculate the impact of ageing on the risk profile as a function of time. Such a case study was developed in Rodionov et al. (2008), where it was proposed to use core damage frequency (CDF) as the average value at one-year intervals calculated for different age points, for example for 10, 20, 30 and 40 years in operation. A three-loop PWR PSA model for a large loss of coolant accident (LLOCA) initiating event was considered for this purpose. The model consisted of four event trees developed for full power operation and hot shutdown reactor states. The set of ‘virtual’ reliability data was prepared on the basis of the results of case studies, available generic data sources and expert opinions. The data included time-dependent reliability models for certain mechanical, electrical and I&C components of low pressure safety injection (LPSI) and containment spray systems (CSS). For most components, the relative increase in failure rate (probability) proved not very significant (see example in Fig. 4.6), while a strong ageing behaviour was considered for one component type – pump motors (see Fig. 4.7). In this case, for the best fitted log-linear model (p-value = 0.98), the relative increase in failure probability was more than three orders of magnitude towards the end of the design lifetime. At the same time, the Weibull model, which fits with a significance level of 0.96, gave a relative increase by a factor of 20 towards the end of the design lifetime.

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‘As bad as old’ preventive maintenance was considered in all cases. Quantification has been calculated for a reference value (no ageing effects considered) and age points of 10, 20, 30 and 40 years.

4.5.2 Effects at the system level The results of calculations performed for the CSS fault tree show that system unavailability increases with time by more than one order of magnitude with regard to the reference value (see Fig. 4.8). Up to the age of 30 years, the main contributor to unavailability (Fig. 4.6) is the failure in the level sensors in the reactor water storage tank (RWST), which provides a signal to switch the CSS to containment sump recirculation. But at the age of 40, the dominant impact on system unavailability is failure of the CSS pump motors. Thus, a rapid and sharp increase in unavailability due to the pump motors

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Css unavailability

9.00E-03 8.00E-03 7.00E-03 6.00E-03 5.00E-03 4.00E-03 3.00E-03 2.00E-03 1.00E-03 0.00E+00

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can be explained by the choice of the log-linear model for the failure rate. As mentioned in the previous section, this model provides more conservative extrapolation results. For comparison purposes, two contributors, with constant failure probability, are presented in Fig. 4.9: human error (20% fractional contribution to the reference value) and CSS pumps fail to start (2% contribution to the reference value). As can be seen from the graph, their contribution to the total system unavailability gradually decreases with time.

4.5.3 Effects at the plant level Figure 4.10 shows the impact of the ageing of selected components on CDF. The result of risk extrapolation is an increase in CDF from 6.58 ¥ 10–8 at

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4.11 CDF as a function of unit age – sensitivity analysis.

10 years to 6.19 ¥ 10–7 at 40 years. In comparison with the reference value (7.18 ¥ 10–8) the increase is by a factor of 8.6 by the end of the designed lifetime. Once the most sensitive components (low pressure safety injection (LPSI) and CSS pump motors and level sensors) are in the minimum cut sets of dominant sequences, the relative contribution of the sequences to the total risk of LLOCA remains approximately the same with age. The same picture can be seen for contributions to the risk associated with the different reactor states and location of the pipe break. Figure 4.11. shows the results of the sensitivity analysis of the reliability model chosen for the most sensitive components, i.e. pump motors. For this component in the reference case, the log-linear model was considered. As mentioned in previous sections (see Fig. 4.7), the Weibull model provides quite different values for extrapolation of failure probability to the end of the design lifetime than the sensitivity analysis.

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The generic conclusion from this analysis is the need to examine the accuracy of several model alternatives before applying any given one to PSA. For the reference case, an analysis of risk importance measures was performed. Figure 4.12 shows the variation of Fussell–Vesely importance for the main contributors to CDF. The nature of the curves is about the same as for the fractional contribution of the failure of a particular component to system unavailability shown in Fig. 4.9. For components not sensitive to ageing, the Fussell–Vesely importance monotonically decreases with time (see LPSI valves or human errors (HE) on RWST sensors). For components sensitive to ageing, the behaviour of the measure could differ according to rate of ageing (see, for example, LPSI pumps and CSS pump motors). In this example, the most dramatic changes in component importance take place between 30 and 40 years for the most sensitive components. The risk increasing factor monotonically decreases with time for all the main contributors to CDF (see example of LPSI valves (not sensitive to ageing) and CSS pump motors (very sensitive to ageing) in Fig. 4.13). However, the most sensitive components remain the same. This behaviour makes the risk importance factor less informative from a decision-making point of view. The basic conclusion is that ageing can alter the risk importance values for particular components and failure modes. This has to be taken into account when applying these measures for component prioritization (for operation or maintenance optimization). The sample case discussed above provides useful insights for a development of a time-dependent PSA study, where ageing effects are properly included.

4.6

Conclusions

Ageing effects may negatively impact plant safety in time. The deterministic assessment of these ageing effects has to refer to the design basis of components 1.80E-01 1.60E-01 1.40E-01 1.20E-01 1.00E-01 8.00E-02 6.00E-02 4.00E-02 2.00E-02 0.00E+00

Fussel–Vesely importance LPSI pump fail to run LPSI valves fail to open HE on RWST sensors CSS pump fail to run RWST level sensors CSS pump's motor fail to start 10

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4.12 Fussel–Vesely importance measure.

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4.13 Risk increasing factor.

and systems and provides information to owners and operators on the need for changes in OLC and/or component replacement. A well-integrated model for PLiM, integrating safety and non-safety programmes, is essential to provide a framework for such assessment and for an effective management of the plant assets in time. Guidelines are proposed for development of those models, adapted to the national context. The proposed probabilistic approach, though more complex from the computational standpoint, provides direct measurement of the effects of ageing on the plant CDF, the dominant accident sequences and contributors to CDF, and the component risk importance measures. Considering ageing effects in PSA, and reliability analysis in general, can help in the selection and prioritization of SSCs and in planning optimum ageing management and maintenance measures as part of a risk-informed decision-making process. The main concerns, which may limit a broad application of the proposed methods, are still related to methodology complexities, data and resource availability. However, the methods presented in this chapter provide some insights and suggest extrapolation techniques for the reliability parameters that may reduce the computational effort and provide reliable, ready-to-use results.

4.7

References

Antonov, A., Chepurko, V., Polyakov, A., Rodionov, A. (2009), ‘Application of generalised linear model for time-dependent trend assessment – a case study for the ageing PSA network’. Reliability Engineering and System Safety, 94 (6), 1021–1029. AREVA (2007), EPR, 2007, www.areva-np.com Atwood, C., LaChance, J.L., Martz, H.F., Anderson, D.J., Englehardt, M., Whitehead, D., Wheeler, T. (2003), Handbook of Parameter Estimation for Probabilistic Risk Assessment, NUREG/CR-6823, US Nuclear Regulatory Commission, Washington DC. © Woodhead Publishing Limited, 2010

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Contri, P. (2007a), ‘Plant life management models with special emphasis to the integration of safety with non-safety related programs’, PLiM Conference, Shanghai, 15–18 October 2007. Contri, P. (2007b), ‘Maintenance issues in long term operation of nuclear power plants’, ENC 2007, Brussels, 19 September 2007. Contri, P., Bieth, M. (2007), ‘A plant life management model as support to plant life extension programmes of nuclear installations – effective integration of the safety programmes into an overall optimization of the operating cost’, joint research center institute for energy. Contri, P., Debarberis, L., Taylor, N. (2007a), ‘Needs in R&D Supporting Nuclear Power Plant Life Management’, PLiM Conference, Shanghai, 15–18 October 2007. Contri, P. et al. (2007b), ‘Maintenance optimisation issues in the EC’ - ISEM 2007, Lansing, 10 September 2007. EU (2007), Work programme 2007, Euratom for nuclear research and training activities, European Commission C (2007) 564 of 26.02.07. IAEA (1992), Safety Related Maintenance in the Framework of the Reliability Centred Maintenance Concept, IAEA-TECDOC-658, IAEA, Vienna. IAEA (1997a), Regulatory Surveillance of Safety Related Maintenance at Nuclear Power Plants, IAEA-TECDOC-960, IAEA, Vienna. IAEA (1997b), Good Practices for Cost Effective Maintenance of Nuclear Power Plants, IAEA-TECDOC-928, IAEA, Vienna. IAEA (2000a), Safety of Nuclear Power Plants: Operation Requirements, Safety Standards Series No. NS-R-2, IAEA, Vienna. IAEA (2000b), Technologies for Improving Current and Future Light Water Reactor Operation and Maintenance: Development on the Basis of Experience, IAEA-TECDOC1175, IAEA, Vienna. IAEA (2000c), Advances in Safety Related Maintenance, IAEA-TECDOC-1138, IAEA, Vienna. IAEA (2001), Applications of Probabilistic Safety Assessment (PSA) for Nuclear Power Plants, IAEA-TECDOC-1200, IAEA, Vienna. IAEA (2002a), Safety Assessment and Verification for Nuclear Power Plants, Safety Guide, Safety Standards Series No. NS-G-1.2, IAEA, Vienna. IAEA (2002b), Maintenance, Surveillance and In-Service Inspection in Nuclear Power Plants; Safety Guide, Safety Standards Series No. NS-G-2.6, IAEA, Vienna. IAEA (2004a), Management of Life Cycle and Ageing at Nuclear Power Plants: Improved I&C Maintenance, IAEA-TECDOC-1402, IAEA, Vienna. IAEA (2004b), Guidance for Optimizing Nuclear Power Plant Maintenance Programmes; IAEA-TECDOC-1383, IAEA, Vienna. IAEA (2006), Principles and Guidelines on Plant Life Management for Long Term Operation of Light Water Reactors, IAEA Technical Reports Series 448, IAEA, Vienna. INPO, (Institute of Nuclear Power Operations) (2004), Equipment Reliability Process Description, INPO AP-913, Atlanta. INSAG-12 (1999), Basic Safety Principles for Nuclear Power Plants, 75-INSAG-3, Rev. 1, IAEA, Vienna. JRC (2006), EUR 21903 EN: ‘Optimization of maintenance programmes at NPPs – benchmarking study on implemented organizational schemes, advanced methods and strategies for maintenance optimization - Summary report’, January. JRC (2008), EUR 23232 EN, ‘A plant life management model including optimized MS&I programme – Safety and economic issues’, JRC EUR report, January.

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Kelly, D., Rodionov, A., Klugel, J-U. (2008), Practical Issues in Component Aging Analysis. Proceedings of PSA’2008 International Topical Meeting, Knoxville, USA, 7–11 September 2008, ANS, La Grange Park. Nitoi, M., Rodionov, A. (2008), Qualitative approach for selection of Systems, Structures and Components to be considered in Ageing PSA, EUR23446EN. Petten: EC DG JRC Institute for Energy. Poghosyan, Sh., Malkhasyan, A., Rodionov, A. (2008), Components Selection for Ageing PSA of Armenian NPP Unit 2. Proceedings of PSA’2008 International Topical Meeting, Knoxville, USA, 7–11 September 2008, ANS, La Grange Park. Rodionov, A., Atwood, C., Kirchsteiger, Ch., Patric, M. (2008), Demonstration of statistical approaches to identify component’s ageing by operational data analysis – A case study for the ageing PSA network. Reliability Engineering and System Safety, 93, 1534–1542. US-NRC (2002), An approach to using probabilistic risk assessment in risk-informed decisions on plant-specific changes to the licensing basis. US NRC Regulatory Guide 1.174. Rev.1. US NRC, Washington, DC. Vaisnys, P., Contri, P., Biéth, M. (2007), Monitoring the effectiveness of maintenance programs through the use of peformance indicators, ENC 2007, 19 September 2007, Brussels. Westinghouse (2007), AP1000, www.AP1000.westinghousenuclear.com.

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5

Assessing the socio-economic impacts of ageing and plant life management (AM-PLiM) programmes for long-term operation (LTO) of nuclear power plants (NPPs)

P h. G. T i pp i n g, Nuclear Energy and Materials Consultant, Switzerland

Abstract: This chapter provides a review of the manpower, fuel supply and economic considerations associated with ageing management (AM) and plant life management (PLiM) programmes in nuclear power plants (NPPs). The chapter underlines the sound economic value of implementing standard operational practices (OPs), ageing surveillance programmes (ASPs), AM and PLiM in a timely manner. These approaches are used not only to assure safe, profitable and reliable operation, but also to put NPPs in a favourable position to continue operation in excess of their original design lifetimes. The business case demands that the costs involved in investing in the repair, refurbishment, replacement of systems, structures and components (SSCs) will have to be amortized over the plant’s operation lifetime. Power uprates (PUs) in NPPs are shown as a way to further increase profits. Longterm operation (LTO) is identified as a sound economical way in which NPP stakeholders can amortize costs over a longer time period and make additional operational profits. Key words: new nuclear power plants (NPPs), economics aspects of AM, ASPs and PLiM for LTO, repair, refurbishment and replacement of SSCs, PU, amortization, full lifecycle, licence renewal, greenhouse gas emissions, radwaste, knowledge management, succession planning, nuclear-based fuel resources.

5.1

Nuclear power as part of the global energy mix: energy demand, environmental issues and manpower

In January 2009, British Nuclear Fuels (BNFL) and British Energy (BE) expressed the wish to develop and construct new NPPs to replace existing ones as they reach their true end-of-life status [1]. If we consider the fact that nuclear power currently accounts for about 19% of electricity generated in the UK, the urgency of the matter is clear; both companies argue that without the contribution of nuclear power, the UK will be unable to meet energy demands, or targets for cutting greenhouse gas emissions (GGEs). A report 117 © Woodhead Publishing Limited, 2010

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from the Royal Commission on Environmental Pollution suggested that the UK should make 60% cuts in CO2 emissions by 2050, and 80% by 2100 if it is to play a significant role in overall efforts to counter global warming [2]. The UK government passed the Climate Change Act in 2008 (updated 2009), committing the UK to reducing carbon dioxide (CO2) emissions by at least 34% by 2020 and by at least 80% by 2050, relative to 1990 levels [3]. Nuclear power has a relatively low carbon footprint (CF), making the development of the nuclear power industry the only practical route to take in order to ensure a ‘lower-carbon’ future for the UK. When it was reported in April 2009 that a list of potential sites for new NPPs had been unveiled by the UK government, the UK Energy and Climate Change Secretary stated that it was an important step towards a new generation of nuclear power stations, as part of the low carbon future of [Great] Britain [4]. The UK government now plans to narrow the ‘electrical energy generation gap’ that will arise when existing NPPs (and coal-fired stations) are progressively taken out of service. Furthermore, the building of new NPPs has the potential to offer thousands of jobs to the UK, and multi-million pound opportunities to British businesses. By the year 2020, the European Union has also set a goal to reduce industrial CO2 emissions by 20%. This is an ambitious plan if safe, clean, reliable and economically competitive energy supplies are to keep pace with demand. The case for maintaining and developing nuclear-generated power has never been stronger, when we recognize that export-driven economies require cheap and ecologically benign energy to create products and jobs. Many of the environmental problems that humanity faces today have their origins in a time about 220 years ago, at the beginning of the fossilfuelled (coal, oil and gas) industrial revolution. Moreover, the Earth’s human population has increased nearly seven-fold since 1900, which has also caused an insatiable demand for food, as well as energy. The steady rise of ‘greenhouse gas’, especially CO2, in the Earth’s atmosphere since the 1900s bears witness to the fact that the processing and burning of fossil-based fuels to produce energy and provide transportation is most probably linked to the extreme effects of climate change being observed today. Additional environmental pollutants such as sulphur dioxide (SO2) also arise when fossil fuels are burned, which causes the formation of acid rain. Furthermore, increased methane gas emissions due to commercial-scale farming has a real potential for increasing the ‘greenhouse gas’ inventory [5]. Methane (CH4) has about 20 times the heat-trapping effect of an equivalent amount of CO2. If the currently observed rate of global warming continues, additional CH4 is likely to be released at increasing rates from underwater sources, such as lakes and oceans, thus further exacerbating the problem of climate change. It is a common moral obligation for both developed and developing

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countries to ensure that present human actions will not leave a chemically and heat-polluted planet as an untenable legacy to future generations. However, though nuclear power has a low CF and operates free from GGE, we cannot ignore the fact that it creates radioactive waste (radwaste). It is therefore of paramount importance to collect, condition and dispose of radwaste in an efficient and safe way. Nuclear power and disposal of radwaste must remain safe technologies in order to maintain a high level of public acceptance. Furthermore, as a mature industry with an ever-evolving future, the nuclear power industry will also create attractive career prospects for young technologists. The challenges presented by the replacement of decommissioned NPPs and increasing the share of nuclear power with new generation NPPs also includes the difficulty of providing sufficient numbers of well-trained engineers, technicians and operating personnel. It is here that NPP owners and operators, national and international organizations, colleges, universities and research institutes must provide environments that actively promote and impart expertise in all aspects of nuclear technology. The vast majority of pioneer and senior generation nuclear power plant designers, technologists, operators and regulators working between 1955 and 2010 have now retired. After the Chernobyl NPP accident in 1986, large numbers of young technology students no longer regarded nuclear science as a career path [6], and many educational institutions dropped nuclear power technology and engineering completely from their teaching syllabuses. The situation has now changed, since many countries now recognize that nuclear power still has a key role to play in the world’s ecological and economic future development. It is thus essential to the future of nuclear power to act now in order to capture the current knowledge concerning, for example, optimum NPP design and principles of maintaining safe operation, inspection and regulation in order that plant personnel are able to make sound judgements in the future. If the world is to develop further industrially, whilst acknowledging the significant GGE, CF and pollution issues associated with the use of fossilbased fuels, convincing arguments exist for the continued use and expansion of commercial nuclear power from both ecological and socio-economic points of view.

5.2

Aspects of current and future nuclear fuel supply and its impact on the viability of nuclear power

It is expected that pollution issues will increasingly influence our way of thinking about all forms of energy use in the future. At least for the next 100–150 years, fission-based (current technology) commercial nuclear power is expected to remain a valuable source of energy, particularly if

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the currently known uranium fuel resources are used efficiently. Continued and expanded use of nuclear power will still be significantly dependent on its public acceptance, and this requires that nuclear power remains a safe technology. At current and projected prices, nuclear fuel is likely to remain a very cost-effective energy carrier. For example, at 45,000 MWd per tonne burn-up, 1 kg of enriched uranium fuel (UO2) will produce about 360 000 kWh of electricity. When conversion, enrichment and fabrication costs are considered, this gives a fuel cost of only about 0.5 US cents per kWh (in 2009). Natural uranium resources suitable for converting into fuel are thought to be sufficient to cover the industries’ requirements for at least 150 years. Indeed, with other fuel cycles (fast breeder reactor concept), increased efficiency and reprocessing of spent fuel, as well as decreasing the degree of enrichment in some ‘high-enriched-uranium’ 235U fuels, it is thought that this could meet demand in excess of 1000 years. Thorium (Th) may also be used as a source of nuclear-based fuel, and, with easy-to-recover reserves of at least 4.4 million tonnes, it could be even more abundant than uranium [7]. Although not fissile itself, the isotope 232Th can absorb slow (thermalized) neutrons to ‘breed’ the isotope 233U, which is fissile and long-lived. An attractive aspect of using the thorium fuel cycle is that less plutonium and other transuranic elements are created when compared with the direct uranium fuel cycle. This has great potential for lowering plutonium inventories, and thus easing many radwaste and nuclear proliferation issues. (Note: The fuels used for nuclear fusion reactors (which are beyond the scope of this book) are isotopes of hydrogen (deuterium (2H) and tritium (3H)) and they are deemed to be practically inexhaustible. Fusion reactor concepts are currently under research and development, and may consequently only become a commercially viable reality towards the latter half of the present century.)

5.3

Economic overview of the nuclear power plant (NPP) lifecycle

NPPs are designed and operated to safely and reliably produce electrical power at a profit. They are particularly suited to the task of providing constant base-load power to the grid. This is particularly relevant if nuclear energy is only a part (e.g. up to about 20–30%) of the overall energy production mix of a country, and alternative sources of energy are available which can be readily adjusted to follow peaks and troughs in demand. In some countries where the majority of electrical power (e.g. about 75%) is nuclear generated, such as in France, the NPPs’ outputs have to be adjusted to follow demand as there is less scope for flexibility. By way of comparison, the UK (with about 19% nuclear-generated power), never designed, licensed or operated

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the first and second generation of NPPs to operate for following varying demands for power, and thus their role continues to be to provide the baseload to the grid [8]. The longer a NPP operates at its full allowable rated power, free from forced outages or safety-related operational restrictions, the more money it creates. Nuclear power plants are expensive to build, but relatively cheap to operate; it is usually only in the last years of their original design lives that they will approach amortization. Therefore, it makes good business sense to first reach the NPP design life, and, having respect for the ability of the NPP concerned to fulfil its licensing conditions (safety aspects), and continue operation. The ‘design life’ of a NPP may be regarded as a relative term, since it is broadly based on usually very conservative engineering assessments of SSCs and how sufficient safety margins thereof may be maintained for as long as possible. Furthermore, there are only a few items in a NPP that truly are life-determining, namely the large, passive SSCs that are practically, technically or economically impossible to replace. A significant focus of OPs, AM, ASPs and PLiM programmes to counter or mitigate ageing degradation (AD) (and to maintain sufficient safety margins) is naturally placed on these SSCs because the longer they are kept fit-for-service, the longer the NPP’s true operational life will be. This leads to a better chance to protect the overall plant investment. Furthermore, SSCs that can be replaced as a matter of routine also benefit from standard OPs, monitoring or maintenance, since they may be kept in service in excess of their nominal design lives if their true condition is known. Replacement costs may then be postponed to a later point in time, or even totally avoided. However, safety must have priority over economic aspects; any safety-related events can potentially become more costly than short-term savings made on delaying repairs or replacements of SSCs. Efforts to minimize AD have a cost, but minimizing contributors that cause premature or unnecessary replacement of SSCs represents avoided costs. The total cost of building and operating a NPP depends on many complex factors. Environmental impact qualification, location of site, access roads, the construction of power lines, land procurement, construction, commissioning, operation, radwaste treatment, radwaste disposal, and eventually decommissioning, are just a few. Other cost factors, apart from securing enough qualified personnel for all aspects of operation, include fuel, licence fees, additional expenditures associated with repairs and replacements of SSCs, and the inherent plant operational costs (e.g. cost of the OPs, ASPs, AM and PLiM programmes). Unusual or unforeseen costs may arise at any time during operation, such as the replacement of steam generators (SGs) and core shrouds, and even reactor pressure vessel (RPV) annealing. They are not only expensive items in themselves, but their replacement will also lead to extended outages and therefore low plant availability and consequently

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lost sales of electricity. For example, SG replacement caused by Alloy 600 tubing integrity issues has become a necessity for many older pressurized water reactors (PWRs). This is a major task and costs around US$150 million, depending on the design of the NPP involved. Such major investment costs are only likely to be recovered if the NPP concerned continues to operate and enters the LTO phase of its life. Re-licensing/licence renewal (LR) (e.g. US practice) or continued operation with periodic safety review (PSR) every 10 years (e.g. European practice) means that any such investment can then be amortized over a longer time period (e.g. 20 years) [9–12]. A typical cost of a LR procedure in the United States is about US$10–20 million. The present net value of LR, if all operating NPPs in the United States operate for 60 years, is approximately US$25 billion [13]. The LR process in the United States takes up to 5 years to perform, but the LR approach is becoming more streamlined as good practices and experience are continually implemented. It is recognized that LR is a relatively cost-effective way to maintain the supply of safe, clean energy, and by April 2009 the US regulator (US-NRC) had already re-licensed 52 NPPs (about half of the US fleet), and eventually up to 85 (from 104) NPPs could benefit from LR.

5.4

Cost drivers of nuclear power plant (NPP) operation

The largest unknown in determining the value of a NPP, particularly in the LTO phase, is the expected market price of electricity. This variable outweighs all other cost drivers, but it is practically impossible to accurately predict the supply–demand–price situation over 20 or more years into the future. However, existing NPPs with effective standard OPs, ASPs, AM and PLiM programmes have several practical and economical advantages over new-build NPPs. Apart from current operational profits, the projected future operation (i.e. LTO) will generate further turnover, income and profits. A NPP in the LTO phase benefits greatly from each extra kilowatt-hour of electricity sold. The overall economic situation is also more certain, since existing and currently operating NPPs are, among other things, free from construction permit delays and limits on heavy industrial capacity to supply new NPPs on time. Capital costs are also less, due to amortization and depreciation. Most utilities and operators have recognized this, and therefore keep their NPPs operating for as long as they remain safe. Financial considerations will decide when NPPs are no longer profitable to operate. However, when all options for SSC modifications, replacements, refurbishing are exhausted, and large, passive items such as containment, buildings and RPVs can no longer fulfil regulatory requirements, the real end of operational life of the NPP will be due to safety considerations. Ideally at this final stage, all

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remaining investments would have been amortized and decommissioning and radwaste disposal costs covered in order to return the site to a greenfield status.

5.5

Basic economic requirements for sustainable operation of nuclear power plants (NPPs)

There are three basic economic considerations that must be reviewed in the industry. These include the necessity of maximizing power generation/ sales (high plant availability and power sold at a market-competitive cost), minimizing operating costs (e.g. focused inspection, replacement of SSCs based on their actual condition and higher levels of fuel burn-up) and assessing the additional costs of operation. When large investments, such as SG replacement are necessary, a cost–benefit analysis must be made. This will take into consideration the projected remaining operating lifetime of the plant and the prospect of achieving complete amortization. Using a holistic approach, realistic economic targets must be identified and then weighed against the expected end-of-life of the plant (i.e. the full lifecycle) and any deductions for fixed and unexpected costs. This methodology is as follows: 1. Take into account current and projected safety issues, their ranking in importance and urgency, and the cost of implementing mandatory requirements for safety upgrade tasks. 2. Identify non-safety-related technical issues that may affect operational flexibility or power output and estimated costs. This may include SG replacement, secondary side equipment, turbine generator replacement or modification, instrumentation and control (I&C), including software, modernization and immediate costs associated with power uprate (PU) (if relevant to the plant). 3. Forecast the price evolution of electricity. 4. To counter the fact that parts may become obsolete, create a strategic inventory of functionally equivalent and qualified spare parts to minimize forced outage or replacement delivery times. 5. Plan for sufficient radwaste treatment and disposal facilities. 6. Consider the costs of decommissioning, in order to return the NPP site back to a greenfield status. (NB: plant dismantling/decommissioning costs are unevenly distributed over time, as they come at the true endof-life of the NPP. Such costs are dealt with by discounting methods, since their values are time-dependent.) Clearly, economic assessment of NPP and LTO operation is greatly influenced by factors that are relatively difficult to quantify or anticipate. However, we can assume that the demand for cheap, environmentally clean and reliable

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supplies of energy will increase with time. This provides a certainty in an otherwise complex economic situation.

5.6

Assessing the costs and economics of nuclear power plant (NPP) operation and the impact of ageing and plant life management (AM-PLiM) programmes for long-term operation (LTO)

Profitable operation of a NPP requires that the total income gained by selling energy (electrical/district heating) will exceed the total operational costs by a margin that covers all foreseeable economic risks. The situation is a difficult and dynamic one, since energy prices may increase or decrease. Furthermore, a NPP at the start of its operation will carry a large capitalcharge debt but will have relatively smaller operation and maintenance (O&M) costs to carry. Over the NPP’s operational life, the capital debt will decrease (e.g. debt repayment, depreciation), whilst the O&M costs will tend to increase through the replacement of SSCs, exceptional maintenance, increased regulatory requirements and monitoring/inspection due to SSC-AD. Furthermore, capital charges and O&M tasks will still cause costs even when the NPP is not generating money through energy sales. To ensure financial success it is therefore necessary to price the energy so as to sufficiently cover all costs. Avoidable costs consist of many factors such as the cost of routine O&M, fuel, future capital expenditure and improvements to the NPP SSCs in certain years. The NPP outage time (forced or planned) must also be factored into the replacement capacity required, charged at the current market price. When no improvements or changes to the NPP are necessary, the avoidable costs will be the sum of O&M and fuel costs. In order to judge the economic viability of NPP operation and PLiM programmes for LTO, three factors must be considered: the cost of continued operation (compared to NPP shutdown and costs of replacement energy), overall net present value (NPV) of any necessary improvement programme, and the internal rate of return (IRR) of the improvement costs. The IRR is the discount rate at which the NPV becomes zero. However, a utilities’ decision to carry out standard or special OPs, AM, ASPs and PLiM will be influenced primarily by safety considerations. It needs a commitment to invest resources for keeping SSCs safe and reliable, as the corporate goal must be to operate the NPP at the best achievable level of safety for at least as long as it takes to amortize such expenditures. Clearly, OPs, AM, ASPs and PLiM will also cause costs, but, according to general experience, freedom from forced outages due to unreliable SSCs will more than counter these expenditures. Regulatory requirements (or political issues) may, however, lead to an earlier than planned-for cessation of NPP operation, © Woodhead Publishing Limited, 2010

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thus causing financial loss. Due to their very nature, however, such aspects are virtually impossible to predict. Nevertheless, safely operated NPPs will also have the best prospects to go into LTO. In deregulated, competitive markets (increasingly the case for electrical power sales), a fully amortized NPP at the start of its new (LTO) licensed period after 35 years of operation is a valuable asset. Such a NPP can sell its power at marginal cost, and thus compete with alternative energy suppliers. The operating profit margin of the NPP will increase accordingly. The majority of countries embrace the option of continuing NPP operations after reaching the original design life, since it makes sound ecological and business sense. However, this option is clearly dependent upon national energy policies and relevant regulations/licensing conditions. Until the mid1990s, the industry, as a whole, tended to focus on achieving NPP design lives and then decommissioning the plants. Now, however, the focus has shifted towards ‘life-after-40’, since OPs, ASPs, AM and PLiM strategies have been shown to pay real dividends from both safety and profitability standpoints. Furthermore, optimized OPs, ASPs, AM and PLiM programmes are favourable precursors to demonstrate to a licensing authority that the plants are operated and maintained to the highest level of safety, increasing the likelihood of being licensed to enter the LTO phase. Another consideration is whether a NPP’s owners will apply for licensing permission to uprate from the originally designed power rating. Power uprate (PU) has a relatively short payback time and can add further value to a NPP. However, NPP-AM, ASPs and PLiM, as well as OPs may require corresponding modification in order to accommodate possible PU-driven SSC-AD effects. Furthermore, it may be necessary to revise shut-down and emergency procedures associated with increased fuel core power and coolant flow rates relative to the pre-PU. All these tasks cost money, but the end result must show that the increased power output has been achieved costeffectively, with no negative effect on the safety level or environment. A candidate NPP for a PU will thus need a comprehensive review of its SSCs in order to assure their fitness-for-service under conditions of increased temperature, pressure or flow. Such increases bring possible environmental impact issues that must also be reviewed to ensure compliance with requirements. For example, radiation levels may increase somewhat, due to the higher power level of the reactor, or the temperature of plant-discharged water to an adjacent river may approach allowed maximum limits. Several trends and factors have favourably altered nuclear power’s economic position over the last 20 years. For example, operating costs of NPPs have steadily fallen as plant availability has increased. The overall steady increase in NPP’s availability, from 75 to 80% in the 1980s to 85 to 90% in 2009, is largely due to the positive effects of continually improving standards of OPs, AM, ASPs and PLiM programmes. Other cost-reducing factors include

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the reliability of the re-licensing process (especially in the United States) and lower than expected financing costs thereof, and increasingly efficient fuel cycles.

5.7

Conclusions

1. There are both financial and safety gains to be made when AM, ASPs and PLiM programmes are implemented along with standard OPs in NPPs since SSCs are more reliable, safety margins are maintained at sufficient levels, potential original design weaknesses are addressed, and both current and LTO is facilitated. 2. All costs associated with the NPP have to be recovered by depreciation and amortization before the end of the NPP’s original design life. When this is reached, LTO will further enhance the competitiveness and profit margins of the NPP. Equally, power uprates will increase the profitability of NPPs. Asset management is therefore a significant incentive for utilities to implement effective OPs, ASPs, AM and PLiM programmes. 3. No matter what type of NPP is under consideration, broad technical aspects remain the same. These include preventive maintenance, time-limited ageing analyses, maintenance of equipment, qualification of OPs, strategies to replace SSCs, strategies to combat original design inadequacies, mitigation or elimination of AD, and the robust implementation of stateof-the-art practices in all areas. A cost–benefit analysis will indicate the overall economic limits on proposed SSC repair, refurbishment, replacement and back-fitting actions, but regulatory safety requirements will dictate their extent. Safety must always have priority over economic considerations. 4. Power uprates (e.g. up to 20% relative to the original rating) are relatively cheap, and so may increase the profitability of NPPs. However, OPs, ASPs, AM and PLiM programmes may have to be modified to account for the new plant power rating status and configuration. 5. Since NPPs have relatively low power generation costs, there is a very strong economic incentive to maintain them with OPs, AM, ASPs and PLiM programmes and thus assure reliable current operation and to eventually facilitate LTO. 6. OPs, AM, ASPs and PLiM programmes have their own costs, but the economic benefit is high for current and long-term operation when OPs, AM, ASPs and PLiM are effective in achieving their respective goals. 7. Operational and business risks must be carefully assessed. Excluding a major accident, there are only two major risks to the overall operational life of a NPP to consider; namely when the non-replaceable SSCs can no longer fulfil their design and safety functions, and when major destruction and damage is caused by natural phenomena (e.g. seismic activity) or by terrorist attack. © Woodhead Publishing Limited, 2010

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Sources of further information and advice

Assurance of Nuclear Fuel Supply: Two Reports by the IAEA Director General. (Source: http://www.carnegieendowment.org/publications/index. cfm?fa=view&id=23256. Published 9 June 2009. Ayers, G., ‘Running nuclear power plants up to 50 years: economics and technology’, Nuclear Europe Worldscan, Vol. 11–12, 1996. Commissioning of Nuclear Power Plants: Training and human resource considerations, IAEA Nuclear Energy Series No. NG-T-2.2, STI/PUB/1334, IAEA, Vienna, 2008. Competency Assessments for Nuclear Industry Personnel, STI/PUB/1236, IAEA, Vienna, 2006. Decommissioning of Nuclear Facilities: Training and human resource considerations, IAEA Nuclear Energy Series No. NG-T-2.3, STI/PUB/1332, IAEA, Vienna, 2008. Efremenkov, V.M., ‘Chemistry and technology of radioactive waste management – the IAEA perspective’, Czechoslovak Journal of Physics, Vol. 53, Supplement 1, Part II, pp. A579–A587, 2007. Managing Human Resources in the Field of Nuclear Energy, IAEA Nuclear Energy Series No. NG-G-2.1, STI/PUB/1397, IAEA, Vienna, 2009. Cost Drivers for the Assessment of Nuclear Power Plant Life Extension. IAEA-TECDOC-1309, Vienna, September 2002. The New Economics of Nuclear Power, World Nuclear Association (WNA) report and press release from 1 December 2005. Economic Performance Indicators for Nuclear Power Plants. IAEA Technical Reports Series TRS-437, STI/DOC/010/437, IAEA, Vienna, 2006. Integrated Approach to Optimize Operation and Maintenance Costs for Operating Nuclear Power Plants, IAEA Technical Document 1509. Vienna, 2006.

5.9

Acknowledgements

The author gratefully acknowledges the helpful suggestions and guidance supplied by Dr B. Raj and Dr T. Jayakumar of the Indira Gandhi Centre for Atomic Research (IGCAR), Kalpakkam, India.

5.10

References

1. ‘Four potential sites for new nuclear power stations have been unveiled by the Nuclear Decommissioning Authority’. Report by R. Prince, Political Correspondent, Telegraph, 23 January 2009. 2. Energy – The Changing Climate, Royal Commission on Environmental Pollution Twenty Second Report, Cm 4794, June 2000. 3. Climate Change Act 2008, Department of Energy and Climate Change: http://www.

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decc.gov.uk/en/content/cms/legislation/cc_act_08/cc_act_08.aspx; accessed January 2010. 4. ‘Potential Nuclear Sites Unveiled’, Report in the Telegraph, citing E. Miliband (Energy and Climate Change Secretary), 15 April 2009. 5. Roosevelt, M., ‘Methane: The really scary greenhouse gas’, Los Angeles Times, 3 February 2009. Source taken from: http://www.huffingtonpost.com/2009/02/23/ methane-the-really-scary_n_169059.html 6. ‘The Chernobyl Accident: Updating of INSAG-1’, IAEA Safety Series No. 75INSAG-7, Vienna, 1992. 7. ‘Our changing Earth’, IAEA-Vienna, Bulletin March 2008, 49–2. 8. Nuttall, W., ‘Why is nuclear power baseload?’, EU Energy Policy Blog, 1 July 2007 (source: http://www.energypolicyblog.com/?p=45). 9. Nuclear Regulatory Commission, Requirements for Renewal of Operating Licenses for Nuclear Power Plants, 10 CFR Part 54, NRC, Washington, DC, 1995. 10. Nuclear Regulatory Commission, Requirements for Monitoring the Effectiveness of Maintenance at Nuclear Power Plants, 10 CFR Part 50.65, NRC, Washington, DC, 1996. 11. Nuclear Energy Institute, Industry Guideline for Implementing the Requirements of 10 CFR Part 54, The License Renewal Rule, Rep. NEI 95–10, NEI, Washington, DC, 2001. 12. Periodic safety review of nuclear power plants, Safety Guide, IAEA Safety Standards Series No. NS-G-2.10, IAEA, Vienna, 2003. 13. Plant life management for long-term operation of light water reactors – Principles and Guidelines, IAEA Technical Report Series No. 448, IAEA, Vienna, 2006.

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Part II Ageing degradation of irradiated materials in nuclear power plant systems, structures and components (SSC): mechanisms, effects and mitigation techniques

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6

Failure prevention and analysis in nuclear power plant systems, structures and components (SSC): a holistic approach

P h. G. T i pp i n g, Nuclear Energy and Materials Consultant, Switzerland

Abstract: Requirements for NPP SSC quality, defence-in-depth (DID) and accident management are discussed within the context of avoiding SSC failures and, should they occur, limiting their consequences. The concept of latent failure conditions (LFCs) in systems is introduced. Questions to ask and aspects to consider when investigating and analysing failures in NPP SSCs are presented. The circumstances already prevailing in the NPP at the time of the failure, and that may have contributed to it, are examined holistically to show how a comprehensive way of understanding the complete event can be arrived at. The development of a systematic approach to guide and assist in the identification of key causative factors and to facilitate an objective SSC failure analysis is explained. Key words: failures in components, root cause identification, latent failure conditions, defence-in-depth, accident management, human factors, holistic approach.

6.1

Introduction

This chapter does not deal directly with how to physically examine failures in nuclear power plant (NPP) systems, structures and components (SSCs), but mainly describes a holistically based approach that may be used to address and analyse such failures. This type of methodology has the potential to provide in-depth information and give indicators and robust answers regarding how and why SSCs fail, and also to reveal any relevant contributing secondary circumstances to the failure, either within or external to the NPP concerned. Accordingly, the chain of events before, during and after the failure can be identified and the overall condition of the NPP may be obtained; this facilitates an appreciation of the sometimes complex human–machine interactions that have led to the SSC failure under consideration. Once root causes are known, corresponding corrective or mitigating actions can be implemented to avoid recurrence of the failure. (Detailed examples of actual failure analysis of materials used in NPP SSCs are given elsewhere in this book). New technology is constantly being added to existing processes (defined 131 © Woodhead Publishing Limited, 2010

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here as industrial plants (e.g. nuclear power, chemical manufacturing, oil refineries)) or other human activities (e.g. transport, construction) usually to improve, modernize or upgrade them in terms of safety, efficiency, reliability, simplification of process control and operation, to facilitate maintenance and to increase profitability. New technology is also added to take advantage of the latest scientific knowledge available concerning materials behaviour (e.g. ageing degradation (AD) aspects) and to combat obsolescence or conceptual ageing in industrial processes. However, this puts increased demands on both system and human reliability that may not always be evident, or sufficiently appreciated, at the time of the new technologies’ implementation [1]. Maximum limits of SSC operating conditions, but within prescribed safety margins, may, nevertheless, be more closely approached, and if SSC-AD effects have not been sufficiently monitored or mitigated, the relative probability of failure in a SSC could increase either insignificantly, tangibly or significantly. It is therefore essential to have in place effective standard plant operational practices (OPs), ageing surveillance programmes (ASPs) and ageing and plant life management programmes (AM and PLiM, respectively) to follow the condition of SSCs with respect to their ability to satisfy their design and safety requirements. The ASPs, AM and PLiM programmes must be checked and modified, if necessary, to address rates and types of AD, and to take into account changes in the NPP’s current operational conditions and configuration (i.e. with the new technology added). Thus, when any type of new technology is incorporated into existing plant, a comprehensive review and assessment of all relevant factors that may impact OPs, ASPs, AM and PLiM programmes or emergency procedures must be carried out to see whether the plant procedures have become outdated, insufficient or redundant to varying degrees. Advances in technology, if not sufficiently analysed for all their possible side-effects, can potentially result in a reduction in the level of human perceived risk, since automation, for example, may isolate personnel from reality. This is particularly so when absolute credence is given to plant parameter signals that may, in themselves, be erroneous or misleading in nature [2]. The NPP operating staff may thus become shielded from the real situation that has developed and therefore be unable to recognize and thus choose, and instigate, the right actions necessary. When failure happens in any industrial plant, transport system or building, for example, and damage to the environment, loss of life and large financial costs occur, the attention of the media, regulators, politicians and the public naturally focuses on who, or what, was responsible. With responsibility and causes established, the issues of compensation, litigation and even punishment in the form of fines or imprisonment will then follow [3]. It is therefore necessary to first define what constitutes failure and then to unravel the circumstances that were present and which contributed, no matter to what degree, to the event. ‘Failure’, within the general context

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of this chapter, may be regarded as a condition of inability of a NPP SSC to fulfil its safety, design/operational function sufficiently, as required by specification or regulatory requirements. However, identification of a SSC that is ‘out of specification’ (but before any event or accident occurs), can be taken as proof that the NPP’s overall OPs (e.g. testing, monitoring, repairs and replacement programmes) and associated programmes are effective. It should be noted that the condition and degree of SSC insufficiency may have developed gradually or may have remained below the detection limits of monitoring, inspection or testing procedures used at the time. The definition of failure may also be extended to also include growing human complacency with time, inability or inadequacy to detect or recognize a developing situation (training level inadequacies), a reduction in questioning attitude or the omission to take appropriate corrective actions once the event occurs. This may be generally classed as deterioration in plant safety culture. Experience shows that major accidents, irrespective in which area of human activity they occur, are usually the accumulation of many small deficiencies, errors or oversights that, individually, would not be capable of causing an accident. Physical failures of NPP SSCs and their consequences may have purely economic impact (e.g. forced plant outages, expensive component replacement [4] or even total loss-of-plant [5]) or safety consequences, or a combination of these to create financial ruin and long-lasting social and environmental issues. Particularly for NPPs, such failures and accidents will cause a loss of political and public acceptance of nuclear power.

6.2

Reducing failure probability and consequences thereof in nuclear power plant (NPP) systems, structures and components (SSCs)

6.2.1 Requirement for SSC quality Due to the very highest standards of manufacture, quality assurance, approval procedures/certification, materials choice, conservative design margins, operation, maintenance, inspection, repair, replacements, OPs, ASPs, AM and PLiM programmes, as well as regulatory requirements, spontaneous failures of major safety-relevant NPP SSCs are exceedingly rare occurrences, but when they happen, they have a large potential to destroy human life, inflict long-term damage on the environment and cause considerable financial losses [2, 4–6]. This ‘low probability of failure, but serious consequences if failure occurs’ scenario is also true for many other industrial plant or human activities, including, for example, chemical manufacturing, aviation, space exploration and mining [7–10]. However, despite using the best known design, manufacturing methods and OPs and principles, and adherence to regulations, NPP SSCs may still be subject

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to erroneous personnel actions or unusual operational stressors that could overcome a design principle or regulatory requirement when rare, unknown or very unusual combinations of circumstances, or demands on the plant, develop in the system. It can thus be stated that although the material quality per se of SSCs, including certification of mechanical, chemical and physical properties thereof, is one very important aspect, any inherent weaknesses, omissions or oversights in design, OPs, ASPs, AM and PLiM programmes, routine monitoring or inspection, or incorrect human actions, for example, may nevertheless contribute to a lessening of the overall reliability and safety of SSCs in service. Unexpected reactions between the materials and the environments they function in may lead to new AD mechanisms, or faster than expected rates of known AD, and this may bring the component into an operating regime whereby maximum limits of existing safety margins are approached and the relative probability of spontaneous failure could increase by some finite degree. Safety margins set the boundaries of acceptable risks at any construction, means of transport or a NPP, for example. That risk level determines what types of safety systems, including emergency procedures and accident management strategies, must be in place in case of accidents.

6.2.2 Defence-in-depth (DID) Nuclear power plant designs feature DID, whereby a series of independent barriers are placed between the radioactive nuclear fuel and the environment. It is clear that DID can only limit the extent and lessen the consequences of a SSC failure, but as a concept, DID itself cannot prevent any actual failure or functional loss of individual SSCs. The salient features of DID are primarily engineered physical barriers (e.g. reactor vessel and containment integrity, core-catcher, emergency core cooling systems, filtered containment venting, fuel cladding, redundant/back-up cooling systems) that, as far as possible, should not depend directly on human action to mitigate or contain the extent of radiotoxic releases or to slow down the rate of accident progression, and secondary measures to be implemented in the framework of the NPP’s accident and emergency management strategies (AEMS). The AEMS cover such items as the control of abnormal situations such as fires and plant emergency procedures, civil evacuation plans and the availability and distribution of prophylactic potassium iodide (KI) tablets to reduce the risk of thyroid gland cancer. (Note: taking KI tablets reduces the risk of thyroid cancer significantly, since the tablets cause the thyroid gland to saturate with respect to non-radioactive iodine, which means that no harmful radioactive iodine can be absorbed by the thyroid gland.)

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6.2.3 Example of the concept of probability The probability (p) scale ranges from 0 to 1; when p equals 0 an event (e.g. failure) will never happen; when p = 1, it certainly will. An underlying principle of DID may be understood as an application of the multiplication law of probability theory, i.e. the individual probabilities of failure in the various DID barriers are multiplied together. Depending on the condition/ fitness-for-service of the individual barriers (or procedures), which should be independent of each other to be able to fulfil their intended function (i.e. no common cause failure (CCF) path present), a usually very low value of probability of p (failure) will be present in the system. For example, when a system has a total of 5 DID barriers, each having a design probability of p = 0.5 of failing, then the probability of them all failing at the same time will be p = 0.03125. Typically, the individual DID barriers in NPPs all have values of p (failure)  0.5. When a system moves away from steady-state, normal operation conditions, seemingly unrelated factors can, however, contribute to some finite increase in the probability of a failure in one or more DID barriers. Nevertheless, the overall p (failure) value will still remain very low, providing the remaining (effective) DID barriers are in optimum condition and the DID concept present in the NPP remains robust.

6.2.4 Accident and emergency management strategies (AEMS) Any number of potential root causes for SSC failure and/or accident may be present within a NPP before an event occurs [11]. This is also true for any other type of industrial plant, transport or construction, but robust designs and DID, also including AEMS are in place as an integral part of the overall goals to ensure that the extent and consequences of any failures are limited as much as possible. Standard (routine) plant OPs, maintenance, monitoring, repairs and replacements of SSCs, plus ASPs, AM and PLiM all complement each other to varying degrees in order that potential inadequacies or failures in SSCs can usually be detected and remedied before they can occur. A vigilant, questioning attitude and good safety culture of the plant personnel will also play a vital role in the overall goal to maintain safety. Nevertheless, failures in NPP SSCs do occur, just as in any other industrial plants (e.g. chemical, mining), constructions (e.g. dams, buildings), equipment (e.g. cranes, hawsers, anchor bolts) or transport (e.g. aircraft, spacecraft and ships). However, failures of NPP SSCs do not necessarily mean that a serious accident will follow. Actions taken to prevent NPP SSC failures from developing into severe or serious accidents are addressed in AEMS, which are basically decision-making tools using logical-based sequential decisions

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to guide plant personnel in an optimum way to mitigate, limit and contain all consequences of plant equipment failures (i.e. within and outside plant), and to bring the NPP into a long-term, stable state. Rapid and effective use of AEMS will not only limit the consequences of the event, but may also limit the financial costs thereof.

6.3

Latent failure conditions (LFCs) and failure terminology

6.3.1 LFCs Latent failures or conditions can be regarded as system errors but may be better envisaged as hazards that have developed or have become unknowingly embedded within a system or process, rather than risks directly associated with human actions. The LFCs arise, for example, from unseen deficiencies residing within such areas as organization culture, management decisions or design of procedures or weaknesses in plant workforce training programmes. These insufficiencies resident in the system may, with the passage of time, translate into error-provoking conditions when certain conditions arise, or they can create dormant weaknesses in the plant’s defences, ready to contribute to the occurrence, or increase the severity of, an event, near-miss incident or accident when combined with active failures. An active (or functional) failure is a malfunction relating to a change in the operating mode of a component or its part. The operating environment (e.g. temperature, pressure, coolant chemistry) or stressors (e.g. vibration levels, radiation and mechanical loading) may be changed either intentionally or unknowingly, and can thus affect the functionality. Compared to the normal full-power operating condition of a NPP, the maintenance condition is, by its very nature, more complicated in character, since it is usually characterized by several, diverse activities being carried out by various individuals all at the same time. For example, NPPs undergo an annual (or 18-month) main revision, which includes refuelling, repairs, replacements and exceptional maintenance, where necessary. As a consequence of this sometimes intense and complex work-schedule and time-pressure to get the NPP back on-line as quickly as possible, human and work-related errors may be potentially less likely to be detected, and then they become unwittingly incorporated into the NPP’s system to contribute, at a later point in time, to a failure. These LFCs are thus hidden conditions that can be precursory to SSC failures. A SSC containing a latent failure can thus exist as a dormant, but unseen, threat to operation and safety of a NPP over an extended period of time. In a wider sense, even inadvertently built-in inadequacies in emergency procedures may only become evident when attempts are used to implement the plant’s AEMS in a real situation. The field

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of LFCs makes a significant contribution to understanding how systems fail, and LFCs are a common feature in sophisticated and extensively automated systems using high-level technology, such as that found in NPPs [12].

6.3.2 Failure terminology By the nature of them, weaknesses or insufficiencies in SSCs, or inadequacies in OPs, ASPs, AM and PLiM programmes and also AEMS, only become evident after the SSC failure event, and the accident, has taken place. This necessitates that a systematic failure analysis is performed, taking into account all available information, prior, during and after the event to obtain the root causes. The following provides some main definitions and terminology that may be used when dealing with NPP SSC failures and related events: ∑

∑

∑ ∑ ∑

∑

∑ ∑

An initiating event is a single event causing the NPP to deviate from its normal operational state. An initiating event for an accident can have on-site or off-site features such as a component failure, a natural phenomenon or a hazardous situation due to intentional or non-intentional human action (i.e. terror attack or aircraft accident, respectively). The diversity principle is where redundant systems or components are available to accomplish the same safety function in such a way that these systems or components have one different feature, such as an operating principle, a manufacturing method or physical parameters. An operator error is a single erroneous action committed while an operator attempts to perform a control action relating to a safety function. A passive failure is the loss of integrity of a component or structure, or the blockage of the flow path of a process, such as coolant supply. A hidden failure is an identified failure which does not activate an alarm and which is not detected in tests or inspections performed according to plans. This raises issues of personnel ability charged with inspection or testing, and the suitability/sensitivity of testing methods used. A random failure is a failure whose occurrence is statistically independent of the failure of other components of a similar type. Small statistical variations in material quality, manufacturing method and tolerances, operating conditions, maintenance and testing sensitivities may eventually cause a component to behave differently compared to other components of a similar type, but from a different batch of production. Quality assurance and component functional equivalence issues may be relevant factors to consider here. A safety system is any system performing safety functions. A common-cause failure (CCF) is the failure of several components or structures as a consequence of the same single initiating event or failure. © Woodhead Publishing Limited, 2010

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A single failure is a random failure plus its consequent effects, which are assumed to occur during either a normal operational condition or in addition to an initiating event and its consequent effects.

6.4

Holistic approach to analysing nuclear power plant systems, structures and components (npp ssc) failure events

The following outlines a holistic approach that can be used to analyse the total structure of failure events in NPP SSCs or, with appropriate modification, any other technological plant, construction or transport system. Specifically, knowing the overall sequence of events leading up to a NPP SSC failure plays an important role in facilitating robust assessment and identification of the root causes of failures. Seemingly small items, such as an ignored minimum amount of liquid weeping from a seal, a fragment of metal lying under a SSC, the necessity to change water or air filters more frequently than previously, gradual decreases in raw and service water intake flow rates or a radiation monitor that increasingly indicates an approach to maximum allowed limits, may all be precursors to more serious events. Hereby, the importance of the NPP’s workforce safety culture and a questioning attitude are highlighted.

6.4.1 Questions to ask A holistic approach to find the root cause of a NPP SSC failure would address the following items, through a series of questions to be raised, such as the following. (a) What was the NPP’s physical state prior to the failure? For example, the NPP may have been recently refurbished or re-fitted with new SSCs, a routine maintenance or exceptional repair may have just been carried out or a power uprating (PU) may have recently been implemented. Are all NPP SSC documents updated, noting and registering repairs, modifications or new operational practices? Was the plant indicating some unusual disturbances, such as increased pump vibration, coolant temperature instabilities or power excursions, before the failure occurred? Were the plant monitoring and control signals functioning correctly? Was SSC unreliability a rare, sporadic or frequent feature of the NPP’s operations to date? If frequent, are the same SSCs involved and is the supplier’s quality assurance sufficient? Have operational reliability problems increased since a change in SSC specification or supplier? Any change in NPP management or training practices? Do the NPP’s OPs, ASPs, AM and PLiM programmes need revision? Did the plant’s management decide to delay (or even cancel) the replacement of the SSC involved, to save money? How did the safety and reliablity assessments for

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the next period of operation of the now-failed SSC become inadequate? Were there any changes in the type, depth or scope of the SSC inspection or the time interval between the failed SSC inspections? All these aspects may be precursors and potential contributors to the overall final failure. (b) What was the NPP’s operational status when the failure under consideration occurred? For example, steady, upset, accident, shut-down for refuelling or maintenance, hot or cold standby, raising or sinking power or full allowable operational power. Were there any unusual circumstances prevailing that could have affected the workload or stress level of NPP personnel? (e.g. exceptional demand for (peak) power, flooding, severe climatic conditions, seismic activity or terror threat). What was the danger to the plant safety? The level of threat to safety is an important aspect of the failure event, e.g. near-miss, low, significant, serious, major. (c) What happened to the NPP after the SSC failure? For example, total loss of plant, radioactive contamination within or external to the plant (i.e. were DID barriers breached and, if so, in what manner), low, reduced or no power generated (also possible loss of emergency power if diesel generators affected – any CCF path?), significance to safety for personnel and environment. Did NPP shutdown and AEMS function smoothly, or were there difficulties in implementing them directly as a result of the failure? If applicable, to what are inadequacies in AEMS attributable? (d) What triggered the SSC failure event or what possible unusual stressors prevailed? The causes may lie in purely mechanical overload failure, physical (e.g. aircraft collision, fires), environmental (e.g. seismic, lightning strike), direct human actions (e.g. wrong valve pressure release setting), or any combinations of these. The final assessment will reflect the role, and to what extent, each factor has played. (e) Identification of the SSC directly involved (e.g. containment, reactor pressure vessel, pressurizer, pumps, piping, seals, control rods, emergency power generators/diesels, instrumentation and control). Was the operational history and reliability of the SSC involved satisfactory until the time of failure? Failure or accident progression mode: was the failure spontaneous in nature or did it progress at slow, medium or rapid speed? Were the methods and tools for SSC monitoring used correctly and at the state-of-theart, science and technology? (Qualified personnel are charged with safetyrelevant inspections, but work and time pressure in an ionizing radiation zone despite ‘as low as reasonably acceptable’ (ALARA) radiological protection procedures and non-optimized planning may affect the quality of results obtained.) Difficulties of access to certain SSCs, due to NPP design features, may compound inspection problems, despite the basic ability of an individual to perform the work according to requirements. (f) Were other SSCs involved (either as joint contributors or themselves affected in a secondary nature by the failure under consideration)? The

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identification of associated SSCs or possible CCF paths will be a vital task to perform, as appropriate, to upgrade the safety in future. For example, valve failures are recurrent problems having potential for causing CCFs. Was the availability of redundant safety and other systems (e.g. DID) nevertheless assured? Has the reliability of other SSCs possibly been affected as a result of the original failure? (i.e. have SSCs adjacent to the failed item been subjected to unusual temperatures, chemical contamination or radiation?). Have actions been taken to verify this? (g) What was the most probable cause of SSC failure? This question cannot usually be answered directly, since, depending on the type of failed SSC, it requires chemical, metallographic, mechanical and physical testing and control of all relevant supporting documentation. (h) How did the plant personnel react? Were the corrective actions taken the best possible to manage the situation? Were existing, validated procedures adhered to, or was it necessary to modify them ad hoc as the failure/accident situation developed? Was it feasible, in the practical situation, to implement the procedures? Was there a proactive approach to deal with the problem? Was there any evidence of lack of appreciation of what was actually happening? Was the personnel situation-awareness consistent with the degree of severity of the problem? Were the emergency procedures (AEMS) adequate and capable to address the situation as it developed and could limit the consequences effectively? (i) Are there any other aspects or features that may conceivably be involved with the sequence of circumstances that resulted in the SSC failure? (see under (d) above).

6.4.2 Acquiring further information and other tasks to do After an initial holistic assessment of the events, which may have already provided deeper insights and information concerning the circumstances leading up to and surrounding the failure event, more specific tasks will have to be performed to find the root cause. A wide range of tools exist to facilitate this. For example, classical metallographic examination will provide information on the failed material’s microstructure, the crack morphology and orientation, and the AD or failure mechanism itself (e.g. transgranular or intergranular stress corrosion cracking, fatigue). Care is needed when removing specimens from failed components, since vital information may be lost if, for example, surface cracks are polished over or the composition of chemical deposits is contaminated. Metallurgical or chemical analysis may be supplemented by mechanical tests (e.g. determination of tensile, hardness and fracture mechanical properties). Results of such tests may then be compared to the SSC supplier’s specification certificates for the SSC in question. Other examinations may include verification of the material’s chemical composition

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(also to be compared with delivery documents) and obtaining an X-ray of a deeply seated flaw, for example. Observing where a crack was initiated will often show the presence of stress raisers in the form of sharp corners (design weakness) or a non-metallic inclusion (material quality). Furthermore, the features of the crack surface will indicate whether corrosion, fatigue, sudden overloading or combinations of these were present. The operating environment (e.g. coolant) and loading history of the item (e.g. number of pressurizations), if applicable, will furnish further supporting information. Analysis of any possible human factor contribution should also be undertaken. Global checks should be made to see whether similarly designed NPPs have experienced the same problems with the SSC under investigation. (Note: Plant-owners associations such as Westinghouse Owners Group (WOG) and Boiling Water Owners Group (BWOG) have a system for rapid notification of problems in place, and regulatory notices also inform plant operators and owners on specific or generic problems.) Similarities or differences in OPs between such NPPs should be identified and their significance examined with respect to identifying any possible contributory aspects.

6.5

Discussion

Failures may be divided conveniently into two areas, namely, ‘materials’ and ‘human’ failures. Materials failures are defined here as those materialrelated events that have occurred in the operation of SSCs (e.g. ruptured pipe, corrosive attack resulting in leakage, loss of electric control signals and isolation), and human failures are defined here as those that have occurred in SSCs as a direct result of human actions (e.g. failure to detect weaknesses in design, errors of commission, inappropriate maintenance or accident management actions taken and failure to report an anomaly, i.e. errors of omission). Discovery of a SSC in a degraded condition, before it has chance to fail, can be regarded as proof of effectiveness of OPs, ASPs, AM and PLiM methodologies. Despite conservative engineering design approaches, good OPs, monitoring, repair and replacements, ASPs and AM and PLiM programmes, the SSCs in NPPs can, nevertheless, undergo AD to the extent that they can no longer fulfil their safety or operational function and fail, particularly if weaknesses are present in the detection, monitoring or inspection schedules, for example. Failure analysis (FA) is performed to understand what happened and to prevent the failure from happening again. The FA must therefore fulfil several goals, all of which will eventually lead to the identification of the root cause of the SSC failure event. When the root cause is unequivocally identified, corresponding strategies and actions can be implemented. Design weaknesses, material unsuitability, OPs and human actions that have either directly contributed to or exacerbated the failure will have to be eradicated.

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In some cases, radical changes have to be undertaken, not only on the SSCs, but also the way they will be inspected, maintained and monitored in the future. Managing operational safety is a major component in overall NPP operations, and avoiding failure in any aspect of operation will contribute to this goal. One element of operational safety is therefore the feedback of operational experiences and implementing lessons learned in the most effective way possible. Irrespective of how it happens, where it takes place and its extent (e.g. NPP, chemical plant, aircraft), spontaneous failure of SSCs in operation, with corresponding loss of life and environmental damage, is direct proof that weaknesses must have existed a priori in, for example, the design, OPs, inspections, monitoring, routine maintenance, human actions, including management strategies, preventative maintenance and, where applicable, AEMS used at the time. Specifically, NPPs are complex social-technical systems, relying on the best available design and featuring SSCs made to the highest standards of materials and a competent, safety-motivated workforce to operate them according to prescriptions. It is also necessary to ensure that management strategies and operational procedures used in NPPs are always at the current state-of-the-art, science and technology. This necessitates continuous implementation of changes as they become necessary. This may entail modifications to SSCs or possible adjustment to inspection schedules and the extent thereof, for example, to satisfy regulatory requirements. Several human and physical-mechanical aspects must therefore unite to achieve the common goal of safe, reliable operation and freedom from spontaneous SSC failure. However, when an adverse conjunction of several diverse causal sequences unite, each one a necessary contributor, but none individually sufficient to breach the system’s defences, a SSC failure and attendant accident may occur. It is evident that a challenge lies in being able to predict or identify such highly unlikely, but potentially serious, contributors to failures and accidents. But herein lies the problem: if the existence of such contributors was known, they would have already been addressed and eliminated in the NPP’s overall OPs, inspection and maintenance activities, or their potential to cause problems correspondingly documented in the ASPs, AM and PLiM programmes before the failure had had a chance to occur. (Note that this observation is also generally valid for aviation, space exploration, shipping or chemical accidents that have occurred.)

6.6

Conclusions

1. A holistic approach is a methodology to assist other tools in analysing why and how SSC failures occur in NPPs. 2. The analysis of failures presents an opportunity to learn from mistakes

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3. 4.

5. 6.

7.

8.

9. 10.

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and to implement the appropriate measures to improve safety and reliability in future NPP operation. Addition of any type of new technology to an existing NPP requires a thorough analysis of its impact and effects on the total plant operations and personnel training and actions. Revision of OPs, ASPs, AM and PLiM programmes and adjustments, repairs or replacements to SSCs must be all documented to enable tracking of future SSC safety performance and reliability, and also to provide data for analysing any future failures. Awareness of maintenance-related LFCs will be increased by a questioning attitude and good safety culture in the NPP’s workforce; any CCF paths that can be identified must be rigorously eliminated. The NPP’s personnel should continue to be trained with regard to appreciating the potential significance of reporting (seemingly irrelevant) observations or small, but unusual, plant operating parameter and signal deviations. Spontaneous failures of SSCs, which may lead to accidents, usually cost much more than regular monitoring, inspection, repair or replacement of them. Appropriate inspection, testing and monitoring procedures must be applied, and updated as the technologies evolve (e.g. improving crack detection rates and obtaining better resolution of crack dimensions with state-of-the-art ultrasonic test instruments). Accident and emergency management strategies (AEMS) are in place to control and limit the consequences if NPP SSCs should fail. Regular training of personnel and control of procedures should be an integral part of plant operational management and operational practices. Emergency exercises, using realistic scenarios, are ideal tools to detect weaknesses in NPP AEMS. The necessary goal to operate NPPs safely is mostly facilitated by robust and tolerant design, DID, standard OPs, ASPs, AM and PLiM programmes. The overall safety culture in a NPP (e.g. questioning attitude, willingness to report anomalies) is an essential facet of operations and potentially can have as much importance regarding reducing the probability of SSC failure as do inspections, monitoring and testing thereof.

Sources of further information

OECD NEA (1995), ‘Chernobyl Ten Years On: Radiological and health impact – 2002 update’. OECD NEA, Paris. IAEA (1996), ‘Ten Years after Chernobyl: What do we really know?’, IAEA/ WHO/EC conference, IAEA, Vienna. IAEA (2006), ‘Environmental Consequences of the Chernobyl Accident and their Remediation: Twenty Years of Experience’, IAEA, Vienna. © Woodhead Publishing Limited, 2010

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Anon. (2002), ‘Potassium iodide for thyroid protection in a nuclear accident or attack’, The Medical Letter, Vol. 44 (W1143C), 11 November, pp. 97–98 (source: http://www.urgencenucleaire.qc.ca/documentation/ potassiumiodide.pdf). Kletz, T. (1998), What Went Wrong: Case Histories in Process Plant Disasters, 4th edn. Gulf Publishing Company, Houston, TX. Lapierre, D. and Moro, J. (2002), Five Past Midnight in Bhopal, Warner Books, New York. IAEA (1991), ‘Safety culture’ Safety Series No. 75-INSAG-4. IAEA, Vienna. IAEA (1991), ‘Reviewing operational experience feedback’, IAEA-TECDOC596. IAEA, Vienna. Becker, W.T. and Shipley, R.J. (2002), ASM Handbook: Volume 11: Failure Analysis and Prevention, 10th edn. ASM International, Materials Park, OH. Colangelo, V. and Heiser, F. (1987), Analysis of Metallurgical Failures. John Wiley & Sons, New York. Moroney, M.J. (1969), Facts from Figures, 3rd edn. Pelican Books, London. Tipping, Ph. (1996), ‘How materials ageing and human factors can lessen safety margins and cause failures’, in Penny, R.K. (ed.), Risk, Economy, and Safety, Failure Minimisation and Analysis, Balkema, Rotterdam, pp. 23–32. Petroski, H. (1999), ‘The Britannia Tubular Bridge: A paradigm of failuredriven design’, in Addis, W. (ed.), Structural and Civil Engineering Design, Ashgate Publishing Ltd, Aldershot, pp. 313–324.

6.8

References

1. Holnagel, E., Human Reliability Analysis Context and Control, Academic Press, London, 1993. 2. Report on the Three Mile Island Nuclear Power Plant Accident: US Nuclear Regulatory Commission NRC Annual Report – 1979, NUREG-0690. 3. Varley, J., Who was to blame for Chernobyl? INSAG’s Second Thoughts, Nuclear Engineering International, Vol. 5, pp. 51–52. 4. US Nuclear Safety Commission, Bulletin 2002-01 ‘Reactor Pressure Vessel Head Degradation and Coolant Pressure Boundary Integrity’, January, 2002. 5. IAEA, Environmental Consequences of the Chernobyl Accident and their Remediation: Twenty Years of Experience. IAEA, Vienna, 2009. 6. Wani, T. and Bessho, Y., Summary of the interim report on the secondary system pipe rupture at Unit 3, Mihama nuclear power plant, in Proceedings Series Material Degradation and Related Managerial Issues at Nuclear Power Plants, Proceedings of a technical meeting organized by the International Atomic Energy Agency, IAEA, Vienna, 15–18 February, 2005. STI/PUB/1260. Printed September 2006. 7. Surviving Bhopal 2002: Toxic Present Toxic Future, report published January 2002 by the Fact-Finding Mission on Bhopal (FFMB). © Woodhead Publishing Limited, 2010

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8. Aircraft accident report 92-11: El Al Flight 1862, Boeing 747-747-258F 4X-AXG Bijlmermeer, Amsterdam. 9. Space shuttle ‘Challenger’ STS 51-L Accident, NASA, updated on 30 March 2009, source: http://history.nasa.gov/sts51l.html. 10. All Mining Fatalities By State, US Department of Labor, Mine Safety and Health Administration, 15 January 2007. 11. Guidelines for the Review of Accident Management Programmes in Nuclear Power Plants, IAEA Services Series No. 9, IAEA-SVS-09, IAEA, Vienna, May 2003. 12. Holnagel, E., Latent Failure Conditions and Safety Barrier Integrity, Joint OECD/ NEA Symposium on Human Factors and Organization in NPP Maintenance Outages: Impact on Safety, 1–8, Stockholm, Sweden, 19–22 June 1995.

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7

Impact of operational loads and creep, fatigue and corrosion interactions on nuclear power plant systems, structures and components (SSC)

M. B a k i r o v, Center of Material Science and Lifetime Management Ltd, Russia

Abstract: The operational lifetime of a nuclear power plant (NPP) is determined mostly by the actual condition of the materials from which the main equipment is made. This chapter addresses the impact of main operational effects (i.e. temperature, stresses, medium present) in causing ageing degradation in the NPP equipment. It also addresses the main mechanisms involved causing damage (e.g. corrosion, fatigue and creep) related to current and long-term operational effects and proposes experimentcalculated approaches to evaluate their impact on the integrity of equipment. Results of work on the analysis of ageing in systems, structures and components (SSCs) in Russian NPPs are presented as specific examples. The most efficient methods of in-situ (specimen-free) non-destructive monitoring and testing to detect ageing degradation are recommended. Key words: mechanisms of metal damage, creep, fatigue, corrosion, thermomechanical loading, ageing, equipment lifetime.

7.1

Introduction

The operational lifetime of a nuclear power plant (NPP) depends significantly on the physical condition of the SSCs and thus their ability to fulfil their design function. The lifetime (fitness-for-service) of nuclear power equipment must be demonstrated in accordance with the relevant standards of NPP equipment and pipelines strength analysis. In Russia, the standard is PNAE G-7-002-86 Strength Calculation Norms, and it is based on characteristics of fracture resistance under static and cyclic loading, taking into account the steel grade, process features of manufacturing and operating conditions. Thus the basic mechanical properties are assumed as guaranteed normative values or values as stated in the delivery certificate of the metal alloy used for the equipment under investigation. This approach to facilitate estimation of the equipment condition can only be absolutely acceptable if the mechanical properties of materials do not change in the course of time/operation. However, NPP operation causes thermo-mechanical loads, the coolant contacts the SSCs 146 © Woodhead Publishing Limited, 2010

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and interacts with them and ionizing radiation is present, all depending on the location of the SSC under consideration. Over time, these operational stressors cause changes in the physical and mechanical properties in the SSCs, namely ageing degradation/effects. Such ageing may impact the structural integrity of the SSCs, and particularly so in those areas where the stressors are severe. Ageing degradation can thus affect designed safety margins and lead to failures, if not detected and mitigated in a timely manner. Globally, there is more than 40 years of experience in commercial NPP operation. This has resulted in a vast accumulation of information on how SSCs react in their respective operating environments and especially how ageing degradation may affect them. Such information provides a basis for establishing the main factors that govern the true lifetime of NPPs (see Fig. 7.1). If very large safety margins were already implemented at the design stage, the equipment should not fail during its operational lifetime, assuming no severe manufactured defects are present. This would be so, since, theoretically, quality control (e.g. crack detection, X-ray) would detect such defects before they could be used in the NPP. Furthermore, if large safety margins were present, they would also counter ageing degradation over the projected lifetime of the NPP. Nevertheless, despite quality assurance and control, there is no equipment that can be guaranteed to be absolutely defect-free. For visual presentation of the causes behind the reduction of equipment operational reliability due to ageing of materials, see Fig. 7.2. If the ordinate axis is the size of a minimum actual (target) defect referred to as atarget crack size in the figure (a defect resulting from manufacturing or that has occurred in the course of operation) and the size of a maximum defect referred to as acritical crack size in the figure (a defect postulated in the design), and time of equipment operation is on the abscissa axis, the essence of equipment Ageing factors

Ageing mechanisms

Consequences

Environment

Embrittlement

Structural damage

Temperature Stress-strain Material

Creep Corrosion Fatigue

Degradation of mechanical properties Cracking Deformation Material loss

7.1 Main ageing factors affecting the lifetime of NPPs.

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aacceptable crack size

Safety margin

Contribution of additional damaging impacts as a result of ultra long-term operation (high- and hyper-cycle fatigue, hydrogenation, medium)

Rupture

atarget crack size

Subcritical crack growth Manufacturing

Design service life Time of the NPP lifecycle

7.2 Diagram for analysis of the lifetime of NPP equipment.

Lifetime extension

Physical life

Crack size, a

Actual growth of the target defect size due to operational influence – fatigue, stress-strain, corrosion, creep (in-service inspection and monitoring)

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Reduction of the critical defect size due to structural damage and mechanical properties degradation (operational history, codes and standards)

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lifetime management activities will consist in following the way in which these defects evolve over operational time. During operation, on the one hand, the size of the actual target defect grows (the lower ascending solid curve), while, on the other hand, the size of the maximum defect decreases (the upper descending solid curve) due to ageing and, as a consequence, the mechanical properties of the equipment’s material, including fracture resistance, decrease as well. The distance between these curves shows the boundaries of equipment strength margin (safety margin in the figure) or, alternatively, the residual strength lifetime. The shape of the curves is determined by experiment-calculated monitoring on the basis of the pertinent industry regulatory documents. Regulatory documents (e.g. ASME Boiler and Pressure Vessel Code, 1995) contain special tables of defect sizes acceptable during operation, and specify methods for their assessment by calculation. The most universal characteristic of material fracture resistance covering calculation in the quasi-elastic and elasto-plastic fields is the J – integral and its critical value JIC. However, its determination is time consuming and for many tough steels used for NPP reactor pressure vessels, for example, the JC value is unknown considering the actual operational level of degradation. For this reason, many regulatory documents use approximate and conservative estimates and require only ordinary mechanical properties (e.g. tensile, Charpy impact test). To assess the kinetics of defect development in operation, it is necessary to perform periodic non-destructive inspection to determine the actual dimensions of defects. Afterwards, it is necessary to acquire data on actual loads experienced by various SSCs. Finally, it is required to perform an analysis of the experimental data regarding the actual defect growth rates related to the various operating conditions. It should be noted that NPP SSCs in the long-term operation (LTO) mode, defined here as NPP operation after the original design life, may experience significant ageing levels due to vibration, high-cycle and hyper-cycle fatigue, chemical reaction with the medium and radiation. The current situation (2009) is that the possible contribution of LTO to increasing ageing levels in SSCs, even if it is deemed to be relatively insignificant, is not yet taken into consideration in regulatory documents (see dashed lines in Fig. 7.2). Accordingly, only the quantitative measures of various parameters of material operational damage mechanisms provide for a confident prediction of NPP safe operation limits. This prediction is impossible without periodical evaluation of the rate of properties degradation at the most loaded areas of the SSC under consideration.

7.2

Nuclear power plant (NPP) equipment materials

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specific values of chemical composition, type of structure and corresponding physical, chemical and mechanical properties. The choice of materials is based on the analysis of chemical and mechanical resistance to ageing, under current and during LTO conditions. For SSCs, including piping, of Russian and other NPPs, various materials are used: heat resistant, high-temperature resistant, corrosion resistant (stainless) and other materials find application. For the fabrication of reactor plant piping, carbon and low alloy steels of grades 10, 20, 20K, VSt3Sp5, 16GS, 17GS, and their foreign analogs AISI C1010, AISI C1020, DIN C10, C22, St 52-3, etc., are used. Their properties depend on the carbon content and alloy additives and heat treatments performed (Anon., 1990; Beskorovainy et al., 1995). These steels are almost free from hot and cold cracking susceptibility caused by welding. Owing to this, piping made of these steels can be welded without any complication by any welding method at both low and high rates of energy input and corresponding rates of heating and cooling down within a broad range of values (Kearns, 1978; Livshits, 1979). Analysis of these steels’ operational performance has shown that the dominating degradation mechanism was erosion-corrosion and this depended on the chemical composition and parameters of the medium, chemical composition and structural properties of the steel and component geometry and features of the circuit in creating turbulence in the coolant. In view of the strict requirements placed on nuclear reactor structural materials, of which the most important are requirements on ultimate tensile and yield stresses, low sensitivity to stress raisers, high hardenability and low sensitivity to temper brittleness, the use of plain carbon (non-alloyed) steels in NPP SSCs is quite limited. In contrast, the reactor pressure vessel (RPV) is typically a low-alloy ferritic steel structure, cladded on the inside with stainless steel. RPVs are heat treated to create a ferritic-bainitic microstructure and maintain high toughness, even if some level of neutron embrittlement has occurred over the service life. For large volume items, such as the RPV and reactor vessel head, heat exchangers and other SSCs, ferritic-pearlitic and various alloyed steels having higher hardenability and better mechanical properties are widely used. Thus, the required mechanical properties can be achieved throughout almost any section thickness. Furthermore, ferritic-pearlitic structured steels are not prone to intergranular corrosion or corrosion cracking in the medium used in pressurized and boiling water reactors. From the point of view of high-temperature strength (having creep-rupture strength as one of its criteria) the heat resistant steels are ranked at the bottom of reactor structural materials. However, taking into account the combination of physical and mechanical properties at low and moderate temperatures, manufacturability and price, their use in NPP SSCs becomes economically feasible.

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Corrosion resistant austenitic steels (e.g. for Russian-designed WWER and RBMK reactors – 08Ch18N10T and for European/US designed pressurized water reactors (PWR) and boiling water reactors (BWR) – AISI 304, 321) are the main structural materials for manufacturing of NPP SSCs operating in contact with primary circuit water. The reason for their extensive use in water-cooled reactors is their high resistance to general surface corrosion. In Russian NPPs, chromium-nickel steels, on the basis of 18% chromium-8% nickel (wt%) classical composition, weld-decay stabilized by titanium, and having generally an austenitic structure are most widely used. Titanium is used as an alloying element in order to increase the hightemperature strength of the steels through the formation of fine-scale strengthening phases of various types (carbides, carbonitrides and intermetallic phases). Additionally, titanium addition is used as a stabilizer in order to reduce the austenitic steel’s tendency to suffer intergranular corrosion by means of carbon binding to form special titanium-based carbides in preference to chromium carbides, thus leaving the corrosion resisting chromium available in the matrix (Beskorovainy et al., 1995). These corrosion resistant austenitic ‘stainless’ steels combine a medium strength (ultimate tensile stress Rm = 550 MPa at room temperature) with relatively high plastic deformation (e.g. relative elongation A5 = 50%). Carbon has no significant effect on the mechanical properties of titanium stabilized 18-8 austenitic steel if present below 0.15 wt% and in the absence of the ferritic a-phase. Furthermore, in order to increase the resistance of non-stabilized austenitic stainless steel to intergranular corrosion, the carbon content has to be low (e.g. < 0.03 wt%). The combined effects of high residual and operational stresses, high aggressiveness of stagnant coolant water in cracks (i.e. ‘chemical hide-out’), heat treatments that have increased the hardness and reduced plasticity, or thermal ageing during operation can all be implicated in the causes of cracking of NPP components made of these steels and having the nature of intergranular corrosion cracking.

7.3

Medium and corrosion

Corrosion degradation of NPP equipment represents the main type of damage sustained in the course of current and long-term operation and, in such a way, it directly affects the remaining lifetime of the NPP. (The large, passive, irreplaceable components dictate the true lifetime of a NPP, but SSCs that are replaceable, but expensive, may impact the profitability of the NPP.) Corrosion types and rates of attack of a given alloy depend on many factors, including the characteristics of the medium the alloy is in (Saji, 2009). The term ‘medium’ here means the coolant water located in ionizing radiation fields of various intensities in a wide range of temperatures, pressures and flow rates/turbulence. © Woodhead Publishing Limited, 2010

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The significance of water-chemical composition parameters of the medium is stipulated for a variety of reasons. First of all, water and water compositions used in the process and cooling circuits of NPPs are universal solvents. When flowing in the primary circuit, the process water and water solutions change their temperature (e.g. in the primary circuit of WWER, and PWR), aggregate state. The term ‘aggregate state’ implies that water passing through the secondary circuit undergoes changes from water-steam mixture in the turbine to water after the condenser. In the primary circuit of a PWR, water does not change its aggregate state – it does not boil. RBMKs and BWRs actually have only one circuit – water is heated in the reactor, boils and steam separated from the water-steam mixture goes to a turbine. The aggregate state is important for consideration if erosive corrosion of carbon steels (e.g. in the pipelines of regenerative heating) is under consideration, where steam contents in the water-steam mixture play a significant role in the erosion wear processes, e.g. in the secondary circuit of WWER and PWR and in the main circulation circuit of RBMK and BWR. In this case, the intensity of the processes of interaction changes between the coolant (or working medium) and the equipment’s material. Intensity of interaction also changes in the case where the amounts of impurities and their composition changes, for instance with chemical dosing and filter efficiency change. The major parameters affecting the NPP SSC corrosion-based lifetime are conductance (i.e. purity) of coolant water, pH and the concentration of chemically aggressive corrosion products in the coolant. The interconnection of these parameters is now explained, using as an example the RBMK-1000 NPP. Due to physically stipulated links to periodically controlled values of all water chemistry parameters, the integrated sum of all iron corrosion products present in RBMK coolant circuit for a year can be used as a parameter for coolant contamination magnitude (Kritsky et al., 2000). In Fig. 7.3 the correlation between the number (ni) of defective RBMK-1000 SSC pieces (modules of steam separators, plugs of instrumentation channels, fuel assemblies, collective doses accumulated during tasks, activity of I-131 in the coolant) and the amount of iron corrosion products (xi) being transferred through the feed water and which have appeared in a reactor is presented. In Fig. 7.3 conventionally three areas can be identified and interpreted as follows: I

the area of high purity of the coolant; owing to that, an effect of coolant characteristics on the NPP equipment lifetime is insignificant, and the equipment lifetime in operation is determined mainly by the original material properties of the equipment and the level of loading fixed by the design; II the area of correlation between the changing of the equipment lifetime characteristics and the coolant characteristics;

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153

^ n i

0.6 0.4

III

0.2 0

II

–0.2 –0.4 –0.6

I

–0.8 –0.8 –0.6 –0.4 –0.2

0

0.2

0.4

0.6

Xi 0.8

Plugged modules of steam separators Activity of I-131 in RBMK coolant Refuels owing to steaming of process tube plugs Refuels on replacement of untight fuel assemblies Collective dose at capital repair Collective dose at usual operation

7.3 Correlation between normalized reliability values of RBMK equipment, collective dozes and normalized values of iron corrosion products transferred through feedwater.

III the area of inadmissible reduction of the NPP equipment lifetime; in this case the NPP personnel implement active measures for replacement of the affected/defective equipment. The structural material or equipment construction are being changed. It is necessary to note that areas I and III are only outlined and are not clearly defined. The reason for this is that it is difficult to maintain the coolant high purity, and the encroachment towards area III means a potential violation of the prescribed regulations. This, in effect, can lead to significant personnel radiation dose aspects and also to economic losses, i.e. increase of repair scope, augmentation of collective doses, etc. If an individual RBMK NPP is considered, it should be noted that the equipment lifetime depends on the reliability of the steam turbine piping (condensers), as well as on the salt content in the cooling water for the condenser tubing. The type and rate of corrosion is also influenced by anodic reactions and degree of oxidation, cathodic reactions and the degree of reduction, efficiency of corrosion inhibition, polarization, presence and types of oxides (adherent or loose), type of galvanic cells, concentration of corrosive substances, structure and condition of metal (e.g. internal tensile stresses). Especially hazardous is the interaction of corrosion with other mechanical stressors such as fatigue cycling or wear which can result in spontaneous failure of equipment. In nuclear power engineering, the following types of corrosion

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are present, to varying degrees and severity (Collins, 1984; Uhlig and Revie, 1985): chemical, galvanic/electrical, crevice, pitting, intergranular, erosion-corrosion, cavitation corrosion, stress-corrosion cracking, hydrogen embrittlement (caused by hydrogen atoms, generated through corrosion processes, diffusing into the metal lattice). A brief review of these corrosion types is presented below. Chemical (equal-rate) corrosion is a widespread type of corrosion when the surface of a metal in contact with a corrosive environment oxidizes evenly. The corrosion rate is assessed by simple laboratory or full-scale in-situ experiments (on test samples) by measuring the loss of weight of coupons, for example. With regard to resistance to equal-rate surface corrosion, the steels and alloys of primary circuit equipment shall belong to materials of resistance groups 1–3, i.e. to have a corrosion rate not more than 0.01 mm/ year (GOST 9.908-85). Regarding NPP equipment, this type of corrosion does not represent any great hazard. There are many methods used to provide protection against this type of corrosion. Electrochemical corrosion takes place when two dissimilar metals are connected through an electrolyte (external medium). Two similar metals, but having different levels of internal stress, may also be affected. The physics of the phenomenon is similar to that in a galvanic cell. Quantitatively, this process is described by Faraday’s law, namely: the weight of reacting metal is equal to K · I · t, where I is current, t is time and K is a constant called the electrochemical equivalent. More precisely, the metal (electrode) acting as the anode suffers a loss in weight (corrodes) and the metal acting as the cathode gains metal or acts in a passive way. Usually accelerated electrochemical corrosion proceeds most intensively at the junction of the two metals. Based on results of tests of many metals, the electrochemical or galvanic series are plotted, usually against a standard reference electrode (e.g. hydrogen standard electrode) (Anon., 1983). There are metals having no galvanic interaction at all since it is possible to select two metals having almost the same (or close) potentials to minimize or exclude the occurrence of corrosive interaction processes between them. Electrochemical corrosion can be countered by the following means: ∑ selection of pairs of similar materials; ∑ electrical insulation of one of the materials (e.g. paint or oxide layer); ∑ provision of a small ratio of cathode surface area and anode surface area (if a large cathode area and small anode area are in electrical contact, the anodic current density is high, which induces high local corrosion rates at the anode; alternatively, a large anode and small cathode area means that the corrosion process is generally under cathodic control (balance of charge must be maintained) and the overall corrosion of the anodic area is low);

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∑ ∑ ∑ ∑ ∑ ∑ ∑

155

application of coatings, including creation of high-quality adherent oxide films; application of inhibitors for reduction of medium aggressiveness; application of cathodic protection; sacrificial electrodes of, e.g., zinc to supply electrons to the cathode and ensure a reduction process to inhibit oxidation; impressed external voltage; maintenance of pH level of coolant at temperature of importance; cleaning of the metal surface to remove deposits that have excessive concentration of salts, chlorides and sulfates. The most efficient measure for the reduction of corrosion damage in important equipment is to use corrosion resistant austenitic steels in the NPP primary circuit.

The steels most widely used in Russian NPPs are austenitic 18% chromium-8% nickel steels stabilized by titanium. The steels most widely used in PWR and BWR reactor plants are austenitic steels of series 300 (e.g. non-stabilized steels 304, 310, 316 and stabilized steels 321, 347, 348). They typically contain 16–26% Cr, 8–22% Ni and 2–3% Mo as the main alloying elements. Crevice corrosion is a localized process of accelerated corrosion in crevices, cracks and under layers of deposits or sediments. Crevice corrosion is usually characterized by a long-term incubation period. It can be reduced by means of periodic removal of contaminating sediments on the surfaces, which are also present in crevices. Sediment removal is usually performed using alkaline solutions or by means of mechanical treatment, elimination of stagnant spaces by mixing and aeration of the medium, implementation of cathodic protection with a polarization up to the value of corrosion potential of the active metal in the crevice. Pitting corrosion is observed in cases when the corrosion rate at some local areas is higher than that on the whole surface. The mechanism is the same as crevice corrosion. Corrosion pits are formed for various reasons, starting from defects of the metal surface, and finishing with the disruption of the protective passive oxide films resulting from temperature and water chemistry conditions. Chemically active (corrodants) sediments can accumulate in corrosion pits, thus provoking their growth. Corrosion pits usually grow in the direction of the force of gravity, since for their active growth they require the presence of relatively concentrated acid or alkaline solution. Pitting corrosion can be reduced by means of maintenance action during NPP outages, elimination of stagnant areas with sediments, reduction of oxygen content in chloride-containing medium and addition of extraneous inhibiting anions (OH- or NO–3). Pitting damage often occurs in NPP equipment manufactured of steels of different grades. The pitting depth is characterized by the pitting factor. It represents a relation of the maximum observed pitting depth to the average

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depth of general surface corrosion. A pitting factor equal to one corresponds to equal-rate corrosion. Intergranular corrosion is a local corrosive destruction along areas on, or adjacent to, metal grain boundaries resulting in loss of strength and plasticity (Chigal, 1969). When a metal crystal structure is deformed, the accumulated deformation energy is higher at the boundaries of the metal grains than in the grains themselves. These high-energy areas of grain boundaries have higher chemical activity than the grains. The small area of inter-grain/grain boundary material acting as an anode, is in contact with a large surface of grains representing a cathode. The corrosion, driven by a large cathode area to small anode area ratio (high anodic current density) proceeds relatively quickly, depending on the medium and other parameters acting. The corrosion attacks and penetrates the metal over time, leading sometimes to catastrophic results. In particular, 18-8 type austenitic steels, of which RBMK piping are made, become subject to intergranular corrosion after sensitization by heating to 500–800 °C. This results in chromium depletion of the metal matrix near to the grain boundaries as a consequence of chromium carbide precipitation in, and adjacent to, the boundaries. Operation of local galvanic cells provokes corrosion at areas with chromium depletion and grains crumble from the boundaries of the matrix. This effect is called weld-decay. In order to minimize the susceptibility of austenitic stainless steels to intergranular corrosion the content of carbon has to be controlled (it should not exceed 0.03%), and stabilizers must be added to prevent chromium depletion near to grain boundaries (for example, addition of molybdenum for 18-8 steel of grade 316). For SSC equipment in operation, the procedure of in-situ recovery heat treatment (i.e. austenitizing of areas prone to intergranular corrosion) has been proven to be very efficient. As a rule, these areas are welded joints of pipelines. Such heat treatment results in the resolution of chromium carbides in the affected areas, thus returning chromium back to the austentic matrix as chromium carbide is dissolved. The technology of in-situ recovery heat treatment was developed in Russia and successfully applied on RBMK austenitic pipelines for the life prolongation of piping that were prone to intergranular stress corrosion cracking. Erosion corrosion occurs due to the action of coolant flow with speeds exceeding the laminar flow parameters at the surface of the metal under consideration (Preece, 1979). At such speeds, the passive film (or corrosion products film) is disrupted and transported away from the surface, thus causing an increase of corrosion rate in these areas. Erosion corrosion causes section-thinning over large areas, and is called erosion-corrosion wear. This type of damage occurs everywhere in operation of all types of NPP equipment manufactured of carbon steels. Initially, it is typically observed on internal surfaces of valves and pumps and, certainly, in pipelines of secondary circuits in areas where changes in flow and turbulence occur.

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This type of corrosion is dangerous since it causes spontaneous catastrophic failures (Anon., 2005). For equipment potentially susceptible to erosion-corrosion wear, the following compensatory actions can be recommended: periodic in-service inspection and measurement of piping wall thickness at potentially affected areas, introduction of experimentally-calculated methods of wear rate prediction, timely performance of repair works at damaged areas, with replacement of metal sections in these areas with carbon steels having a higher content of chromium (≥ 0.5%), since such contents of chromium in steel provide for an effective reduction of the rate of erosion-corrosion wear, implementation of online monitoring with support of calculative codes, e.g. Checkworks, Wathec, Dasy, Comsy, etc. (Shah et al., 1997; IAEA-TECDOC1260). Liquid-particle erosion is another type of degradation. This type of erosion affects the blades of NPP turbines. In steam turbines, the working medium is superheated steam at high pressure. This saturated steam passes through the turbine blades to turn them, and then, after performing the thermomechanical work, it cools down and transforms into low pressure steam. During the expansion process, its temperature decreases below the dew point and as a result, water droplets of 100–200 mm diameter are formed. The front edges of rotating blades collide with these water drops with impinging speeds close to the linear speed of the blades. Owing to this, cavities are caused on the turbine blade edges, and this has consequences for the necessary aerodynamic characteristics and balance of the blades. At very moist steam conditions, containing, at the last stages of the low pressure turbine, more than 10–15% of water (wt%), the drop cavitation erosion can cause significant damage of blade edges so that their replacement is necessary. Cavitation corrosion is observed in turbines, pumps and pipelines owing to the collapse of steam bubbles near to the surface in high pressure areas of the system (Bogachev et al., 1974). At the disappearance of steam bubbles, shock waves of high pressure, of up to 1000 MPa, can occur, thus locally deforming metal and destroying protective oxide film on the surface. As a consequence, the corrosion processes are accelerated locally. Corrosion fatigue is characterized by the cracking of a metal under the action of a periodically applied tensile stress in a corrosively active environment. If this stress does not exceed the critical value, called the endurance limit strength (sometimes fatigue limit strength), the metal not exposed to the corrosive medium will not break under any duration of operation. The true endurance limit strength is not usually achieved in a corrosion medium, because the metal breaks after a certain number of cycles, regardless of the stress. If the metal subjected to constant tensile stress in a specific corrosion medium or coolant cracks after loading during a certain time, then this degradation

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and failure is due to stress-corrosion cracking. Stresses in equipment can be caused by residual stresses after machining, heat treatment, grinding and welding, for example, or can have external origins. The cracks can be either intergranular or transgranular in nature, depending on the properties of metal involved and corrosion medium acting. Degradations of this type have fundamental differences from intergranular corrosion that do not depend on whether or not the metal is in a stressed condition.

7.4

Stress-corrosion cracking

Stress-corrosion cracking represents a very important type of degradation mechanism since it may occur in many various alloys (Scully, 1990). This type of damage is characterized by the formation of multiple cracks in a metal under the simultaneous action of a tensile stress and a corrosion medium. Most of the metal surface is free from any traces of damage, but inside the matrix, a system of intergranular and intragranular cracks may be propagating over time. The stress levels typical for stress-corrosion cracking are considerably lower than the yield stress of the material, so the damage can be caused by both applied and residual stresses. Low stress levels usually mean a longer time for crack incubation and growth. Apparently there is a certain threshold stress, below which the corrosion cracking does not take place. For NPPs with RBMK reactor plants, this type of damage is of great importance since the pipelines of the forced circulation coolant circuit of the installation are made from welded austenitic stainless steel of the type 08Ch18N10T (the foreign analog is AISI 304). This steel is not prone to other types of corrosion, but it is sensitive to stress-corrosion cracking in the presence of chloride-ions Cl and oxygen O2. The severity of the corrosion cracking depends on stress, composition of alloy, environment, temperature and condition of the surface, as well as the oxide film condition. Obviously, the spreading of cracks proceeds irregularly. Growth of cracks at statically loaded sections takes place at the interface between the mechanical deformation at the crack tip and the chemical corrosion processes in the crack tip area (Bakirov et al., 2003). The maximum value of the stress intensity factor under plane strain conditions in the corrosion medium, at which crack does not grow, is designated as KI SCC and it is called the threshold value of the stress intensity factor for corrosion cracking. Depending on the corrosive medium present and tensile loading, austenitic steel, is especially sensitive to corrosion cracking when it is in a sensitized state. Uhlig and Revie (1989) have obtained data on the influence of heat treatment duration and temperature on the susceptibility of austenitic steel to intergranular corrosion. Experiments showed that several minutes of heating

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at temperatures about 700 °C are equivalent to several hours of heating at lower or higher temperatures. Intergranular corrosion is caused by the slow cooling of the alloy and its transition through a range of sensitizing temperatures during heat treatment, as well as by welding causing also the alloy to pass through the sensitizing range. With electric arc welding, the damage rate rises with the increase of heating time, especially when welding massive section components. Hazardous (sensitizing) temperatures are attained several millimeters from the welding area where the metal is heated to the melting point or higher. Hence, at contact with an aggressive medium, the destruction of the welded joint does not occur in the weld itself, but in the adjacent heat affected zone. It is well known that stress-corrosion cracking in austenitic stainless steels is highly dependent on the content of chlorides and oxygen in water. As an example, Fig. 7.4 shows at what ratios of chlorine ions and oxygen concentration the corrosion cracking takes place, and at what ratios it does not occur (Uhlig and Revie, 1989). In Russia, special research on corrosion cracking of austenitic steel in the initial and sensitized conditions and of austenitic steel welded joints was carried out (Nikitin, 1980). The tests were performed on tubular samples at temperature of 300 °C. The samples were loaded by axial tensile load, and the initial and accumulated deformation was measured. The working medium selected for the tests was a solution containing 100 mg/kg of Cl ions and 100 mg/kg of dissolved O2. As usual

Concentration of O2, mg/l

100

10 Cracking 1

0.1

0.01 0.1

No cracking

1 10 100 Concentration of Cl, mg/l

1000

7.4 Effect of chlorides and oxygen contents in water on stresscorrosion cracking of 18-8 austenitic stainless steel working in a contact with steam phase and periodically dampened by water having pH = 10.6, 50 mg/l concentration of PO43–, at a temperature of 242–260 °C.

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for corrosion cracking conditions, destruction of the samples occurred with very small plastic deformation, typically some hundredth or tenth parts of a percent. As an example, Fig. 7.5 presents diagrams of corrosion cracking of samples as a function of time before destruction from the applied stress (Nikitin, 1980). As the presented data show, the steel 08Ch18N10T (AISI 304) in the initial condition (austenitizing and high-temperature tempering) is characterized by small sensitivity to corrosion cracking, in spite of high aggressiveness of the medium. The conventional limit of corrosion cracking on the basis of 105 hours is 370 MPa, which is considerably higher than the yield stress (256 MPa) and is close to the ultimate stress (439 MPa). The criterion reflecting the resistance against corrosion cracking can thus be written as follows:

RCC – RP 0.2 RB – RP 0.2

7.1

where RCC is the conventional limit of corrosion cracking; RP0.2 is yield stress; and RB is ultimate stress. The criterion [7.1] for the described case results in a rather large value – about 60% of which is specific for high resistance to corrosion cracking (Nikitin, 1980). Sensitizing heat treatment of the steel (650 °C for 1 hour) considerably increases the sensitivity of the steel to corrosion cracking. In this state, the conventional limit of corrosion cracking on the basis of 105 hours is 300 MPa. Although higher than the yield stress of the steel, it is not significantly so. The heat affected zones of the welded joint are more sensitive to corrosion cracking. Kinetic dependence crosses the yield stress level at t = 3000 hours, s, MPa 700 600 500

1

400

2

300 3 200

100 100

101 102 103 104 1 – initial state; 2 – sensitized material; 3 – welded joints

t, h

7.5 Diagram of 08Ch18N10T steel corrosion cracking in water solution (with 100 mg/kg concentration of Cl– and O2) at temperature 300 °C.

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and the conventional limit of corrosion cracking, on the basis of 105 hours, shows a value of 200 MPa. Metallographic analysis showed that cracks, as a rule, formed in the heat affected zone of the welded joint. During the initial stage, the cracks usually grow along the lines normal to the surface, and cross grains and their boundaries. However, as they go deeper into the metal, they begin to develop in an intergranular way, mainly demonstrating a clear tendency to splitting. In general it can be stated that corrosion cracking of welded joint heat affected zones has, in principal, intergranular nature. Figure 7.6 presents data on the rate of corrosion cracking development based on results of the referenced work (Anon., 1986). In contrast to the above-stated results, the data has been processed on the basis of fracture mechanics parameters and provides for the evaluation of the rate of crack development. However, in this case, it is not clear how long the incubation period is before formation of the main crack. So, for practical calculations, research continues with sub-division of the destruction process into at least two stages: crack formation stage and 10–3

Rate of crack growth, ¥ 25.4 mm

Curve A

10–4

Curve B

10–5

10–6 0

10 20 30 40 Stress intensity factor, ¥ 1.1 MPa·m1/2

50

Curve a – material subjected to provoking thermal treatment in the oven (oxygen content is 0.2 ppm) Curve B – material sensitized at welding (oxygen content is 0.2 ppm)

7.6 Rates of crack growth at stress-corrosion cracking for austenitic stainless steel.

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crack development stage. For separation of the stages, special experiment methods have to be used. The most promising of the methods appears to be the acoustic emission method. Taking into account the great practical need for the development of robust engineering methods for corrosion cracking kinetics evaluation, it is hoped that in the next few years (or decades) this research will continue and result in practical calculations. In summary, there is a large accumulated practical experience and knowledge about the conditions for the occurrence of corrosion cracking for NPP equipment materials (effect of the medium, properties of materials, admissible stresses) (Azbukin, 1983; Azbukin et al., 1997; Bieniussa and Reck, 1997; Bogoyavlensky, 1984; Pogodin and Bogoyavlensky, 1970; see Scott, 2000). With reference to PWRs, the most typical corrosion damage in nickelbased alloys (Alloy 600, X750), corrosion-resistant steels (steels 304 and 316) and high-strength corrosion-resistant steels (A 286, A 410) have been reviewed (see Scott, 2000). The main causes of cracking are joint action of high residual and operational stresses, high aggressiveness of water coolant in crevices, use of heat treatment increasing the hardness and reducing the plasticity, or thermal ageing due to long-term operation. Steels 304 and 316 are successfully used in the deoxygenated water coolant in PWR primary circuit but are nevertheless subjected to corrosion cracking due to initial contamination of the surface of components, accumulation of sludge admixtures (e.g. chlorides, sulphates) and air ingress into dead-end crevices (e.g. canopy seals) at refuelling. A common problem in PWRs is the cracking and leaking of coolant at RPV head penetration areas, where the seals of casings for control rod drive mechanisms are located. The issues of power equipment corrosion damage are covered widely and in detail in monographs (Pogodin and Bogoyavlensky, 1970; Azbukin, 1983; Bogoyavlensky, 1984) and reviews (Scott, 1987; Hanninen and Torronen, 1987) for the period from 1970 to 1987. The incidences of corrosion damage to structural elements, taken from experience of NPP operation in Russia, analysis of their causes and implemented mitigative actions are given in Azbukin et al. (1997). The NPP items most susceptible to corrosion cracking are as follows: steam generator tubing, primary circuit pipelines and parts of the reactor pressure vessel internals. As regards the tubing, since the steam generators have a considerable reserve of thermal power, the loss of tightness of single tubes is not regarded as an emergency situation. In a period of planned preventive maintenance, the damaged tubes are usually plugged. In-service damage of pipelines, caused by corrosion cracking, mainly occur in boiling water NPPs (i.e. RBMK and BWR). The common aspect for all cases of pipeline degradation made of steel 18-8 is that the intergranular cracks occurred either in heat-affected areas, in which the metal was sensitized at

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welding, or resulted from a sensitizing heat treatment of pipeline components during equipment manufacture. In NPP operation abroad (e.g. USA, Germany and Japan) this problem is ranked as a top priority (Danko and Stahlkorf, 1984). When analysing the operational experience of 21 light water cooled reactors in Germany, it was noted that the cracking mechanism appeared in pipeline systems made of stabilized austenitic corrosion-resistant steels. Review of corrosion damage of German light water reactors over 20 years is given in Bieniussa and Reck (1997). The analysis of pipeline failures at 14 German PWRs revealed that during 20 years of operation only a small amount of corrosion damage occurred in the base metal and then only in small diameter pipelines of niobium-stabilized austenitic steel. The cause was chloride-induced transcrystalline cracking, owing to the absence of circulation of the medium in cracks (i.e. chemical hide-out). During 20 years of operation of German BWRs there were detected about 30 instances of corrosion damage related to the mechanism of intergranular stress-corrosion cracking. The damage occurred in pipelines of small (up to 25 mm) and large (250 mm and more) diameter made of steels 1.4550 and 1.4551 (Ti-stabilized) mainly in the areas of welded joints (on large size pipes). The damage was caused by corrosion caused by a combination of residual stresses, material sensitization and oxygen in the coolant plus manufacturing defects or shortcomings such as undercuts and displacement of butt-edges in welded joints. The damage was found in residual heat removal systems, coolant cleanup system, sealing and flushing water supply systems. The depth of cracks in large diameter pipelines attained 3/4 of the wall thickness; however, no leakage occurred (Erve et al., 1997). Based on the experience of German light water reactor operation, it is evident that the role of titanium or niobium stabilization alloy additions in the prevention of intergranular cracking has been overestimated (Jungclaus et al., 1999). Russian stabilized austenitic steel 08Ch18N10T has high corrosion resistance in high-temperature oxygen-containing water coolant. Nevertheless, at several Russian NPPs with RBMK reactor plants, formation of intergranular stress corrosion cracks in welded joints of 300 mm nominal diameter pipelines and in welds of the bottoms of distributing pipe-group headers, manufactured of the same steel, was also found after operation times of about 100 000 hours. This corresponds to about half of the design lifetime of RBMK NPP operation (Karzov, 1998). In order to extend the service life of equipment made of alloys prone to corrosion cracking, the most efficient procedure is the in-situ technique for the reduction of mechanical plastic deformation of welded joints. This procedure is based on the redistribution of the residual stresses at pipeline cross-sections in the area of a welded joint. As a result of external plastic deformation, the stresses in the weld root area, being in contact with coolant,

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change from tensile stresses to compressive stresses. This procedure is successfully applied on pipelines in the United States, Japan and Russia for the reduction of operational damage related to the intergranular stress corrosion mechanism.

7.5

Evaluation of impact of thermo-mechanical loading on strength of equipment materials

7.5.1 Material degradation mechanisms (time-dependent mechanical properties) Long-term operational thermo-mechanical effects can be considered on the micro and macro levels. Evaluation of phenomena occurring in material on the micro level facilitates understanding of the origin of the material reaction to external influences. For these processes, investigation structural methods of research are used. The obtained experimental results are practically not subjected to any normative regulation. Phenomena taking place at the macro level already allow the use of normative calculation methods of continuum mechanics and engineering approaches for the calculation of strength and durability. In these calculations, the assumed material properties are the properties averaged over the whole volume of the equipment material under consideration. These design approaches make a basis for the substantiation of equipment durability in nuclear power engineering. At the same time, as has been demonstrated by many years of practical analysis of NPP equipment operational damage, understanding of processes ongoing in local micro areas provides for determination of real damage susceptibility of equipment. Based on this, the research into material behaviour in the intermediate scale condition between the micro and macro levels, so called mesoscale level, seem to be promising. Thus, macroscopic mechanical behaviour of material depends on the aggregate total of effects taking place on different scale levels. Despite the external difference in behaviour of material of a sample (or a component) under fatigue and creep, these phenomena have a common feature representing the accumulation of irreversible damage to the material over time. Thus, for example, at high-cycle fatigue, the stresses are elastic and the deformation is reversible, but after a certain number of cycles, a crack originates and destruction occurs. The nature of susceptibility to damage lies in structural heterogeneity and anisotropy of the structure of materials. Owing to this, a local irregularity in the stressed-deformed state occurs. Eventually all types of damage irreversibility occur at the microscale at rather small externally applied loads and deformations, and detailed accounting of these types of irreversibility is hardly possible. Therefore, the susceptibility to damage of the material can be taken into account using models allowing the

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conversion of microscopic effects into the class of macroscopic, measurable design values. For example, susceptibility to damage can be assessed by means of a scalar quantity reflecting, generally speaking, the decrease of sample cross-section area, deemed equivalent to degradation of material properties. Damage equal to zero means that the material is intact (not damaged), damage equal to one means that the material is destroyed. In this case, the physics of the phenomenon are not specified, but the governing equations can be written and solved and, eventually, useful results for practical application are obtained. The process of equipment lifetime exhaustion in the course of operation has the following stages: accumulation of damage until the formation of a microcrack; growth of the microcrack into a macrocrack; growth of the macrocrack to dimensions of a detectable through-wall crack. In this case, the limiting state is taken as the moment of microcrack formation (the first stage of the process of lifetime exhaustion). This state is characterized by limiting the susceptibility to damage of the material to numerically equal one. As the practice of NPP equipment operation shows, the limiting state of crack formation at thermo-mechanical cyclic loading can occur in zones of stress concentration as a result of the impact of power and temperature loads, as well as in zones outside the stress concentration areas. The cyclic damage of metal to be calculated is determined both by use of values of specified deformations (conventional elastic stresses), and the ‘history’ of loading (quantity and sequence of operational and test modes). Quantitative measures of susceptibility to damage can be assumed as the current number of loading cycles in relation to the limiting number of cycles at the given value of the load of a cycle. For example, according to the Russian standard (PNAE G-7-002-86 Strength Calculation Norms), calculation of the admissible number of cycles of specified amplitude of stresses or, on the contrary, admissible amplitude of stresses for the set forth number of cycles, shall be made using fatigue diagrams whose equations for NPP equipment steels are known. If several types of cyclic loads are dealt with, the linear law of damage accumulation is used. Thus the number of cycles of the specific type is related not to the limiting, but to the admissible number of cycles of the same type. The obtained sum-total should not exceed one. This statement can be recorded as a cyclic strength condition:



K S N i = a ≤ [aN ]

i =1

N oi



7.2

where Ni is number of cycles of i type during the operation time; K is total number of cycle types; Noi is limiting number of cycles of i type, at which microcrack forms in material; a is accumulated susceptibility to damage and [aN] = 1 – limiting susceptibility to damage.

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The most evident quantitative and at the same time physical measure of susceptibility to damage is the plastic deformation. Its irreversibility allows its interpretation as an overall damage of metal integrating all types of metal microdamage that have occurred in the course of deformation. For example, in calculations of creep-rupture strength, in the same way as in calculations of fatigue, the conditions of non-exceeding of admissible deformations are introduced. When several loading modes are considered, the linear law is used as well. Making analysis of the residual lifetime of equipment under thermal loading, the accumulation of damage is a function of the operating time. Similar to fatigue, the time of operation in this mode is related to the admissible time of operation in this mode only. The amount of damage accumulated in various modes should not exceed one. The procedure of strength and durability calculations under the effects of fatigue (high-cycle and low-cycle) and creep consists in the determination of strength and durability parameters, at which the accumulated damage attains the limiting values. In this case, the environmental effect participates in the acceleration of reaching the point in time where limiting states occur, owing to the intensification of material damage processes. Deformation and destruction calculation models chosen for evaluation of residual lifetime of specific equipment need an exact definition of boundary conditions. Such boundary conditions need a basis from experimental research that accurately simulate operating conditions for the equipment. In order to reduce the uncertainty of calculated estimates related to simulation of inservice loading models for the next operating period (i.e. for the additional time relative to the design service life) and to lessen uncertainty related to the choice of actual mechanical properties of aged metal, new methods and procedures of in-situ monitoring have to be introduced. Such an integrated approach assures more objective information about the real operational susceptibility to damage. However, in NPP operation, there has not yet been established a unified methodology of non-destructive testing application intended for ageing monitoring. Efforts in this field are underway in many countries (Anon., 1999, 2002; Bakirov et al., 2002, 2006). Due to the fact that non-destructive methods of inspection assure acquisition of data directly from real equipment, with a periodicity required for analysis of ageing rates in local areas subjected to maximum loads, it can be stated that results of such in-situ inspection represent ‘bridges’ interconnecting the theoretical durability forecasts with a real construction. Recommendations, based on specific examples of in-service ageing analysis, on the use of non-destructive testing methods for inspection of thermocyclic damage of NPP equipment metal will be given later in this chapter. The existing differences in calculation procedures set forth in normative regulations of various countries require a comparative analysis and elaboration of recommendations on their application. This should be facilitated by

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collaboration to improve efficiency and provide improved estimates on equipment strength and reliability.

7.5.2 Fatigue NPP operating experience shows that most cases of failure of components or structures in service can be attributed to the fatigue damage mechanism. Therefore, the residual lifetime of NPP equipment is often determined concerning this damage type. The word ‘fatigue’ refers to the behaviour of materials under the effect of cyclic strains or stresses. Anon. (2000) gave the following definition of fatigue: ‘the process of progressive localized permanent structural change occurring in a material subjected to conditions that produce fluctuating stresses and strain at some point or points and that may culminate in cracks or complete fracture after a sufficient number of fluctuations’. August Wohler originally undertook a systematic investigation of fatigue (Stephens et al., 2001). Using stress versus life (S-N) diagrams, Wohler in the middle of nineteenth century showed how fatigue life decreased with higher stress amplitudes and that below a certain stress amplitude the test specimens did not fracture. Thus he introduced the concept of the S-N diagram (Wohler’s diagram) and the fatigue limit. He pointed out that for fatigue, the range of stress is more important than the maximum stress. Fatigue life is such a number of cycles of stress or strain of a specified character that a given specimen sustains before occurrence of the failure due to fatigue (Terentev, 2002). The fatigue strength is a hypothetical value of stress corresponding to the exact number of cycles when failure happens, as determined from S-N diagrams. The fatigue limit is the limiting value of the median fatigue strength as the number of cycles becomes very large. The above ASTM definitions are all based on median lives or 50% survival. The endurance limit is not defined by ASTM but is often implied as being analogous to the fatigue limit. The progression of fatigue damage, in general, can be classified into four stages: early fatigue damage, fatigue microscopic crack initiation, fatigue crack propagation and fracture (see Fig. 7.7). The original approach to build fatigue damage allowance into design concepts involved characterizing the total fatigue life to failure of initially uncracked test pieces in terms of the number of applications of cyclic stress range (the S-N curve in ‘high-cycle fatigue’ referred to the number of cycles 105 < N < 108) or a cyclic strain range (‘low-cycle fatigue’, N < 5 · 104). Test methods are described in detail in the ASTM Standards E466–E468. When conducting fatigue calculations there are two main approaches (see Fig. 7.8): the S-N approach and the fracture mechanics approach (Alkazraji, 2008). The S-N approach is generally used at the design stage for pressure

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Stress

S Fatigue Fatigue crack final   crack fracture line    growth     region fatigue crack nucleation line

Number of cycles

N

7.7 S-N scheme of fatigue crack nucleation, growth and final fracture. Fatigue problem S-N approach

S-N approach

Bases on curves of fatigue life for unknown materials, and also effects on discontinuities such as welds

Assumes a pre-existing defect in the material or structure

Paris equation Look-up S-N curves for different welds

Initial and final defect sizes

7.8 S-N and fracture mechanics approaches for fatigue calculations.

vessels and is described in codes such as the British Standard Specification. This code provides fatigue curves for different classes of welded joints. To conduct a defect assessment of features (i.e. a pre-existing defect) such as cracks and laminations, a fracture mechanics approach is required. Many components of power equipment suffer the effects of cyclic loads with the number of cycles N > 105. At these loads, the material can be damaged owing to fatigue or corrosion fatigue. Most often such damage is caused by vibrational loads or fluctuation of temperature. An example of this is the formation in vessels of thermal fatigue cracks located at steam–water interfaces in the coolant-level fluctuation area or at areas of mixing of coolant streams having considerable difference in temperature. Damage of this kind has been observed on heating steam distribution pipes of separator-superheater at high local humidity downstream to the separator (moisture inflows). At high-cycle fatigue, the temperature variations are not very significant, and the number of cycles is comparable on the basis of usual fatigue tests or exceeds it. Under these conditions, the process can be considered as isothermal in nature and no difference should be made between the thermal fatigue and isothermal mechanical fatigue. Just as in the case of mechanical

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fatigue, here the durability will depend on the dimensions of components, presence of stress raisers, cycle asymmetry, frequency of stress variation, quality of surface treatment, temperature, corrosion and other technological factors (Stephens et al., 2001). The statistical nature of fatigue destruction leads to a considerable scattering of the number of cycles before failure, even when testing laboratory samples in air, using the same testing machine, identical samples manufactured from the same melt and strictest observance of experimental conditions. Inter-melt variations and additional factors can increase the difference in number of cycles before failure by several orders of magnitude, especially when small amplitudes of stresses are used. Therefore, the fatigue curves of materials of the same grade, obtained in the course of different research campaigns at different testing facilities, using test samples of material taken from different supplies can, as a rule, significantly differ from each other. The practice of lifetime estimation by calculation for different equipment allows some general recommendations to be made on the selection of material fatigue curves. The most reliable results can be provided only by use of full probability maps of material fatigue, specially developed on multiple numbers of samples. It is desirable that dimensions of samples, quality of surface treatment, temperature and frequency of tests and composition of the environment conform to operating conditions of components under investigation. For equipment that is already in operation, it is practically impossible to do so. In the Russian case, in complete absence of experiment data, the recommendations of the standard (PNAE G-7-002-86 Strength Calculation Norms) shall be applied. In this case, the calculation of lifetime and, correspondingly, plotting of fatigue curves shall be made for two probabilities of failure, namely, p = 0.5 and close to zero (p ª 0.001). When constructing the first curve, the coefficients of stress margin ns and cycle number margin nN are not included into the formulas recommended by the standard (ns = nN = l), and it is supposed that this curve corresponds to failure probability 0.5. A curve of admissible amplitudes (ns = 2, nN = 10) is used for estimation of the lifetime at a failure probability close to zero. It was supposed that such a curve gives minimum values of endurance limit with deviation from the average one, equal to three mean square deviations (assuming a normal law of distribution, the fractile equal to 3 corresponds to the probability p ª 0.001). Consideration of the joint impact of various factors on material fatigue resistance of power installation components (scale factor, availability of welded joints, condition of surface, corrosion, stress concentration, etc.) requires an individual review. When test samples are subjected to testing in conditions similar to the maximum extent to the real operating conditions, the real situation impact

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is taken into account automatically. For the determination of fatigue characteristics by calculation, the joint effect of all factors is often taken into account through multiplication with corresponding coefficients obtained experimentally and reflecting effects of each individual factor. With regard to NPP power equipment components, it is not always acceptable. For example, when significant corrosion impact of the medium is already present in the initial condition, the surface condition makes little effect. There is some experimental data on the joint impact of several factors. Karpenko (1985) presents results of research of the joint impact of stress concentration, scale factor and corrosion on the fatigue behaviour of various steels. It was determined that with the increase of diameter of the sample the endurance limit reduces in air and increases in the corrosion medium, and the presence of stress raisers abruptly reduce the limit of endurance in the air, but not very considerably, in the corrosion medium. The author explains these effects by the cooling action of the corrosion medium being more intensive on larger samples, as well as by removing the bottom of the undercut (concentrator) and the appearance of simultaneously developing cracks causing a lessening in magnitude of the main stress acting. The data for evaluation of corrosion impact of the coolant medium used in NPPs may have a tendency to underestimate the corrosion resistance value. It is explained by the fact that experimental research of moisture action is usually made using tap water without corresponding analysis of chemical composition, and free access of air oxygen takes place. Components of power equipment operate, as a rule, in coolant which is carefully monitored for its composition. The performed experiments showed that even if there was an impact of medium on fatigue resistance, this impact was considerably smaller than that stated in the data of standards. The same is also confirmed by experimental data given in the work by Mahutov (1985). Thus the conclusion can be made that it is almost impossible to develop theoretically a material fatigue curve that correctly accounts for the joint impact of various factors on the basis of the high number of cycles for operating conditions of power equipment components. The preference shall always be given to the experimental development of these curves. As a rule, the approach described below is usually used for calculations. If there is no recommendation on material resistance to fatigue, the design curves are built according to recommendations of the PNAE G-7-002-86 Strength Calculation Norms as per the above-stated interpretation, i.e. for 0.001 and 0.5 probabilities of failure. Additional effects of corrosion can be accounted for by a reduction of fatigue curve ordinates in accordance with the coefficient:

bK = s–1NK/s–1NB

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7.3

Impact of operational loads and creep, fatigue and corrosion

171

where s–1NK and s–1NB are the endurance limits for the set forth number of cycles N in corrosion medium and in air respectively (see Fig. 7.9). It is recommended to apply the following bK values, equal to minimum values as per the data of Mahutov (1985): 0.95 for stainless steels; 0.75 for low-carbon steels. Following acquisition of experimental data, the stated recommendations can be specified more precisely. The same approach is also used in thermal power engineering (OST 108.031.08-85; OST 108.031.10-85). Since the parameters of equipment in organic fuel power plants are higher and a lot of equipment operate in creep conditions, the low-cycle fatigue and creep damage are summed up following the recommendations of OST 108.031.08-85 and OST 108.031.10-85. Some experimental and calculation results, related to building of fatigue diagrams in corrosion conditions for austenitic stainless steels, are presented at the end of this section (Anon., 1995; Filatov, 2004, 2007; Kochik et al., 2005; O’Donnell et al., 2005). The data are presented without any details related to executed calculations and experiments. Comparative analysis shows satisfactory compliance between results of tests on fatigue of austenitic stainless steels in a heat-treated state (the tests were performed in air conditions and in reactor water conditions) and fatigue curves, which were calculated using empirical equations. Average mechanical characteristics of steels and experimentally determined reduction of fatigue strength due to tests in reactor water were used in the equations. Figure 7.10 shows the results of testing some stainless steel grades in air environment at 20 °C and with a controlled strain symmetrical cycle (O’Donnell et al., 2005). A number of force controlled tests were carried 1 1

Coefficient bK

0.8

2

0.6

3

0.4 0.2 0 0

50

100 150 200 250 Temperature, °C 1 – water of high purity; 2 – distilled water; 3 – 3% NaCl solution

300

7.9 Effect of corrosion on the endurance limit depending on environment and surface temperature.

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e a, % 10.00 1

1.00 3

4

6

2

5 0.10 8

7

N

0.01 101 1 4 6 7 8

– – – – –

102

103

104

105

106

107

108

Langer curve; 2 – ANL model; 3 – japanese model; design curve in current use; 5 – design cuve proposed by ASME; curve on fatigue failure criterion, symmetrical cycle, ns = nN = 1; design curve on fatigue failure criterion (str)max = R20 p02, ns = 2, nN = 20; design curve on fatigue crack initiation criterion (str)max = R20 p02, ns = 2, nN = 10

7.10 Results of experiments in air environment (austenitic stainless steels of series 3XX, at 20 °C) as compared with fatigue curves based on different models.

out for the high-cycle region. In Fig. 7.10 the following curves are built: 1 – Langer curve; 2 – ANL model; 3 – Japanese model; the design curve in current use (4) and proposed (5) by ASME; and the curves calculated from equations: 6 – curve on fatigue failure criterion, symmetrical cycle, ns = nN 20 = 1; 7 – design curve on fatigue failure criterion (str)max = R p02 , ns = 2, nN 20 = 20; 8 – design curve on fatigue crack initiation criterion (str)max = R p02 , ns = 2, nN = 10. From Fig. 7.10 it is evident that all the models of fatigue curves, despite some differences in premises being taken for their construction, lay in the scatter band of the experimental data. Thus, serious deviations of calculation results based on various norms are not expected to exist. It also should be noted that irradiated austenitic steels demonstrate a lower cyclic strength in reactor water compared to tests in air environment, reduction of deformation rates in water tests is attended by a reduction of durability likewise, as is also observed for non-irradiated steels.

7.5.3 Creep Creep in its simplest form represents a gradual accumulation of plastic deformation in a component under stress at high temperature during a certain © Woodhead Publishing Limited, 2010

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period of time (Evans and Wilshire, 1999). Usually it is considered that for metals the creep becomes apparent at temperatures above 0.3–0.5 °C from the melting temperature Tm (or solidus temperature for steels and alloys). However, research on deformation and destruction mechanisms showed that the creep processes start at various levels of homological temperatures (homological temperature is a temperature normalized to melting temperature). In Table 7.1 there are presented data on the minimum temperatures, at which the creep processes start to be observed, depending on the type of bonding in the lattice and on the type of crystal lattice present in the material. According to the table, the minimum temperature of austenitic steel creep is 0.3 Tm, and that of chrome-molybdenum and chrome-molybdenum-vanadium steels is 0.35–0.4 Tm, whereas in engineering practice, it is assumed that creep shall be taken into account for austenitic steels at temperatures about 100 °C, and more, above the temperatures corresponding to carbon and low alloy steels. This difference is related to the fact that the time of failure in creep conditions at temperatures 0.1–0.3 Tm exceeds, by several orders of magnitude, the service life of the component being considered for use in power equipment design and operation. It should be mentioned that, at relatively low homological temperatures, creep can impact the growth of cracks, strength and plasticity at varying levels of loading. As of today, these effects have been studied insufficiently and in practice are not taken into account for the resolution of engineering tasks. As has already been noted, the strength of materials working in creep Table 7.1 Minimum temperature of creep for various groups of materials Group designation

Type of bonding, type of lattice

Minimum temperature of creep

Metals and their alloys Metallic bonding, face-centered cubic

0.1Tm – Monel400 (0.1 – 0.2)Tm – Ni 0.2Tm – Cu, Al, Pb, RR58, DS-Nickel (Ni + 2%ThO2), Nimonic-80A, Inconel-X750 0.25Tm – nichromes 0.3Tm – Ir, Rh, austenitic steels

Transition metals Metallic bonding, and their alloys body-centered cubic

0.2Tm – Ta 0.25Tm – Nb 0.3Tm – W, Mo, Cr (0.35–0.4)Tm – a-Fe, Cr-Mo, Cr-Mo-V and 9–12% Cr steels

Metals and their alloys Metallic bonding, face-centered close-packed

0.3Tm – Re, Be, Ti, Mg, Magnox-L80, Magnox-ZR55 0.4Tm – Cd

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conditions decreases in the course of operation owing to changes in microstructure. As a result, pores and microcracks occur (mainly at grain boundaries). Pores can merge together and form microcracks and the latter, in turn, can grow into macrocracks and finally lead to failure of the equipment. For reactor plant equipment that comes into contact with water coolant, the level of temperatures are such that the creep processes for real service life periods should only be accounted for in items made of zirconium and titanium alloys. The calculation is to be made in accordance with recommendations of the PNAE G-7-002-86 Strength Calculation Norms, and checking the calculation of which covers further calculations of static and cyclic creeprupture strength Calculation of static creep-rupture strength in modes differing by reduced stress or design temperature uses the linear summation of accumulated static creep-rupture damage according to the formula below:



K S ti ≤ 1 i

[t ]

7.4



where ti is accumulated time of work at i – mode; and [t]i is admissible accumulated time of work at the same mode. For calculations for creep-rupture cyclic strength, the duration of units’ operation at temperatures exceeding the Tt value shall be taken into consideration. Tt refers to the temperature above which the characteristics of creep-rupture strength, plasticity and creep shall be taken into account. According to the PNAE G-7-002-86 Strength Calculation Norms, the temperature Tt, is assumed equal to: ∑ 20 °C for aluminum and titanium alloys; ∑ 250 °C for zirconium alloys; ∑ 350 °C for carbon, alloyed, silicon-manganese and high-chromium steels; ∑ 450 °C for corrosion resistant austenitic class of steels, temperature resistant chrome-molybdenum-vanadium steels and nickel-iron alloys. For calculation of creep rupture strength, the impact of non-metallic inclusions is to be taken into account. Whereas in nuclear power engineering the damage in creep conditions is considered only for a limited number of components of power units with WWER and RBMK reactor plants, and this damage is only really relevant for equipment in high-temperature reactors with liquid-metal and gaseous coolant (in Russia these are fast reactors BOR-60, BN-350, BN-600), in thermal power engineering, this type of damage is the principal one. Therefore, in thermal power engineering, great experience of practical application of durability prediction methods for components operating in creep conditions has been © Woodhead Publishing Limited, 2010

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accumulated. Specifically, there are about 100 experimental installations where tests of real components of boilers and pipelines are performed under field loads and temperatures up to 800 °C. It should be particularly noted that many results of predictive durability calculations have been proven in practice (Coffin, 1976). In Russia, the works in the field of durability estimation for SSCs operating in creep conditions should be noted, which were carried out by the following organizations: Central Research Institute of Structure Materials ‘Prometey’, Institute of Physics and Power Engineering, Afrikantov Experimental Machine Building Design Bureau, Engineering Strength Center of N.A. Dollezhahl Research and Development Institute of Power Engineering (NIKIET). By order of Concern ‘Rosenergoatom’, these organizations have issued the regulation RD EO 1.1.2.09.6714-2007. Most often (especially in the initial stages), the calculations are made using approved regulatory procedures and the linear cumulative damage rule. Approaches used for development of regulations are similar to those considered above, though there are some differences (ASME Boiler and Pressure Vessel Code, 1980). The joint effect of fatigue and creep is taken into account by summing up the damage:



N + Ê1.25 s c ˆ [N ] ÁË s c /t ˜¯

m

≤ Dcreep



7.5

where N is the set forth number of cycles; [N] is the admissible number of cycles; sc is the maximum local design stress, determined accounting for the creep; sc/t is the conventional limit of creep rupture strength at tension; and Dcreep is cumulative damage. It is recommended to use:

sc/t = 1.5 [s]

In addition to this, m is the exponent of creep rupture strength. The values of cumulative damage Dcreep are presented in Fig. 7.11. If Dcreep < 1 the crack does not initiate, and one can perform calculation of fatigue damage of the construction. During recent years, calculation procedures for the determination of material damage rate, based on results of metallographic research, has been intensively developed worldwide. Relevant research shows that accumulation of damage at the grain boundaries occurs at the level of operational stresses below the yield stress value for the metal; and damage inside the grains occurs if the stress level exceeds the yield stress. Apparently, the most interest is paid to procedures based on non-destructive tests of replica models. However, these procedures have fundamental limitations since they are based on structural analysis of a superficial layer only.

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Understanding and mitigating ageing in nuclear power plants 1 0.9 0.8 0.7 Dc

0.6 0.5 0

0.1

0.2

0.3

0.4

0.5

0.6

0.7

0.8

0.9

1

[1.25(sC/sC/t)]m

7.11 Cumulative damage as a function of damage caused by creep.

More exact evaluations are made through the analysis of metal samples. The sections of material selected for analysis should be representative with regard to the damage rate of the corresponding components. Deterioration in the material, revealed by metallography, is estimated in quantitative units, which are used in state equations for the calculation of residual lifetime. Thus estimation of the degree of reliability depends, to a great extent, on the quality of the equation used and the quantitative evaluation of accumulated damage detected by metallography. At the present time, this model works only for materials prone to creep at grain boundaries. Thus it is not applicable for martensitic steels. When necessary, the calculation methods are supplemented by hightemperature tests under stress or creep tests to destruction. These can be tests of samples or real components at constant (working) stress, but at higher temperature in order to get results in 1500–3000 hours under load. In the practical case, for ‘aged’ boiler and pipeline components, the most exact and reliable results are obtained after tests of real elements (for example, from samples of pipes of heating surfaces) cut from areas most affected by damage.

7.6

Equipment condition monitoring, prediction and testing

Implementation of works on proof of prolongation of the projected lifetime of the NPP equipment requires forecasting of the equipment material behaviour during the projected additional service life. For this purpose, in laboratory conditions, it is necessary to simulate the accumulation of static and cyclic damage caused by joint action of several damaging factors of various intensity. The most important factors which affect the mechanical behaviour of the

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construction material are: temperature, speed of a deformation process, cyclicity parameters, features of the stressed state, absolute dimensions of working parts and environment (Mahutov and Permyakov, 2005). As practice shows, the accuracy required in the analysis of the residual lifetime is attained through application of combined research methods of the analysis of an operational loading history, nominal and local stresses, deformations, such as: ∑

carrying out the modelling in bench-top conditions, with application of small-sized and full-scale test specimens (models) and reproducing the loading conditions as close to the real ones as possible; ∑ execution of periodical in-situ investigation of NPP components’ metal in zones having maximum operational loading.

In such research the following approaches must be preferably used: ∑

on-line continuous measuring in each time period ti of forces F(ti), local deformations e (ti) and stresses s (ti) using the methods of tensometry, roentgenography, holography, magnetometry and acoustic emission; ∑ on-line continuous measurement of vibrations, pressure pulsations and temperatures; ∑ periodical in-situ inspection of changes of the metal’s mechanical properties and microstructure using specimen-free non-destructive methods. In order to facilitate realization of the mentioned approaches, it is necessary to use special automated (computerized) instrumental systems (hardware) which allow restoration of the history of the actual operational loading F(t i), s(ti), e(ti) and to evaluate the accumulated damage sustained, as well as the actual residual lifetime. Availability of such integral experiment-calculated information on forces, stresses, deformations and so on is the basis for building of the ultimate (critical) loading curves for the construction material of specific equipment:

Fcrit = f {(sred, ered)max, t, N, t}

7.6

where Fcrit is the critical combination of mechanical and thermal forces for the chosen loading modes at definite temperature t, number of cycles N, time t; sred is local reduced stress; and ered is local reduced deformation. Fcrit depends on the values of local reduced stress (sred)max and local reduced deformation (ered)max which are determined through the equation of non-isothermal fatigue (low- and high-cycle) curve taking into consideration the material mechanical properties:



{(s red )max , (ered )max } = f

Ê N , s uts , Y, Su , t sy, n Ê ÁË

7.7 ˆ ˜¯

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where N is number of loading cycles; suts is ultimate tensile stress; Y is cross-section at fracture; Su is true fracture strength; t is temperature; sy is yield stress; and n is tensile strain hardening exponent. The described approach for modelling of materials ageing was used in the work on assessment of the microstructure and mechanical properties degradation of WWER-440 and RBMK-1000 austenitic piping (Mamaeva and Bakirov, 2003; Mamaeva et al., 2006). Hereinafter, the work being executed is briefly presented as being representative. Several identical full-scale test samples were fabricated from the archive pipelines metal (18-8 austenitic stainless steel). They had the form of rings having a length of 300 mm and a circumferential welded joint in the middle. The full-scale test samples were welded in accordance with a standard technology which was used for austenitic pipeline welding 30 years ago, i.e. at the time the NPP was commissioned. After the samples were welded they were subjected to investigations by specimen-free inspection methods intended for assessment of the chemical composition, mechanical properties, microstructure and residual stresses in order to check out the identity of the samples and to obtain comprehensive data related to the initial state of the investigated metal before ageing. One full-scale test sample was cut to manufacture standard test samples which were subjected to usual destructive tensile testing in order to collect statistical data concerning the initial values of the mechanical properties; much more detailed analysis of the metal microstructure was carried out as well. At the next stage, the analysis of real operational loads of investigated NPP pipelines was performed. It allowed determination of the loading modes simulating the impact due to thermally-forced operational loading extended to a 50-year term of NPP operation. Full-scale samples were subjected to testing in a specially designed laboratory test bench – a resonance testing machine, which allows carrying out a reversed cyclic bend loading of the samples at the definite temperature being selected (see Fig. 7.12). The samples were equipped with high-temperature resistance strain gauges along the welded joint perimeter, inductive sensors to detect bend displacements, and also by thermocouples. A control of the loading process was fully automatic and realized through a computer providing a possibility of on-line monitoring. Calculation of loading bend stresses was done in accordance with the requirements of PNAE G-7-002-86 Strength Calculation Norms. Additionally, a finite element method calculation of the test bench was carried out using as input data the selected loads which had to be applied to the full-scale test sample to simulate the ageing deformation. Calculation of fatigue curves was executed using actual mechanical properties obtained after testing of the standard samples cut out from the original material. For the purpose of execution of the accelerated thermal ageing, the temperature of testing was set in accordance with recommendations of the work performed by Asano

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Impact of operational loads and creep, fatigue and corrosion Analysis of the initial metal propelties

Assessment of the loading parameters

179

Design of the test bench Loading motor

full-scale sample {F}

Loading grip Stage 2 FEM calculation of deformation loading

Stage 1 Fabrication of the full-scale samples and preliminary test

Elaboration of the experiment-calculated model of loading e, % 0.26

N = 4000 t = 20 °C

N = 24000 t = 450 °C

Stage 3 Preparation of bending resonance test machine using computer control

Change of the metal properties and microstructure 50 years of operation

0.10 suts, sy, y, Kc etc.

0 Simulation of construction and starting initiaI forces {F }

Simulation of operational forces {F }

Stage 4 Thermo-deformation cyclic loading of the full-scale sample

Stage 5 Examination of the metal properties after ageing

7.12 Stages of the thermo-deformation ageing and research of the full-scale test sample.

et al. (1996). Duration of the thermal ageing was determined based on the theory of activation processes. Accumulated actual fatigue damage resulting from the applied symmetric cyclic loading appeared to be considerably lower than the value corresponding to a crack formation. Nevertheless, the value of the fatigue damage in the considered experiment exceeds one if it is calculated by using the normative fatigue curves prescribed for application by the PNAE G-7-002-86 Strength Calculation Norms and constructed using the normative margin coefficients on the stress ns = 2 and on the number of cycles nN = 10. When testing was completed, the aged specimens were cut into pieces in order to manufacture standard test samples intended for the comprehensive material science investigations. The data obtained were compared to the data

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related to the initial un-aged state of the metal. At all stages of the experiment there were also conducted comparative specimen-free investigations in order to find out the possibility of ageing assessment by using various nondestructive methods such as hardness, electric conductivity and magnetic permeability. The analysis of the experimental data shows that due to thermal-forced ageing of the austenitic steel metal, its strength properties – Rm and Rp0.2 – as well as hardness have increased 10–15% on average, Charpy impact strength KCV decreased by 20–25%, plastic fracture toughness JIC of the weld metal decreased by 40%, low-cycle fatigue and fatigue crack growth properties had changed insignificantly. Ferrite content (d-ferrite) of the weld metal had increased by 1.5–2% on average. The results of the investigation have also shown that the most sensitive indirect parameters suitable for detection of the pipeline metal condition under cyclic loading are the electric conductance and the field gradient of a residual magnetization. It was found that the electric conductance is decreasing during cyclic loading, and the field gradient of a residual magnetization is, as a rule, increasing during cyclic damage accumulation. Consequent decrease of the electric conductance of the base metal after additional thermal ageing for 1500 and 3000 hours have been indicated as well. The possibility of application of electro-magnetic instrumental facilities for qualitative assessment of the austenitic steel condition and damage to welded joints, as well as to reveal surface cracks, have been demonstrated. Parallel with the study of the physical-mechanical properties of the austenitic test samples subjected to artificial ageing, detailed metal structure investigations were also carried out. The analysis allowed changes in the metal microstructure to be revealed, which in turn stipulates the mechanical properties of piping welded joints in the course of long-term operation. The changes observed in the weld metal consist in decay of the d-ferrite with chromium carbides M23C6 evolving, and, in some cases, with formation of the brittle a¢-phase. Regarding the heat affected zone, the changes were in increased evolution of carbides at grain boundaries. Also it was demonstrated that microstructure change due to long-term operation can cause favourable conditions for the propagation of a stress corrosion crack not only in a heat affected zone, but also in the weld metal, and this fact should be taken into account in the calculation of the residual lifetime of the piping. In accordance with the results of the experimental work, it can be stated that ageing of the austenitic pipeline metal after 50 years of operation in artificially modelled operational conditions does not result in changing of the basic mechanical properties outside of the allowable limits prescribed by the adopted normative standards. This statement is valid only if construction rules and operational conditions of the SSCs are fully complied with. However, in some cases when dealing with the task of ageing analysis,

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it is necessary to define not only the actual mechanical properties, but also to find out if the processes which are in compliance with the SSCs design, are adhered to, but may not be acceptable. The most relevant example is to demonstrate the absence of metal creep processes of Russian-designed BN-600 NPPs piping having a liquid metal coolant. For this purpose for the BN-600 NPP, after its 30-year operation, an appropriate inspection procedure based on the hardness creep test (see Li et al., 1991) has been elaborated. There are a number of possible mechanisms of plasticity some of which may contribute to indentation creep, such as plasticity (dislocation glide, slip plane), recovery (dislocation climb) and diffusion (grain boundary, volume). All of these creep mechanisms are schematically illustrated in Fig. 7.13 (see Li et al., 1991). A high level of similarity of creep mechanisms observed in both indentation and tensile tests is confirmed by the similarity of the primary creep curves. It comes from the analysis of the primary creep curves (see Fig. 7.14) being constructed for the 18-8 steel using the testing method of ball indenter indentation and the standard tensile method (Markovets, 1979). On the primary creep curves obtained in ball indentation tests (10 mm diameter) there are observed the same regions as on the ‘classical’ creep curves obtained in tensile tests, i.e. the primary region of unsteady creep (see region I in the figure, i.e. primary creep with the decreasing creep rate), the next following straight region of steady-state creep (see region II in the figure, i.e. secondary creep with the minimum creep rate which remains approximately constant) and the failure region (see region III in the figure, i.e. accelerating creep with the

D

H = 4P2 pd

d

Boundary elasticplastic zone

sz

sJ

sx sJ

Lattice diffusion

sx sz Z

Slip plane

Dislocations move mainly by climb

Boundary diffusion Dislocations nodes

Dislocations move mainly by glide

7.13 Schematic illustration of the various mechanisms contributing to indentation creep.

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d, %

d, %

1.0

1.0 III

0.8

III

1

0.8

0.6

0.6 II

0.4 0.2 0.0 0

1

II 2

I

0.4 2

0.2 I 3 200 400 600 800 1000 t, h (a)

0.0 0

3 200 400 600 800 1000 t, h (b)

7.14 Primary creep curves for the 18-8 austenitic steel at 700 °C obtained by the ball indentation method (a) and the tensile method (b) at different values of the applied stresses: 1 – s = 100 MPa, H = 193 MPa; 2 – s = 60 MPa, H = 60 MPa; 3 – s = 50 MPa, H = 50 MPa.

increasing creep rate leading to fracture). All three regions described above are observed only at high loads P applied to the ball indenter in the case of indentation tests. The rate of uniform creep is calculated by the formula:

v=

Dd Dt

7.8

where Dd is the increment of deformation; and Dt is the increment of time. By the results of calculation of the rate of uniform creep on the second straight region there were built graphs of v as a function of a contact stress H in the projected surface area of the indentation and a tensile stress s in the sample at different H and s values. The graphs were made for the 18-8 steel used for manufacturing of the BN-600 secondary circuit piping having liquid sodium coolant as a working medium and operating at a temperature of about 550 °C. The graphs are presented in Fig. 7.15 in a logarithmic scale both for v and H, s. The obtained trends are well approximated by equation [7.9] with tensile tests and by equation [7.10] with indentation tests:

v = A · sn



v = A1 · s n1

7.9 7.10

where A, A1, n, n1 are experimental coefficients. Solution of both equations [7.9] and [7.10] under consideration of equal creep rates gives the following equation:

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H, s, MPa 600 400 300 200

550°C 1

5

2

650

4

3

6

100 80 60

700°C

7 8

40 30 20 –5 10

600°C

2

3 4

6 8 10–4

2

3 4

Vn %/h

7.15 Logarithmic diagrams of creep for 18-8 austenitic steel obtained by the ball indentation method (solid lines) and the tensile method (dashed lines) at different temperature values: 550 °C (lines 1, 2); 600 °C (lines 3, 4); 650 °C (lines 5, 6), 700 °C (lines 7, 8).



ÊA ˆ s = Á 1˜ Ë A¯

1/n

· H n1 /n



7.11

Processing of the experimental data for the different temperatures 550, 600, 650 and 700 °C gives the final equation:

s = k · H

7.12

In order to specify the coefficient of proportionality k in the final equation, the task of calculation of the stress-distribution picture of a ball indenter contact to the spherical contact surface area of the indentation was solved. The task was resolved into a plasto-elastic statement taking into consideration thermal deformation and material properties changing due to temperature increase, as well as due to joint action of metal strengthening and softening. Special monitors were designed and manufactured (see Fig. 7.16) for the task of in-situ application of the described method, and they were rigidly mounted on a pipe to create a constant load on the indenter ball (10 mm diameter). The monitor is mounted through studs which are in turn welded to the material under test by contact welding, since this technology does not damage the piping metal. The monitors are to be installed in zones of maximum (usually these are pipe bends) and minimum (straight sections) loading. Comparison of monitors’ records in maximum and minimum loaded zones allows detection of the process of creep, as well as its rate, if creep exists. The in-situ creep monitoring records are used as the initial data intended for substantiation of the service life prolongation of the investigated piping if one is dealing with damage due to the creep mechanism.

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Loading screw High-temperature force sensor Casing Force spring

Indenter Monitor base Mounting stud

Inspected pipe

7.16 General view of the monitor for in-situ inspection of creep damage.

7.7

Conclusions and future trends

Until now, there has been accumulated a rich theoretical, experimental and practical scope of knowledge providing for a confident calculation of strength and durability of NPP equipment. It is clear that any theoretical models need mechanical characteristics obtained experimentally and which are indispensable for carrying out calculations. However, calculations used at the design stage have to be separated from those calculations related to follow up NPPs in operation, namely, to consider current and extended operational reliability of components, due to changes in properties through ‘ageing’. At the design stage, there is a possibility to determine the necessary mechanical properties on full-scale standard samples. At the operational stage of a NPP the removal of samples is practically impossible. Thus the task of material ageing (degradation) rate determination by sample-free methods appears paramount and urgent when dealing with the task of NPP equipment service life prolongation. Various methods are known, but the most promising appears to be the method of component surface kinetic indentation. Research in this field will help to eliminate the problems pertaining to the determination of mechanical properties of service-exposed material and thus to make a great

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contribution to solving the lifetime extension problem by providing reliable data on the actual metal condition.

7.8

Acknowledgements

Acknowledgements are due to Dr V. M. Filatov for supplying valuable data on fatigue damage of NPP equipment steels, Professor E. M. Morozov for remarks given throughout the whole chapter, as well as the specialists of the Center of Material Science and Lifetime Management Ltd – V. I. Levchuk, A. A. Romanova, S. V. Chubarov – for help in preparation and printing of materials.

7.9

References and further reading

Alkazraji D (2008), Pipeline Engineering (Quick Guide), Cambridge, Woodhead Publishing. Anon. (1983), Condensed Metric Practice Guide for Corrosion, Annual Book of ASTM Standards, V.03.02, American Society for Testing and Materials, Philadelphia, PA. Anon. (1986), ‘Norms for boilers and pressure vessels. Evaluation of defects of pipelines made of austenitic steel’, Theoretical Bases for Engineering Calculations, 3, 146–171. Anon. (1990), Adhesives and Sealants, Engineered Materials Handbook, Materials Park, OH, ASM International. Anon. (1995), ASME BPVC, Sec. III, Div. I – Appendices, New York, ASME. Anon. (1999), Proceedings of Joint European Commission – IAEA specialists meeting ‘NDT methods for monitoring degradation’, Petten. Anon. (2000), Standard Terminology Relating to Fatigue and Fracture, Testing, ASTM, West Conshohocken, PA. Anon. (2002), Proceedings of the international symposium ‘Mechanisms of material degradation and non-destructive evaluation in light water reactors’, Osaka. Anon. (2005), Secondary Piping Rupture Accident at Mihama Power Station, Unit 3, of the Kansai Electric Power Co., Inc. (Final Report), Tokyo, The Nuclear and Industrial Safety Agency. Asano M, Hattori S and Suzuki I (1996), ‘Effect of long-term thermal aging on the material properties of austenitic stainless steel welded joints’, Proceedings of the 4th International Conference on Nuclear Engineering, ASME, 5, 183–188. ASME Boiler and Pressure Vessel Code, (1980), Section III, Case No. 47, New York, American Society of Mechanical Engineers. ASME Boiler and Pressure Vessel Code, (1995), Section XI, Rules for In-service Inspection of Nuclear Power Plant Components. IWB-3000, Appendix A, C, H. New York, American Society of Mechanical Engineers. Azbukin V (1983), Corrosion Resistant Steels and Alloys for NPP Equipment and Piping, Kiev, Nauka dumka. Azbukin V, Gorynin V and Pavlov V (1997), Perspective Corrosion Resistant Materials for NPP Equipment and Piping, St. Petersburg, CNII KM ‘Prometey’. Bakirov M, Potapov V and Massoud J (2002), ‘A phenomenological method of mechanical properties definition of reactor pressure vessels (RPV) steels VVER according to the

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ball indentation diagram’, Proceedings of the 5th international symposium ‘Contribution of materials investigation to the resolution of problems encountered in pressurized water reactors’, Fontevraud, 2, 777–779. Bakirov M, Kudryavcev E, Sarichev G and Tutnov I (2003), Diagnostics of Corrosive Damages of Metal Structures and Pipelines of Nuclear Energy Objects, Moscow, Radecon. Bakirov M, Potapov V, Iljin V and Nemitov B (2006), ‘Development and introduction of the technology of physical-mechanical properties inspection of nuclear reactor pressure vessels during the whole lifetime period’, Proceedings of the 6th international symposium ‘Contribution of materials investigation to improve the safety and performance of LWRs’, Fontevraud, 2, 781–791. Beskorovainy N, Kalin B, Platonov P and Chernov I (1995), Structural Materials of Nuclear Reactors, Moscow, Energoatomizdat. Bieniussa K and Reck H (1997), ‘Evaluation of piping damage in german nuclear power plants’, NED, 171, 15–32. Bogachev I, Korobeinikov V, Litvinov V and Poleva V (1974), ‘Aspects of surface fracture of metals during cavitation’, Materials Science, 2, 130–133. Bogoyavlensky V (1984), Corrosion of Steels at Water Cooled NPPs, Moscow, Energoizdat. British Standard Specification for Unfired Fusion Welded Pressure Vessels, 2000, London, BSI. Chigal V (1969), Intergranular Corrosion of Stainless Steels, Leningrad, Hymia. Coffin L, Jr (1976), Fatigue and Elevated Temperature, Philadelphia, PA, American Society for Testing and Materials STP 520. Collins J (1984), Material Damage in Structures. Analysis, Prediction, Prevention, Moscow, Mir. Danko J and Stahlkorf K (1984), ‘A review of the boiling reactor owners group research program on pipe cracking’, Proc. Amer. Power Conf., Chicago, 46, 654–658. Erve M, Wesseling U, Kilian R, Hardt R, Brümmer G, Maier V and Ilg U (1997), ‘Cracking in stabilized austenitic stainless steel piping of German boiling water reactors – characteristic features and root cause’, NED, 171, 113–123. Evans R and Wilshire B (1999), Introduction to Creep, Poole, Oakdale Printing Company Ltd. Filatov V (2004), ‘Prediction of austenitic stainless steel mechanical characteristics at irradiation’, Proceedings of the eighth International Conference on Material Issues in Design, Manufacturing and Operation of Nuclear Power Plants Equipment, St. Petersburg-Sosnovy Bor, 2, 189–202. Filatov V (2007), ‘Neutron exposure fatigue of structural materials’, Voprosi Materialovedenia, 3 (51), 253–264. GOST 9.908-85 Metals and alloys. Methods for determination of corrosion and corrosion resistance parameters, Moscow, Energoizdat. Hanninen H and Torronen K (1987), ‘Environment sensitive cracking in pressure boundary materials of light water reactors’, Int. J. Pres. Ves and Piping, 30, 253–291. IAEA-TECDOC-1260, Material Degradation and Related Managerial Issues of Nuclear Power Plants, IAEA, Vienna, 2006. Jungclaus D, Michel A and Schulz H (1999), ‘Operating experience with pressurized components of German light-water reactors’, NED, 192, 331–335. Karpenko G (1985), Working Ability of Construction Materials in Aggressive Environment, Kiev, Nauka dumka.

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Karzov G (1998), ‘Nature of the damage of KMPC downcomer pipelines on RBMK reactors in operation and ways of its overcoming’, Proceedings of 5th international conference ‘Material Issues in Design, Manufacturing and operation of NPP Equipment’, St. Petersburg, 2, 38–52. Kearns W (1978), Welding Handbook, Volume 2 – Welding Processes, Miami, FL, American Welding Society. Kochik I, Postler M and Zamboch M (2005), ‘Effect of netron irradiation on microstructure and mechanical properties of VVER-type reactor vessel internals’, 5th International Symposium on ‘Contribution of Metals Investigation to Improve the Safety and Performance of LWR’, Fontevraud 5, 1, 362–370. Kritsky V, Berezina V and Styazhkin P (2000), ‘Effect of coolant quality on operational reliability of equipment elements of RBMK-1000 NPPs’, Teploenergetika, 7, 2–9. Li W, Henshall J and Hooper K (1991), ‘The mechanisms of indentation creep’, Acta Metal, Mater, 12, 3099–3110. Livshits L (1979), Physical Metallurgy for Welders, Moscow, Mashinostroenye. Mahutov N (1985), Vibrations and Longevity of Marine Power Equipment, Leningrad, Sudostroenie. Mahutov N and Permyakov V (2005), The Lifetime of Safe Operation of Vessels and Pipelines, Novosibirsk, Nauka. Mamaeva E and Bakirov M (2003), ‘Assessment of the microstructure and properties degradation for WWER-440 austenitic piping welded joints after modeling of long operating life’, ESReDA 26th Seminar on Lifetime Management of Industrial Systems. Mamaeva E, Bakirov M, Chuvaev S and Fedorova O (2006), ‘Microstructure and mechanical properties of welded joints of NPP pipelines after long-term operation’, Metallovedenie i Termicheskaya Obrabotka Metallov, 7, 36–42. Markovets M (1979), Definition of the Metal Mechanical Properties by the Hardness, Moscow, Mashinostroenye. Nikitin V (1980), ‘Determination of metals resistance against corrosion cracking’, Physical-Chemical Mechanics of Materials, 1, 13–23. O’Donnell W J, O’Donnell W John and O’Donnell T P (2005), Proposed New Fatigue Design Curves for Austenitic Stainless Steels, Alloy 600 and 800, PVP 2005-71409, 2005 ASME PVP Division Conference, Denver, CO, p. 23. OST 108.031.08-85 Stationary boiler and pipelines of steam and hot water. Strength calculation norms. General provisions on substantiation of wall thickness, St Petersburg, NPO CKTI. OST 108.031.10-85 Stationary boiler and pipelines of steam and hot water. Strength calculation norms. Determination of strength coefficients, St Petersburg, NPO CKTI. PNAE G-7-002-86 Strength Calculation Norms for Nuclear Power Plant Equipment and Piping, Moscow, Energoatomizdat, 1989. Pogodin V and Bogoyavlensky V (1970), Intergranular Corrosion and Corrosion Cracking of Stainless Steels in Water Mediums, Moscow, Atomizdat. Preece C (1979), Erosion, Murray Hill, NJ, Bell Laboratories, Inc. RD EO 1.1.2.09.6714-2007 Procedure on strength calculation of main elements of a reactor installation on fast neutrons with sodium coolant. Saji G (2009), ‘Degradation of aged plants by corrosion: ‘Long cell action’ in unresolved corrosion issues’, NED, 239, 1591–1613. Scott M (1987), ‘Environment-assisted cracking in austenitic components’, Int. J. Pres. Ves and Piping, 65, 255–264. © Woodhead Publishing Limited, 2010

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Scott M (2000), ‘Stress corrosion cracking in pressurized water reactors – interpretation, modeling and remedies’, Corrosion, 56, 771–782. Scully J (1990), The Fundamentals of Corrosion, 3rd edn, Oxford, Pergamon Press. Shah V N, Ware A G and Porter A M (1997), Technical Report NUREG/CR-6456 ‘Review of industry efforts to manage pressurized water reactor feedwater nozzle, piping, and feedring cracking and wall thinning’, Washington, DC, Nuclear Regulatory Commission. Stephens R I, Fuchs H O, Fatemi A and Stephens R R (2001), Metal Fatigue in Engineering, New York, John Wiley & Sons. Terentev V (2002), Fatigue of Metallic Materials, Moscow, Nauka. Uhlig H and Revie R (1985), Corrosion and Corrosion Control, New York, John Wiley & Sons. Uhlig H and Revie R (1989), Corrosion and the Fight Against It, Leningrad, Hymia.

© Woodhead Publishing Limited, 2010

8

Microstructure evolution of irradiated structural materials in nuclear power plants

M. H e r n á n d e z - M a y o r a l, CIEMAT, Spain and M. J. C a t u r l a, University of Alicante, Spain

Abstract: This chapter describes the microstructure evolution in materials irradiated with fast neutrons in the pressure vessel and its internal components of nuclear power plants. In reactor pressure vessel steels, precipitates and matrix features, that include a large range of defect types and compositions, are observed. In austenitic steels, commonly used for internal components, Frank loops and ‘black dots’ are observed while at higher temperatures, voids, bubbles and precipitates appear. Measures to restore the mechanical properties of irradiated materials are enumerated. Finally, the latest research methodologies applied to understanding microstructure evolution of irradiated materials are described. Key words: austenitic stainless steel, ferritic/martensitic materials, neutron irradiation, transmission electron microscopy (TEM), dislocation structure.

8.1

Introduction

One of the key components in a nuclear power plant (NPP) is the reactor pressure vessel (RPV), since it contains the active fuel core and is a Safety Class I primary boundary pressure-retaining component. Consequently, special attention must be taken in the selection and quality of the structural materials used for its fabrication (Tipping, 1996). Likewise, the integrity of those materials used in the reactor internals must be preserved during the life of the NPP. The structural materials employed in current power plants, both in the pressure vessel and in the core internals, are conventional materials, low alloy steels in the case of the pressure vessel (clad internally with stainless steel) and austenitic stainless steels and nickel-base alloys in the case of internal structures. These materials undergo operational conditions that are not particularly aggressive, regarding water chemistry, temperature or stress level. However, exposure to high energy neutrons coming from the fission of fuel, together with the overall nuclear reactor core environment (e.g. gamma radiation, neutrons, radiolysis of water), affects significantly the physical and mechanical properties of these materials, resulting in ageing degradation effects during the life of the NPP that could result in early failure of these structural components if not detected or mitigated. Such events or 189 © Woodhead Publishing Limited, 2010

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necessary replacements are potentially expensive, impacting not only safety but also economic viability of NPPs. The RPV is a component that receives a relatively low neutron flux (neutron dose rate), with a total neutron fluence at the end-of-life of the reactor (approximately 40 full power years) from 5 ¥ 1022 m–2 up to 1.6 ¥ 1024 m–2, equivalent to 0.0075 to 0.24 displacements per atom (dpa). It is, nevertheless, an irreplaceable component (plant life-limiting) and a safety barrier in the overall defence-in-depth (DID) concept, layout and principle of a NPP (IAEA NP-T-3.11, 2009). The mechanical property changes in NPP systems, structures and components (SSCs) during the operational life of the NPP are clearly of considerable importance in the safe operation of the reactor, and play a major role in impacting the fitness-for-service and thus total achievable operational lifetime. The RPVs of pressurized water reactors (PWRs) and boiling water reactors (BWRs), for example, are made from ferritic low alloy steel (e.g. A533B type), forged and heat treated to a ferritic-bainitic microstructure and possessing an initially very high fracture toughness. The neutron irradiation experienced by the RPVs leads to an increase in the tensile yield strength or hardening and a concomitant reduction in ductility. This is particularly relevant in the RPV’s beltline region, opposite the core, where the neutron flux, hence fluence with time, is highest. When impact toughness is measured through a Charpy V test, a shift of the impact energy curve to higher temperatures is observed for irradiated materials, as well as a reduction in the upper shelf energy (USE), as shown schematically in Fig. 8.1. The transition temperature, often taken as the intersection between the Charpy curve and the 41 J line, although other definitions are also used (see Was, 2007, for a detailed description), characterizes the change from ductile to brittle behaviour. Due to neutron irradiation, this transition occurs at higher temperatures, and this effect is known as neutron embrittlement.

Charpy V energy

Unirradiated DUSE

DT41

Irradiated

Temperature

8.1 Schematic representation of Charpy V notch impact toughness as a function of temperature for unirradiated and irradiated ferritic steels.

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The ductile-to-brittle transition temperature (DBTT) can also be determined by other fracture-mechanical-based assessments such as the Master Curve methodology (IAEA-TRS-429, 2005; Wallin, 1984, 1985; ASTM, 2005). Other factors to take into account concerning the overall ageing degradation processes of reactor materials are the in-service temperature, which in the case of PWRs and BWRs is around 300 °C, operational stressors and the reactor specific water environment, which could contain impurities such as sulphate, chloride and peroxide that have potential to cause corrosion (Tipping, 1996). This is particularly crucial for the case of RPV internal components and their materials, in contact with the water coolant. Moreover, these components are subject to different levels of tensile stresses (internal or applied) and are exposed to much higher neutron fluxes (fluences with time) than the RPV, i.e. up to 5 ¥ 1026 m–2 (80 dpa) at the end-of-life of the reactor (COSU CT94-074, 1997). The main ageing degradation mechanism under irradiation associated with these internal components (usually made from various grades of austenitic stainless steel) is known as irradiation assisted stress corrosion cracking (IASCC) and is described in detail in Chapter 9. Austenitic stainless steel IASCC is of concern in the environment of oxygenated water of BWRs and to a lesser extent in PWRs since the critical neutron fluence for intergranular cracking is higher in non-oxidizing environments (Bruemmer et al., 1997; IAEA-TECDOC-1557, 2007; Scott, 1994). One of the factors that contributes to IASCC is the segregation of impurities or alloying elements due to the presence of defects produced by neutron irradiation. This radiation induced segregation (RIS) results in enrichment or depletion of certain elements at grain boundaries consequently changing their mechanical and corrosion-resistant properties. Chromium (Cr) depletion from grain boundaries due to RIS is considered one of the factors for the intergranular corrosion and intergranular stress corrosion cracking observed in BWR internal components (Scott, 1994). However, this is not the only change induced by neutron irradiation in internal components. Besides microchemical changes such as RIS, microstructural changes can cause radiation hardening and embrittlement to occur, just like in the case of ferritic RPV steels. Radiation hardening associated with the localization of plastic deformation seems to be another parameter that controls the loss of ductility leading to intergranular cracking of these components. Detection of degradation in these internal structures (e.g. core barrels, jet pumps, fixtures, brackets) is very difficult due to ease-of-access problems or other design features and so, accordingly, repairs and replacements can be potentially difficult to perform. The embrittlement of the RPV, as well as intergranular cracking and stress corrosion cracking-related degradation in stainless steels of the internal structures are important technological and economic problems that affect the operation, reliability and overall operational life of NPPs worldwide.

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Changes in mechanical and physical properties of materials such as neutron embrittlement or IASCC are a consequence of changes in the microstructure during irradiation, since ultimately it is the microstructure which determines materials behaviour. It is, therefore, crucial to understand how neutron irradiation affects the microstructure. This is, however, not an easy task due to the numerous parameters involved related to the material, such as alloy type and composition, prior heat treatment history, grain size, impurity content, precipitates or initial dislocation density, and also to the irradiation and environmental conditions: neutron energy, fluence, flux, helium/dpa ratio, temperature, contact with water coolant or stress levels, for example. When fast energetic neutrons (E > 1 MeV) bombard a material, changes at an atomic scale take place. The incident neutrons create point defects or clusters of defects and give rise to nuclear reactions that affect the crystalline structure, the microstructure and the chemical composition. The degree of change (e.g. number of displaced atoms) depends on the neutron energy spectrum and temperature of irradiation. The lower the temperature e.g. < 100 °C, the higher the level of remaining damage for an equivalent irradiation. These changes at an atomic level determine the mechanical and physical/ chemical properties of the components and their evolution during the life of the reactor. Consequently, identifying these processes is fundamental to understand the degradation phenomena that affect the components subject to neutron irradiation; with understanding, the appropriate mitigation methods may be applied. This is the main subject of this chapter. Microstructural characterization of irradiated materials is yielding new insights into the mechanisms that cause radiation damage in structural materials. Specialized techniques are required to study the damage produced by irradiation with sufficient resolution to reach the adequate size-scale typically expected for each phenomenon. Microstructural tools provide qualitative and quantitative information on irradiation-induced precipitation, matrix damage and grain boundary segregation and deformation modes. Transmission electron microscopy (TEM) is possibly the most widely used technique to study the microstructure of irradiated materials (Jenkins and Kirk, 2001). It provides information about the density of defects, size of these defects down to the nanometre level, and their character (dislocation loop, stacking fault tetrahedra, etc.). Other very important experimental techniques are atom probe tomography (APT), small-angle neutron scattering (SANS), positron annihilation spectroscopy (PAS), atomic force microscopy (AFM), secondary ion mass spectroscopy (SIMS) and auger electron spectroscopy (AES). The progress gained in the use of these experimental tools during the last few years is increasingly providing more detailed information about the microstructure of irradiated materials (Zinkle et al., 2009). (See Chapter 13 for details of these and other techniques.) Nevertheless, the combination of data

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from different techniques is necessary to obtain a complete characterization at all levels since no individual microstructural technique can provide a precise description. Therefore, an iterative process is required, using data obtained by different tools to characterize the microstructure that is consistent with all the microstructural information. An assessment of the mechanisms relevant for degradation will determine the selection of techniques to be employed. As mentioned above, there are a large number of parameters that influence the microstructural evolution in irradiated materials. Despite the numerous amounts of information gathered over the years from radiation experiments, many unknowns still remain regarding the role of the different parameters (temperature, fluence, impurity content of alloy, etc.). Modelling has always played a major role in understanding the different processes, and an important example is the work on void swelling from the 1970s and 1980s (see Trinkaus and Singh, 2003 for a review on this subject). During the last few years a new approach in modelling has emerged. With the development of first principle models aided by high performance computing, the objective now is to build computational tools capable of predicting materials performance from a fundamental understanding of physical processes (Phillips, 2001). These models could play a major role in forecasting the behaviour of structural materials under operation in NPPs. The development of such tools requires a solid experimental programme that can be used to validate every aspect of the models. These tailored experiments help in the selection of relevant mechanisms that should be introduced or considered in the models. However, neutron irradiation experiments are costly and difficult to perform. Other radiation sources such as protons, ions or electrons together with accurate models are becoming increasingly useful to obtain relevant information about radiation effects. These experiments allow for systematic studies of critical variables such as fluence, flux, irradiation temperature or material composition to obtain basic understanding of defect characteristics produced under irradiation. Fundamental experiments have been performed since the early 1960s (Meechan et al., 1960) and provide information such as defect mobilities, stability with temperature or defect type. However, the damage produced by different types of irradiation sources differs significantly. For example, irradiation with electrons or with light ions results in the production of mostly Frenkel pairs (pairs of vacancies and self-interstitial atoms, in short, interstitials), while irradiation with heavier ions, as well as neutrons, results in the formation of cascades, where not only point defects are created but also clusters of vacancies and interstitials. The distribution in space of the defects will also depend on the type of irradiation source and energy spectrum. Therefore, the results obtained with such alternative sources can only be extrapolated to ‘neutron-equivalent’ conditions with the help of a validated model. Finally, it is important to point out that a link must be established between

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the changes in the microstructure and the macroscopic changes in materials properties. This is by no means a trivial task and requires a strong collaboration between both experimental and theoretical groups and a significant research effort. In this section, the main studies of microstructure evolution of irradiated RPV ferritic steel and austenitic stainless steel internal components will be reviewed. Sections 8.2 and 8.3 describe the main structural materials and components that are prone to suffer degradation due to neutron irradiation in NPP environments, as well as the operational conditions. Section 8.4 explains those changes in the microstructure of RPV steels and internals materials under irradiation paying special attention to the effects of fluence and irradiation temperature, and their connection to the different degradation mechanisms. Some of the mitigation paths to restore the initial microstructure are described in Section 8.5. Finally, the most recent approaches to develop radiation resistant structural materials from a fundamental understanding of radiation effects are described in Section 8.6.

8.2

Structures and materials affected

The core of a NPP, constituted by the fuel elements, and its internal components is where the fuel fission and release of energy and neutrons takes place and, therefore, in addition to being the main source of heat, it is where the greater flux of neutrons occurs. The pressure vessel that surrounds the core is also exposed to significant levels of neutron radiation that could alter its mechanical properties. In this section the materials most commonly used for the RPV and internals are described.

8.2.1 RPV steels The RPV is one of the most important barriers between the fuel core of the reactor and the outside environment. Different concepts and designs of NPP exist worldwide that give rise to differences in dimensions and materials employed in the construction of RPVs. Western designs can be categorized into two main types of reactors, PWR and BWR, with different designs (power ratings) depending on the company constructing or running them. Regarding Eastern PWR designs (i.e. VVER, sometimes called WWER, depending on the strict use of the Russian abbreviation or adopted English thereof), they are classified according to the output power rating as VVER-440 (440 MeV) and VVER-1000 (1000 MeV). The RPVs are made of forged low alloy ferritic steels, with a suitable heat treatment to give a bainitic-ferritic structure of high initial toughness. A typical RPV is cylindrical with a hemispherical bottom head and a flanged and gasketed head, which is bolted to the main cylindrical body. In BWRs, the bottom has penetrations for control rods

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and instrumentation access; in PWRs, it is the RPV closure head which has penetrations for the control rods and instrumentations. Typical RPVs have the following dimensions: wall thickness depending on designs from 150 mm up to around 250 mm, 12 m height, from almost 3.5 m up to 5 m inner diameter and weight about 400 metric tonnes. A full and comparative description of different RPV designs can be found in IAEA NP-T-3.11 (2009). It is worth mentioning that BWR RPVs generally have a larger diameter and thus receive lower neutron flux, while PWRs, and in particular VVERs, have smaller RPV diameter, which results in a higher neutron flux on the region surrounding the core due to the smaller water gap. In all of them, the part of the vessel of primary concern with regard to age-related degradation is the core beltline typically located in the intermediate lower shells. Independent of the material employed for the RPV, it is common to distinguish among them depending on the procedure of fabrication: plate, forging and welds. The materials employed for its construction should be certified to be in accordance with the design and regulatory requirements for pressure retaining equipment of the highest safety class. Low alloy ferritic steels that show a high fracture resistance, combined with an acceptable ductility, are the materials selected for the pressure vessel. The body of the vessel is of low alloy carbon steel, commonly cladded with a minimum of about 3 to 10 mm of austenitic stainless steel to minimize corrosion at the inside surfaces in contact with the coolant. The most common RPV materials used in western NPPs are manganese-molybdenum-nickel (MnMoNi) steel, which is similar to the astm a533b1 specification, its forging equivalent astm a508-3 or a manganese-molybdenum-chromium (MnMoCr) forging steel a508-2, while in the VVER design, the selection for rpv construction was chromium-molybdenum-vanadium (CrMoV) steel under the specification 15Ch2mfa. Table 8.1 shows the ranges of compositions of those steels most commonly used to fabricate RPVs. The main difference is the addition of chromium (Cr) and vanadium (V) in those steels used in the VVER reactors, and the higher Ni level in VVER-1000. Since the 1970s, much data has been continually obtained concerning changes in materials properties caused by radiation damage. Now it is known that the susceptibility to embrittlement of RPV steels is strongly affected by copper (Cu), nickel (Ni) and phosphorous (P), and since then, therefore, chemical restrictions for RPV materials have been applied both in Western and Eastern-designed NPPs. The levels of these elements are restricted to below 0.1wt% for Cu and 0.020wt% for P. Also the Ni content is restricted in VVER-1000 to below 1.5wt% due to the enhanced radiation embrittlement in high Ni steels (Hawthorne et al., 2000). The initial microstructure of these materials is one of the parameters that will influence their behaviour under irradiation. Ferritic, tempered martensitic or bainitic microstructures are usually found in RPV steels,

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Steel

Cu

Ni

Mn

Mo

Cr

Si

P

S

C

V

Reference Phythian and English, 1993; IAEA report 2009 Carter et al., 2001; Miller et al., 2007; Fukuya et al., 2003 IAEA report 2009 Hawthorne et al., 2000, Carter et al., 2001; IAEA report 2009 Brillaud and Hedin,1992

IAEA report 2009

Kuleshova et al., 2002; Miller et al., 2000a,b;

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Base Type Not Not 1.15–1.50 045–0.60 Not 0.15–0.30 £0.035 £0.040 £0.25 – A302B specified specified specified earlier RPVs in BWR/PWR A533 GrB 0.06–0.16 0.40–0.70 1.15–1.50 0.40–0.60 0.08–0.11 0.15–0.40 0.008– 0.007– 0.12–0.23 – or cl 1 0.014 or 0.017 or (Plate) £0.035 £0.040 A508-3 0.02–0.04 0.50–1.00 1.2–1.5 0.49–0.70 0.09–0.45 0.15–0.40 0.008– Max 0.19–0.27 Max (Forging) Max 0.10 0.015 0.018 0.05 French 0.040– 0.64–0.84 1.23–1.56 0.39–0.57 0.13–0.26 0.16–0.33 0.005– 0.004– 0.138– – reactors 0.090 0.017 0.014 0.181 16MnD5 or 18MnD5 German £0.12 0.60–1.20 0.50–1.00 Max 0.60 0.25–0.50 0.15–0.35 Max Max 0.17–0.23 Max reactors 0.012 0.008 0.02 22NiMoCr 37 (Forging) 15Ch2MFA Max 0.30 Max 0.40 0.30– 0.60– 2.50–3.00 0.17–0.37 0.009– Max 0.13–0.18 0.25– VVER-440 0.60 0.80 0.0375 0.025 0.35 max

196

Table 8.1 Typical RPV steel compositions (wt%) used in Western and VVER (Russian) designs (from various sources)

Hawthorne et al., 2000 Kuleshova et al., 2002; Miller et al., 2000a,b; Hawthorne et al., 2000 Phythian and English, 1993; Hawthorne et al., 2000

Weld KS–01 0.37 1.23 1.64 0.70 0.47 0.18 0.017 0.012 0.06 – Linde 80 0.3 0.58 1.63 0.39 0.1 0.61 0.017 0.012 0.08 – flux weld Linde 1092 0.21 1.00 1.23 0.54 0.05 0.21 0.0014 0.009 0.13 – flux weld French 0.03–0.13 0.07–0.78 1.31–1.88 0.44–0.60 0.01–0.03 0.29–0.51 0.003– 0.005– 0.045– – reactors 0.019 0.020 0.078 SV-10ChMFT 0.15–0.21 0.09–0.29 0.97–1.03 0.43–0.50 1.37–1.58 0.15–0.35 0.018– 0.012– 0.05– 0.19– VVER-440 0.039 0.013 0.07 0.23 SV-10Ch 0.05–0.08 1.17–1.88 0.72–0.94 0.55–0.70 1.70–1.88 0.14– 0.010– 0.006– 0.05– 0.01– GNMAA 0.41 0.011 0.012 0.12 0.03 VVER-1000

Miller et al., 2003a,b Carter et al., 2001 Carter et al., 2001 Brillaud and Hedin, 1992 Kuleshova et al., 2002 Kuleshova et al., 2002

Microstructure evolution of irradiated structural materials

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15Ch2 Max 0.10 Max 0.40 0.30– 0.60– 2.50–3.00 0.17–0.37 max Max 0.13–0.18 0.25– MFA-A 0.60 0.80 0.012 0.015 0.35 VVER-440 15Ch2 Max 0.10 1.00–1.5 0.3–0.6 0.50–0.70 1.8–2.30 0.17–0.37 Max Max 0.13–0.18 0.10– NMFA-A 0.01 0.12 0.12 VVER 1000

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containing different types of precipitates with size ranging from 10 nm to micrometre. They include cementite (Fe3C), Mo2C, M23C6, M7C3, V(CN), etc., depending on the alloy type and heat treatment of the steel (Miller et al., 2007). For example, SA533-1 has a mixed ferrite tempered martensite and bainite with a grain size of 20–30 mm. It shows discrete aligned inclusions between 2 and 34 mm, carbides on both grain and lath boundaries and other large precipitates (approximately 100 nm) Mn rich, Mo rich or Mn/Si rich. Irradiation does not seem to have a detectable effect on these precipitates (Carter et al., 2001). For VVER-440 the basic matrix of the alloy is a mixture of tempered granular and lath bainite with a small amount of ferrite with large differences in dislocation density between base and weld alloys. In both base and weld, Cr-rich M7C3, M23C precipitates of approximately 200 nm located preferentially at grain boundaries can be found (Miller et al., 2000a; Ko�ik et al., 2000 and Kuleshova et al., 2002), as well as intergranular fine (<20 nm) vanadium carbides (VC).

8.2.2 Internal structures The internal structures of water-cooled reactors are essentially made of austenitic stainless steel and nickel-base alloys. These structures play an important role as they support the reactor core, they channel the water flow inside the vessel and they support and guide the instrumentation necessary for controlling and monitoring the reactor. One should note that there are design differences in the core of BWR and PWR reactors. In a BWR the internal structures include not only those structures that support the core, as in a PWR, but also those parts related to separating the water flow in the core of the reactor (core shroud), the feedwater spargers, the jet pump assemblies and the steam separator and dryer assemblies (Ware and Shah, 1992). Table 8.2 lists those stainless steels most commonly used in Western reactors and Russian-designed VVER reactors, while Table 8.3 lists the internal components fabricated with these materials. As shown in these tables, in both BWR and PWR internals the austenitic stainless steels most commonly employed are of the types 304 and 316. In German reactors Nb or Ti stabilized steels are employed as type 316Ti, with German nomenclature 1.4751. Those stabilized steels are also the ones used in VVERs. High nickel content alloys are rarely used in VVER reactors, while titanium stabilized stainless steel (SS) is used almost everywhere (COSU CT94-074, 1997). Internal components are of diverse shapes, dimensions and compositions, from the plates of baffle formers, tubes for control rods, cylindrical core shroud or bolts at the junctions of the baffle formers. A full description can be found, for example, in COSU CT94-074 (1997). As shown in Table 8.2, chromium contents higher than 15% are used in these materials to protect against intergranular cracking. The general resistance

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Steel

C

Mn

Si

304 316 X6CrNiNb 18–10 (1.4550) Germany CW 1.4571 Germany (type 316 Ti) 08Kh18N10T

0.04– £2.0 £1.0 £0.03 £0.045 8.0– 11.0 18–20 – – – £0.05 0.08 (304L <0.04) 0.03– £2.0 £1.0 £0.03 £0.045 10.0–14.0 16–18 2.25–3.0 – – £0.20 £1.0 0.08 £0.04 £2.0 £1.0 £0.02 £0.035 9.0–12.0 17–19 – £0.65 – £0.20 –

(Nuclear (304L grade £0.0018) £0.10)

£0.06

£2.0

£0.08

1.0–2.0

£1.0

S

P

Ni

Cr

Mo

£0.02

£0.035

10.5–13.5 16.5–18.5 2.0–2.5

£0.02

£0.035

9.0–11.0

Nb

–

17–19

Ti

£0.7

Co

£0.20

Cu

–

N

B

£0.08

–

£0.08

–

–

–

£0.6

Microstructure evolution of irradiated structural materials

Table 8.2 Typical austenitic stainless steels compositions (wt%) used for reactor vessel internal components from AMES report, IAEATECDOC-1557 (2007)

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Table 8.3 Materials used for internal structures and components in Western and VVER (Russian) designs (IAEA-TECDOC-1557, 2007) NPP type

General remarks

US-type reactors (ASTM, ASME)

304 316CW 347/348

French-type reactors (RCC-M, AFNOR)

Plates of 304L with controlled nitrogen content: core barrel, core baffles and formers 316CW for threaded structural fasteners and bolts

German-type reactors Niobium stabilized austenitic SS: (KTA 3104) ∑ X6CrNiNb 18-10 (1.4550) ∑ X6CrNiMoTi 17-12-2 (1.4571) similar to AISI type 316 but Titanium stabilized, employed for fasteners Inconel X-750: support pins of the control rod mechanisms, and in former German design for baffle bolts and fuel aligment UK-type reactors

304 SA: core support structures 316CW for threaded structural fasteners and bolts Similar to French design

VVER-440 and VVER-1000

Titanium stabilized austenitic stainless steels 08Ch18N10T (equivalent to A-321), ChN35VT (VD) ∑ Core barrel ∑ Core shroud ∑ Block of guides tubes Welding materials: Sv-04Ch19N11M3, EA-400/10T

of austenitic stainless steels to water corrosion and neutron irradiation embrittlement makes these materials the most suitable for this application. In order to stabilize the austenitic phase, carbon (C), nitrogen (N), Ni or Mn are added, which are austenite-forming elements. Precipitates that can appear at grain boundaries are thus chromium-carbides of the type M23C6 that can severely (and locally) deplete the grain boundary region with respect to Cr, therefore reducing the resistance to corrosion there. Consequently, the C content is kept low in these steels: specifications require less than 0.08% C and in low carbon steels (‘Nuclear Grade NG’) a maximum of 0.03% C. Other second phases can also be present in the initial microstructure of austenitic steels such as ferrite, martensite, nitrides, borides, sulphides, oxides and intermetallics. Martensite formation can be caused by the presence of certain elements such as C, N and Ni or by plastic deformation. Impurities such as N will increase the strength of the alloy and levels can be as high as 0.1% in some steels (Bailat, 1999; Bruemmer et al., 2001). Nickel is used for austenite stability with specifications between 8 and 14% depending on the alloy specification given in the 300-series. Mn and Si are also present at levels below 2% and 1% respectively. Mo, Nb and Ti form carbides at

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temperatures higher than can Cr, and thus have a ‘stabilizing’ effect in this respect, removing carbon before chromium has the chance to form chromium carbides. Mo is added in 316 for high temperature creep and pitting corrosion resistance (Bailat, 1999). There is also a large number of impurities present on these materials, phosphorous (P), sulphur (S), boron (B), cobalt (Co) and copper (Cu) among others. Boron, particularly, seems to have an important effect in IASCC possibly due to the generation of helium through transmutation reactions with neutrons (Bruemmer et al., 2001). Important differences exist between the 300-series stainless steels, regarding initial composition and microstructure. It is important to note that small differences in concentrations of alloying elements such as C, Ni or Cr can have significant consequences in terms of phase stability or stacking fault energy, which for austenitic steels is generally less than 100 mJ/m2. Differences can also be observed regarding grain size, dislocation densities or precipitates. Therefore, a detailed characterization of the initial microstructure of these materials is necessary to understand the subsequent evolution under irradiation. The microstructure will be very dependent on the fabrication conditions. Grain sizes can range between 12 and 120 microns with typical sizes between 30 and 80 mm (Bruemmer et al., 1996). The concentration of carbon in solid solution versus carbides will also depend strongly on fabrication conditions. Typically two types of austenitic steels are used: solution-annealed (SA) and cold worked (CW), which have very different dislocation densities with very low densities in SA (~ 1012 m–2), and very high densities in CW (~1016 m–2). As explained below, this will have significant consequences for the evolution of the microstructure during irradiation. Due to the low stacking fault energy of these materials, deformation twins can also be present in the initial microstructure, with different densities depending on the processing treatment.

8.3

Environmental and other stressors

As mentioned in the introduction, the significant environmental stressor that will affect the performance and properties of both the RPV and its internal components is the exposure to fast neutrons. However, other issues must also be taken into account, particularly for the case of internals such as the water coolant or localized stresses. One of the problems in understanding the relationship between the environmental conditions and the microstructure evolution is that, besides the fact that these are very complex structures as described in the section above, the information gathered over many years of research comes from different sources with different irradiation conditions. This is the case both for RPV and internal components. Table 8.4 shows, as an example, the irradiation parameters of different reactors used to study microstructure evolution.

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Table 8.4 Irradiation conditions and parameters of different neutron sources from facilities used to study microstructure evolution of structural materials (from COSU CT94-074, 1997) Reactor

Location

Temperature

Neutron flux, E > 1 MeV

BOR-60 BR2 OSIRIS HFR LWR-15 VVER-2

Dimitrovgrad, Russia Mol, Belgium Saclay, France Petten, the Netherlands Rez, Czech Republic Rheinsberg, Germany

330 °C 150 °C 250–1000 °C 250–400 °C 200–300 °C 255 °C

~ 1 ¥ 1019 n/m2s 3 ¥ 1018 n/m2s 2 ¥ 1018 n/m2s 2.5 ¥ 1018 n/m2s 2.3 ¥ 1017 n/m2s 0.15–5.4 ¥ 1016 n/m2s

Although this chapter will centre discussions on light water reactors (LWR), much of the information existing in the literature regarding microstructure evolution comes from fast breeder (FB) type nuclear power reactors, which are cooled by liquid sodium, for example. The main differences between LWR and FB reactors are in the operational temperature of the reactors, with lower temperature for LWR, and the neutron flux, with two orders of magnitude lower neutron flux in LWRs.The difference in neutron flux and energy spectra between FB and LWR could have a significant effect on defect evolution, particularly for the transient regime of swelling (Garner and Greenwood, 2003; Okita et al., 2002).

8.3.1 RPV steels Depending on the location and reactor design, the range of fluence received by RPVs is relatively wide, from as low as 5 ¥ 1022 m–2 (~0.0075 dpa) to as high as 1.6 ¥ 1024 m–2 (~0.24 dpa) (IAEA NP-T-3.11, 2009). In PWRs the neutron flux (>1 MeV) received by the pressure vessel is between 1.2 and 4 ¥ 1014 m–2 s–1 which results in a neutron fluence at the design end-of-life of about 4 ¥ 1023 m–2 (~0.06 dpa). The maximum design operating temperature in PWRs is around 340 °C, but they operate normally at 280–325 °C. Their RPV is designed for a pressure of 17 MPa and the operational pressure is typically ~15 MPa. In BWRs, both neutron fluence and flux are lower than in PWRs since the RPV has larger diameter in the former and, therefore, more water gap exists between the core and the wall of the vessel. Both the design temperature and pressure are also lower than in PWRs, 302 °C and 8 MPa respectively, and operate at temperatures between 280 and 288 °C and a pressure of ~7 MPa. The VVER reactors are of the PWR type, with design lifetime of about 30 or 40 years for VVER- 440 and 40 for VVER-1000, but their smaller diameter means that they receive higher neutron fluences relative to Western-design PWRs. The expected maximum fluence at the end-of-life is 1.6 ¥ 1024 m–2 (E > 0.5 MeV) for a neutron flux of 1.5 ¥ 1015 m–2  s–1 in 440 design, and 3.7 ¥ 1023 m–2 for a neutron flux of 4 ¥ 1014 m–2  s–1 (E >

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0.5 MeV) in VVER-1000 design. Their respective working temperatures are: 264/299 °C VVER- 440 and 289/322 °C in VVER 1000.

8.3.2 Internal structures The temperature of internal structures ranges between 300 and 380 °C, with temperatures higher in PWRs than in BWRs. The location of these components will determine the level of radiation received, from very low fluences, < 6 ¥ 1024 m–2 (< 1 dpa) up to 5 ¥ 1026 m–2 (80 dpa) at the endof-life of the component (COSU CT94-074, 1997). In the case of PWRs the lower core structures are the most critical components not only because of their role as supporting structures but also due to the high levels of neutron radiation they are exposed to. These lower core structures are made mostly of 304L SS and CW316 SS, as explained in the previous section and they are: the core barrel with doses between 6 ¥ 1024 m–2 and 6 ¥ 1025 m–2 (1 and 10 dpa), the baffle plates, between 2 ¥ 1026 m–2 and 5 ¥ 1026 m–2 (36 and 82 dpa) and the formers, around 3 ¥ 1026 m–2 (48 dpa). Only a few components are made of other nickel-based alloys such as Inconel 600 and X750. However, these components do not receive high dose levels and their degradation problems are not attributed to the effect of irradiation (Scott, 1994; IAEA-TECDOC-1557, 2007; COSU CT94-074, 1997). Note that due to the different sizes, the levels of radiation in BWR internal components are much lower than in PWRs, ~3 dpa and up to 100 dpa for the lifetime of the reactor, respectively (Garner and Greenwood, 2003). Another important issue to take into account in the case of austenitic stainless steels is the production of helium (He) during neutron irradiation. It is well known that He is one of the main factors responsible for swelling of irradiated materials (Mansur and Coghlan, 1983) and it can also contribute to embrittlement due to helium bubble formation at grain boundaries (Ullmaier, 1984). Helium can be produced by the transmutation reaction of boron (B) to form He and Li. The other source of He production is through (n, a) reactions with metallic elements, particularly with Ni. Those He atoms produced by neutron reactions can interact with the vacancies also created by the irradiation. Helium binds very strongly with vacancies, increasing the stability of vacancy clusters in three-dimensional structures (voids or bubbles) that, without the presence of He, would be unstable or would collapse to other microstructural features. The helium production rate expected for austenitic steels in PWR is of the order of 6 appm (atomic parts per million) per dpa (Foster et al., 1995). Notice that in LWR the rate of helium production is expected to be higher than in FB reactors (see Table 8.4). Void swelling, however, can also occur without the presence of He, since other gas atoms could be present, such as hydrogen, that also contributes to stabilize vacancy clusters and therefore to void swelling. The phenomenon of void swelling is treated in detail in Chapter 10. © Woodhead Publishing Limited, 2010

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Besides neutron irradiation, there are other environmental issues that must be taken into account in the case of internal components. These components are in contact with the primary water coolant, which will lead to degradation due to stress corrosion cracking (SCC) as discussed in the next section. Note that the water chemistry is different in the case of BWRs and PWRs. Moreover, each component is subject to different stress levels and conditions.

8.4

Changes in the microstructure and degradation mechanisms

The interaction of high-energy neutrons with a material will produce damage in terms of atom displacements from their lattice sites, creating a supersaturation of vacancies and interstitials. Along the path of the neutron through the lattice, several atoms will be displaced with high enough energy to produce the subsequent displacement of other atoms, as shown schematically in Fig. 8.2. The first atom in the matrix displaced by the neutron is known as the primary knock-on atom (PKA), and the group of displacements that it produces is called the collision cascade. Atoms will be displaced from their lattice sites as long as the energy transferred is higher than the threshold displacement energy (Eth) which for Fe, as an example, is of the order of 40 eV. Within a few picoseconds from the PKA initiation, the number of displaced atoms in the collision cascade will reach a maximum. This is shown in Fig. 8.3 as an illustration of this process. This figure shows the evolution of a collision cascade produced by a 30 keV Fe PKA in body-centred cubic (bcc) Fe as obtained from molecular dynamics simulations (see Chapter 15 for a description of modeling tools). Figures 8.3a and 8.3b show the atoms displaced in the collision cascade for the maximum number of displacements and the final number of defects respectively. The size (radius) of these collision cascades is on the order of a few nanometres. Figure 8.3c shows the total number of Frenkel pairs (pairs of vacancies and self-interstitials) produced in the collision cascade as a function of time. Initially, depending on temperature, a large number of the displaced atoms return to lattice Recoil cascades

n

8.2 Schematic representation of cascade damage.

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Microstructure evolution of irradiated structural materials

(a)

205

(b)

Number of Frenkel pairs

2000

1500

1000

500

0 0

2 ¥ 10–12

4 ¥ 10–12 Time (s) (c)

6 ¥ 10–12

8 ¥ 10–12

8.3 Collision cascade of a 30 keV PKA event in bcc Fe as obtained from molecular dynamics simulations. Dark dots show location of self-interstitial atoms while light dots are vacancy sites for (a) maximum number of displacements and (b) final number of displacements in a collision cascade. Part (c) shows the number of Frenkel pairs produced as a function of time.

sites, therefore much of the damage is recovered, but some defects remain in the lattice, with vacancies located in the centre of the cascade region and interstitials in the surroundings, reaching a constant number of defects within a few picoseconds. These remaining defects can then interact with the initial microstructure (grain boundaries, dislocations, impurities, solute atoms, etc.) giving rise to microchemical and microstructural changes. In

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complex steels, such as those described here, mixing between the different alloying elements will occur, and interstitials will be not only of the base element (iron) but also of alloying elements (for example Cr or Ni in the case of austenitic steels). The extremely short timescales for the production of these collision cascades together with their associated small volume, make the resolution of this process experimentally unaccessible. Nevertheless, it is possible to observe at least the largest of those defects produced by the cascade with, for example, in-situ TEM experiments in pure iron (Yao et al., 2008). Smaller defects could also be observed experimentally using SANS or PAS (for the case of vacancy type of defects). However, we must take into account that all experimental observations are done in timescales significantly larger than the timescale for a collision cascade, and therefore many intermediate processes could feasibly occur from the time of the collision cascade until the time of the observation, such as cluster growth through interaction between clusters. Notice in Fig. 8.3 how, after the collision cascade, not only point defects are created, but also defects in clusters. TEM resolution allows for the observation of defects sizes of ~1.5–2 nm and higher, which, for example, in Fe corresponds to a loop with ~120–200 defects. The size of those clusters created in the collision cascade depends on the energy of the recoil as well as the material characteristics (stacking fault energy, lattice structure, etc.). Figure 8.4 shows two examples of cascade damage produced in bcc Fe (Fig. 8.4a) and face-centred cubic (fcc) Cu (Fig. 8.4b)

(a)

(b)

8.4 Defects in Fe (a) and Cu (b) produced by 30 keV atom recoils after the collision cascade as obtained from molecular dynamics simulations. Dark dots show location of self-interstitial atoms while light dots are vacancy sites.

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from molecular dynamics simulations of 30 keV recoils. In this particular case, self-interstitial clusters are observed for both cases. However, in Fe, most of the vacancies are isolated and in Cu they form large vacancy clusters in the centre of the cascade region. Although, as shown above, after a few picoseconds the number of defects produced in the cascade reaches a steady value, on a longer timescale this number will change, since some of these defects will migrate and reach sinks such as dislocations, grain boundaries or the debris of other cascades. In metals, at any given temperature, interstitial atoms diffuse relatively faster than vacancies therefore, at low temperatures defect migration will be mostly of interstitials. They could then contribute to the growth of interstitial clusters produced in the cascade, or could annihilate at vacancies, either isolated or in clusters. Moreover, some interstitial clusters are so highly mobile that they diffuse quasi-athermally. At higher temperatures, vacancies will also migrate contributing to the growth of cavities or the annihilation of interstitials. The diffusion of these point defects is also responsible for the phenomenon of radiation-induced segregation, since impurities and solutes will rapidly diffuse, aided by these defects, as described below. The evolution of the microstructure of irradiated materials will then be very dependent on the initial damage created in the cascades (size of clusters, their mobility, number of point defects and their distribution). In the following we will describe the microstructures observed in neutron irradiated RPV and austenitic steels, to provide a description of those fundamental processes that could be responsible for the development of such microstructures.

8.4.1 RPV steels Neutron irradiation embrittlement is the main degradation mechanism for RPV ferritic steels and welds. Its understanding requires detailed characterization of those changes in the microstructure induced by the irradiation, as well as its dependence on material and irradiation parameters. Point defects produced by neutron irradiation are responsible for changes at the nanoscale in the alloy’s matrix. On one hand, nano-features are formed that lead to hardening (Odette and Wirth, 1997), but, on the other hand, non-hardening embrittlement can also occur due to changes of the local chemistry at grain boundaries (Phythian and English, 1993; Odette, 1994). These changes are due to the redistribution of solutes and impurities in the matrix from the enhanced diffusion of these elements in the presence of defects created by the irradiation. As a result, enhancement or depletion of some of these elements near surfaces or grain boundaries can occur, changing material properties such as strength or resistance to oxidation. This phenomenon is known as radiation induced segregation (RIS). (For a detailed description of this process see Was, 2007, p. 433.) But, most importantly for these steels,

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point defects can enhance the diffusion of matrix elements and induce the formation of precipitates. Due to RIS, solute concentrations can reach local concentrations that are well above the solubility limit (of the element or impurity under consideration) and therefore precipitate out from solid solution into a different phase that would otherwise not be stable under thermodynamic equilibrium (Was, 2007). All these microstructural and microchemical changes depend on material and irradiation parameters, such as composition, initial microstructure, flux, total fluence and irradiation temperature. The microstructural features observed in RPV steels after irradiation, and considered responsible for hardening and embrittlement, can be of very different nature and can be divided into well-formed precipitates and matrix features (Odette and Wirth, 1997). Table 8.5 shows a summary of the main defects observed in irradiated RPV steels and some of their characteristics. It is worth pointing out that the boundary between the different features is not clearly defined, and in some types an overlap exists. The classification presented here, in accordance with the literature, can sometimes appear arbitrary and it is possible that some of these nano-features are of the same type, but appearing as different in the microstructural studies due to differences in irradiation conditions or in alloy composition (Odette and Wirth, 1997; IAEA NP-T-3.11, 2009). Our current understanding of the nucleation of these nanometer-size features results from both physically-based models (Wirth and Odette, 1999) as well as complex microstructural and microchemical characterization methods (APT, SANS, TEM, PAS). Furthermore, most of the information comes from basic experiments dealing with model alloys and controlled irradiation conditions. In particular, in order to fully understand the role of individual alloying elements, model alloys with systematic variations in the alloying additions are studied (Meslin et al., 2010; Bergner et al., 2010; Lambrecht et al., 2008; Hernández-Mayoral et al., 2010). Matrix features include a large range of defect types and compositions. For example, clusters of point defects can be identified, mainly nanovoids, that is, vacancy type of defects, and also present are interstitial clusters in the form of interstitial loops. Other matrix features are point defect clusters together with solute elements that have not resulted in a well-formed precipitate. These are mostly vacancy-solute complexes. Among well-formed precipitates one can distinguish between copper-rich precipitates (CRP), manganese-nickelrich precipitates (MNP) and the so-called ‘late blooming phases’ (LBP). All these defects can act as pinning sites for dislocation glide, increasing the tensile yield strength and consequently contributing to matrix hardening that can be detected by hardness tests, in contrast to embrittlement due to segregation at grain boundaries (non-hardening embrittlement), which is detected by Charpy or fracture mechanical tests. For Western-designed RPV steels and welds, the most relevant impurity element is Cu, since it can come out of solid solution as matrix-coherent

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Table 8.5 Nano-sized microstructural features observed in irradiated ferritic RPV steels (from various sources) Radiation-induced nanofeatures

Remarks and description

Vacancies-CuMnNiSi: dilute solute atmospheres or enriched areas containing Mn, Si, Ni, Cu,etc., that can be associated with vacancies, or microvoids. Could be precursors of wellformed precipitates. Size < 2 nm, whatever the fluence or chemical composition and mainly contain Fe atoms. Number density: 1024 m–3.

209

Matrix Vacancy-solute-complexes features Auger et al., 1995; Auger et al., 1994; Odette and Wirth, 1997; Carter et al., 2001; Kasada et al., 2005

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Well-formed Copper-rich precipitates (CRP) Their nucleation and growth rates are accelerated by radiation enhanced diffusion. precipitates Beaven et al., 1986; Miller They do not appear in low Cu steels, even at high fluences, 1025 n/m2. et al., 2000a,b, Carter et al., Average size ranges: 0.5–1.5 nm, and number density ~1024 m–3. High neutron fluence, 2001; Asoka-Kumar et al., high Cu content and low flux are factors promoting coarser CRPs. 2002; Ulbricht et al., 2005, Composition of the CRPs: Bergner et al., 2008 ∑ Cu-rich core, sometimes Cu-Ni-rich core Cu>50%-MnSiNi, Cu level depends on Cu nominal content Fe content in the precipitates under discussion, due to the different criterion for definition of precipitate, boundary precipitate/matrix ∑ Enriched in Mn, Ni, Si and P, depending on the alloy content of these elements ∑ Mn-Ni-Si-P at the interface CRP-matrix and Mn, Ni, Si area extended or larger than the Cu-rich area. MnNi-rich precipitates (MNP)* Phases appearing in steels with Cu > 0.1%, but later (at higher fluences) than CRPs. Auger et al., 1995; Odette Cu acts as catalyser of MNPs, i.e., the presence of Cu is required for the clustering of and Lucas 1998; Miller et al., MnNiSi. 2003a,b; Meslin 2007, IAEA MnNiSi (>50%)-Cu. report, 2009 Size and number density similar to CRPs. Late Blooming Phases MnNiSi-rich phases have been observed in Cu-free steels. These MNPs are nearly pure (LBPs)* Glade et al., 2006, and they are slow to nucleate, i.e. they appear at higher fluences than CRPs, they have Miller et al., 2006 been named ‘late blooming phases’ (LBP). Precipitates, distinguished from MNPs by the definition that they do not require significant Cu to form. Found by APT in a VVER steel low Cu, high Ni at neutron flux 7 ¥ 1011 and high fluence 14.1 ¥ 1023 n/m2 y 11.5 ¥ 1023 n/m2 E>0.5MeV.

210

Table 8.5 Continued Remarks and description

Voids Interstitial dislocation loops Ko�ik et al., 2000, Kasada et al., 2005, Maussner et al., 1999, Gurovich et al., 1999 Other nanoclusters Odette and Wirth, 1997; Miller et al., 2003a,b

Cluster number density increases as fluence increases, but size and microstructure do not evolve with increasing fluence. They are classified as ‘unstable matrix features’ (UMF)’ and ‘stable matrix features’ (SMF)’. For steels with Cu > 0.05wt%: CuMnNiSi atoms are not randomly distributed after irradiation and also spatial correlation; as fluence increases correlations increase and finally MnCuNiSi clusters are observed but still not well-formed precipitates. Threshold dose from which these defects are detected seems to decrease as Cu content increases. Not reported for RPV steels, but they can appear in pure iron or low Cu FeCu model alloys. Only observed at high fluences. Size and number density increase with increasing fluence.

Phosphides are observed for high impurity P content alloy as well as carbonitrides: they are formed due to irradiation (Miller et al., 2003a,b) P clusters with P about 54 at% and number density 1 ¥ 10e23 m–3 (MnNiCuMoFe). Do not need Cu to be present.

* Note: the boundary between the different features is not well defined and sometimes an overlap exists.

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copper-rich precipitates assisted by the presence of vacancy-type point defects. The formation and volume fraction of Cu-enriched precipitates would depend, in the first place, on the amount of copper supersaturation present in the matrix (heat treatment aspects). Usually, two situations can be distinguished, namely RPV steels and welds with Cu content below and above 0.1 wt%.1 At high Cu concentrations (> 0.1 wt%) CRPs are the dominant defects (Odette and Lucas, 1998). The excess concentration of vacancies produced under irradiation enhances copper diffusion and accelerates its precipitation out of the matrix. One should note that since the solubility limit of Cu in iron is very low, 0.007% at 300 °C (Auger et al., 2000), there is a supersaturation of Cu in solution arising from the welding or processing heat treatment and subsequent relatively rapid cooling. Other solutes, Ni, Mn and Si, also experience enhanced diffusion by the excess point defects. With or without irradiation, diffusion of solutes, such as copper, takes place by thermal jumps into adjacent vacancies, but the number of these jumps in a given time interval is much higher during irradiation, depending on the amount/concentration of vacancies present at the prevailing temperature. Subsequent repetition of this vacancy-solute exchange process results in random diffusion of copper. When a diffusing Cu atom encounters another Cu atom (or cluster of Cu atoms), they bind with one another. Small Cu clusters can re-dissolve, but at a sufficient size, the Cu clusters form matrix-coherent precipitates. This precipitation results in the formation of Cu-rich precipitates (CRP) that continue to grow by radiation-enhanced diffusion up to the point when Cu is mostly depleted from the matrix. The structure of CRPs has been studied by different techniques, APT (Auger et al., 2000; Miller et al., 2007) and SANS (Solt et al., 1993; Phythian and English, 1993), among others. In general it is observed that CRPs have a copper rich core with levels higher than the rest of the elements that conform the precipitate, as shown in Figs 8.5(a) and (b) of APT measurements from Miller et al., (2007). Characterization with PAS on model alloys with high Cu content has shown the evolution of Cu precipitation with dose (Xu et al., 2006, 2008). Cu precipitated first, followed by the growth of microvoids near these precipitates, and then the evolution of precipitates and the shrinkage of microvoids started at the same time. This fact lead the authors to conclude that Cu atoms were located on the microvoid surfaces when microvoids grow near Cu precipitates, in agreement with other authors (Nagai et al., 2001). They proposed a threestage process in the early formation, i.e. low fluences, of CRPs: (i) the precipitates nucleate by vacancy migration; (ii) microvoids form and grow 1

It should be noted that the selection of this value is arbitrary, and some authors consider the frontier below 0.07 or even 0.05wt%Cu. We selected 0.1 because it is the maximum value required in specifications of RPV steels since 1970, when experimental data suggested that Cu was harmful for RPV steels performance.

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Cu

Ni



P

Mo

Mn

Cr

Si

C

5 nm

(a) 25

Cu Mn Ni Si

Concentration, at%

20

15

10

5

0

0

1

2 3 Distance, nm (b)

4

5

8.5 (a) Atom maps of individual elements in CRPs and (b) average radial concentration profiles from the centre of mass of the CRP (from Miller et al., 2007).

at these precipitates; (iii) the aggregation of Cu atoms is promoted at these microvoids. An additional fourth stage is identified at higher fluences where Cu precipitation is complete, but microvoids still form and grow (Xu et al., 2008). As mentioned above, concentrations of other elements besides Cu, such as Mn, Ni and Si, have also been detected at CRPs (Odette and Lucas, 1998). In some cases these precipitates contain more Mn, Ni and Si than Cu and so are referred to as MNPs (Mn-Ni-rich precipitates). These MNPs are favoured at lower base matrix Cu contents. The segregation of Mn, Ni and Si elements

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that initially are not in supersaturation in the matrix can be explained by a reduction in the interface energy (Odette and Lucas, 1998, 2001; Liu et al., 1997). Mn and Ni are found to be distributed over a more extended area than Cu and sometimes they are found at the interface between the matrix and the precipitate (Miller and Russell, 2007). The structures found seem to imply that the enrichment of Mn, Ni and Si is independent of the existence of Cu (Odette and Lucas, 1998). It is possible to explain the experimental observations considering that Cu aggregation occurs first and then Ni, Mn and Si segregate during further irradiation, or that simultaneous clustering of Cu, Ni, Mn and Si occurs followed by redistribution of the solutes in the clustered region. In low Cu steels the formation of MNPs at high fluences and low fluxes is a concern for LWRs since they could limit the lifetime of RPV steels (Odette and Lucas, 2001). These so-called ‘late blooming phases’ were first predicted by empirical models and later observed experimentally, and is an example of the advantage of developing physically based models. In addition to the important role of Mn and Ni on embrittlement, it is worth to point out that P is sometimes found in CRPs and that there is increasing evidence of the contribution of this element to irradiation embrittlement, which could be of significance to VVERs where the P level is higher. Matrix features (MF) can be defined as dislocation obstacles that produce hardening other than well-formed precipitates. They can be of particular relevance in low copper steels (<0.1% Cu) and can be separated into thermally unstable matrix defects (UMD) and those that are thermally stable (SMD) (Odette and Wirth, 1997). The former ones are considered to be small vacancysolute complexes formed directly in displacement cascades (Miller et al., 2000a; Auger et al., 2000). The second type of matrix features include stable nanovoid-solute complexes that could also be growing slowly. Recent evidence (Nagai et al., 2001) seems to indicate that these stable matrix features also include dilute solute atmospheres (DSA) that may contain vacancy-solute clusters, or nanovoid complexes at their cores, and Ni, Mn and other solutes. These would be consistent with the proposal of some authors that damaged regions associated with the displacement cascades act as nucleation sites in low Cu steels (<0.1%) (Auger et al., 2000). As mentioned above, different irradiation and materials parameters can have a significant effect on the evolution of these nano-features. Number density of nanoclusters increases as fluence increases (Auger et al., 1995) while size remains unchanged. Figure 8.6 shows the clusters produced by neutron irradiation as a function of dose for two IAEA reference RPV steels as measured by SANS. Neutron flux rate and energy spectrum appears to have a strong effect on size, composition and number density of nano-features, but not on volume fraction (Bergner et al., 2008). Under certain conditions, namely relatively high fluences, interstitial loops have also been observed in RPV steels (Ko�ik et al., 2000; Kuleshova et al.,

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0.6

0.4

Sans measurements JRQ

0.2 JFL 0.0 0

25 50 75 Fluence F, 1018 n/cm2 (E > 1 MeV)

100

8.6 Neutron irradiation-induced clusters in RPV steels measured with SANS (from Ulbricht et al., 2005).

100 nm

8.7 TEM image of dislocation loops observed in irradiated VVER steels (after Kuleshova et al., 2002).

2002). Figure 8.7 shows an example of a TEM image of dislocation loops in VVER steels from Kuleshova et al., (2002). Dedicated experiments in model alloys (Ebrahimi et al., 1988; Hoelzer and Ebrahimi, 1995; HernándezMayoral and Gómez-Briceño, 2010) have shown that interstitial dislocation loops are affected by the presence of solutes or alloying elements, Cu, Mn and Ni. However, for the typical conditions of RPV steels, interstitial loops are not considered to play a significant role in radiation hardening (Odette and Wirth, 1997), although they can be of concern for high fluences, i.e. longterm operation conditions, and low Cu steels (Kuleshova et al., 2002). In summary, the causes, character and consequences of matrix features as well as LBP in low Cu content steels (<0.1% Cu) are not as well understood

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as in the case of CRPs or MNPs, and are still the subject of study by the research community. The effect of the different variables involved is still under study, in particular the effect of some alloying elements such as Ni and Mn, and the effect of very low flux together with high fluence that are relevant conditions for LWRs in case of extension of life programmes. Finally, for some steels, non-hardening embrittlement can be caused by radiation-enhanced solute segregation of elements to grain boundaries. It is well known that phosphorus can segregate to grain boundaries in RPV steels (base metal, weld metal, and heat affected zone) during service at elevated temperatures. In fact, classical temper embrittlement is an example of this. The presence of P in the grain boundaries can then lower their cohesive energy, causing the material to fail through intergranular embrittlement. Thus this type of embrittlement, typically defined as irradiation-assisted temper embrittlement (McMahon et al., 1981) is manifested as an intergranular fracture rather than the usual transgranular cleavage fracture. In this case, the Charpy or other toughness indication of the transition temperature shift (TTS) can occur even if the yield strength or hardness does not increase appreciably; in principle, combinations of irradiation hardening and grain boundary embrittlement can interact synergistically to produce very large TTS (PWRMRP-19, 2000). This phenomenon can have an effect at fluences near end-of-life or in case of long-term operation. It is also of concern in eastern RPVs where annealing is employed to restore the damaged microstructure and where, depending on the annealing temperature, P segregation could occur (PWRMRP-19, 2000).

8.4.2 Austenitic steels for core internals The main degradation mechanisms in austenitic stainless steels are hardening, swelling, embrittlement and stress corrosion cracking (Maziasz, 1993). Again, all these phenomena are related to changes in the microstructure. Before irradiation, austenitic (i.e. face-centred cubic crystal structure – fcc) steels have dislocations with Burgers vector b = (a0/2)<110> in concentrations that range from very low (~ 1012 m–2 ) in solution-annealed materials to very high (~ 1016 m–2) in cold-worked materials (Maziasz, 1993). The effect of radiation in the microstructure of austenitic stainless steels can be divided into two regimes: below and above 300 °C. This is the temperature at which vacancy clusters dissociate emitting vacancies. As a consequence, the type of defects observed in these two temperature regimes are very different. Incidentally, the operational temperature of light water reactors lies at this temperature ~ 300 °C. Note that most of the data existing in the literature refers to high temperature conditions and a limited number of experiments exist for those conditions in LWRs. Reviews on austenitic steels microstructural evolution under neutron irradiation can be found in Maziasz and McHargue (1987),

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Maziasz (1993), Zinkle et al. (1993) and Garner (1993) as well as Rowcliffe et al. (1998) and Bruemmer et al. (2001). Low temperature regime (< 300 °C) In the low temperature regime, two types of defects are reported in the literature: Frank loops and ‘black dots’. Figure 8.8 shows an example of Frank loops and ‘black dots’ observed in 304 SS irradiated at 330 °C for a dose of 0.8 dpa (after Pokor et al., 2004). The loops observed at low temperature and low dose are Frank (faulted) loops with Burgers vector b = (a0/3) <111> (Maziasz, 1993). A Frank loop can be formed by the collapse of a platelet of vacancies or interstitial atoms, such as those formed during irradiation. In the case of austenitic stainless steels large Frank loops have been identified as being interstitial in character. Figure 8.9 shows the atomic structure of a Frank loop formed by interstitials (Fig. 8.9a). This defect includes a stacking fault, as can be seen in the figure (after Hull and Bacon, 1984). Consequently, these loops cannot glide and will not move under an applied stress or temperature, and are therefore considered as sessile. However, the stacking fault can be removed either by dislocation reactions or by growth. If the loops grow to large enough sizes they will unfault by nucleating Shockley partials inside themselves and thus will reduce the energy due to the presence of the stacking fault. In a fcc lattice a perfect edge dislocation consists of two extra (110) half planes. These two planes do not need to be adjacent to each other and each of them are dislocations with a Burgers vector smaller than (a0/2) <110>, called partial dislocations. The separation of the two partials results in a stacking fault; therefore, the boundary of a stacking fault

10 nm

10 nm (a)

(b)

8.8 Example of Frank loops (a) and ‘black dots’ (b) observed in 304 SS irradiated at 330 °C for a dose of 0.8 dpa (after Pokor et al., 2004).

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C B A C B A (a) C B A C B A (b)

8.9 Frank interstitial loop (a) and prismatic interstitial loop (b) (see Hull and Bacon, 1984).

inside a crystal is a partial dislocation. Those partial dislocations formed by slip are called Shockley partials (see Hull and Bacon, 1984). Temperature or stress will also drive or favour the unfaulting process. The unfaulting reaction will produce perfect loops with Burgers vector b = (a0/2) <110> which are now glissile, that is, they are able to move easily along the glide plane. Figure 8.9b shows the atomic structure of a perfect interstitial loop. The interaction between multiple glissile loops can result in the formation of a dislocation network (Maziasz, 1993). Frank loops are considered to be responsible for hardening in austenitic stainless steels. The second type of defects observed, the so-called ‘black dots’, are those defects that, due to their small size, are difficult to identify. There is some controversy about the nature and type of these defects. Some authors assume that they are simply small Frank loops that cannot be resolved as such due to their small dimensions. Therefore, since larger Frank loops (more than 10 nm) have been identified as being of the interstitial type (Zinkle et al., 1993) these ‘black dots’ are also supposed to be interstitial loops. Pokor et al. (2004) claim that, in fact, ‘black dots’ and Frank loops are one single defect, observed under different conditions. Alternatively, other authors argue that a vacancy component of the damage must exist, since there is an equal production of vacancies and interstitials in the irradiation, and that these must be the ‘black dots’ (Edwards et al., 2003). There is some limited evidence, coming from positron annihilation experiments, that points towards the existence of a population of very small, under TEM resolution, vacancy

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clusters (Fukushima and Shimomura, 1993). In pure fcc metals such as Cu, the predominant defect observed under irradiation is the stacking fault tetrahedra (SFT) (Singh and Zinkle, 1993) in agreement with the fact that SFTs are the most stable geometries for vacancy clusters for low stacking fault energy materials (Zinkle et al., 1987). An SFT is a tetrahedron of intrinsic stacking faults on {111} planes with stair-rod dislocations along the edges of the tetrahedron (Hull and Bacon, 1984), as shown in Fig. 8.10. This structure can be formed by the collapse of a triangular platelet of vacancies. Surprisingly, in commercial alloys the density of SFTs is very low, less than 1% (Zinkle et al., 1993; Maziasz, 1993) or have not been observed. Only in high purity alloys or pure ternary alloys are SFTs clearly observed (Horiki and Kiritani, 1994). Moreover, some authors (Edwards et al., 2003) have pointed out that small Frank loops or partially dissociated Frank loops can be mistaken for SFTs. These last ones could be vacancy type defects that have not completed the transition to SFTs, maybe due to the presence of solute atoms or impurities. In summary, the nature and character of these ‘black dots’ observed in austenitic steels at low dose and low temperature still needs to be resolved. The average size and density of Frank loops at this low temperature regime, as a function of dose, is shown in Fig. 8.11 as reported by Edwards et al. (2003) and Bailat (1999) for different 304 and 316 SS. The density increases rapidly with dose reaching a saturation level at ~ 1 dpa or even sooner, while loop size continues to increase, levelling off after 2 dpa. As the dose increases, Frank loops become more prominent and increase in size. At low dose, 0.15 dpa, the mean loop size is ~ 2 nm, increasing with dose up to approximately 8 nm (Bailat, 1999). Note that interstitials are highly mobile

8.10 Stacking fault tetrahedra.

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Loop density, m–3

1024

1023

1022

Edwards-304B Edwards-304E Bailat-304CP Bailat-304HP Edwards-316K Edwards-316P Bailat-316CP Bailat-316HP 0

5 10 Neutron irradiation dose, dpa (a)

15

Average loop size, nm

15

10

Edwards-304B Edwards-304E Bailat-304CP Bailat-304HP Edwards-316K Edwards-316P Bailat-316CP Bailat-316HP

5

0

0

5 10 Neutron irradiation dose, dpa (b)

15

8.11 Density of loops (a) and average size of loops (b) as a function of dose for different 304 and 316 SS as reported by Bailat (1999) and Edwards et al. (2003).

and can contribute to the growth of these loops. This growth will continue until a saturation is reached when there is cascade overlap and vacancies and interstitials annihilate at the same rate at these interstitial loops. As seen in Fig. 8.11, at this low temperature, there is not much difference in the microstructure with alloy composition (Maziasz, 1993). However,

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there is a significant difference for different pre-treatments of the sample, that is, cold-worked (CW) vs. solution-annealed (SA). In CW steels there is a significant recovery of the damage produced by irradiation. The high dislocation density of these materials delays the build-up of damage due to the preferential recombination of self-interstitials at dislocations. At 5–10 dpa larger Frank loops are not observed in 20–25% CW 316. The density of these loops increases with temperature and seems to reach a maximum at around 330 °C for irradiation at 7 dpa (Maziasz, 1993). Some impurities can also have an effect on the microstructure evolution. The presence of P increases the dislocation density and decreases their size since P can bind to self-interstitials, decreasing their mobility (Watanabe et al., 1988; Yoshida et al., 1992). A complex interaction between different impurities has been observed that cannot be resolved unless a systematic study is performed. An increase in matrix hardness (and tensile yield stress) is expected from the interaction of these defects and coherent precipitates produced by the irradiation with dislocations. But the interaction of moving dislocations can also result in the annihilation of the obstacles such that subsequent dislocations will move more easily along the same path. This dislocation channeling results in microscopic flow localization (Bruemmer et al., 1997). The presence of these defect-free channels was first observed at high doses and temperatures, but they are also present at low temperatures. This flow localization could also contribute to IASCC (Lucas, 1997), particularly for those conditions where RIS cannot explain this phenomenon. High temperature regime (> 300 °C) For temperatures around 300 °C and higher, the features observed in irradiated austenitic stainless steels are quite different from those at lower temperatures. Besides Frank loops, which are generally larger than at lower temperatures (between 20 and 200 nm), voids, bubbles and precipitates can be observed (Maziasz, 1993). Already at a temperature of 375 °C both Frank loops and cavities can be observed in 304 and 316 SS irradiated in fast breeder reactors, as shown in Fig. 8.12 from Pokor et al. (2004), after a dose of 10 dpa. As the temperature increases, the concentration of ‘black-dots’ decreases and at temperatures of ~350 °C no significant fraction of black-dots is observed (Bruemmer et al., 1997). Figure 8.13 from Zinkel et al., (1993) illustrates the expected behaviour of damage as a function of temperature. At high temperatures (above 300–350 °C) the microstructure is dominated by cavities, large Frank loops and network dislocations. An important difference that drives the microstructure evolution at these temperatures is the loss of thermal stability of vacancy clusters. These clusters emit vacancies that can migrate and will annihilate at different sinks. In particular, the decrease in concentration of ‘black-dots’ in this regime could

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Microstructure evolution of irradiated structural materials

50 nm

221

50 nm

(a)

(b)

8.12 Frank loops (a) and cavities (b) observed in 304 SS irradiated at 375 °C for a dose of 10 dpa (after Pokor et al., 2004).

Low

High

Very high

1

Density (relative units)

10–1

Network dislocations

Black dots

Faulted loops 10–2

Cavities

10–3

0

200

400 600 Temperature (°C)

800

8.13 Schematic representation of defect evolution as a function of temperature (from Zinkle et al., 1993).

be due to recombination between mobile vacancies and small self-interstitial clusters. Moreover, at these temperatures, larger Frank loops can unfault, thus becoming mobile and contributing to the dislocation network. On the

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other hand, in the presence of gases, such as He, that can stabilize cavities, the emission of these vacancies can result in the growth of bubbles and voids, that are ultimately responsible of void swelling (Trinkaus and Singh, 2003). Helium binds strongly with vacancies and preferentially occupies substitutional sites. Other vacancies or He atoms can then be trapped forming bubbles and voids. These voids can grow by the addition of vacancies at these high temperatures which could lead to a significant change in volume of the material and concommittant swelling. However, helium is not the only element responsible for void swelling. Hydrogen can also play a significant role in the stabilization of bubbles and voids. These phenomena are treated in detail in Chapter 10 and have been the topic of numerous studies, both experimentally and theoretically. Table 8.6 shows a summary of those microscopic features observed in irradiated austenitic stainless steels for both low and high temperature regimes. Impurities and solutes play a significant role in the evolution of the microstructure of austenitic stainless steels. On one hand, the presence or absence of SFTs seems to be very dependent on the impurity content. An explanation of this behaviour has been given by Zinkle et al. (1993). Moreover, RIS in these steels is considered as one of the mechanisms responsible for irradiation assisted stress corrosion cracking (IASCC), which is described in detail in Chapter 9. It can result in enrichment of trace or minority alloying or impurity elements at the grain boundary, such as Si or P, or depletion of Cr, which could influence the behaviour of the material under stress. Figure 8.14 shows the concentration of Cr, Ni, Si and P across a grain boundary for a stainless steel irradiated at 300 °C up to several dpa (after Bruemmer et al., 1999). In irradiated stainless steels there is always a depletion of Cr at grain boundaries, changing the resistance of these materials to corrosion. Table 8.6 Summary of microscopic features observed in neutron-irradiated austenitic stainless steels Defect type

Description

‘Black dots’

Only at low temperatures (below 300–350 °C) and low dose (< 1 dpa). Controversy about their nature: vacancy or interstitial type and their characteristics, small SFTs or loops.

Frank loops

Seen at all doses and temperatures in solution-annealed steels. Not observed in cold-worked steels at low temperature. Size increases with dose and temperature.

SFTs

Only in high purity alloys and at low concentrations

Bubbles

Only at high temperatures (> 300 °C)

Voids

Only at high temperatures (> 300 °C)

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22

Cr or Ni concentration, wt%

18

Cr

16

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14 12

Ni

2

Si

1

10 8

P

6 –20

–15

3

–10 –5 0 5 10 15 Distance from grain boundary, nm

Si or P concentration, wt%

5

20

0 20

8.14 Concentration of Cr, Ni, Si and P across a grain boundary of a 300 series stainless steel irradiated in a light water reactor to several dpa at 300 °C (after Bruemmer et al., 1999).

RIS increases with dose up to ~ 5 dpa when it seems to saturate. Other impurities present in engineering alloys such as Si or P may also undergo RIS. In the case of Si and P, both are enriched at the grain boundaries as shown in Fig. 8.14. The enrichment or depletion of different elements depends on their migration mechanisms and velocities. Cr and Ni diffuse with the aid of vacancies, produced by the irradiation with Cr diffusing faster than Ni in austenitic steels. As vacancies reach the grain boundaries, Cr is more likely to diffuse away from the interface. For the case of Si or P, it is considered that they migrate with the aid of self-interstitials, therefore as self-interstitials reach the boundaries so do these ‘undersized’ atoms. The depletion of Cr from grain boundaries has been studied in detail by several authors due to the consequences for corrosion resistance and IASCC (see review by Bruemmer et al., 1999). However, there are still some controversial results regarding the enrichment or depletion of certain alloying and impurity elements at grain boundaries under LWR conditions, particularly for high doses. Helium can also segregate to grain boundaries and consequently bubbles and voids could be produced at these interfaces. It is important to note that RIS will depend on the initial grain boundary composition and microstructure. For example, if boron is present at the grain boundaries, this could lead to the enhanced formation of He (Bajaj et al., 1995) and subsequently the generation of gas cavities with the detrimental effects that this could have on grain boundary cohesion. On the other hand, the high dislocation densities of cold-worked

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materials could reduce the amount of RIS at least at low doses, due to enhanced recombination at dislocations.

8.5

Mitigation paths

The research studies undertaken by the nuclear fission community over the years have allowed measures to be developed to mitigate some of the degradation mechanisms mentioned above. For instance in the case of RPVs, reduction in Cu and P levels or the use of forging instead of plate removing welds from the high flux positions are some of the paths taken to reduce the effect of irradiation. When the embrittlement has reached levels close to the design safety margin limits of the component, there is the option of thermal annealling. Increasing the temperature will thermally activate some processes such as defect mobility and the emission of defects from defect clusters, enhancing recombination, removing unstable matrix features and decreasing the total defect concentration, restoring, to varying degrees, the original mechanical properties of the material. The recovery of the mechanical properties will depend on the chemical composition of the alloy. The selection of annealing time and temperature is of vital importance for the efficiency of this mitigation mechanism in RPV structural materials (Eason et al., 1998). A general tendency is to use annealing temperatures of about 150 °C higher than the irradiation temperature. This will also cause copper-rich precipitates to grow and eventually lose their coherency with the matrix. This will reduce their efficiency to dislocation pinning. However, temperatures should not be high enough to encourage other mechanisms of degradation to occur such as ‘temper embrittlement’. Treatment at temperatures higher than 460 °C for at least 100 hours restores the mechanical properties with an acceptable value of residual embrittlement. Information extracted from surveillance capsules can be very helpful to understand the correct heat treatment parameters and conditions (Tipping, 1996). A better understanding of the processes occurring in the irradiated material during annealing at a fundamental level would also provide more accurate temperatures and times. However, annealing cannot be used for some components, such as many internal structures. Another mitigation option is to reduce the level of neutron flux on the RPV beltline wall, which will slow down the rate of embrittlement in that area, giving a longer operational time before toughness limits are reached (Tipping, 1996). The rate of fluence accumulation can be lowered through the use of dummy fuel elements (stainless steel) and implementation of a low neutron leakage core geometry (fuel elements in core with ‘insideout’ configuration). However, this can result in a slight loss of core power output. For the case of water reactor internal components, COSU CT94-074

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(1997) points to a series of measures that can be taken to prevent or mitigate irradiation damage in these components. On one hand, it is noted that 316 grade steels perform slightly better under irradiation than 304 steels since hardening occurs later in the former. The stabilization of the initial dislocation network could prevent dislocation loop formation and this could be achieved either by adding titanium in solid solution in the presence of carbon or by using a steel with a fine dispersion of precipitates. On the other hand, accelerating the formation of network dislocations from dislocation loops or enhancing the strength of grain boundaries would also improve their mechanical behaviour under irradiation. However, more research is needed to develop practical solutions to these problems.

8.6

Application of research and operational experience to the practical solution of problems

The development of advanced NPPs presents new challenges for materials. On one hand, these reactor concepts are usually designed to operate at higher temperatures. While in a commercial LWR temperatures no higher than 300 °C are reached, values as high as 1000 °C are expected in VHTR (very high temperature reactors). These materials will have to sustain higher irradiation doses and the plant lifetime expectancy should be at least 60 years, instead of the current 30–40 years for original design lifetimes. Current NPPs are also candidates for long-term operation (LTO) with relicensing or open-ended licence conditions with periodic safety review approaches. Moreover, the materials in new NPPs, particularly for internals, should be compatible with the new coolants being considered (Yvon and Carré, 2009). The current candidate materials for these applications are ferritic/martensitic steels (9–12%Cr), nickel-based alloys, oxide dispersion strengthened ferritic/ martensitic steels and ceramics (SiC, TiC, etc.). The behaviour of these different materials under irradiation, corrosion under representative environments, and high temperature coupled with operational stressors must be studied to determine their applicability and suitability. These are likely to be costly and long-term experiments that require multiple neutron irradiations as well as laboratories for testing radioactive material under a large number of conditions. Moreover, as mentioned above and pointed out by several authors (Bruemmer et al., 1999, Garner and Toloczho, 1997), most of the data existing in the literature has been obtained for conditions different from those in LWRs. Therefore, it is becoming increasingly important to use other radiation sources, such as proton or ion irradiation, to obtain fundamental information on materials behaviour under irradiation. The advantage of proton vs. neutron irradiation is time, cost and no sample radioactivation is caused when using protons (Was et al., 1999). Accordingly, no special hot-cells or radiological protection measures are

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required, which makes for ease of testing and lower costs. Consequently it is easier to perform systematic studies of different parameters such as temperature, fluence or material characteristics. However, the differences in the damage produced by proton or ion irradiation with respect to that created by neutron irradiation must be understood. That is, a fundamental knowledge of defect production and evolution is needed for the development of materials more resistant to radiation. Robust models must be developed to be able to extrapolate to different experimental conditions. And such models must be validated through tailored experiments. Modelling has always played a very significant role in the process of understanding radiation effects, particularly for void swelling (Mansur et al., 1986; Trinkaus and Singh, 1993), hardening of RPV steels (Wirth and Odette, 1999) and RIS in austenitic steels (Allen and Was, 1998). In the last few years, with the development of high performance computing, powerful first principle computational tools have emerged, that are being applied to produce very accurate calculations of defects in metals (Domain and Becquart 2002; Fu et al., 2005). This has resulted in a new approach in modelling irradiated materials: the development of models from first principles. Unlike in previous years, the challenge now resides in developing a predictive model with no adjustable parameters, and no parameters extracted from experiments. Information from the most accurate models is transferred to less accurate but more efficient tools in order to reach the time and length scales of experiments and to go from microscopic defects to macroscopic properties in what is known as multiscale modeling (Phillips, 2001). This is quite a challenge that requires the collaboration of different research groups as well as experimental validation. This is treated in detail in Chapter 15. The combined effort of experimental researchers and modellers to develop predictive tools should help in finding the answers to many of the existing open questions. This will require more experimental results in the particular conditions of LWRs, especially for high fluences. A clear and quantitative connection between the different microscopic features (precipitates, loops, voids, gas bubbles) to macroscopic observations (hardening, IASCC, RIS, shift in Charpy ductile to brittle transition temperature (DBTT)) is needed. For the case of IASCC in austenitic steels, there are some explanations available for BWR conditions, since Cr depletion due to RIS will have a significant effect in the (relative) oxidizing water of these reactors compared to the ‘oxygen-free’ water in PWRs. It should be noted that hydrogen water chemistry (HWC) or proprietary noble metal chemical additions affect the electrochemical potential of stainless steels in BWRs by depressing the effect of dissolved oxygen. However, the consequences of RIS on IASCC in the PWR environment is not clear.

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Acknowledgements

We want to thank Dr Frank Bergner (FZD) for his thorough review of this chapter and very helpful comments. We also want to thank Dr Marta Serrano (CIEMAT) for her valuable information on RPV and internal components.

8.8

Definitions

Neutron flux: number of neutrons crossing an area per unit of time, measured in neutrons/m2s in the International System of Units. Neutron fluence: total number of neutrons crossing an area integrated over time (neutrons/m2). dpa: displacements per atom, number of times an atom is displaced from its lattice site.

8.9

Sources of further information and advice

Akamatsu M, Van Duysen J C, Pareige P, Auger P (1995) ‘Experimental evidence of several contributions to the radiation damage in ferritic alloys’, J. Nucl. Mat. 225, 192–195. Asoka-Kumar P, Wirth B D, Sterne P A, Howell R H (2002) ‘Composition and magnetic character of nanometre-size Cu precipitates in reactor pressure vessel steels: implications for nuclear power plant life extension’, Phil. Mag. Lett. 82, 609–615. Auger P, Pareige P, Akamatsu M, Van Duysen J C (1994) ‘Microstructural characterization of atom clusters in irradiated pressure vessel steels and model alloys’, J. Nucl. Mat. 211, 194–201. Beaven P A, Frisius F, Kampmann R and Wagner R (1986), ‘Analysis of defect microstructures in irradiated ferritic alloys’, in Atomic Transport and Defects in Metals by Neutron Scattering, C Janot, W Petry, D Richter, T Springer (eds), Springer Proceedings in Physics, Vol. 10, SpringerVerlag, Berlin, p. 228. Becquart C S (2005) ‘RPV steel microstructure evolution under irradiation: a multiscale approach’, Nucl. Instrum. and Meth. B 228, 111–121. Brillaud C, Hedin F (1992) ‘In-service evaluation of French pressurized water reactor vessel steel’, Effects of Radiation on Materials: 15th International Symposium, ASTM STP 1125, R E Stoller, A S Kumar, D S Gelles (eds), American Society for Testing and Materials, Philadelphia, PA, pp. 23–49. Cammelli S, Degueldre C, Kuri G, Bertsch J, Lützenkirchen-Hecht D, Frahm R (2009) ‘Study of atomic clusters in neutron irradiated reactor pressure vessel surveillance samples by extended X-ray absorption fine structure spectroscopy’, J. Nucl. Mat. 385, 319–324.

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Cumblidge S E, Motta A T, Catchen G L, Brauer G, Böhmert J (2003) ‘Evidence for neutron irradiation-induced metallic precipitates in model alloys and pressure vessel weld steel’, J. Nucl. Mat. 320, 245–257. Diaz de la Rubia T, Zbib H M, Khraishi T A, Wirth B D, Victoria M, Caturla M J (2000), ‘Multiscale modelling of plastic flow localization in irradiated materials’, Nature 406, 871. Fisher S B, Buswell J T (1987) ‘A model for PWR pressure-vessel embrittlement’, Int. J. Pressure Vessels and Piping, 27, 91–135. Fukuya K, Ohno K, Nakata H, Dumbill S, Hyde J M (2003) ‘Microstructural evolution in medium copper low alloy steels irradiated in a pressurized water reactor and a material test reactor’, J. Nucl. Mat. 312, 163–173. Gan J, Was G S, Stoller R E (2001), ‘Modeling of microstructure evolution in austenitic stainless steels irradiated under light water reactor condition’, J. Nucl. Mat. 299, 53–67. Garner F A (1994) ‘Irradiation performance of cladding and structural steels in liquid metal reactors’, Materials Science and Technology, Volume 10A, Nuclear Materials, R W Cahn, P Haasen, E J Kramer (eds), VCH, Weinheim. Glade S C, Wirth B D, Odette G R (2006) ‘Positron annihilation spectroscopy and small angle neutron scattering characterization of nanostructural features in high-nickel model reactor pressure vessel steels’, J. Nucl. Mat. 351, 197–208. Gurovich B A, Kuleshova E A, Lavrenchuk O V, Prikhodko K E, Strombakh Y A (1999) ‘The principal structural changes proceeding in Russian pressure vessel steels as a result of neutron irradiation, recovery annealing and re-irradiation’, J. Nucl. Mat. 264, 333–353. IAEA-TECDOC-1442 (2005) ‘Guidelines for prediction of irradiation embrittlement of operating VVER-440 reactor pressure vessels’, IAEA, Vienna. IAEA-TECDOC-1470 (2005) ‘Assessment and management of ageing of major nuclear power plant components important to safety: BWR pressure vessels’, IAEA Vienna. Jenkins M L (1994) ‘Characterization of radiation damage microstructures by TEM’, J. Nucl. Mater., 216, 125–156. Kasada R, Kudo T, Kimura A, Matsui H, Narui M (2005) ‘Effects of neutron dose, dose rate and irradiation temperature on the irradiation embrittlement of a low-copper reactor pressure vessel steel’, J. ASTM Int. 2 (3) Paper ID JAI12399. Kuri G, Cammelli S, Degueldre C, Bertsch J, Gavillet D (2009) ‘Neutron induced damage in reactor pressure vessel steel: an X-ray absorption fine structure study’, J. Nucl. Mat. 385, 312–318. Maussner G, Scharf L, Langer R, Gurovich B (1999) ‘Microstructure alterations in the base material, heat affected zone and weld metal of a 440-VVER

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reactor pressure vessel caused by high fluence irradiation during long term operation; material: 15CH2MFA~0.15C-2.5Cr-0.7Mo-0.3V’, Nucl. Eng. Design 193, 359–376. Miller M K, Russell K F, Sokolov M A, Nanstad R K (2003) ‘Atom probe tomography characterization of radiation-sensitive KS-01 weld’, J. Nucl. Mat. 320, 177–183. Miller M K, Wirth B D, Odette G R (2003) ‘Precipitation in neutron-irradiated Fe-Cu and Fe-Cu-Mn model alloys: a comparison of APT and SANS data’, Mater. Sci. Eng. A353, 133–139. Miller M K, Sokolov M A, Nanstad R K, Russell K F (2006) ‘APT characterization of high nickel RPV steels’, J. Nucl. Mat. 351, 187– 196. Miller M K, Russell K F, Ko�ik J, Keilova E (2000) ‘Embrittlement of low copper VVER 440 surveillance samples neutron irradiated to high fluences’, J. Nucl. Mat. 282, 83–88. Odette G R, Mader E V, Lucas G E, Phythian W J, English C A (1993) ‘The effect of flux on the irradiation hardening of pressure vessel steels’, Effects of Radiation on Materials: 16th International Symposium, ASTM STP 1175, A S Kumar, D S Gelles, R K Nanstad, E A Little (eds), American Society for Testing and Materials, Philadelphia, PA. Odette G R, Wirth B D, Bacon D J, Ghoniem N M (2001) ‘Multiscalemultiphysics modeling of radiation damaged materials: embrittlement of pressure vessel steels’, MRS Bulletin, 26, 176–181. Odette G R (1995) ‘Radiation induced microstructural evolution in reactor pressure vessel steels’, Microstructure of Irradiated Materials, MRS Symp. Proc. 373, I. Robertson L E Rehn, S J Zinkle, W J Phythian (eds) Pittsburgh, PA, MRS, pp. 137–148. Planman T, Pelli R, Törrönen K (1994) ‘Irradiation embrittlement mitigation’, AMES Report No. 1, Espoo. Quinot P, Desfontaines G (1999) ‘The main components of the European pressurized water reactor’, Nucl. Eng. Design 187, 121–133. Ranganath S, Sandusky D W, Chapman T L, Gordon G M, Kiss E (1990) ‘Proactive approaches to assure the structural integrity of boiling water reactor components’, Nucl. Eng. Design 124, 53–70. PWRMRP (2000) Review of Phosphorus Segregation and Intergranular Embrittlement in Reactor Pressure Vessel Steels (PWRMRP-19): PWR Materials Reliability Project (PWRMRP), EPRI, Palo Alto, CA, TR114783. Schaublin R, Yao Z, Baluc N, Victoria N (2005) ‘Irradiation-induced stacking fault tetrahedra in fcc metals’, Phil. Mag. 85, 769–777. Slugen V, Kögel G, Sperr P, Triftshäuser W (2002) ‘Positron annihilation studies of neutron irradiated and thermally treated reactor pressure vessel steels‘ J. of Nucl. Mat. 302, 89–95.

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Soneda N, Díaz de la Rubia T D (2001) ‘Migration kinetics of the self-interstitial atom and its clusters in bcc Fe’, Philos. Mag. A 81 (2), 331–343. Stoenescu R, Schaublin R, Gavillet D, Baluc N (2007)‘Welding-induced mechanical properties in austenitic stainless steels before and after neutron irradiation’, J. Nucl. Mat. 360, 255–264. Ulbricht A, Bergner F, Dewhurst C D, Heinemann A (2006) ‘Small-angle neutron scattering study of post-irradiation annealed neutron irradiated pressure vessel steels’, J. Nucl. Mat. 353, 27–34. Yu Q, Was G S, Wang L M, Odelte R, Alexander D E (2001) ‘Hardening and microstructure of model reactor pressure vessel steel alloys using proton irradiation’, Microstructural Processes in Irradiated Materials, MRS Symp. Proc. 650, G E Lucas, L Snead, M A Kirk, Jr, R G Elliman Pittsburgh, PA, MRS. Zubuchen C, Viehrig H-W, Weiss F-P (2009) ‘Master curve and unified curve applicability to highly neutron irradiated Wester type reactor pressure vessel steels’, Nucl. Eng. Design 239, 1246–1253.

8.10

References

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Lambrecht M, Malerba L, Almazouzi A (2008) ‘Influence of different chemical elements on irradiation-induced hardening embrittlement of RPV steels’, J Nucl. Mat., 378, 282–290. Liu C L, Odette G R, Wirth B D, Lucas G E (1997) ‘A lattice Monte Carlo simulation of nanophase compositions and structures in irradiated pressure vessel Fe-Cu-Ni-Mn-Si steels’, Mater. Sci. Eng. A – Struct. 238, 202–209. Lucas G E (1997) ‘Irradiation-induced changes in the mechanical properties and microstructures of solution annealed austenitic stainless steel at low to intermediate temperatures’, Mat. Res. Soc. Symp. Proc. 439, 425. Mansur L K, Coghlan W A (1983) ‘Mechanisms of helium interaction with radiation effects in metals and alloys: a review’, J. Nucl. Mat. 119, 1–25. Mansur L K, Lee E H, Maziasz P J, Rowcliffe A P (1986) ‘Control of helium effects in irradiated materials based on theory and experiment’, J. Nucl Mat. 633, 141–143. Maussner G, Scharf L, Langer R, Gurovich B (1999) ‘Microstructure alterations in the base material, heat affected zone and weld metal of a 440-VVER reactor pressure vessel caused by high fluence irradiation during long term operation; material: 15CH2MFA~0.15C-2.5Cr-0.7Mo-0.3V’, Nucl. Eng. Design 193, 359–376. Maziasz P J (1993), ‘Overview of microstructural evolution in neutron-irradiated austenitic stainless steels’, J. Nucl. Mat. 205, 118–145. Maziasz P J, McHargue C J (1987) ‘Microstructural evolution in annealed austenitic steels during neutron irradiation’, Int. Metals Rev. 32, 190. McMahon Jr, C J, Vitek V, Kameda J (1981), ‘Mechanics and mechanisms of intergranular fracture’, in Developments in Fracture Mechanics, 2nd edn G G Chell (ed.), Applied Science Publishers, Englewood Cliffs, NJ. Meechan C J, Sosin A, Brinkman J A (1960) ‘Thermally activated point defect migration in copper’, Phys. Rev. 120, 411–419. Meslin E (2007) ‘Mécanismes de fragilisation sous irradiation aux neutrons d’alliages modèles ferritiques et d’un acier de cuve: Amas de défauts’, Thesis, University of Roven. Meslin E, Radiguet B, Pareige P, Barbu A (2010) ‘Kinetic of solute clustering in neutron irradiated ferritic model alloys and a French pressure vessel steel investigated by atom probe tomography’, J. Nucl. Mat. 399, 137–145. Miller M K, Russell K F (2007) ‘Embrittlement of RPV steels: an Atom probe tomography perspective’, J. Nucl. Mat. 371, 145–160. Miller M K, Pareige P, Burke M G (2000a) ‘Understanding Pressure Vessel Steels: An Atom Probe Perspective,’ Materials Characterization 44, 235–254. Miller M K, Russell K F, Ko�ik J, Keilova E (2000b) ‘Embrittlement of low copper VVER 440 surveillance samples neutron irradiated to high fluences’, J. Nucl. Mat. 282, 83–88. Miller M K, Russell K F, Sokolov M A, Nanstad R K (2003a) ‘Atom probe tomography characterization of radiation-sensitive KS-01 weld’, J. Nucl. Mat. 320, 177–183. Miller M K, Wirth B D, Odette G R (2003b) ‘Precipitation in neutron-irradiated Fe-Cu and Fe-Cu-Mn model alloys: a comparison of APT and SANS data’, Mater. Sci. Eng. A353, 133–139. Miller M K, Sokolov M A, Nanstad R K, Russell K F (2006) ‘APT characterization of high nickel RPV steels’, J. Nucl. Mat. 351, 187–196. Miller M K, Russell K F, Sokolov M A, Nanstad R K (2007) ‘APT characterization of irradiated high nickel RPV steels’ J. Nucl. Mat. 361, 248–261. Nagai Y, Tang Z, Hasegawa M, Kanai T, Saneyasu M (2001) ‘Irradiation-induced Cu

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aggregations in Fe: an origin of embrittlement of reactor pressure vessel steels’, Phys. Rev. B, 63, 134110. Odette G R (1994), ‘On the ductile to brittle transition in martensitic stainless steels – mechanisms, models and structural implications’, J. Nucl. Mat. 212–215, 45–51. Odette G R, Lucas G E (1998) ‘Recent progress in understanding reactor pressure vessel steel embrittlement’, Radiation Effects and Defects in Solids 144 (1–4), 189–231. Odette G R, Lucas G E (2001) ‘Embrittlement of nuclear reactor pressure vessels’, J. Metals, 53 (7), 18–22. Odette G R, Wirth B D (1997) ‘A computational microscopy study of nanostructural evolution in irradiated pressure vessel steels’, J. Nucl. Mat. 251, 157–171. Okita T, Sato T, Sekimura N, Garner F A, Greenwood L R (2002) ‘The primary origin of dose rate effects on microstructural evolution of austenitic alloys during neutron irradiation’, J. Nucl. Mat. 307–311, 322–326. Phillips R (2001) Crystals, Defects and Microstructures: Modeling Across Scales, Cambridge, Cambridge University Press. Phythian W J, English C A (1993) ‘Microstructural evolution in reactor pressure vessel steels’, J. Nucl. Mat. 205, 162–177. Pokor C, Brechet Y, Dubuisson P, Massoud J-P, Barbu A (2004) ‘Irradiation damage in 304 and 316 stainless steels: experimental investigation and modeling. Part I: evolution of the microstructure’, J. Nucl. Mat. 326, 19–29. PWRMRP-19 (2000) Review of Phosphorus Segregation and Intergranular Embrittlement in Reactor Pressure Vessel Steels (PWRMRP-19): PWR Materials Reliability Project (PWRMRP), EPRI, Palo Alto, CA: TR-114783. Rowcliffe A F, Zinkle S J, Stubbins J F, Edwards D J, Alexander D J (1998) ‘Austenitic stainless steels and high strength copper alloys for fusion components’, J. Nucl. Mat. 258–263, 183–192. Scott P (1994) ‘A review of irradiation assisted stress corrosion cracking’, J. Nucl. Mat. 211, 101. Singh B N, Zinkle S J (1993) ‘Defect accumulation in pure fcc metals in the transient regime: a review’, J. Nucl. Mat. 206, 212. Solt G, Frisius F, Waeber W B, Tipping Ph (1993) ‘Irradiation induced precipitation in model alloys with systematic variations of Cu, N and P content: a small angle neutron scattering study’, Effects of Radiation on Materials, 16th Int. Symp., ASTM STP 1175, ed. A S Kumar et al. ASTM, West Conshohocken, PA, pp. 444–462. Tipping Ph. (1996) ‘Lifetime and ageing management of nuclear power plants: brief overview of some light water reactor component ageing degradation problems and ways of mitigation’, Int. J. Pres. Ves. & Piping, 66, 17–25. Trinkaus H, Singh B N (2003) ‘Helium accumulation in metals during irradiation – where do we stand?’, J. Nucl. Mat. 323, 229. Ulbricht A, Böhmert J, Viehrig H-W (2005) ‘Microstructural and mechanical characterization of radiation effects in model reactor pressure vessel steels’, J. ASTM Int. 2, 151–164. Ullmaier H (1984) ‘The influences of Helium on the bolb properties of fusion-reactor structural materials’, Nucl. Fusion 24, 1039. Wallin K (1984) ‘The scatter in KIC-results’, Eng. Fract. Mech. 19, 1085. Wallin K (1985) ‘The size effect in KJC results’, Eng. Fract. Mech. 22, 149. Ware A G, Shah V N (1992) ‘Age-related degradation of boiling water reactor vessel internals’, Nucl. Eng. Design 133, 49. Was G S (2007) ‘Fundamentals of Radiation Materials Science’, Springer, New York.

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Was G S, Allen T R, Busby J T, Gan J, Damcott D, Carter D, Atzmon M, Kenik E A (1999) ‘Microchemistry and microstructure of proton-irradiated austenitic alloys: toward an understanding of irradiation effects in LWR core components’, J. Nucl. Mat. 270, 96–114. Watanabe H, Aoki A, Murakami H, Muroga T, Yoshida N (1988) ‘Effects of phosphorus on defect behavior, solute segregation and void swelling in electron irradiated Fe-CrNi alloys’, J. Nucl. Mat. 155–157, 815. Wirth B D, Odette G R (1999) ‘Kinetic lattice Monte Carlo simulations of cascade aging in dilute iron-copper alloys’, Microstructural Processes in Irradiated Materials, MRS Symp. Proc. 540, 5. J Zinkle, G E Lucas, R C Ewing, J S Williams (eds) Warrendale, PA, MRS, pp. 637–642. Xu Q, Yoshiie T, Sato K (2006) ‘Dose dependence of Cu precipitate formation in Fe-Cu model alloys irradiated with fission neutrons’, Phys. Rev. B 73, 134115. Xu Q, Yoshüe T, Sato K (2008) ‘Formation of Cu precipitates and vacancy clusters in neutron-irradiated Fe-Cu alloys’, Phil. Mag. Lett. 88, 353–362. Yao Z, Hernández-Mayoral M, Jenkins M, Kirk M (2008) ‘Heavy-ion irradiations of Fe and FeCr Model alloys. Part 1: Damage evolution in thin foils at lower doses’, Phil. Mag. 88, 2851–2880. Yoshida N, Xu Q, Watanabe H, Muroga T, Kiritani M (1992) ‘Low dose fission neutron irradiation on P- and T-modified austenitic alloys with improved temperature control’, J. Nucl. Mater. 191–194, 1114. Yvon P, Carré F (2009) ‘Structural materials challenges for advanced reactor systems’, J. Nucl. Mat. 385, 217–222. Zinkle S J, Seitzman L E, Wolfer W G (1987) ‘Stability of vacancy clusters in metals: I. Energy calculations for pure metals’, Phil. Mag. 55, 111. Zinkle S J, Maziasz P J, Stoller R E (1993) ‘Dose dependence of the microstructural evolution in neutron-irradiated austenitic stainless steel’, J. Nucl. Mat. 206, 266. Zinkle S J, Ice G E, Miller M K, Pennycook S J, Wang X-L (2009) ‘Advances in microstructural characterization’, J. Nucl. Mat. 386–388, 8–14.

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9

Stress corrosion cracking (SCC) of austenitic stainless steels in high temperature light water reactor (LWR) environments

P. L. A n d r e s e n, GE Global Research Center, USA

Abstract: This chapter focuses on stress corrosion cracking (SCC) of austenitic stainless steels in high temperature light water reactor (LWR) environments, and emphasizes common grades of wrought austenitic stainless steel over cast, ferritic or martensitic stainless steels. Corrosion fatigue and environmental effects on fracture are related forms of degradation, and are also addressed. This chapter provides insight into the SCC dependencies and the underlying processes and mechanisms. There are 20–30 key, interdependent variables, and this complexity makes a purely empirical approach intractable. Research has transformed this complexity to one in which the origin of the effects and their interaction is reasonably well understood and quantified. Key words: stress corrosion cracking, crack growth rate, high temperature water, light water reactors, stainless steel, nickel base alloys, sensitization, stress intensity factor, corrosion potential, irradiation effects, prediction.

9.1

Introduction

Many incidents of stress corrosion cracking (SCC) in austenitic stainless steels have occurred in boiling and pressurized water reactors (BWRs and PWRs), and laboratory studies of this phenomena extend over 50 years [1–28]. The corrosion and SCC behavior of structural materials in high temperature water have many unique elements, and extrapolation of data and intuition based on low temperature (<100 °C) response is of limited use, especially when considering low temperature, aggressive forms of SCC such as in halide cracking, including in concentrated MgCl2 environments at 100–154 °C. By contrast, SCC in stainless steels in relevant high temperature LWR environments represents a lower SCC susceptibility or growth rate that ideally satisfies the need for long life in nuclear power plant (NPP) applications. The pressure boundary design codes (e.g., ASME Section III [29]) address only cyclic loading, and reflect an evolution of their historical emphasis on fatigue and fracture; SCC is simply addressed by noting that the design must ensure it will not occur. Thus, the concepts of SCC immunity and thresholds (conditions below which SCC does not occur) are attractive, although careful 236 © Woodhead Publishing Limited, 2010

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studies show that immunity rarely exists [10–12]. High resistance to initiation can be achieved by careful attention to design issues such as low stress; no crevices or sharp corners; proper selection of materials, heat treatment and fabrication/welding/grinding; good control of environment including during transient operation; minimizing grinding or other sources of surface cold work; polishing/treating surfaces to minimize surface roughness and create compressive stresses, etc. However, in large welded structures ‘initiation circumventing’ phenomena (broadly, anything that results in much faster initiation than expected from smooth surface data) have contributed to extensive SCC in operating plants. While continued efforts to improve initiation are important, complete reliance on a thin 1–50 mm skin in welded structural components is unwise, and ‘inherent resistance’ to SCC crack advance remains a key consideration. A crack growth rate approach is supported by the vastly better experimental techniques for quantifying crack growth than crack initiation. SCC is often schematically shown as a confluence of stress, material and environment factors (Fig. 9.1), with an implication that there is a very small central region of susceptibility. It would be wise to view such diagrams as representing an iso-susceptibility boundary, with the central overlap representing high susceptibility or growth rate, and increasingly larger circles overlap representing boundaries of lower susceptibility or growth rate. Figure 9.1 also shows the underlying crack tip processes that more directly control SCC growth, as well as how these processes are affected by radiation. Improved crack detection and monitoring techniques, coupled with improved testing techniques, permit SCC growth rate response to be observed into regimes that were previously hidden ‘below ground level’, and had given the impression of immunity and masked the inter-relationships of related phenomena. The modern view is that all variables follow a well-behaved continuum, with growth rates often dropping by, e.g., 100¥ or 1000¥ – but not to zero – for changes in corrosion potential or other parameters. Examples include corrosion potential (SCC occurs in deaerated water), degree of sensitization or neutron fluence level (SCC occurs in unsensitized stainless steel), water purity (SCC occurs in ultra high purity water), crack tip stress intensity factor, K (SCC has been readily observed in at ~5 MPa÷m), etc. Thus it is much more realistic and appropriate to consider a continuum in SCC response as a function of material, environment and stressing parameters, with the need to define boundaries of adequately low susceptibility or growth rate, rather than viewing very low susceptibility as immunity. The objective of this chapter is to provide a broad framework for understanding and interpreting SCC (especially the crack growth rate response) of austenitic stainless steels; to define linkage and commonality of SCC in structural materials; to define the dependencies (e.g., corrosion potential, water purity, stress intensity factor (K), neutron fluence, sensitization, etc.);

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Stress

Environment

Oxide rupture rate at crack tip Microstructure

Passivation rate at crack tip (a) Solution renewal rate to crack tip Stress

Df anionic   transport

Oxide rupture rate at crack tip

Environment Microstructure

g-field –

Crack tip f (A) , pH

Hardening Relaxation

Passivation rate at N-fluence crack tip G.B. denudation Segregation (b)

9.1 SCC is often shown schematically as an overlap of stress, environment and microstructure, although the small central area should be viewed as a region of high SCC susceptibility. The underlying factors include mass transport, dynamic strain that damages the protective oxide, and the repassivation process. This structure provides a basis for anticipating the effects of irradiation. The complexity of SCC is reflected in the large number of influential variables and the associated requirement that all 20 to 40 in a given system be adequately controlled.

to propose controlling processes and SCC mechanisms; to describe SCC mitigation approaches; and to summarize the state of ability to predict SCC. This understanding facilitates appropriate mitigation measures to be taken for NPP ageing and plant-life management (AM and PLiM, respectively) to achieve safe, reliable and long-term performance/operation in the system, structure or component (SSC) involved.

9.2

Historical problems and structures affected

At the start of commercial nuclear power over 50 years ago, it was reasonable to select stainless steels for use in high purity, unaggressive water chemistries

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at moderate stresses, and this decision was supported by field experience in other industries and accelerated testing. However, in retrospect, such accelerated tests did not have sufficient sensitivity to ensure high SCC resistance over long operating times. Earlier and more incidents of SCC developed in boiling water reactors (BWRs) because of the oxidizing conditions that exist in the coolant, although the incidence of SCC in pressurized water reactors (PWRs), with very low dissolved oxygen, has also increased in the last 25 years. The common grades of austenitic stainless steel are types 304, 316, 321 and 347 stainless steel (Table 9.1). These are generically called 18-8 stainless steels because they have approximately 18% Cr and 8% Ni, along with 0.04–0.07% C, ~1% Mn and ~0.5% Si. The low carbon specification (e.g., 304L and 316L) is <0.03% C, and 0.015–0.020% C is typical. Type 316 stainless steel includes 2–3% Mo, while Type 321 incorporates ~0.5% Ti and Type 347 ~0.5% Nb. C, N and Ni stabilize the austenite phase, and reducing their concentration or adding Cr and Mo promotes formation of the ferrite phase. Stainless steel weld metals, such as Type 308/308L, are designed to form 3–12% ferrite, which is necessary to avoid hot cracking during weld solidification. Ferritic and martensitic stainless steels typically have < 1% Ni. Common nickel alloys (which are all austenitic) include Alloy 600 (~15.5% Cr, ~8% Fe), Alloy 690 (~30% Cr, ~9% Fe) and Alloy 800 (~30% Ni, ~20% Cr, bal Fe). Nickel alloy weld metals, such as Alloy 182 (~15% Cr) and Alloy 82 (~20% Cr), are used both to weld nickel alloys as well as for most dissimilar metals welds. Alloys 152 and 52 are the higher Cr (~30%) weld metals. In BWRs, water boils on the fuel cladding surfaces, and the steam, after going through the steam separators and dryer, directly drives the turbine. The steam is condensed and then the water is demineralized and returned as feedwater to the pressure vessel. Water is circulated upward through the Table 9.1 Typical composition of common grades of austenitic stainless steel, (wt%) AISI Grade UNS

Fe

Cr

Ni

Mn*

Si*

C*

Other

304 304L 316 316L 321 347

Bal Bal Bal Bal Bal Bal

18–20 18–20 16–18 16–18 17–19 17–19

8–12 8–12 10–14 10–14 9–12 9–13

~1.2 ~1.2 ~1.2 ~1.2 ~1.2 ~1.2

~0.5 ~0.5 ~0.5 ~0.5 ~0.5 ~0.5

~0.045 ~0.020 ~0.045 ~0.020 ~0.045 ~0.045

– – 2–3 Mo 2–3 Mo ~0.5 Ti* ~0.5 Nb*

S30400 S30403 S31600 S31603 S32100 S34700

* Mn is 2% maximum; Si is 1% maximum C is 0.08% maximum in non-L grades, typically 0.045% in modern heats, 0.065% in older heats; L-grades are typically 0.015 – 0.020% Ti in 321 is 5¥(C+N) minimum, 0.7% maximum Nb in 347 is 10 ¥ C minmum, 1.0% maximum

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core and steam separators, and down through the annulus between the core shroud and the pressure vessel. An average water molecule recirculates 7–10 times before leaving the reactor as steam. The accumulation of impurities in the reactor water is controlled by the reactor water clean-up system. Radiation–primarily by neutrons–produces radiolysis of water that simplistically generates hydrogen (H2) and hydrogen peroxide (H2O2). Most of the H2 partitions to the steam phase, leaving behind a net oxidizing environment in the reactor water. Nuclear core reactivity is controlled by moving control rods, containing neutron absorbers such as B or Hf, that penetrate through the bottom head of the reactor pressure vessel. Normal water chemistry (NWC) is considered to have 100–200 parts per billion by weight (ppb, or parts per 109) dissolved O2 and ~10 ppb dissolved H2, although the conditions in the core are different due to radiolysis. All US and some international BWRs inject H2 (hydrogen water chemistry or HWC) to mitigate SCC by reducing the O2 and H2O2 – the H2 level in the reactor water can vary from ~40 to ~250 ppb. Most US BWRs employ NobleChem™, which creates an electrocatalytic surface on all wetted components. This requires much less H2 injection for SCC mitigation, typically resulting in reactor water with 35–40 ppb H 2. This also avoids an increase in radioactive N16, which shifts from soluble (e.g., NO3–) to volatile (e.g., NO or NH3) at higher levels of H2, and increases the radiation level in the piping and turbine (called turbine shine). N16 forms at very low concentrations from transmutation of O16 in the core, and has a half-life of 7.1 seconds. PWRs circulate pressurized primary coolant water through the core, where it is heated from ~290 °C to ~323 °C. It then flows through the inside of the steam generator tubing, with boiling occurring on the outside in the secondary water, which produces steam that drives the turbine. The primary water (that flows through the core) contains boric acid (H3BO3) and lithium hydroxide (LiOH), with ~2 bar of dissolved H2 (~3 parts per million (ppm) H2 or ~35 cc H2 per kg H2O) to minimize radiolysis. The reactivity during the fuel cycle is controlled by adjusting/tapering the boron (B) level from ~1500 ppm to < 50 ppm at the end of cycle. Li (as LiOH) is adjusted to maintain an approximately constant pH at temperature (pH300C ~7.1), and so is typically adjusted/tapered down from ~3 ppm to ~0.3 ppm over the fuel cycle (the target pH300C has varied by plant and over time). Control rods that penetrate from the upper head are used primarily as a safety system and for shutdown. Over-pressure in the primary system is maintained using a pressurizer – a smaller pressure vessel connected to the primary system in which water is electrically heated to ~343 °C, where the vapor pressure is about 151.7 bar, about 34.5 bar above the vapor pressure of 323 °C water. The secondary system uses high purity, low O2 feedwater with O2 scavengers such as hydrazine or morpholine. Trace level impurities (ppb levels) that

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accumulate during steam generator operation are ‘blown down’ by releasing water from the bottom of the steam generator. In BWRs, stainless steels are extensively used for piping, pressure vessel cladding and structures inside the pressure vessel, including the core shroud (which separates the up-flow through the core from the downward flow in the annulus), the core plate (which supports the bottom of the fuel), the top guide (which aligns the top of the fuel bundles), the shroud dome (the dome is above the core, and supports the steam separators), the steam separators, the steam dryer (above the separators), etc. The earliest incidents of SCC in BWRs occurred in stainless steel fuel cladding [30–33] (before zirconium alloys were used). Radiation produces atom displacement damage that causes grain boundary segregation (including Cr depletion) and radiation hardening, which significantly increase susceptibility to SCC (as discussed later). SCC then developed in creviced, cold-worked and/or furnace sensitized components, and these factors were eliminated from designs. Subsequently, SCC occurred in small, then increasingly larger diameter, stainless steels pipes that were weld sensitized; these were primarily type 304 stainless steel, which typically had carbon levels > 0.06%. Most problems occurred in piping external to the pressure vessel (e.g., recirculation and reactor water clean-up piping), and considerable efforts were devoted to understanding and mitigating SCC in these piping systems. Most piping was replaced by a low carbon grade which incorporated nitrogen to retain strength. This was designated as ‘LN’ (low carbon, added nitrogen) or ‘NG’ (nuclear grade) stainless steel. Subsequently, SCC developed in many core internals, initially in highly stressed components (such as absorber tubes that contain the B4C neutron absorber material, which swells over time), core spray piping welds, and core shroud welds (Fig. 9.2). Related structures (e.g., the top guide) and materials (e.g., Alloy 182 nickel-base weld metals) also exhibited cracking. The control blades and absorber tubes are replaceable, but SCC can permit dissolution of the B4C and potentially affect control rod insertion. Various approaches have been taken to manage SCC in core internals, including hydrogen injection (H2 water chemistry), electrocatalysis (NobleChem™) [34–37], mechanical restraints (tie rods) that limit motion of the shroud in the unlikely event of through-wall circumferential cracking, etc. In PWRs, the incidents of SCC were slower to appear, but have increased steadily over time and include stainless steel fuel cladding, baffle former bolts, pressurizer heater sleeves, canopy seals in the control rod drives, steam generator safe ends, etc. Some debate continues regarding whether SCC occurs primarily under unusual circumstances, such as upset water chemistry (e.g., O2 in make-up water), crevice/boiling situations (like pressurizer heat sleeves), severely ground (surface cold-worked) piping welds, very high O2 conditions such as in unvented control rod drives (where air is pressurized

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Internals holddown bracket

Steam dryer and brackets Core spray piping and spargers

Shroud head bolts Feedwater nozzle weld butters

Core spray safe end to nozzle weld butters

Top guide rim welds and beam

Shroud head and separators

Core shroud weld Jet pump riser brace

SRM dry tubes and in-core housings

Recirc. inlet and outlet nozzle to safe end weld butters

Core plate rim welds Recirc. inlet safe ends

Access hole cover and shroud support shelf

CRD stub tube to housing welds (SS)

Shroud-to-shroud support weld Recirc. piping welds

Jet pump riser elbow Alloy 182/ 600

304, 304L, or 347SS

9.2 Schematic of a BWR showing materials and areas where cracking has occurred.

and trapped during start-up), etc. However, with a large number of incidents and the widespread observation of SCC of cold-worked stainless steel in laboratory tests in PWR primary water, it is likely that SCC of stainless steels is a generic issue whose incidence is likely to continue to rise somewhat with operating time. The presence of ‘cold work’ from shrinkage strains in the weld heat affected zone is a universal issue, and is addressed later. Examples of intergranular SCC in BWR and PWR components are shown in Fig. 9.3. The emphasis in this chapter is on SCC in stainless steels, but SCC has also been observed in many nickel-base alloys and weld metals, including in BWRs (Alloy 600 shroud head bolts, Alloy X-750 jet pump beams, and various Alloy 182 weldments) and PWRs (Alloy 600 steam generator tubing, Alloy 600 upper head and lower head penetrations, Alloy 182 and 82 weld metals in various locations), etc., so some comments will be directed at showing the areas of common dependencies and mechanisms in stainless steels and nickel alloys.

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1.3 cm Crack

Iascc pwr Baffle Bolt

2 1

3

9.3 Examples of intergranular SCC in stainless steel: BWR pipe welds (left), BWR control rod sheath (middle) and PWR baffle former bolts (right).

9.3

Stress corrosion cracking (SCC) dependencies – introduction

Increasingly sophisticated measurement techniques and patient observation have eroded the historical concepts of immunity and thresholds to SCC in high temperature water [11–13, 19, 21], and have revealed a continuum in SCC response among the relevant materials, environments and stresses. High resolution crack depth measurements monitor SCC growth rates into regimes that were previously hidden ‘below ground level’, which had previously given the impression of immunity and uniqueness. The modern view is that all variables follow a well-behaved continuum, with growth rates often changing dramatically for changes in corrosion potential (Figs 9.4 and 9.5), water purity (Figs 9.4–9.7), stress intensity factor (Fig. 9.8) and other parameters. If a low growth rate is achieved, another variable may become more aggressive and cause the growth rate to increase to a readily observable level; demonstration of the underlying interconnection requires high resolution measurements. The interdependency is widespread, with the effect of a given variable generally dependent on the state of all other variables (e.g., Figs 9.4–9.8). It is important to understand that corrosion potential represents a mixed electrochemical potential, and does not refer to the ‘potential or rate of corrosion’. As oxidants (e.g., O2 or H2O2) are

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added to water, the corrosion potential of iron- and nickel-base materials increases, although the corrosion rate often does not. When not specified in more detail, corrosion potential refers to the mixed electrochemical potential on the relevant surface, usually an iron- or nickel-base structural material. The absence of SCC growth rate thresholds (i.e., where SCC ceases) is demonstrated for sensitization, corrosion potential and water purity in Figs 9.4–9.7, which show data (and predictions) for sensitized and unsensitized stainless steels and nickel alloys in moderate to high purity 288 °C water. Figures 9.4 and 9.5 include data from the SKI/EPRI round robin [38] (the smaller symbols at about +150 to +200 mVshe, or mV referenced to the standard hydrogen reference electrode), and also include data (larger symbols in Fig. 9.4b) for carefully controlled changes in potential and at low potential. Similar data exist for the non-threshold response vs. temperature, buffered 304 STAINLESS STEEL 1

25 mm CT Specimen 10

–6

Furnace sensitized; 15 C/cm 288 °C water; 0.1-0.3 mS/cm Constant load; 25 Ksi√in

5

2

10 8 6

Crack propagation rate, mm/s

42.5 min/h

10–7

11

14

14.2 min/h Theoretical curves

a aa 9

mS/cm 0.1

10–8 Hydrogen 12 water chemistry b b 10–9

–600

0.3 0.2

3 7 2

4

Normal water chemistry (ex-core)

–400 –200 0 +200 Corrosion potential, mVshe (a)

+400

9.4 SCC growth rate vs. corrosion potential for stainless steels tested in 288 °C high purity water containing 2000 ppb O2 and 95–3000 ppb H 2.

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1.E-05

Crack growth rate, mm/s

1.E-06

Sensitized 304 Stainless Steel 30 MPa√m, 288 °C water 0.06–0.4 mS/cm, 0–25 ppb SO4 filled triangle = constant load open squares = “gentle” cyclic

¨200 ppb O2 ¨500 ppb O2 ¨2000 ppb O2

Stress corrosion cracking of austenitic stainless steels

Screened Round Robin data – highest quality data – corrected corr. potential – growth rates corrected   to 30 MPa√m

42.5 28.3

1.E-07

14.2 min/h GE PLEDGE Predictions 30 MPa√m

1.E-08

0.5

0.25

2000 ppb O2 Ann. 304SS 200 ppb O2 0.1

0.06 mS/cm

0.06 mS/cm Industry mean 30 MPa√m 1.E-09 –0.6 –0.5 –0.4 –0.3 –0.2 –0.1 0.0 0.1 Corrosion potential, Vshe (b)

9.4 Continued

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0.2

0.3

0.4

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Understanding and mitigating ageing in nuclear power plants 1.E-05 Sensitized 304 Stainless Steel 30 MPa√m, 288 °C water 0.06–0.4 mS/cm, 0–25 ppb SO4 SKI Round Robin Data filled triangle = constant load open squares = “gentle” cyclic

¨200 ppb O2 ¨500 ppb O2 ¨2000 ppb O2

246

Crack growth rate, mm/s

1.E-06 CW A600

316L (A14128, square) 304L (Grand Gulf, circle) non-sensitized SS 50%RA 140 C(black) 10%RA 140C (grey) 1.E-07

42.5 28.3 14.2 min/h

CW A600 GE PLEDGE Predictions 30 MPa√m Sens SS

0.5

2000 ppb O2 Ann. 304SS 200 ppb O2

0.25

1.E-08 0.1

0.06 mS/cm

GE Pledge Predictions for Unsensitized Stainless Steel (upper curve for 20% CW) 1.E-09 –0.6 –0.5 –0.4

–0.3 –0.2 –0.1 0.0 0.1 Corrosion potential, Vshe

0.2

0.3

0.4

(c)

9.4 Continued

water chemistry, presence vs. absence of grain boundary carbides, cold work, etc., as discussed later. The behavioral characteristics that reflect similarities and continua in SCC response include: ∑

BWR and PWR primary water chemistry, including pure water and the buffered B/Li chemistry in PWR primary water ∑ type or grade of stainless steel ∑ stainless steels vs. nickel alloys ∑ corrosion potential, controlled primarily by oxidants (dissolved O2 and H2O2), reductants (H2) and pH, and the effects of impurities such as chloride and sulphate ∑ temperature

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Stress corrosion cracking of austenitic stainless steels

247

∑ intergranular morphology, even in unsensitized materials ∑ grain boundary Cr depletion vs. grain boundary carbides ∑ cold work from bulk deformation, surface cold work, weld residual strain, etc. ∑ neutron irradiation ∑ stress intensity factor (K) ∑ general corrosion rate ∑ grain boundary silicon ∑ effects of environment on fracture ∑ crack initiation.

1.E-05 Sensitized 304 Stainless Steel 30 MPa√m, 288 °C water 0.06–0.4 mS/cm, 0–25 ppb SO4 SKI Round Robin Data filled triangle = constant load open squares = “gentle” cyclic

4 dpa 304 SS

Crack growth rate, mm/s

1.E-06 316L (A14128, square) 304L (Grand Gulf, circle) non-sensitized SS 50%RA 140 C(black) 10%RA 140C (grey) 1.E-07

CW A600 GE PLEDGE Predictions 0.5 mS/cm 30 MPa√m 0.25 Sens SS 0.1

1.E-08

CW A600

2000 ppb O2 Ann. 304SS 200 ppb O2

0.06

GE Pledge Predictions for Unsensitized Stainless Steel (upper curve for 20% CW) 1.E-09 –0.6 –0.5 –0.4 –0.3 –0.2 –0.1 0.0 0.1 Corrosion potential, Vshe (a)

0.2

0.3 0.4

9.5 SCC growth rate vs. corrosion potential in 288 °C high purity water for stainless steels, irradiated stainless steel and various nickel-base alloys.

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Understanding and mitigating ageing in nuclear power plants

1.E-06

Sensitized 304 Stainless Steel 30 MPa√m, 288 °C water 0.06–0.4 mS/cm, 0–25 ppb SO4 SKI Round Robin Data filled triangle = constant load open squares = ‘gentle’ cyclic

¥ 750 20% CW HTH AH

316L (A14128, square) 304L (Grand Gulf, circle) non-sensitized SS 50%RA 140 C(black) 10%RA 140C (grey)

1.E-07

¨200 ppb O2 ¨500 ppb O2 ¨2000 ppb O2

1.E-05

Crack growth rate, mm/s

248

718 CW A600 42.5 28.3 14.2 min/h

CW A600

20% CW ¥ 750 AH

GE PLEDGE Predictions 30 MPa√m Sens SS

0.5

2000 ppb O2 Ann. 304SS 200 ppb O2

0.25 1.E-08 0.1

¥750 718 HTH

0.06 mS/cm

GE Pledge Predictions for Unsensitized Stainless Steel (upper curve for 20% CW) 1.E-09 –0.6 –0.5 –0.4

–0.3 –0.2 –0.1 0.0 0.1 Corrosion potential, Vshe (b)

9.5 Continued

© Woodhead Publishing Limited, 2010

0.2

0.3

0.4

249

¨2000 ppb O2

Alloy 182, Alloy 600 & St. Steel 30 MPa√m, 288 °C Water Sens 304 SS Round Robin (open) Circles = constant load Triangles = ‘gentle’ cyclic SO4 & Cl data in pink

¨500 ppb O2

1.E-05

¨200 ppb O2

Stress corrosion cracking of austenitic stainless steels

Includes: • All alloys tested • All heat treatments • No cold work data • Medium screening

1.E-06

Crack growth rate, mm/s

42.5 28.3 1.E-07

14.2 min/h Expected 16¥ peak in CGR for alloy 182 vs. H2 in pure water

1.E-08

CGR data in pure water in Ar or low H2

1.E-09 –0.6 –0.5 –0.4

0.25 0.1 0.06 mS/cm

–0.3 –0.2 –0.1 0.0 0.1 Corrosion potential, Vshe (c)

SS Pledge Prediction * Normal YS

0.2

0.3

0.4

9.5 Continued

9.3.1 Need for SCC insights and modelling based on data and fundamental understanding The number of factors (and sub-factors) and their complexity (interdependency) make a purely empirical approach to understanding and quantifying SCC intractable. Without interdependencies, evaluating five levels for each of 20 variables would require 5 ¥ 20 = 100 tests. But with interdependency, the number leaps to 520 = 1014 experiments. Of course, all experiments would need to be definitive and reproducible, which are notorious problems in SCC experiments. Another approach is to emphasize fundamental understanding and modeling of SCC. The challenge is that there are hundreds of processes that might be

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Understanding and mitigating ageing in nuclear power plants

Crack growth rate, mm/s

10–6

10–7

316L Stainless Steel 25 mm CT Specimen Constant load 288 °C water Test conditions: 0 C/cm2 EPR ª 27.5 MPa√m 200 ppb O2

10–8

Predicted curves from Pledge Code for typical range in ECP 10–9

100 10–1 Solution conductivity, mS/cm (a)

101

10–7

Crack growth rate, mm/s

250

10–8 Theoretical prediction for typical range in corrosion potential 10–9 304 Stainless steel Water 288 °C; 200 ppb O2 ~25 Ksi√in; 15 C/cm2 (27.5 MPa√m) 10–10

10–1 100 Solution conductivity, mS/cm (b)

101

9.6 Predicted and observed SCC growth rates of stainless steel in 288 °C BWR water as a function of water purity, e.g., from additions of chloride or sulphate.

© Woodhead Publishing Limited, 2010

Stress corrosion cracking of austenitic stainless steels Sens. 304 Stainless Steel Deaerated, 289 °C water fc ª – 0.6 Vshe SSRT 10–6/s

10–2

in/h

Crack growth rate, mm/s

10–4

251

10–5 10–3

Prediction 1 ¥ Œ· app

TG + Ductile IG + Ductile

10–6

10–2

102

10–1

100 H2SO4, ppm (a)

10–4

101

102

Sensitized Type 304 Stainless Steel Factor of improvement for change in conductivity from 0.5 to 0.1 mS/cm

Factor of improvement

K = 5 MPa√m

101

K = 25 MPa√m

SSRT = 1 ¥ 10–6/s

100

–0.5 –0.4 –0.3 –0.2 –0.1 0 0.1 Corrosion potential, Vshe (b)

0.2

0.3

9.7 (a) Elevated growth rate can occur in deaerated water if sufficient H2SO4 is added. (b) The factor of improvement for a change in water purity depends on corrosion potential and stressing.

controlling SCC, and while formulations exist to describe those processes, often the constants cannot be accurately estimated in high temperature water. Historically, electrochemists have tried to describe SCC as a predominantly

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Understanding and mitigating ageing in nuclear power plants

10

–5

4

Stress intensity, ksi√in 6 8 10 20 30 40 60 80 Sens. 304 Stainless Steel 288 °C Water

Crack growth rate, mm/s

10–6

10–4

10–7

NRC disposition line

Theory 15 C/cm2­, –50 m Vshe 0.5 mS/cm Theory 15 C/cm2­, –50 m Vshe 0.2 mS/cm Theory 15 C/cm2­, –200 Æ – 500 m Vshe 0.2 mS/cm

10–8

10–9

10–10

10–3

4

10–5

10–6

Crack growth rate, in/h

252

10–7

6 8 10 20 30 40 60 80 Stress intensity, MPa√m

9.8 Predicted and observed SCC growth rate vs. stress intensity factor for various materials and water chemistry conditions in 288 °C water.

electrochemical phenomenon, but with little success. Mechanical engineers have focused on stress and stress intensity factor and cycling/vibration, also with limited success. Nuclear engineers have modeled radiation damage; materials scientists considered phases and nucleation and diffusion, etc. All of these factors are important, but only within a more complex context of the whole of SCC. The most effective approaches have evaluated existing data and dependencies, hypothesized the underlying processes (e.g., the effects of water chemistry or mechanics), performed critical experiments to validate the hypothesis, then modeled and quantified the critical parameters. From an engineering perspective, SCC is often, and simplistically, characterized as a confluence of stress, environment and metallurgy (Fig. 9.1). But this is only an engineering, macroscopic, external (to the crack) perspective. Crack advance must instead be understood from the perspective of the crack tip

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Stress corrosion cracking of austenitic stainless steels

253

system, and then the role of more readily measurable, external, engineering parameters on the crack tip must be described. Historically, the occurrence of SCC in each material or environment has been considered unique, but as increasingly careful studies are performed, perhaps 80% of the dependencies and characteristics are common among iron and nickel-base structural materials. This section will emphasize the description of individual effects, as itemized above – primarily for stainless steels, but with some discussion of the dependencies and mechanisms shared with nickel alloys. The next section will provide a mechanistic approach for interpreting these parameters, including their interdependencies, from the perspective of the crack tip system.

9.3.2 Distinctions between BWRs and PWRs The primary differences between BWR and PWR primary conditions are associated with: coordinated changes vs. time in the B and Li level in PWR primary water that increases the pH at temperature in pure water from 5.65 (BWR) to ª7.2 (PWR); the H2 fugacity (from ª40 to 3000 ppb H2); and temperature (274/288 °C vs. 288/323/343 °C for the PWR core inlet, core outlet and pressurizer). Of these differentiating factors, temperature is the most important in stainless steels, whereas for nickel alloys, both temperature and H2 are important. Irrespective of whether BWRs operate at high (electrochemical) corrosion potential or low corrosion potential on the exposed surfaces (that is, with or without H2 addition and/or NobleChem™), the crack tip itself is always deaerated and at a low corrosion potential. The details of secondary (steam generator) chemistry in PWRs (a few early BWRs also had steam generators) is beyond the scope of this chapter, and the comments below should be viewed as a broad, introductory perspective. Once-through and recirculating steam generator designs are both in use, and there are typically three or four steam generators per PWR. The boiling process leaves behind non-volatile species, and steam generators do not have a clean-up system (such as the reactor water clean-up system in BWRs) but rather depend on periodic ‘blow-down’ to purge the system of impurities. Improved control of feedwater chemistry has been key to controlling corrosion; this includes control of impurities, dissolved oxygen and other oxidants, and additives to control pH and scavenge oxidants. The control of pH was once done using phosphate, but essentially all commercial steam generators now use all volatile treatment (AVT) comprising species like hydrazine, morpholine, EDTA, etc., most of which also scavenge oxidants. Lead (Pb) and sulfur (S) are major concerns, as is the corrosion product ‘sludge’ that builds up on the tube sheet and can induce stresses and strains on the tubing as well as an occluded (crevice) chemistry. The heat exchanger tubes are the primary SCC problem in steam generators, and most Alloy 600 tubing

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254

Understanding and mitigating ageing in nuclear power plants

has been changed from mill annealed (most susceptible) or thermally treated Alloy 600 to Alloy 690. Thermal treatments for both Alloy 600 and Alloy 690 are typically ~705 °C for ~12 hours, which produced grain boundary carbides without Cr depletion. Canadian (CANDU) and German (PWR) plants generally use Alloy 800, which at 30Ni-20Cr has a composition between a stainless steel and Alloy 600, and has given very good service. (CANDU reactors are heavy water moderated, but the secondary side is similar to that in LWRs.)

9.3.3 BWR and PWR primary water chemistry The presence of the B and Li chemistry in PWR primary water sometimes leads to the assumption that the SCC growth rate behavior is fundamentally different from that in BWR pure water. Certainly the presence of oxidants and high (electrochemical) corrosion potential in BWRs produces an elevated growth rate compared to deaerated water (Figs 9.4 and 9.5). However, the crack tip, where the SCC process takes place, is deaerated in all cases. Also, the vast majority of BWRs now operate at low corrosion potential using surface catalysis and a small amount of H2 [34–37], as discussed in the Section 9.8. The net H2 at the surface is typically 35–40 ppb H2, and the corrosion potential of structural materials is close to the potential in deaerated water, which is predominantly controlled by the H2/H2O reaction. Thus, a major distinction in chemistry between BWR and PWR primary water comes down to the role of B and Li in PWRs, whose concentrations are typically ~1500 ppm B as H3BO3 and 3 ppm Li as LiOH at the start of the fuel cycle, tapering to < 50 ppm B (sometimes <10 ppm) and ~0.3 ppm Li at the end of cycle. B and Li levels are coordinated to maintain an approximately constant pH300C of ~7.0–7.3. Laboratory studies involving dozens of duringthe-test changes in B and Li concentrations [39–43] showed no effect on crack growth rate in stainless steels or nickel alloys, even when changing from deaerated pure water (Fig. 9.9). This is consistent with expectations based on the Pourbaix diagram (Fig. 9.10), because shifts in pH produce a change in corrosion potential (controlled by the H2/H2O reaction) that is exactly parallel to the metal-metal oxide reactions for Ni, Fe and Cr (of these, only Ni/NiO is close to the H2/H2O reaction). In the mid-pH range, the solubility of metal cations (e.g., Ni2+) or anions (e.g., HNiO2–) is limited; at more extreme pHs, a significant effect of pH might be expected. In the presence of oxidants, the relatively concentrated B and Li chemistry plays a large role, producing growth rates that are perhaps 10¥ higher than in pure water (Fig. 9.11). Thus, oxidants should be avoided in PWR primary water. Indeed, the common practice of adding of H2O2 during PWR shutdown (typically at <130 °C) to react with the dissolved H2 should be evaluated to determine the extent to which cracking may be accelerated.

© Woodhead Publishing Limited, 2010

Stress corrosion cracking of austenitic stainless steels

To 1100 ppm B, 2 ppm Li @2675h

–8

pH325C constant at ~7.25

c283 – 0.5TCT of A600 CRDM, 325C 27.5 MPa√m, 30 cc/kg H2, Varying B/Li

11.19

Pt potential

11.14 1200

1700

2200 Test time, h (a)

Ct potential 2700

0

–0.2 –0.4 –0.6 –0.8 –1

3200

SCC#3 – c315 – Alloy 600, CRDM Tube, 93510

11.23

0.2

0.4

Outlet Conductivity ∏ 0.01

11.225

0.2

11.205

To 600 ppm B, 2.2 ppm Li @2325h

Crack length, mm

11.21

To Constant K @2011h

11.22 11.215

11.2

11.195

Conductivity, mS/cm or Potential, Vshe

3.5 ¥ 10 mm/s

0.4

4.3 ¥ 10 mm/s

0 6.6 ¥ 10–9 mm/s

–0.2 –0.4

c315 – 0.5TCT of A600 CRDM, 325C 17.5 MPa√m, 18 cc/kg H2, Pure Water

–9

–0.8

11.19

Pt potential

11.185 2000

2200

–0.6

2400 2600

2800 3000 3200 3400 Test time, h (b)

Ct potential

–1 3600 3800 4000

9.9 Crack length vs. time for Alloy 600 tested in 325 °C water containing H2 under constant K conditions. Various changes in B, Li and pHT, including additions to pure water, produced no discernible change in crack growth rate.

© Woodhead Publishing Limited, 2010

Conductivity, mS/cm or Potential, Vshe

11.24

0.6

To 60 ppm B, 0.3 ppm Li @3315h

11.29

0.8

Conductivity ¥ 0.01 To 3200 ppm B, 7 ppm Li @1880h

Crack length, mm

11.34

To 1100 ppm B, 2 ppm Li

11.39

SCC#2 – c283 – Alloy 600, CRDM Tube, 93510

To Constant K @1201h

11.44

255

Understanding and mitigating ageing in nuclear power plants 2.50

(Ni++)

(Ni(OH)3–)

2.00 (Ni(OH)2aq)

(NiOH+) 1.50

Potential (Volts VS. SHE)

(b)

1.00

NiO2 Ni

0.50

++

H

2 /H 2O

(a)

0.00

Fe 2O

3

–0.50

Ni3O4

Fe

≠H2

3O 4

Fe

NiO

–1.00

≠pH

Ni –1.50

–2.00

Ni(OH)3–

0

4

pH

8

12

16

(a)

Dissolved H2, cc/kg

256

75 70 65 60 55 50 45 40 35 30 25 20 15 10 5 0 260

Film reduction and/or no film formation Film formation Experimentally measured Ni/NiO transition

Experimentally measured Ni/NiO transition

Ni regime NiO regime

280

300 320 340 Temperature (°C) (b)

360

380

9.10 (a) Ni–H2O Pourbaix diagram at 300 °C. (b) Right Ni/NiO phase boundary as a function of H2 fugacity and temperature [44].

© Woodhead Publishing Limited, 2010

Stress corrosion cracking of austenitic stainless steels

28.5 28.4 28.3

2.5 ¥ 10–6 mm/s

1 ¥ 10–8 mm/s

0.1 0 –0.1 –0.2 –0.3 –0.4 –0.5 –0.6

28.2 6 ¥ 10–6 mm/s

CT Corrosion Potential

–0.7

Potential, Vshe or Conductivity, mS/cm

Crack length, mm

28.6

33 MPa√m, R = 0.7 0.01 Hz + 900 s hold

pH288C ~ 6.79 To 95 ppb H2 @2518h

28.7

c85 – 1T CT of Sens. SS – AJ9139 95 ppb H2, 1200 ppm B, 2.2 ppm Li, 288C

To 200 ppb O2 @2426h

28.8

257

–0.8 28.1 2375 2400 2425 2450 2475 2500 2525 2550 2575 2600 2625 2650 Time, h

9.11 SCC crack length vs. time for sensitized stainless steel. Adding 200 ppb O2 at 2279 h causes a marked increase in growth rate in water with 1200 ppm B and 2.2 ppm Li as LiOH at 288 °C.

9.4

Stress corrosion cracking (SCC) dependencies – materials and water chemistry

9.4.1 Type or grade of stainless steel A brief introduction to the common grades of stainless steels was given in Section 9.2. A lot of emphasis has been made on differences among grades of stainless steel, but the primary differences relate to whether sensitization and grain boundary Cr depletion occurs, which is described in more detail in the next section. Among common grades of austenitic stainless steel, there is little distinction in their SCC growth response if they are in the same metallurgical condition. That is, if there is a similar Cr depletion profile, cracks in all types of stainless steels will grow at similar rates. If unsensitized stainless steels are cold worked to the same yield strength, they all grow cracks at nearly identical rates [13–21]. This can be seen in prior and future figures (e.g., Figs 9.4, 9.8 and 9.20). Austenitic stainless steels or nickel alloys with low carbon and/or additions of Nb or Ti (also Mo) have less tendency to sensitize, and thus L-grade (low carbon) or NG-grade (low carbon, high nitrogen) or Nb/Ti-stabilized stainless steels can show less SCC susceptibility than standard Type 304 stainless steel. However, sensitization is still possible in all of these grades. They have similar work hardening characteristics, so cold working produces

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Understanding and mitigating ageing in nuclear power plants

a similar increase in tensile yield strength. Stainless steels and Alloy 600 of a similar condition (level of sensitization or cold work) all exhibit about the same SCC growth rates in BWR water – there is very little difference as a function of Mo, Ti or Nb [13–21]. The fact that stainless steels and nickel alloys (e.g., Alloy 600 and Alloy 182/82 weld metal) exhibit similar crack growth rates and response vs. corrosion potential and water purity (Figs 9.4 and 9.5) suggests that a similar crack advance mechanism is operative.

9.4.2 Stainless steels vs. nickel alloys Subsequent sections will identify some unique elements of SCC in austenitic stainless steel vs. nickel alloys. In broad terms, their SCC growth rate dependencies are 80% similar. Figure 9.12 shows that the crack growth rate in 288 °C water is close to identical in stainless steels and nickel alloys, and the electrochemical response (e.g., corrosion potential and corrosion rate vs. dissolved O2) is the same. Since the metal ion solubility is also comparable, it is reasonable to expect that the crack chemistry will be essentially identical. It is not surprising that the effect on SCC of aqueous impurities, such as Cl– and SO4=, is also similar for these two classes of material. A primary difference is related to the higher temperature dependence of SCC in nickel alloys, which typically has an activation energy of ~135 kJ/ mol [41, 45–46] vs. ~80–100 kJ/mol for stainless steels [15, 16, 19, 47], which may be associated with the higher diffusivity and creep in the nickel alloys. A second difference is that the corrosion potential in deaerated water is close to the Ni/NiO phase boundary (Fig. 9.10a), and far from the Fe/ Fe3O4 phase boundary; it is effectively impossible to add sufficient H2 to drop the corrosion potential near the Fe/Fe3O4 phase boundary. For nickel alloys, being near the phase stability for nickel might be expected to be important. Indeed, the growth rate of nickel alloys peaks near the Ni/NiO phase boundary, with a peak height of 2.5–3¥ for Alloy 600 and 8–20¥ for Alloy 182/82 weld metal [41–43]. While few comparative crack initiation tests have been performed on stainless steels and nickel alloys, field experience indicates that crack initiation occurs more readily in nickel alloys in PWR primary water, e.g., at 320–340 °C than at 274 °C (BWRs) or than in stainless steel. This is probably because of the higher creep rate in nickel alloys at elevated temperatures or in stainless steels. While thermal creep at these temperatures would not be a significant problem at constant load, it becomes more important as crack advances and the strain field is redistributed and creates dynamic strain.

9.4.3 Corrosion potential and water chemistry There are two distinct mechanisms for the effects of corrosion potential on SCC growth rate, one related to the oxidants in the bulk water that © Woodhead Publishing Limited, 2010

Stress corrosion cracking of austenitic stainless steels

12.15 1000

0.2 Outlet Conductivity 1.8 ¥ 10 mm/s 1.9 ¥ 10–8 mm/s

2.7 ¥ 10–7 mm/s 1200

0.1

CT Potential 1400

1600

1800 Time, h (a)

–7

2000

–0.3 –0.4 –0.5

2200

2400

0.2

CT Potential

0.1

23.5 4700

2.1 ¥ 10 mm/s

2.0 ¥ 10 mm/s

3.2 ¥ 10–8 mm/s

–7

To 2000 ppb O2 @5107h

To 95 ppb H2 @4816h

Outlet Conductivity To const 31.1 MPa√m To 200 ppb @3449h O2 @4174h

Crack length, mm

23.7

23.55

–0.6

0.3 Pt Potential

23.6

–0.1 –0.2

23.75

23.65

0

–7

0 –0.1 –0.2 –0.3

Annealed + 20%CW Alloy 600 30 MPa√m, 2000 ppb O2, Pure Water

–0.4 –0.5 –0.6

4800

4900

5000 Time, h (b)

5100

5200

Potential, Vshe or Conductivity, mS/cm

12.25

0.3

To 2000 ppb O2 @2108h

12.35

To 6% H2 in Ar @1245h

12.55 To static load @367h Pt/Rh coated @960h

Crack length, mm

12.65

12.45

0.4

SSC of c126 – 316L SS 20% Cold Work 27.5 MPa√m, 2000 ppb O2, Pure Water

Potential, Vshe or Conductivity, mS/cm

12.75

259

–0.7 5300

9.12 Crack length vs. time for an unsensitized, 20% cold-worked 316L stainless steel (a) and Alloy 600 (b) showing similar growth rates and effects of corrosion potential.

create a gradient in corrosion potential and altered crack chemistry, and one related to the presence of reductants (especially H2) at the crack tip. Oxidants increase the corrosion potential on the external surface and create an electrochemical (differential aeration) cell vs. the region inside the crack (where O2 is consumed, and therefore the corrosion potential is low). This

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Understanding and mitigating ageing in nuclear power plants

electrochemical cell produces a shift in the crack chemistry. In deaerated water, changes in dissolved H2 or shifts in pH (from B/Li, ammonia, etc.) also shift the corrosion potential, but these species are not consumed in the crack, and thus the corrosion potential is essentially identical throughout the crack, and no unique crack chemistry forms. Altered crack chemistry can also form from temperature gradients (especially boiling). Oxidants markedly elevate the corrosion potential on the free surface above that in argon (Ar) or nitrogen (N2) deaerated water, e.g., to between lines (a) (H2/H2O) and (b) (H2O/O2) in Fig. 9.10a. The elevation in corrosion potential occurs only on the external surface and near the crack mouth; it does not persist far into the crack because oxidants are rapidly consumed once any convective flow decays within the crack. Oxidants have an indirect effect on SCC; the potential gradient that forms near the crack mouth (Fig. 9.13) produces a change in crack chemistry that affects the entire crack. Figure 9.13 shows that electromigration, which concentrates anions in the crack, is balanced by back-diffusion of species out of the crack (if convection is present, it generally overwhelms the contributions of electromigration or diffusion). The dynamic equilibrium between electromigration and diffusion occurs a short distance into most cracks, and after the dynamic equilibrium in concentration is initially formed, the species move deeper into the crack by ordinary diffusion. The gradient in corrosion potential is determined by both the external corrosion potential and the corrosion potential within the crack. The corrosion Cl–

Time evolution after increase in ECP of –0.5 to 0.1 Vshe

fC = +0.1

O 2 , fC fC = –0.5 Vshe

Df

e–

e– Cl–, SO42–, OH–

1 Ni Æ Ni2+ Æ H+ 2 e– Microcell 1

Ni Æ Ni2+ + 2e– + H2O Æ NiO + 2H+

Zn2+ Æ e 3 Microcell

–

4 O2 + 2H2O + 4e– = 4OH–

3 H2 Æ 2H+ + 2e–

2 2H+ + 2e– Æ H2 JA = –DADCA – zmCAFDf + CAV flux = diffusion + f-driven + convection

9.13 Schematic of crack chemistry transport processes in high temperature water with O2.

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Stress corrosion cracking of austenitic stainless steels

261

potential in the crack is controlled by the pH and the concentration of H2 in the crack, but it remains along the H2/H2O line that controls the potential in deaerated water (Fig. 9.10a). The change in pH in the crack is generally more extreme than the ~1.5 unit change in high temperature pH from the addition of B/Li to deaerated water. If most impurities are species like Cl– and SO4=, acidification occurs. This produces a somewhat higher corrosion potential in the (deaerated) crack because of the effect of pH on line (a) in Fig. 9.10a, which reduces the potential gradient somewhat. The potential gradient does not always cause acidification of the crack. Acidification requires the presence of anions other than OH– (such as SO4= or Cl–) to balance the charge of the H+ cation. In high purity water containing oxidants, the pH in the crack shifts in the alkaline direction (higher pH), with the increase in OH– concentration balanced by an increase in metal ion solubility. An alkaline shift also occurs if species such as sodium hydroxide (NaOH) or potassium hydroxide (KOH) are the predominant impurities. As the crack tip moves in the alkaline direction (higher in pH), the corrosion potential in the deaerated crack drops because of the pH change (line (a) in Fig. 9.10a), and the potential gradient increases. However, the balancing cations (e.g., Na+) also tend to be rejected from the crack by the potential gradient. The effect of oxidants on crack growth is quite similar for sensitized and unsensitized, cold-worked stainless steel (Fig. 9.4), and for Alloy 600 and Alloy 182 weld metal (and other nickel alloys and irradiated stainless steel, Fig. 9.5). Detailed examples of the crack length vs. time response for changes in corrosion potential are shown in Fig. 9.12 for stainless steel and Alloy 600 – many other similar examples are shown in References [10–23]. The oxidant concentration itself is not the controlling factor. Electromigration is controlled by the potential gradient, not by the oxidant concentration (or gradient) per se. The relationship between dissolved O2 and corrosion potential is not linear, but rather follows a complex relationship (Fig. 9.14). Above ~2–20 ppb O2 (in tests with well-controlled chemistry), the corrosion potential follows a roughly logarithmic relationship, as predicted by the Nernst equation. At lower oxidant concentrations, the potential falls rapidly to that associated with deaerated water. The reason for the rapid drop is that O2 becomes mass transport limited, with O2 consumed on the surface (by corrosion or reaction with H2) faster than it can diffuse though the nearsurface, stagnant boundary layer. The O2 concentration at which the potential drops depends on the O2 reaction rate (including corrosion rate and reaction with H2) and the convection (flow) rate. As the oxidant level and corrosion potential decrease, the growth rate also decreases but does not cease. Another mechanism of corrosion potential effect on crack growth is associated with the presence of H2 in deaerated water (note that this mechanism can also be operative if the bulk water has stoichiometrically more H2 than

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Laboratory and in-reactor data for 304/316 stainless steel 288 °C water Unirradiated condition Conductivity range .1 – .5 mS/cm Variety of flow rates Both bright and oxidized surfaces

101 102 Dissolved oxygen content, ppb

103

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9.14 Effect of dissolved oxygen on corrosion potential in 288 °C water. Note that modern measurements made with improved water purity and higher refresh rates tend to represent the data where the large drop in corrosion potential occurs at the lowest O2 concentration. Also, improved, more accurate reference electrodes show the corrosion potential to be perhaps 50 mV higher than the upper line.

O2). The presence of H2 always decreases corrosion potential; H2 decreases the corrosion potential in aerated or deaerated water, although by a smaller magnitude than the increase from oxidants. For a tenfold increase in H2 in deaerated water, the corrosion potential of Pt, stainless steel and nickel alloys decreases by 55 mV at 281 °C and 60 mV at 332 °C (according to the Nernst equation, which depends on temperature in Kelvin). Unlike oxidants, H2 is not consumed in the crack – some H2 might be created by corrosion reactions, or some might be lost by transport through the metal, but these are small factors. Generally, it is reasonable to assume that the H2 level at the crack tip is the same as the bulk H2 concentration (or the H2 that remains after reaction with oxidants in the bulk water). An increase in H2 causes the H2/H2O line to shift vertically downward on a Pourbaix diagram (Fig. 9.10a), and for stainless steels and nickel alloys the corrosion potential in deaerated water is exactly thermodynamic and identical to the response on catalysts like platinum (Pt) [10–23, 39–43]. Among Fe, Cr and Ni, only Ni/NiO has a metal-metal oxide equilibrium near the H2/H2O line. Thus, it is reasonable to expect that the crack growth

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rate response of nickel alloys might be affected, but stainless steels not. References [10, 13–19, 21] give examples of the lack of effect of H2 on the SCC growth rate of stainless steel, although it should be noted that effects of H2 are observed in crack initiation under corrosion fatigue initiation [48, 49], corrosion fatigue crack growth [50] and slow strain rate testing [51, 52], perhaps because the corrosion rate is higher at low corrosion potential [10, 14]. For nickel alloys, a peak in crack growth rate is observed at the Ni/NiO phase boundary (Fig. 9.10a), which occurs at different H2 levels as a function of temperature (Fig. 9.10b). The growth rate response is symmetrical and the growth rate similar on the NiO and Ni-metal sides of the phase boundary, making it unlikely that there is a change in crack growth mechanism. Morton et al. [44, 53–56] found that the peak height is higher for Alloy 82 weld metal and Alloy X-750 than for Alloy 600 (Fig. 9.15a), but that the peak height and shape were unchanged over the 260–360 °C range they evaluated. Andresen et al. [39–43] have confirmed the effect on Alloy 600 and Alloy 182 weld metal (Fig. 9.15b), but have seen a stronger effect of H2 (higher peak) of 15–20¥. Bruemmer et al. [57] also observed a higher peak for Alloy 182. This is likely due to their use of constant K (no cycling) vs. Morton’s use of periodic partial unloading throughout the test, because cycling biases low growth rate data upward more than high growth rate data. Figure 9.5c shows the effect of H2 in the context of oxidants. At high corrosion potential the presence of oxidants causes a relatively high crack growth rate. As the oxidant level decreases, the growth rate drops. If H 2 is added to deaerated water, the corrosion potential begins to shift down. For example, 100% O2 bubbled in pure water at standard temperature and pressure (STP) gives ~42 ppm dissolved O2, but even 0.1 ppm O2 can increase the corrosion potential by >500 mV compared to deaerated water. 100% H2 gives 1.58 ppm dissolved H2, and shifting by 100¥ (e.g., from 10 ppb to 1000 ppb H2) only causes a 114 mV decrease in potential at ~300 °C. As H2 is changed and the Ni/NiO phase boundary crossed, there is a peak in the crack growth rate of nickel alloys, as shown in Fig. 9.15. While it has not been verified in deaerated pure water, the data of Morton et al. showing that the peak height is unchanged from 260–360 °C. The observation described above shows that the crack growth rate is unaffected by B/Li additions to deaerated water [39, 40] and provides a strong basis for anticipating that H2 affects SCC growth of nickel alloys in pure water. Since most BWR structural materials are exposed to 274 °C water, the Ni/ NiO phase boundary and associated peak in crack growth rate occurs at ~200 ppb H2 (Fig. 9.10b), which is the regime in which ‘medium HWC’ BWRs operate.

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3.0 K = 60 ksi√in, 338 °C K = 25 ksi√in, 338 °C K = 60 ksi√in, 316 °C

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Effect of H2 on Crack Growth Rate of 182 Weld Metal – c380, 325C 6.E–08

Observed Predicted

Crack growth rate, mm/s

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Predictions scaled to CGR at 10.4 cc/kg H2. The decrease in CGR was larger than predicted based on an 16X peak vs. H2.

4.E–08 3.E–08 2.E–08 1.E–08

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9.15 Crack growth rate of Alloy 600 (a) [44] and Alloy 182 (b) [42, 43] in high temperature water as a function of potential/H2. The peak growth rate occurs very close to the Ni/NiO phase boundary, with a characteristic height and width associated with the material and condition.

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9.4.4 Temperature The effect of temperature is somewhat complex when considered in the range of ~25–350 °C. Focusing on operating temperatures (274–288 °C for BWRs and 288–343 °C for PWRs) and low corrosion potential conditions, stainless steels and nickel alloys show a significant temperature dependency, although the effect on crack growth is stronger on nickel alloys (at least Alloy 600 and its weld metals) at ~ 135 kJ/mole [41, 45, 46] vs. ~80–100 kJ/ mole for stainless steel [15, 16, 19, 47]. For 135 kJ/mole, the difference in growth rate is ~13¥ from 288 to 343 °C, while for 90 kJ/mole, the difference is 5.6¥. The lower activation energy for stainless steel translates to a lower incidence of cracking of these components in PWRs compared to Alloy 600 and its weld metals, but the limited incidence of cracking in stainless steel PWR components is certain to increase with time. For an activation energy of 135 kJ/mol, a temperature change from 274 and 343 °C increases the growth rate by ~28¥, and explains why good SCC mitigation is achieved by reducing corrosion potential of materials in BWRs (most structural materials operate at 274 °C), while problems have developed in the low corrosion potential PWR primary water, especially in nickel alloys. Temperature also affects the dissolved H2 concentration in the coolant, where the peak in growth rate occurs in nickel alloys (at the Ni/NiO phase boundary) and the dissolved H2 associated with the phase boundary varies with temperature (Fig. 9.10b). While Attanasio and Morton [44] reported that the peak in crack growth rate may not always occur exactly at the Ni/NiO phase boundary, the data are too few and too scattered to convincingly support a deviation from the Ni/NiO phase boundary, especially since the deviation was sometimes slightly positive and sometimes slightly negative. The effect of temperature over wider ranges (to lower temperatures) is more complex. When the environment is more aggressive (resulting in higher growth rates), there is a tendency for the increase in growth rate vs. temperature to be monotonic (Fig. 9.16), but if the water chemistry is unaggressive (e.g., less oxidizing), a peak in growth rate is generally observed at ~175 °C (Fig. 9.16). An interest in this wider temperature regime is usually related to plant start-up and shutdown, but these involve many other changes in addition to temperature. In BWRs, the water often has higher impurity levels and is more oxidizing due to prior exposure to air, from radiolysis and/or because H2 is not injected during shutdown or start-up. There are also differences in pressure loading, vibration, and differential thermal strains, e.g., in the ferrite containing stainless steel weld metal vs. the base metal (all austenite) due to the different coefficients of thermal expansion. In PWRs, H2O2 is generally added during cool down to consume most of the H2 and minimize transport of radioactive crud; during heat-up, the water is more oxidizing due to exposure to air during shutdown.

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SCC of high strength, cold worked austenitic steels in pure water, oxygen 0.1–300 ppm, K = 50 MNm–3/2 10–5

X50MnCr185, Rp0.2 = 1000 MPa X6MnCrN1818, Rp0.2 = 1500 MPa

–7

10–8

10–9

300 250 200

°C

25 10–3

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Air Sat’d 0.224 mS/cm HCl 0.077 mS/cm HCl

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(a)

100

Open = temp. ≠ Close = temp. Ø

10–10

10–12 1.6

150

in/h

Da , (m/s) Dt Stress corrosion crack growth rate,

© Woodhead Publishing Limited, 2010

10

Effect of Temperature Constant K = 33 MPa√m + R = 0.5, 0.01 Hz every 1000s 200 ppb O2, 0.27 mS/cm H2SO4

0

2.0 2.4 2.8 3.2 3.6 Reciprocal temperature, 1/T (1/K)·103

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Sens Sens Sens Sens

304 St.St. AJ9139, C49 Alloy 600 NX8608, C46 304 St.St. 2P4932, C52 304 St.St. 71635, C53

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10–5

0.0035

(b)

9.16 Effect of temperature on stainless steel in aggressive (a) and more typical (purer) water chemistry (b). Even in less aggressive environments, an increase in growth rate is observed as temperature is increased.

Understanding and mitigating ageing in nuclear power plants

10–6

20

Crack growth rate, mm/s

10–5

Temperature, T, (°C) 200 160 100 60

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Stress corrosion cracking of austenitic stainless steels

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9.4.5 Intergranular crack morphology, even in unsensitized materials While an intergranular (IG) crack morphology is not surprising for sensitized (Cr-depleted grain boundaries) stainless steels and nickel alloys, it might not be expected in unsensitized materials. Figure 9.17 shows the IG crack morphology in unsensitized, cold-worked Type 316L stainless steel. There is no evidence of a role of grain boundary segregants (e.g., S, P, B, N and C), since similar (or faster) IG growth rates are often observed in high purity alloys, and no significant effect of heavily impurity doped alloys was observed [58]. The preference for cracks to follow the grain boundary can be strongly promoted by grain boundary chemistry (e.g., Cr depletion or Si enrichment), but is also promoted by preferential deformation under loading in the grain boundary, which is a common explanation for intergranular SCC in stainless steels with no measurable chemical difference in the grain boundary (especially Cr depletion). One measure of the preference for cracking in the grain boundary is the tendency for the crack morphology to shift from IG to

9.17 Intergranular SCC morphology in unsensitized, 20% cold-worked Type 316L stainless steel tested in 288 °C BWR water.

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transgranular (TG) with an increase in DK and frequency, the applied strain rate, or even Kmax in tests without cycling. In heavily sensitized stainless steel in aggressive (oxidizing and impure) water, IG cracking can be sustained to relatively high DK or frequency, whereas very resistant materials in unaggressive environments can shift to TG morphology even at very low DK or frequency. Indeed, materials that are very resistant to SCC may exhibit TG cracking without cycling. No fundamental difference should be attributed to a TG vs. an IG crack path – an IG morphology reflects a chemical or strain rate preference between the grain boundary vs. slip planes or twin boundaries. The tendency to dismiss TG observations originates primarily from the early use of highly accelerated tests, e.g., employing slow strain rate or cyclic loading.

9.4.6 Grain boundary cr depletion vs. grain boundary carbides Sensitization resulting from grain boundary chromium (Cr) depletion occurs in stainless steels and nickel alloys because various types of Cr carbides form in the temperature range of ~500–750 °C. Chromium carbide formation ties up significant amounts of Cr in and near the grain boundaries, thus rendering the matrix in the boundary poor in Cr, relative to the bulk composition, and susceptible to corrosion and SCC. The most common Cr carbide in stainless steel is Cr23C6, although other forms exist. Carbon (C) atoms diffuse much faster than Cr atoms, and the growth of the carbide produces a Cr depletion profile near the grain boundary. Since grain boundary diffusion is faster than matrix diffusion, the entire grain boundary is depleted in Cr, even though the carbides are discontinuous. Depending on the carbide formed, it becomes stable below about 1000 °C, but the nucleation rate is fastest at a peak temperature between about 650 and 750 °C (depending on composition of the steel), and is more sluggish at higher or lower temperatures. Below about 550 °C, nucleation is sluggish, although in cold-worked stainless steels, carbides can form at temperatures as low as 400 °C in a few hours. Grain boundaries are the preferred sites for nucleation, although in cold-worked materials extensive intragranular nucleation occurs. Cr borides can also form and cause sensitization, usually in materials with low C or elevated B. Time-temperature-sensitization (TTS) diagrams are used to show the time to nucleation vs. temperature. Once the carbide nucleates, it can grow over a wide temperature range. Two key factors are the Cr activity at the carbide interface, and the Cr diffusion kinetics. Above 900–1000 °C, the Cr activity at the carbide interface is higher than the Cr activity in the matrix, and the carbide dissolves. At 700–800 °C, the Cr activity at the carbide boundary is somewhat below the matrix level, and the level of Cr depletion in the grain boundary is moderate (e.g.,

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12–15% Cr). The Cr depletion profile is broad because of the more rapid diffusion. At lower temperatures (e.g., 500 °C), the Cr activity at the carbide interface is very low, so the grain boundary Cr concentration is lower (e.g., 5–8% Cr) and the Cr profile is narrow. At LWR operating temperatures of ~ 300 °C, carbides continue to grow, but very slowly – this is called low temperature sensitization. For many stainless steels, a short high temperature (e.g., 1050 °C for 30 min) heat treatment dissolves the Cr carbides; this is called a solution anneal. When all of the free carbon is consumed, carbides stop growing but the Cr continues to diffuse, eventually eliminating the Cr depletion profile – a process called healing. This is difficult to achieve in stainless steels because of its lower Cr diffusivity, but is used commercially in nickel alloys (e.g., in thermally treated steam generator tubing) to create grain boundary carbides without Cr depletion. Sensitization is often measured using standard corrosion or electrochemical tests, such as the EPR test (electrochemical potentiokinetic repassivation). The EPR test reflects the area (width) along the grain boundary where the Cr concentration is below ~15% Cr–it does not measure the minimum Cr level in the grain boundary, which controls SCC susceptibility [59, 60]. Thus, EPR data can be very misleading, because very high readings can result from higher temperature sensitization heat treatments that have very wide but relatively shallow Cr depletion, while very low readings can result from low temperature treatments that give narrow but deep Cr depletion profiles. Sensitization is mitigated by controlling the alloy composition and restricting exposure to temperatures where it occurs most readily. Low C (L-grade) alloys, such as types 304L and 316L stainless steels, are common, and the presence of ~2.5% Mo in 316(L) further inhibits carbide nucleation. Current melting practice typically results in a 0.04–0.05% C steel, down from the 0.07–0.08% C range, which was common 40 years ago. The specification for L-grade steels is <0.03% C, but in practice a C content of 0.015–0.020% is common. The addition of titanium (Ti) or niobium (Nb) (in types 321 and 347 stainless steel, respectively) helps to reduce or eliminate sensitization by reacting with the free C in TiC or NbC carbides, which are more stable and form at higher temperatures than Cr carbides, and limit the free carbon available. These MC carbides are primarily intragranular and, because they form at higher temperature, do not result in grain boundary Cr depletion. However, during welding, the area adjacent to the weld fusion line heats above ~ 1200 °C, and MC carbides can partially dissolve. Because cooling to 800 °C in welds is rapid compared to solidification in an ingot, the MC carbides do not have a chance to tie up the C, and the cooling rate below 800 °C can permit the formation of grain boundary Cr carbides and Cr depletion. IG corrosion can occur in the heat affected zone and is known as knife line attack. IG SCC will obviously also be accelerated.

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Historically, SCC in nickel alloys has been considered to be fundamentally different in PWRs than in BWRs because grain boundary carbides were advantageous in PWRs (e.g., in steam generator tubing and other Alloy 600 components) but detrimental in BWRs (e.g., furnace and weld sensitization in stainless steels and Alloy 600). However, controlled studies on stainless steels with grain boundary carbides, but without Cr depletion, show that the carbides were beneficial in BWR chemistries (Fig. 9.18). Cr depletion is the detrimental factor, primarily under oxidizing (external) water chemistry conditions where the crack tip chemistry/pH is aggressive; sensitization is a much smaller factor in PWR or BWR environments where the materials are at low corrosion potential.

9.5

Stress corrosion cracking (SCC) dependencies – cold work, stress intensity factor and irradiation

9.5.1 Cold work from bulk deformation, surface cold work, weld residual strain, etc. Components that had bulk cold work cracked relatively early in BWR service, and subsequent designs restricted cold work to less than a few 1

0.4 Comparison of 20% CW SSs 27.5 MPs√m, 2000 ppb O2, Pure water

0.3

Crack length, mm

0.8

0.2

0.7

Alloy 600 2 ¥ 10–7 mm/s

Last 60% of test at 200 ppb O2

0.6

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347 SS 2.3 ¥ 10–7 mm/s

0.5 0.4

–0.1 304L SS 1.5 ¥ 10–7 mm/s

316L SS 0.3 1.6 ¥ 10–7 mm/s

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304 GB Carbides 2 ¥ 10–8 mm/s

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–0.4 –0.5

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9.18 Crack length vs. time in 288 °C water for 20% cold-worked stainless steels (and Alloy 600) with and without grain boundary carbides. Grain boundary carbides were formed at 621 °C for 24 h, then Cr depletion was eliminated by an equilibration heat treatment at 950 °C for 5 h, followed by water quenching.

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Conductivity, mS/cm or potential, Vshe

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271

percent. However, surface cold work from machining, grinding and related processes is more difficult to control or quantify, and has been responsible for accelerated initiation in many LWR components. Importantly, shrinkage from weld solidification produces residual strains as well as residual stresses near the weld. The development of electron backscattered diffraction (EBSD) permitted measurement of residual strains on a very fine scale [61–64]; EBSD was extensively calibrated using both tensile and double-cone compression specimens of many materials. It was discovered that peak weld residual strains of 20–30% (equivalent room temperature tensile strain) are often present near welds. These peak values occur at the root of the weld (e.g., at the inside diameter of welded pipes) and decrease toward zero within a few millimeters from the fusion line (Fig. 9.19). Lower residual strains were observed near the final (outside) welding passes, and they also tapered towards zero away from the fusion line. The importance of the weld residual strains on SCC was directly measured on weld heat affected zone aligned specimens (e.g., Fig. 9.20), which exhibited high crack growth rates consistent with bulk cold-worked specimens. The historical inattention to weld residual strain in sensitized stainless steel piping

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Various BWR stainless steel weld HAZs

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9.19 Equivalent room temperature tensile strain in the heat affected zone adjacent to the weld fusion line determined from electron backscattered diffraction. The peak strain always occurs near the fusion line and at the root of the weld, with peak values of 20–30% strain being common both in stainless steel and Alloy 600 HAZs.

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Understanding and mitigating ageing in nuclear power plants 0.4

11.9

0.35 0.3

11.7 3.2 ¥ 10–7 mm/s

Corrosion potential of Pt

0.25

To 200 ppb O2 @565h

11.6 11.5 11.4 11.3 11.2

Corrosion potential of CT

0.2 0.15 0.1

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Crack length, mm

2.2 ¥ 10–7 mm/s

0.5T CT of HAZ Aligned 348 Weld 288C, 2 ppm O2, 27.5 MPa√m

0 800

9.20 SCC response for a Nb-stabilized, high N-bar stainless steel specimen aligned along the weld HAZ. This and similar observations confirm the detrimental role of shrinkage strains adjacent to weld HAZs.

is not surprising because the peak in sensitization in Type 304/316 stainless steel does not occur near the weld fusion line, but typically 5–15 mm away where the thermal profiles favor nucleation and growth of grain boundary carbides. At the 5–15 mm distance, the residual strain decreases to quite low levels, and ignoring its contribution represents a small oversight. As SCC has developed in unsensitized stainless steel weld heat affected zones, the role of residual strain has become very important, and indeed most cracks occur close to the weld fusion line. The effect of bulk cold work on SCC has been studied in many materials (Fig. 9.21) [7–22], and is best characterized in terms of a basic effect of yield strength on SCC growth rate. Factors that increase the yield strength (e.g., cold work, precipitation hardening and irradiation) tend to increase SCC growth rates, all other factors being equal. The crack growth rates are accelerated at high corrosion potential, but remain relatively high in deaerated or in PWR water (Figs 9.4, 9.5 and 9.12). Indeed, Fig. 9.4 shows that lower crack growth rates can be achieved at low potential on sensitized materials than on cold-worked materials. While the incidence of SCC in stainless steel welds has not been as high in PWRs as in BWRs, the growth rates observed in the laboratory at low corrosion potential suggest that it will become more evident over time. High yield strength materials (which can develop from high cold work,

© Woodhead Publishing Limited, 2010

Stress corrosion cracking of austenitic stainless steels 1.e–07

Crack growth rate, mm/s

Unsensitized 304, 304L & 316L SS & A600 288 °C High purity water, 95 or 1580 ppb H2 CT Tests at 27.5 – 30 MPs√m  = High martensite SS Predicted  = Alloy 600 Response

273

Very high martensite

2 Sensitized Points for Comparison

1.e–08

Very low or no martensite

Annealed   Cold Worked at –55 °C or +140 °C 1.e–09

0

1.E–06

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300 400 500 Yield strength, MPa (a)

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Crack growth rate, mm/s

Unsensitized 304, 304L & 316L SS & A600 288 °C High purity water, 2000 ppb O2 CT Tests at 27.5 – 30 MPa√m  = High martensite SS  = Alloy 600 Predicted

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Response

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9.21 Effect of yield strength and martensite content on the SCC growth rate on stainless steel and Alloy 600 in 288 °C, high purity water (ª0.06 mS/cm outlet) containing 95 or 1580 ppb H2 (a) or 2 ppm O2 (b).

neutron irradiation, etc.) can exhibit both high growth rates and a more limited effect of corrosion potential at high stress intensity factor or under ‘gentle’ cyclic loading conditions (Fig. 9.22), which typically enhance the SCC growth rate only very slightly. In some high yield strength materials, rapid crack advance is observed during the reloading portion of the waveform, with rates

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Understanding and mitigating ageing in nuclear power plants 0.3 Outlet conductivity

13.75

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2.2 ¥ 10–7 mm/s

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To 1.58 ppm H2 @3396h

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13.65

To R = 0.7, 0.001 Hz + 85,400s hold @1808h To 2000 ppb O2 @3037h

0.1 13.7

–0.3 –0.4 CT potential

c157 – 0.5T CT of 316LSS, 50%CW 32 – 34.5 MPa√m, Pure water

13.45

PT potential

K slowly rises from 32 to 34.5 MPa√m in this graph 13.4 3100

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3150 3200

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3300 3350 3400 Time, h (a)

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16.55 16.5 16.45 5500

Average rate: 5 ¥ 10–7 mm/s

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c157 – 0.5T CT of 316LSS, 50%CW 45–49 MPa√m, Pure water

To To To 85

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9.22 Crack length vs. time for a 0.5TCT specimen of unsensitized 316L stainless steel ‘cool’ worked to 50% showing the effect of gentle unloading cycles on environmental crack advance on stainless steel whose yield strength is elevated by cold work.

of 10–4 mm/s and higher. Whether this ‘gentle’ cyclic phenomenon at a load ratio, R = 0.7 (R is Kmin/Kmax) would be observed under higher frequency, lower amplitude conditions still needs to be evaluated. High growth rates are

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also observed as the stress intensity factor increases at constant K. Figure 9.23 shows the very high growth rates and, while there is some reduction as the corrosion potential is decreased, the rate rapidly increases as stress intensity factor increases.

3.4 ¥ 10–7 mm/s

4 ¥ 10–7 mm/s Pt potential

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4 ¥ 10 mm/s

1600

–0.2 –0.3 –0.4

0.3 –4

3 ¥ 10–5 mm/s

3900

–0.1

0.4

1.5 ¥ 10–6 mm/s

3860 3880 Time, h (b)

0

–0.6 1800

c181 – 0.5T CT of 316L 50%CW 2000 ppb O2, Pure water

61 MPa√m

Crack length, mm

19

1500

CT potential

To constant load & 95 ppb H2 @3851h

19.5

53 MPa√m To 2 ppm O2 @2389h NobleChem @1837h To 0.001 Hz + 9000s hold @3037h

Pt potential

1300 1400 Test time, h (a)

3920

0.2 Specimen failed @3927h Actual K at failure = 124 MPa√m

1000

0.1

–0.5

CT potential 12 900

0.2

0.1 0 –0.1 –0.2 –0.3 –0.4 –0.5

–0.6 3940

9.23 Crack length vs. time for 0.5TCT specimens of unsensitized types 304L and 316L stainless steel ‘cool’ worked to 50% showing the effect of high yield strength and high stress intensity on environmental crack advance and fracture toughness.

© Woodhead Publishing Limited, 2010

Conductivity, mS/cm or potential, Vshe

13

1.1 ¥ 10–6 mm/s

Outlet conductivity

0.3

Conductivity, mS/cm or potential, Vshe

14

Stop cycling @1255h K = 45 MPa√m

15

To 6% H2 in Ar @829h To R = 0.7, 0.001 Hz + 9000s hold @890h

Crack length, mm

16

5.2 ¥ 10–6 mm/s

48 MPa√m

17

End of test @1758h K = 64 MPa√m

0.4 c182 – 0.5T CT of 304L ~50%CW 2000 ppb O2, Pure water

276

Understanding and mitigating ageing in nuclear power plants

At lower temperatures (<150 °C), martensite in stainless steel can have a significant effect on cracking. In high temperature water, the effect of martensite was evaluated by rolling or forging at –55 ºC where very high levels of deformation-induced martensite can form. However, Fig. 9.21 shows that there is no consequential difference in crack growth rate at a given yield strength in materials with high or low (or no) martensite, including in Alloy 600, where martensite does not form. To evaluate a possible role of hydrogen on SCC of stainless steels, H2 permeation was measured [7–22, 63–68]. Hydrogen exists in the coolant as a dissolved gas (H2), but dissociates to H0, adsorbs on the surface, then permeates the metal in atomic form. Hydrogen permeation is controlled by the coolant H2 fugacity, which is non-zero even in water containing no H2 (Fig. 9.24). Once hydrogen had permeated into the 4.6 mm ID tube, reducing the H2 fugacity in the water readily produced dissociation of H2 and permeation back into the water. Since the hydrogen permeation rate is high compared to the H2 generation rate from corrosion, radiolytic proton injection, or transmutation, it is very unlikely that a high H 2 fugacity (i.e., well above the coolant H2 fugacity) is generated in metals exposed to hot water. Measurements of high hydrogen concentration in metals [25, 69] are a reflection of hydrogen storage locations, e.g., in heavily cold-worked materials (as atoms), or at radiation-induced voids (as atoms on the surface and H2 in the void) – not of a high hydrogen fugacity that can have extraordinary effects on sub-critical crack advance. No difference in the growth rate of stainless steels was observed when changing from N2 deaerated water to various levels of H2 [13, 15–19]. Similarly, the effect of electrocatalytic species (e.g., Pt) on the surface has no accelerating effect on the crack growth rate in deaerated water [34–37]. Under dynamic strain conditions, such as corrosion fatigue, some detrimental effect of lower corrosion potential from dissolved H2 is observed on crack initiation [48, 49] and crack growth rates of stainless steels [50] as noted earlier.

9.5.2 Irradiation It is not the intent of this chapter to discuss the details of irradiation assisted SCC, which are presented in references [7, 24–28]. Irradiation assisted is an apt term – the primary effects of irradiation on SCC (Figs 9.1 and 9.5a) are associated with radiation hardening, radiation induced segregation (especially Cr depletion and Si enrichment), radiation creep relaxation and radiolysis (insofar as it elevates the corrosion potential of the metal). Figure 9.5a highlights the effect of radiation hardening – the large-triangle data point at low corrosion potential shows that the growth rate is similar to cold-worked stainless steel of similar yield strength. The large-triangle

© Woodhead Publishing Limited, 2010

Stress corrosion cracking of austenitic stainless steels 10000

0.2 Outlet conductivity To 1580 ppb H2 @2459h

Closed end tube of 304LSS 95 ppb H2, Pure water

8000

4000 3000 2000

0

85 microns/h

–0.1 –0.2 Repeat data 1000h later –0.3 –0.4 CT potential

30 microns/h

Pt potential

–0.6 –0.7

1000 0 2350

2370

2390

2410

2430

2450 2470 Time, h (a)

2490 2510

–0.8 2530 2550

0.3

10000 Outlet conductivity

9000

0.2

Pt potential

8000

0.1

CT potential

7000 H2 pressure, microns

–0.5

315 °C water –90 microns/h

6000 5000

0 –0.1

304LSS closed end tube 2000 ppb O2, Pure water

–0.2 –0.3

4000 3000

–0.4

H2 Pressure ¥ 10

2000

–0.5 –0.6

1000 0 4090

4140

4190

Time, h (b)

4240

4290

Conductivity, mS/cm or potential, Vshe

5000

To 95 ppb H2 @1332h

H2 pressure, microns

7000

0.1 Conductivity, mS/cm or potential, Vshe

9000

6000

277

–0.7 4340

9.24 Hydrogen permeation vs. time and coolant H2 fugacity in unsensitized Type 304L stainless steel.

data at high potential shows the effect of both hardening and segregation, in that the growth rate is higher than for the unirradiated data that is either cold worked or sensitized. The Cr depletion profile from irradiation is a few nanometers (nm) wide, ~100¥ narrower than exists from thermal sensitization. But a wide range of Cr depletion widths were evaluated using complex heat

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Understanding and mitigating ageing in nuclear power plants

treatments and it was shown that SCC depends on the minimum Cr level, not the width [59, 60] of the Cr profile near grain boundaries. Radiation creep relaxation is often beneficial because components under constant displacement loading experience load relaxation vs. time, so that if cracks do not nucleate early in life (nor are reloaded later in life), the probability of SCC can decrease with neutron fluence because the load drops below a third of its original value within a few displacements per atom (dpa). Above ~ 320 °C and above a few dpa, radiation swelling can occur. While the PWR core outlet temperature is only ~323 °C, gamma heating in stainless steel components such as baffle plates and baffle former bolts can produce an elevation in their internal temperature by about 30 °C. Because core baffle plates are typically fabricated from annealed stainless steel, while the baffle former bolts are typically fabricated from ~15% cold-worked stainless steel, and because swelling is delayed in cold-worked stainless steel, differential swelling can occur preferentially in the plates. Early in life (within ~5 dpa) the bolts undergo radiation-induced relaxation by > 80%, but subsequently differential swelling in the plates can reload the bolts and promote SCC – the balance between reloading and relaxation is complicated.

9.5.3 Stress intensity factor An understanding of the effects of the (crack tip) stress intensity factor (K) is key both to rationalize the historical development of cracks and to disposition their future growth. K is proportional to stress (s) times the square root of crack length (a) (K = B s÷a, where B is a factor that accounts for crack and component geometry). Quantifying the effects of K on SCC growth rates has had a long and troubled history. Simple transgranular (TG) fatigue precracking followed by static (esp. bolt) loading is a very poor approach for quantifying the effects of K, as well as most other parameters. Decreasing K in large steps is also likely to produce crack arrest in cases where it would not otherwise occur. Modern techniques [20, 21, 38, 70] employ transitioning to IG SCC, and re-transitioning at each K level or employing a –dK/da technique where the K decreases only as crack growth occurs. Using such techniques (Figs 9.6, 9.25 and 9.26), a consistent K dependency can be measured, with no evidence of crack arrest (KISCC), at least down to ~5 MPa÷m. A similar K dependency has been observed for irradiated stainless steel [66], and for nickel alloys in BWR water [7, 21, 65]. Effects of K can easily be mis-measured and misinterpreted. In constant displacement specimens there is a strong tendency toward low or no growth rates at low K (i.e., much lower than would be observed with good techniques), while reasonable growth rates are often measured at high K. The net effect is to increase the K dependency because the lower K crack growth rate data are biased low. Somewhat conversely, specimens with no side grooves tend to

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Stress corrosion cracking of austenitic stainless steels

279

711 mm SCH80 Pipe-ranganath analysis Buchalet/Bamford Method Circumferential crack 55

Upper limit

Mean

33

+ 30

Residual stress Ksi

22

142

11

Lower limit

0

40

Stress intensity, MPa√m

35

5

10

Uper limit

Lower limit

+8.1 1.423 Crack depth

Pressure stress 41 MPa deadweight stress   7 MPa Thermal stress 41 MPa

15 20 25 Crack depth, mm (a)

30

35

15

304SS, 2-sided weld, EPRo = 5 C/cm2 0.3, 0.2, 0.1 mS/cm, +0.175 Vshe Symmetrical sres profile 3 ¥ 1019 n/cm2-y

30 25 20 15

sres with +69 MPa above nominal (shroud.mj3)

0.3 mS/cm

10

0.2 0.1

5 dK/da for 0.3 mS/cm

10 5 0 0

dK/da, MPa√m per mm

Sress intensity, MPa√m

44

0 10

20 Crack depth, mm (b)

30

9.25 Residual stress through wall (a, inset diagram) and resultant stress intensity factor vs. through-wall crack depth (a) for a typical large diameter pipe weld, with all welding passes made from the outside diameter. (b) Stress intensity factor vs. crack depth for a BWR core shroud that is welded by alternate passes from the inside and the outside diameter. © Woodhead Publishing Limited, 2010

Understanding and mitigating ageing in nuclear power plants

58.1

Stress intensity factor

58 6020

11.88 11.83

To constant K @816h At 19.8 MPa√m @2146h

Crack length, mm

11.93

11.78 11.73

40 35 30

6025

25

6030

6035 6040 Test time, h (a)

6045

6050

20 6055

20

12.03 11.98

45

1.5 ¥ 10–7 mm/s

58.05

50

5.7 ¥ 10–8 mm/s

1168 11.63 2200

2700

Stress intensity factor 3.4 ¥ 10–8 mm/s

2.5 ¥ 10–8 mm/s

18 16 14 12 10 8 6

c209 – 0.5TCT 316L 20% CW 20 MPa√m, 2000 ppm O2, Pure water 3200

3700 Test time, h (b)

4200

4700

4

Stress intensity factor, MPa√m

58.15

1.0 ¥ 10–5 mm/s

End Varying-K at @6044h

58.2

To varying-K at @6029h 86.6 MPa√m per mm

58.25

55 3.5 ¥ 10–7 mm/s

Unload @3006h

Crack length, mm

58.3

To 27.5 MPa√m, R = 0.5, 0.001 Hz @5897h

58.4 58.35

60

c317 – 2TCT of 316L +20% RA at 140C 27.5 MPa√m, 2 ppm O2, 100 ppb SO4

Stress intensity factor, MPa√m

58.45

Re-initiated varying-K drop at 21.7 MPa√m per mm @2452h

280

2 0 5200

9.26 Effect of +dK/da and –dK/da at values relevant to plant components. Growth rates up to 1000¥ faster than constant K conditions have been observed under +dK/da conditions. By contrast, under representative –dK/da conditions, the crack growth rate is sustained and the observed rates are similar to those obtained at constant K.

© Woodhead Publishing Limited, 2010

Stress corrosion cracking of austenitic stainless steels

281

exhibit extensive crack branching above moderate K, and this decreases the local K at each branched crack tip. This biases the higher K data downward, and yields a lower K dependency. In general, the agglomeration of many sets of data that involve different heats and test techniques tends to dilute any trend, including K, giving a shallower K dependency. The issue of crack branching can be important in the analysis of plant components. As the K increases, extensive branching may occur, and this can lead to a significant deviation from a fixed power-law dependency to a plateau-like behavior (fixed growth rate vs. K). However, if there are reasons why the crack plane is constrained – e.g., the narrow plane where the weld residual strain is high, adjacent to the fusion line – then plateau behavior may not exist. Another important consideration is the change in K as the crack depth increases. In plant components, K changes almost completely because of crack advance since, e.g., changes in pressurization are small and represent a small fraction of the total stress. Unfortunately, most laboratory tests have been performed at constant K or constant load, and make no effort to simulate the actual profile in K vs. crack depth. Tests to evaluate +dK/da have shown elevated growth rates [71, 72] because there is a positive feedback mechanism that causes an increase in growth rate (da/dt), which in turn causes an increase in K vs. time (dK/dt = da/dt · dK/da). Growth rates up to 1000¥ higher than constant K data have been observed under realistic +dK/da conditions. Under –dK/da conditions, the possibility that cracks might arrest was an attractive prospect, but –dK/da data have shown sustained cracking in ~25 cases, some involving a more than twofold reduction in K. Stainless steels and nickel alloys show the same response under dK/da conditions.

9.6

Stress corrosion cracking (SCC) dependencies – miscellaneous

9.6.1 General corrosion rate The general corrosion rates of stainless steels and nickel alloys (i.e., Alloy 600 and Alloys 182/82 weld metals) are very similar (Fig. 9.27). Corrosion rate can be important for many reasons, including: ∑

the loss of a thin polished or compressive layer or a passivation film that was designed to mitigate SCC initiation; ∑ release of corrosion product into the water; ∑ altering the rate of consumption of oxidants as they diffuse into the crack; ∑ shifting the corrosion potential when the oxidant concentration is low (a higher corrosion rate will consume more O2 and cause the corrosion potential to drop, e.g., at 10 ppb vs. 1 ppb, Fig. 9.14); and

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Understanding and mitigating ageing in nuclear power plants

1000

Electrode potential, mV(she)

800

304 SS polarization measurement in deaerated, high purity water (< 2 ppb O2) ast 288 °C

600

After preoxidation in 200 ppb O2 for 3 weeks

400

After preoxidation in 150 ppb H2 for 3 weeks

200 After preoxidation in 1 ppm H2O2 for 3 weeks

0

After preoxidation in deaerated water for 3 weeks

–200 –400 –600 0.01

1000

Electrode potential, mV(she)

800

0.1

1 10 Current density, mA/cm2 (a)

100

1000

Alloy 600 polarization measurement in deaerated, high purity water (< 2 ppb O2) at 288 °C and 150 cc/min

600 400 After preoxidation in deaerated water for 3 weeks

200 After preoxidation in 200 ppb O2 for 3 weeks

0

After preoxidation in 150 ppb H2 for 3 weeks

–200 –400 –600 0.01

0.1

1 10 Current density, mA/cm2 (b)

100

1000

9.27 Corrosion rate vs. prior exposure for stainless steel (a) and Alloy 600 (b) in 288 °C pure water. Exposure in a given dissolved gas chemistry was for several weeks, then a rapid change was made to deaerated water to make the polarization measurements.

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Stress corrosion cracking of austenitic stainless steels

∑

283

enhancing the environmental effect, especially on surfaces, e.g., for corrosion fatigue.

While seemingly contradictory, the corrosion rate at ‘high corrosion potential’ is lower than at ‘low corrosion potential’ because the passive oxide films that form at low potential are not as protective. The data shown in Fig. 9.27 are short-term data obtained by exposing the material in a specific environment for ~2 weeks, then changing to Ar-deaerated water so that O2 and H2 do not interfere with electrochemical measurements of the metal corrosion reaction. Long-term corrosion tests show lower corrosion rates because the rate appears to decay parabolically with time. The similarity in general corrosion behavior means that the corrosion potential response vs. O2 and H2 is very similar among these alloys, as is the short crack and long crack chemistry development. The higher corrosion rate at low potential may explain why the corrosion fatigue crack initiation life of stainless steels and nickel alloys is reduced at low vs. high corrosion potential [48, 49].

9.6.2 Grain boundary silicon At sufficiently high concentrations in the grain boundary, Si can enhance SCC because it oxidizes to SiO2 at all relevant potentials and is soluble in high temperature water. The primary concern is for irradiated materials, where Si can segregate to levels above 20 at%, and can cause high crack growth rates, and no effect of corrosion potential or stress intensity factor on crack growth rate. This has been observed in stainless steels (Fig. 9.28) and nickel alloys with elevated bulk Si levels [73].

9.6.3 Effects of environment on fracture Environmental effects on fracture is an emerging area that can be subdivided into J-R tearing resistance and fracture toughness (e.g., KIC-Env) [74, 75]. Both stainless steels and nickel alloys can exhibit environmental effects on fracture in light water reactor environments, and very large reductions have been observed at lower temperatures in Alloy 82 weld metal (Fig. 9.29). Exposure to high temperature water allows H2 gas molecules to dissociate and permeate as hydrogen atoms throughout the material, although the effects of hydrogen are generally more pronounced after cooling to <130 °C. Pre-saturation of the metal with hydrogen by exposure to high temperature water can also create the opportunity for rapid failure, as distinct from slower tearing fracture, where (without pre-exposure) there is a rate-limiting step associated with hydrogen dissociation and permeation into the metal. To date, several laboratories have observed sudden fracture while performing SCC testing, with KIC values as low as ~75 MPa÷m (Fig. 9.30).

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Understanding and mitigating ageing in nuclear power plants

16.1 15.7 Pt potential

Outlet conductivity

12.5

200

250

300

350 400 450 Test time, h (a)

500

550

600

c270 – 0.5TCT of A182 +3%Si + 26%RA 27.5 MPa√m, 2 ppm O2, Pure water

11.6

3.4 ¥ 10–7 mm/s

Pt potential

11.5

CT potential

11.4 700

750

–0.3 –0.4 –0.5

–0.6 650

0.5 0.4

0.1

800

850

0

To 2000 ppb O2 @997h

11.7

–0.2

0.2 Outlet conductivity

11.8

–0.1

0.3

To 95 ppb H2 @633h

Crack length, mm

11.9

To constant K @277h

12.1 12

0

2.2 ¥ 10–6 mm/s

4 ¥ 10–6 mm/s

12.1 150

0.1

900 950 1000 Test time, h (b)

–0.1 –0.2 –0.3 –0.4 –0.5 1050 1100 1150

Conductivity, mS/cm or potential, Vshe

12.9

To 6% H2 in Ar @312h

13.3

To constant K @226h

13.7

To 0.001 Hz + 9000s hold @167h

Crack length, mm

14.1

1.0 ¥ 10–6 mm/s

Conductivity, mS/cm or potential, Vshe

0.2

CT potential

14.9 14.5

0.3

Varying-K drop to 15 MPa√m @630h

15.3

0.4

c183 – 0.5T CT of 5Si-SS 15%CW 29.7 MPa√m, 2000 ppb O2, Pure water

Being Varying-K decrease @492h

284

–0.6 1200

9.28 Crack length vs. time for a 0.5TCT specimen of an unsensitized model ‘stainless steel’ containing 5% Si ‘cool’ worked to 22% reduction in area (a) and a wrought Alloy 182 with 3% Si with 26% RA (b).

9.6.4 Crack initiation This chapter has emphasized SCC growth, partly because it can be quantified more accurately than crack initiation, and partly because the study and use of initiation data poses a variety of conceptual and pragmatic concerns. These

© Woodhead Publishing Limited, 2010

Stress corrosion cracking of austenitic stainless steels T= 340

Air (A2a,b) 29 (150 cc/kg)

54 °C Water 54 °C Water

15

68

150 cc/kg

338 °C Water

Welded in 100%Ar

285

Alloy 82H (T-S) Weld A2a,b

245

327

Air (C4a) 2 (150 cc/kg)

54 °C Water

54 °C Water 50 13 93 °C Water

3 (150 cc/kg)

Welded in 100%Ar

Weld C4a

121 °C Water 150 24

454

Air (C4c) 54 °C Water

4 (150 cc/kg)

54 °C Water

50 36

54 °C Water

150 cc/kg

93 °C Water 150 23 149 °C Water

150 cc/kg

338 °C Water

150 cc/kg

144 Welded in 100%Ar

Weld C4c

230 310

26

Air (C2) 54 °C Water

4 (150 cc/kg)

54 °C Water

4 (50 cc/kg) 0

Welded in Ar-5%H2

500 JIC, kJ/m2

Weld C2 1000

9.29 J-R data of Brown and Mills on various Ni alloy 82H weld metal showing a large reduction in fracture resistance in water compared to air [76].

include challenges associated with the absence of continuous monitoring (even of failure); use of sufficient replicates and surface area so that the results are statistically significant; and uncertainty in defining the average, distribution or extremes of surface condition in plant components that is key to assuring the applicability and relevance of the data. Conceptually, there is no agreement on the meaning of crack initiation (e.g., a depth of 1 mm, of 50 mm, of 1 grain size, when conveniently detectable by ultrasonic inspection, when long crack growth rate response is achieved, etc.). Most examinations of cracked plant components show significant surface cold work, and indeed

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Understanding and mitigating ageing in nuclear power plants

most surfaces show heavy damage when examined in detail (Fig. 9.31a), and how well understood or reproducible that the surfaces are has a strong bearing on the initiation behavior.

17.6

17 16.8 16.6 16.4

c192 – 0.5T CT of 304L 70%CW ~41 MPa√m, 2000 ppb O2, Pure water 1.6 ¥ 10–5 mm/s

Outlet conductivity

2 ¥ 10–6 mm/s

16.2 6715

0.25 End of Test @6749h

17.2

0.3

CT potential Start Varying-K at 43.3 MPa/m per mm @6687h

Crack length, mm

17.4

0.35

Actual KIC = 75.5 MPa√m

Pt potential

6720

6725

6730 6735 Test time, h (a)

6740

0.2 0.15 0.1 0.05

Conductivity, mS/cm or potential, Vshe

17.8

0 6750

6745

18.3

0.4 Pt potential CT potential

14.3

K rising because test is controlled by specimen c275

13.3 1.7 ¥ 10 mm/s

12.3 3900

This tandem specimen grew rapidly over 2 week holiday and failed @5914h

15.3

0.2

c276 – 0.5TCT A182 Weld Metal, 288 °C 27.5 MPa√m, 2 ppm O2, Pure water

–7

To Constant K @4515h

16.3

To 30 ppb SO4 @2961h To R = 0.8, 0.001 Hz + 85,400s hold @3900h

Crack length, mm

17.3

2.4 ¥ 10–7 mm/s

0

–0.2

–0.4 KIC ~ 90 MPa√m

Conductivity, mS/cm or potential, Vshe

Outlet conductivity

–0.6 4400

4900 Test time, h (b)

5400

5900

9.30 Examples of sudden failure of (a) cold worked stainless steel and (b) Alloy 182 weld metal tested in 288 °C water at increasing K until failure occurred. The load and crack depth at failure are very well defined, and the resulting KIC (not necessarily obtained valid conditions) is relatively low [75]. The fracture surfaces are illustrated in (c) and (d) and show the demarcation between scc and fast fracture. © Woodhead Publishing Limited, 2010

Stress corrosion cracking of austenitic stainless steels

287

(c)

(d)

9.30 Continued

Crack initiation can be tracked to relatively small dimensions, both on smooth specimens and specimens with defects such as a CT specimen with a controlled radius, blunt notch. The transition from short crack growth to long crack growth (Fig. 9.31b) generally occurs in the 20–50 mm range of

© Woodhead Publishing Limited, 2010

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Understanding and mitigating ageing in nuclear power plants

40.0 Transverse cross-section

Plastic strain (%)

30.0

20.0

10.0

0.0 0

Crack length from notch, microns

400

50

100 150 Distance from surface (microns) (a)

350

Sensitized 304 SS 25 mm CT C37 (1y) Kmax = 33 MPa√m, R = 0.7, 0.01 Hz 200 ppb 02, 0.5 mS/cm HCI

300

Cracking from blunt notch, wet machined with a 220 grit wheel

200

250

Blunt notch CT specimen 250 200

200 ppb O2 0.5 mS/cm HCl

150 100

6% H2 in Ar 0.5 mS/cm HCl

50

50 0 80

0 130 100

120

140

160

180 200 Time, h (b)

220

180 240

260

280

9.31 (a) Example of surface cold work present on many surfaces when section obliquely polished and evaluated by EBSD. (b) Transition from short crack to long crack growth in a carefully wet ground blunt notch CT specimen of sensitized stainless steel.

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Stress corrosion cracking of austenitic stainless steels

289

crack depth, with the growth rate being slower when the crack is shorter – apparently related to nucleation of isolated cracks and their coalescence to a single crack. Repeated interruption of slow strain rate specimens revealed a strong effect of surface preparation, temperature (Fig. 9.32a) and impurities (Fig. 9.32b). Many crack initiation tests use very aggressive loading, temperature, surface preparation and/or water chemistry conditions, sometimes so aggressive that

7

Pure water

5 4 55 mS/cm H2SO4 All

3

A40 Na2CO3

A20 H2SO4

A25,A22 HCl+H2SO4 3.3+0.53 ¥ 10–7/S A26, H2SO4

1

4

5

6 pH288°C (a)

7

Weld + 400 °C/10d 40 Weld + 500 °C/24h 600 °C/24h 600 °C/24h + shot peen Open symbols-% strain Closed symbols-% life

40

Initiation, % strain

NaCl, A39 NaHSO4 A27

Neutral

2

A30 Na2SO4

30

8

100

80

60 20 40 10

0 100

20

150

200 Temperature, °C (b)

250

0 300

9.32 Strain to crack initiation for sensitized stainless steel tested by slow strain rate at 3.3 ¥ 10–7 s–1 in 288 °C water and repeatedly interrupted to detect initiation.

© Woodhead Publishing Limited, 2010

Initiation, % of life

Crack initiation %e

6

0

A37 NaOH

10 mS/cm Impurity 304SS Weld + 500 °C/24h Cert 3.3 ¥ 10–7/S 200 ppb Oxygen A13

290

Understanding and mitigating ageing in nuclear power plants

they call into question the relevance of the data. For example, in BWR tests it is common to use graphite wool as a crevice agent, but the wool releases large amounts of impurities. U-bend and crevice bent beam specimens usually are interrupted for SCC evaluation, which makes the time to initiation ambiguous. Test designs that allow for controlled, active load and continuous monitoring are more attractive, such as the ‘Keno’ test which can test up to 150 specimens in one autoclave [76]. Failure of any specimen is revealed by a numbered indicator ball. Good statistical confidence can be obtained (Fig. 9.33). The effect of many variables on crack initiation and growth are similar, e.g., temperature, corrosion potential, water purity, irradiation, etc. This suggests that a similar mechanism may be responsible for the early stages of the formation of a mechanically distinct geometry that will tend to grow in preference to its surroundings and the later stages when a long crack exists. The sole reliance on a ‘paper thin’ layer whose characteristics and SCC behavior is difficult to quantify with confidence and relevance is not wise. While every effort should be made to control component design and surface characteristics to delay initiation, the concept of inherent resistance in the bulk material to crack advance is a safer focus.

324 MPa 324 MPa 47 Ksi 47 Ksi creviced

99

%

95 90 80 70 60 50 40 30

193 MPa 28 Ksi creviced

193 MPa 28 Ksi

Stress, 20X

20 Stress w/ crevice, 10X

10

Crevice @193 MPa, 6X

Crevice @324 MPa, 3X

5 3 2

Sensitized 304 SS 288 °C, 80 wppm O2, 0.87 mS/cm H2SO4

1

10

100 Time to failure, h

1000

9.33 Effect of stress and crevicing on the crack ‘initiation’ of sensitized stainless steel in actively loaded smooth specimens whose failure is detected on-line [77].

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Stress corrosion cracking of austenitic stainless steels

9.7

291

Mechanism of stress corrosion cracking (SCC)

Despite the many common characteristics across environment, temperature and material (e.g., Figs 9.1–9.6), SCC tends to be viewed narrowly in terms of specific observations, with the inherent implication that different mechanisms and dependencies are involved. However, the crack tip is deaerated and at low potential in LWR environments, and a well-behaved continuum exists for changes in corrosion potential, water purity, temperature, stress intensity factor, material and condition, cold work/yield strength, etc. (Figs 9.1–9.6). Thus, it is reasonable to propose that a common crack advance mechanism applies to all of these materials, with some special factors that must be accounted for in specific materials or environments.

9.7.1 Underlying crack advance mechanism and primary sub-processes When the protective oxide film is removed, engineering alloys are highly reactive in aqueous environments, not unlike sodium in water (whose oxide/ hydroxide, NaOH, is highly soluble and therefore non-protective). Thus, most models of SCC advance (or, more generically, environmentally assisted cracking, EAC, which encompasses the spectrum of corrosion fatigue ´ slow strain rate ´ constant load loading) in hot water involve strain induced disruption of the protective oxide film at the crack tip. Following such disruption, rapid metal oxidation and subsequent repassivation proceeds, a process that can be quantitatively linked to crack advance (Fig. 9.34) [7–9]. Even in hydrogen embrittlement or film cleavage models of crack advance, oxide rupture plays an essential role, e.g., in enhancing the kinetics of (atomic) H0 formation and avoiding the restricted kinetics of H0 transport through oxide films. Hydrogen embrittlement is sometimes invoked as a primary mechanism of crack advance in high temperature water, but there is a significant array of contrary evidence [13–21]. Hydrogen in the metal has been shown to be proportional to the H2 fugacity in the coolant, and hydrogen permeation into and back out of closed (permeation) tubes is also controlled by the square root of the coolant H2 fugacity. All iron and nickel-base structural materials show a similar effect of corrosion potential (e.g., Figs 9.1–9.5), and their growth rates rise dramatically as the oxidant level rises (usually with the H2 level decreasing), whether the crack tip pH moves acidic or basic. Efforts to achieve even small amounts of crack growth under static load in 300 °C gaseous H2 requires high pressure, and generally only achieves a modest decrease in fracture toughness (vs. growth at a constant and reasonable K). For temperatures below 150 °C, the role of hydrogen embrittlement can become much more pronounced, and crack growth rates in water and gaseous

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VS J H2O

VT

Ê ˆ Penetration distance Á = M · Qf ˜ Ë nrF ¯

Oxidation charge density, Qf

J M+ dc

Qf

Average crack tip penetration

Qf

Q Q VT = M · f = M · f · e nrF t f nrF e f Crack side penetration VS = M · ip nrF Oxide nucleation and growth

Bare-surface dissolution

Time

tf Oxide rupture e tf = f e

9.34 Schematic of the corrosion (oxidation charge density) vs. time for metal that is unstrained (lower curves, ‘crack side penetration’) or strained to produce damage to the protective oxide (upper curves), with subsequent repassivation. While shown for dissolution the model equally applies to chemical oxidation in which the anodes and cathodes are separated.

H2 can be similar. Of course, hydrogen might have a role in enhancing dislocation mobility at 300 °C, but dislocations are more mobile at 300 °C, and hydrogen is ubiquitous and permeates in atomic form readily in metals exposed to high temperature water. In high temperature water environments, the slip–film rupture–oxidation (S/FR/O) model of crack advance (Fig. 9.34) is the most widely accepted model, and it has been extensively developed by Ford and Andresen for light water reactor environments [7–9]. It ascribes fundamental importance to dynamic strain that occurs at the crack tip (not, e.g., to stress intensity factor) and to the kinetics of film repair, which are related to local material chemistry and crack tip solution chemistry. The nature of passivity in hot water is different than at lower temperatures, and the oxides are relatively thick (0.1–1 mm), even on highly corrosion resistant materials such as Alloy

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690 (30% Cr) and Alloy 22 (22% Cr, 13% Mo). Thus, while the crack tip solution chemistry and local (e.g., grain boundary) material composition are very important and predictable, no theoretical basis has been developed to predict the resulting repassivation response. Thus, investigators [7–9, 77] have relied on measurements of repassivation response. Since the repassivation response generally follows a power law decay of fixed slope ‘n’, the S/FR/O model can be reduced to a very simple formulation that derives from the periodicity of strain (oxide ‘rupture’) and the repassivation response: –n



it = i0 ÈÍ t ˘˙ Ît 0 ˚



i0 t 0n Vt = M e n z r F (1 – n ) e fn ct



V t = A (ect )n

Methods of direct, in-situ measurements of crack tip strain rate have not been devised, especially since the deformation behavior is highly inhomogeneous, in terms of specific slip planes within one grain, specific grain and grain boundary response, and vs. time. Thus, correlations to engineering/macroscopic stressors, such as stress intensity factor (K), stress intensity factor amplitude (DK), frequency, and applied strain rate (for slow strain rate tests) have been developed. Continuum mechanics formulations, such as proposed by Gao and Hwang [78], have also been used (e.g., by Shoji [77] and Young et al. [79]), but they rely on many assumptions and approximations, and have not been carefully validated against SCC observations, e.g., as a function of K, yield strength, rising/falling K profiles, growth rate (i.e., varying water chemistry), etc. While relying on measurements of dissolution currents during repassivation to provide conceptual and quantitative data in many EAC systems, the S/ FR/O model does not conceptually require separated anodic dissolution and cathodic reactions, since the process of film rupture and reformation is equally applicable to gaseous environments such as steam. Indeed, the continuum in EAC from high temperature water (e.g., 200–360 °C) to steam (e.g., >400 °C) can be conceptually explained by the shift from a dissolution dominated crack advance process at lower temperatures (where there is a ‘catalytic’ benefit associated with separated anodic and cathodic reactions) to oxidation dominated crack advance at higher temperatures (where direct chemical oxidation can occur readily). The sub-mechanisms associated with the effects of water chemistry – including dissolved O2 and H2, corrosion potential, aqueous impurities, etc. – were discussed in Sections 9.3.3 and 9.4.3. Anodic and cathodic

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reactions near the crack mouth create the potential gradient and associated crack chemistry. Since these reactions are remote and uncoupled from the crack tip reactions associated with crack advance, this sub-mechanism can be isolated and solved independently. The sub-mechanisms related to grain boundary compositional inhomogeneities, e.g., from sensitization or radiation segregation, can similarly be isolated and addressed independently of the overall crack advance mechanism. The sub-mechanism related to repassivation is addressed by measurement – no useful models exist to describe film formation, the extent of corrosion during repassivation, etc., as a function of crack tip material and crack tip water chemistry. Simulating the crack tip conditions and measuring the repassivation kinetics is not trivial, but has been done. The sub-mechanism related to crack tip deformation kinetics is the most difficult, as measurements of deformation rate at growing crack tips are very difficult, and continuum mechanics formulations and finite element models are inadequate for providing a robust description of the dislocation kinetics at a crack tip. However, they do provide guidance regarding the effect of temperature, stress, stress intensity factor, stress intensity factor amplitude, etc. For example, under inert conditions, fatigue crack growth occurs by reversed slip, and the amount of reversed slip per increment in crack advance can be accurately estimated. Reversed slip can readily be associated with a crack tip strain rate, which permits good estimation of the crack tip deformation under cyclic conditions. Similarly, under slow strain rate test conditions, there is a sound basis for estimating a crack tip strain from the applied strain rate and the number of cracks that form. The biggest challenge is for constant K conditions, where sustained crack tip deformation is maintained by virtue of sustained crack advance, which requires that the strain field at the crack tip be redistributed. Nonetheless, there are empirical formulations that have been developed and provide reasonably accurate SCC prediction [7–9] of the spectrum of loading conditions (cyclic, slow strain rate, constant) and the inter-related effects of loading and material and water chemistry effects (Fig. 9.35).

9.8

Stress corrosion cracking (SCC) mitigation

Some of the mitigation approaches should be clear from the preceding sections; mitigation can be viewed in terms of moving or shrinking the size of one or more of the circles in Fig. 9.1. Factors that minimize SCC initiation include low-stress designs, absence of sharp corners or crevices, control of surface cold work during fabrication and welding, and improved surface characteristics. Many of these techniques are only applicable to new components, but processes have been developed that reduce one of more of the stressors, namely stress, cold work, surface roughness and surface

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Stress corrosion cracking of austenitic stainless steels 10–3

Crack propagation rate, cm/s

10–4 10–5

295

Theoretical and observed crack propagation rate vs. strain rate sensitized 304 stainless steel epr 15 C/cm2 288 °C Constant load --- SSRT --- Corrosion fatigue

10–6 10–7

Theory 8 ppm O2 0.5 mS/cm

10–8

Theory (SSC) deaerated 0.1 mS/cm

10–9 10–10 10–10

Theory (fatigue) deaerated 0.1 mS/cm

10–9

10–8

10–7 10–6 10–5 10–4 Crack tip strain rate, 1/s

10–3

10–2

10–1

9.35 SCC growth rate of sensitized stainless steel vs. crack tip strain rate spanning constant K, slow strain rate and corrosion fatigue loading. Because the curves diverge, the change in growth rate (factor of improvement) associated with the indicated water chemistry change varies with loading condition [7].

degradation. While surface stresses can be redistributed using various types of shot, laser and water-jet peening processes, decreasing the surface cold work, surface roughness and surface degradation requires removal of some of the surface. Perhaps the most promising technique addresses all of these factors [80]. It is more common to consider techniques that mitigate SCC growth rates, because it is essentially impossible to prove that components that have been in service for years or decades have not already developed small cracks, or to ensure that new components will not crack. Because extensive cracking of stainless steel was observed first in BWRs, the most fully developed techniques for SCC mitigation apply to that system. Billions of US dollars were spent replacing piping systems with low-carbon and/or elevated nitrogen stainless steels that were resistant to weld sensitization. But replacement of pressure vessel internals is far more costly, and the most efficient mitigation techniques involve optimized water chemistry, since they provide systemwide benefits. Hydrogen water chemistry (HWC) [81] was the first water chemistry mitigation technique developed. It involved injecting H2 into the feedwater, and was effective because the annulus of most BWRs provided an optimal range of gamma radiation for recombination of oxidants and H2. This significantly

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reduced the corrosion potential in the external piping and in the lower plenum (below the core) – its effectiveness in the core region, where there is a high neutron flux, is less clear. Adequate mitigation generally requires a ‘moderate’ H2 injection rate that makes the bulk water reduce enough to promote reduction of NO3 to volatile forms (NO2, NO and NH3). Because there is always some transmutation of O16 to N16, moderate HWC causes a 5–7-fold increase in the radiation levels in the steam lines and at the turbine. Since the half-life of N16 is only 7.2 seconds, it quickly decays if it remains soluble as NO3 or NO2. Because H2 partitions to the steam phase, the rate and cost of the continuously injected H2 is relatively high. These drawbacks led to the development of electrocatalytic surface technologies, most notably NobleChem™ and OnLine NobleChem™. By introducing very low levels of platinum ion (usually in the form of Na2Pt(OH)6) into the reactor water – either at ~130 °C for NobleChem™ or during fullpower operation for OnLine NobleChem™) – all wetted surfaces become sufficiently catalytic that oxidants (e.g., O2 and H2O2) and H2 (a reducant) will fully react on the surface. This creates a very low corrosion potential even when the bulk water retains significant levels of oxidant – provided there is a stoichiometric excess of H2 vs. oxidants (Fig. 9.36). Since water is the product, only a small amount of H2 has to be injected to achieve the 2:1 molar ratio associated with its reaction with O2 (2H2 + O2 Æ 2H2O), and essentially no increase in radiation level occurs. Other techniques that have been considered or applied include stress mitigation (e.g., with last-pass heat sink welding, or mechanical stress improvement process) and full annealing of welds to eliminate weld residual stresses and strains. An important factor in evaluating mitigation methods is recognizing that the factor of improvement is dependent on the conditions of test or component. Figure 9.36b shows that the benefit associated with water chemistry changes depends on loading and corrosion potential, and Fig. 9.35 shows that the divergence of the lines translates to a different benefit depending on the loading conditions of the specimen or component.

9.9

Prediction of stress corrosion cracking (SCC) and irradiation assisted stress corrosion cracking (IASCC)

There are a number of narrow applicability, empirical models of crack advance, e.g., for SCC of stainless steels in BWR water, SCC of irradiated SS in BWRs, SCC of Alloy 600 in PWR primary water, SCC of Alloy 182 weld metal in PWR primary water, etc. However, the limits of such empirical approaches were recognized 30 years ago [7–9], because with 20–30 primary variables, with the effect of most variables being interdependent on the others, the matrix of experiments required to identify the dependencies exceeds

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Stress corrosion cracking of austenitic stainless steels 400

nmca-18

300 1.0 ppm O2, 288 °C, 150 cc/min.

200

nmca-18 (ust) nmca-19 nmca-19 (ust)

100 ECP, mV (she)

297

nmca-18: 60 ppb Pt+20 ppb Rh, 200 ppb O2, 48 hours, 50 °C nmca-19: 60 ppb Pt+20 ppb Rh, 200 ppb O2, 48 hours, 88 °C

0 –100 –200 –300 –400 –500 –600

0

1

2 3 4 Molar ratio of H2/O2 in water (a)

5

6

O2

0 Bulk water with O2 and H2 Metal with catalyzed surface

2H2 + O2 Æ H2O on surface Stagnant water layer (b)

9.36 (a) Corrosion potential of an electrocatalytic surface vs. H2 to O2 molar ratio. The corrosion potential drops significantly once stoichiometric excess H2 conditions are achieved at the surface. (b) Schematic of an electrocatalytic surface showing rapid reaction of O2 and H2 at the surface, mass transport through the stagnant boundary layer, and the bulk levels of O2 and H2. If H2 is in excess, all of the O2 arriving at the surface is reacted, and the corrosion potential drops significantly.

1014. So the conceptual model described in Section 9.7 was created with engineering inputs, such as stress intensity factor, neutron fluence, corrosion potential (on the external surface), bulk water chemistry, grain boundary Cr depletion, etc. While many common measurements are imperfect – for example, solution conductivity does not account for specific anion effects – better inputs can be used if available.

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The model evaluated critical processes to properly model the individual effects – such as corrosion potential and water purity, which combine to affect the crack tip chemistry – and to understand the repassivation response in deaerated water of varying pH and anion concentrations. The overall concept is to develop, evaluate and quantify a model based on the crack tip system, then compare it to controlled laboratory data and then to well-defined plant data. The model provided excellent agreement with a diversity of laboratory data, e.g., for the effects of corrosion potential, water purity, stress intensity factor, cold work, sensitization, fluence, etc., as can be seen by the predicted lines in Figs 9.4 – 9.8 and 9.21. Insofar as well-documented field cracking data existed, reasonably accurate predictions can be made for the time when cracks appear (e.g., in recirculation piping, Fig. 9.37) as a function of distinguishing characteristics of plant operation – in this case, the impurity concentrations that are reflected in the average plant coolant conductivity. Difference in BWR water purity dominated the behavior of other components, including the incidence of cracking in shroud head bolts (Fig. 9.38). The predicted response based on average conductivity matched the plant response in most cases, but there were three outliers, where the incidence of cracking was high at low average conductivity. Investigation revealed that these plants were unique in having very high coolant conductivity early in life, followed by low conductivity subsequently. The early exposure caused more extensive cracking, and these cracks were then at higher K where growth could occur more readily after the plant conductivity improved. The predicted curve in these cases was based on year-by-year conductivity data rather than an overall lifetime average. Cracking in core internal structures, such as the core shroud, were accurately predicted years in advance of their detection in-plant (Fig. 9.39). Cracking in some cases was associated with significant irradiation damage, and in other cases cracking occurred in areas of low neutron fluence. An important factor in welded, bolted or other constant displacement loaded structures is that radiation creep causes stress relaxation at the same time as it causes radiation segregation and hardening. Although no residual stress measurements vs. neutron fluence have been made to verify the quantitative role of radiation creep, there is a strong basis for its effect. SCC in highly irradiated structures, like stainless steel control rods, has also been predicted [7–9, 24, 25, 82], although a range of responses is expected since swelling of the neutron-absorbing B4C pellets plays a large role. Studies on irradiated stainless steel from BWR control rods in the laboratory were performed by Jacobs as described in references [24, 25]. Figure 9.40 shows the predicted vs. observed response for these tests in one range of stress. Other examples of predicted response in sensitized stainless

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Fraction of wall thickness penetrated

Stress corrosion cracking of austenitic stainless steels

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28≤ dia. schedule 80 304 piping theoretical vs observed igscc penetration Theoretical Residual predication stress Mean Upper limit

0.5 0.4 0.3

Browns Ferry–1 0.326 ms/cm

Hatch–2 0.401 mS/cm

N.d.t. resolution limit

0.2

Big Rock Point 0.15 ms/cm

vermont Yankee 0.216 ms/cm

0.1

20

40

60

80

100 120 140 160 On–line, months (a)

180

200 240

260

10≤–28≤ dia. sch 80 304 piping in BWR environment Observed data with, in parentheses, the no. of cracks quoted Theoretical relationships based on “pledge” code, assuming Ecorr = OmVshe

Conductivity (Geo. mean), mS/cm

0.5

0.4

(42) (2)

0.3

0.2

(1) Curve # sresid. 1 Mean 2 Upper

0.1





Degree of sensitizN 15C/cm2 15C/cm2

Upper 20C/cm2

3

20

(19)

(50)

(6)

(6) (3) (19) (19) (2) (2)

1 2

10

Mean condition

(50)

3

Worst condition

40 60 80 100 120 140 On-line time (months) for wall penetration, a/t = 0.25 (b)

160

9.37 Predicted and observed response of sensitized stainless steel piping as a function of plant water purity (solution conductivity).

steels, unsensitized stainless steel, irradiated stainless steel and other materials are discussed in References [7–9, 24, 25, 82].

9.10

Future trends

Advanced LWR designs – such as Advanced BWR (ABWR), Economic Simplified BWR (ESBWR), AP1000 (Advanced PWR), European Passive

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% components with IGSCC divided by on-line months

1

Creviced Alloy 600 Shroud Head Bolts

0.9

UT inspections at BWR plants with detection at 10% of wall

0.8

BWR's with high mS/cm excursions not reflected in plant average

0.7 0.6 0.5

Predicted from stress distribution Eq. EPR 13 C/cm2 fc 150 mVshe Uniform Stress a0 50 mm

0.4 0.3 0.2

Prediction for high conductivity Early in Plant Life

Prediction for Avg. Conductivity

0.1

0.5

0.1 0 0

0.2 0.3 0.4 Avg. plant conductivity, mS/cm

0.6

9.38 Predicted and observed effect of average plant water purity (solution conductivity) on incidence of SCC in Alloy 600 shroud head bolts in BWRs [7]. The three outlier plants had high coolant conductivity early in life, then good water purity (low conductivity) subsequently, so that the average conductivity did not accurately characterize their response. When predictions were made on a yearby-year basis, the agreement was good.

Reactor (EPR), and others – incorporate simplifications, improved materials, better inspectability, etc. Nonetheless, in April 2008 there were 439 nuclear power plants operating in 31 countries (104 in the United States, 59 in France, 55 in Japan, 31 in the Russian Federation and 20 in Korea, among others), and the need to maintain their safe and reliable operation must remain a priority. The development of improved materials, fabrication and joining techniques, control of microstructure and surface conditions, etc., also remains a priority. High nickel alloys such as Alloy 690 are unattractive in the core because nickel activates and complicates inspection and servicing during outages. Welding creates residual stresses and strains (and possible weld defects), and these have been the primary causative factor in SCC in LWRs. The importance of microstructure is highlighted in observations of high growth rates in cold-worked Alloy 690 [83–85]. Despite efforts to carefully fabricate components, most surfaces have significant surface damage, which can include both surface cold work and compositional changes from heat treatment. The desire to achieve component reliability over 40 or 60+ years of plant operation will require sustained, diligent efforts.

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Stress corrosion cracking of austenitic stainless steels Core Shroud Analysis #J30/31e

25

32 mm thick 304SS, 2-sided weld 0.15 mS/cm, EPR0 = 0 C/cm2 Symmetrical sres profile Stepped thru-wall flux 5 ¥ 1019 n/cm2-y at ID +0.20 Vshe

20

Crack depth, mm

301

sres with + 103 MPa above nominal

15 sres with +69 MPa above nominal

10

Indication #4: prior UT current average current maximum

5

0

0

Boat sample

100

200 Time, months

300

400

9.39 Predicted and observed SCC in stainless steel core shrouds in BWRs [7, 24, 25].

Fluence, n/cm2 (E > 1 MeV)

1022

3

25 ¥ 1021

Theoretical V Observed Relationships, Double Ligament Specimens 304 Stst, 32 ppm O2 Water, 288 °C

5

18 30 43 48 44 46 42 40 61 62 60 7

59

50-70 KSI

(345 – 483 MPa)

30-50 KSI

(207 – 345 MPa)

71

69

70

1020

(483) 70 1019 101

102

Theory

103

(207) 30

104

Applied tensile stress KSI (MPa)

105

Time to failure, h

9.40 Predicted and observed SCC response of irradiated stainless steel tested under constant load conditions in oxygen-saturated (~40 ppm) water at 288 °C.

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Sources of further information and advice

There are a number of publications and proceedings focused on SCC of stainless steels and other structural materials (i.e., nickel alloys, weld metals and pressure vessel steels) that are excellent introductory sources of information [1–5], which include publications that cover a broader spectrum of environmental cracking phenemona. The IAEA, NRC and EPRI are organizations that are sources of extensive information, and the latter two organizations have undertaken major efforts in proactive materials degradation management. The US Nuclear Regulator Commission website (www.nrc. gov) has extensive links to plant incidents and laboratory reports. Important conferences include the Env Deg and Fontainvraad conferences, along with nuclear sessions such as the annual NACE Corrosion conference. An annual gathering of technical experts involved in environmental cracking in LWRs – the International Cooperative Group on EAC – meets every year (contact the author). Some company, national laboratory and contractor reports are available from their websites (e.g., www.epri.com) or by direct contact, e.g., with Areva, GE, Westinghouse/Toshiba and Hitachi. Journals in which SCC data in high temperature water environments are published include Corrosion, J. of Nuclear Materials, Corrosion Science and others.

9.12

References

1. Proc. 1st–13th Int. Symp. on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, NACE/ANS/TMS, 1983–2007. 2. Proc. 1st–6th Int. Symp. Fontevraud, Contribution of Materials Investigation to the Resolution of Problems Encountered in Pressurized Water Reactors, 1986–2006. 3. Proc., Chemistry and Electrochemistry of Corrosion and SCC: A Symposium Honoring the Contributions of R.W. Staehle, Ed. by R.H. Jones, TMS, Feb. 2001. 4. Proc. Parkins Symp. on Fundamental Aspects of Stress Corrosion Cracking, ed. by S.M. Bruemmer et al., AIME, 1992. 5. Stress Corrosion Cracking and Hydrogen Embrittlement of Iron-Base Alloys, Firminy, France, June 1973, Ed. by R.W. Staehle, J. Hochmann, R.D. McCright and J.E. Slatern, NACE, Houston, TX, 1977. 6. Fundamental Aspects of Stress Corrosion Cracking, Ed. by R.W. Staehle, A.J. Forty, and D. VanRooyen, Ohio State Univ., NACE, Houston, TX 1967. 7. F.P. Ford and P.L. Andresen, ‘Corrosion in Nuclear Systems: Environmentally Assisted Cracking in Light Water Reactors’, in Corrosion Mechanisms, Ed. by P. Marcus and J. Ouder, Marcel Dekker, pp. 501–546, 1994. 8. F.P. Ford and P.L. Andresen, Proc. Third International Symposium on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, Ed. by G.J. Theus and J.R. Weeks, The Metallurgical Society of AIME, 1988, p. 789. 9. P.L. Andresen and F.P. Ford, Mat. Sci. Eng., Vol. A1103, 1988, p. 167. 10. P.L. Andresen and G.S. Was, ‘SCC of Unirradiated Stainless Steels and Nickel Alloys in Hot Water’, 17th International Corrosion Congress, Las Vegas, NACE, Houston, TX, 2008.

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11. P.L. Andresen, ‘Perspective and Direction of Stress Corrosion Cracking in Hot Water’, Proc. Tenth Int. Symp. on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, NACE, 2001. 12. P.L. Andresen, T.M. Angeliu and L.M. Young, ‘Immunity, Thresholds, and Other SCC Fiction’, Proc. Staehle Symp. on Chemistry and Electrochemistry of Corrosion and SCC, TMS, Feb. 2001. 13. P.L. Andresen, T.M. Angeliu, L.M. Young, W.R. Catlin and R.M. Horn, ‘Mechanisms and Kinetics of SCC in Stainless Steels’, Proc. Tenth Int. Symp. on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, NACE, 2001. 14. P.L. Andresen, L.M. Young, W.R. Catlin and R.M. Horn, ‘Stress Corrosion Crack Growth Rate Behavior of Various Grades of Cold Worked Stainless Steel in High Temperature Water’, Corrosion/02, Paper 02511, NACE, 2002. 15. P.L. Andresen, P.E. Emigh and L.M. Young, ‘Mechanistic and Kinetic Role of Yield Strength/Cold Work/Martensite, H2, Temperature, and Composition on SCC of Stainless Steels’, Proc. Int. Symp. on Mechanisms of Material Degradation in Non-Destructive Evaluation in Light Water Reactors, Osaka, Japan, May 2002, published by Inst. of Nuclear Safety System, Japan, 2002. 16. P.L. Andresen, P.W. Emigh, M.M. Morra and R.M. Horn, ‘Effects of Yield Strength, Corrosion Potential, Stress Intensity Factor, Silicon and Grain Boundary Character on the SCC of Stainless Steels’, Proc. 11th Int. Symp. on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, ANS, 2003. 17. P.L. Andresen, P.W. Emigh and L.M. Young, ‘Mechanistic and Kinetic Role of Yield Strength/Cold Work/Martensite, H2, Temperature, and Composition on SCC of Stainless Steels’, Invited overview, Proc. of 10th Anniversary Symposium of Institute for Nuclear Systems Safety, Osaka, Japan, May 2002. 18. P.L. Andresen, ‘Factors Influencing SCC and IASCC of Stainless Steels in High Temperature Water’, PVP, Vol. 479, ASME, 2004. 19. P.L. Andresen and M.M. Morra, ‘IGSCC of Non-sensitized Stainless Steels in High Temperature Water’, J. of Nuclear Materials, Vol. 383, Issues 1–2, December 2008, pp. 97–111. 20. P.L. Andresen, ‘Perspective and Direction of Stress Corrosion Cracking in Hot Water’, Proc. Tenth Int. Symp. on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, NACE, 2001. 21. P.L. Andresen and M.M. Morra, ‘SCC of Stainless Steels and Ni Alloys in High Temperature Water’, Corrosion, Vol. 64, 2008, pp. 15–29. 22. P.L. Andresen, ‘Critical Processes to Model in Predicting SCC Response in Hot Water’, Paper 05470, Corrosion/05, NACE, Houston, TX, 2005. 23. P.L. Andresen and L.M. Young, ‘Characterization of the Roles of Electrochemistry, Convection and Crack Chemistry in Stress Corrosion Cracking’, Proc. Seventh International Symposium on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, NACE, 1995, pp. 579–596. 24. P.L. Andresen, F.P. Ford, S.M. Murphy and J.M. Perks, Proc. Fourth International Symposium on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, Ed. by D. Cubicciotti and G.J. Theus, NACE, 1990, pp. 1–83. 25. P.L. Andresen, ‘Irradiation Assisted Stress Corrosion Cracking’, in Stress Corrosion Cracking: Materials Performance and Evaluation, Ed. by R.H. Jones, ASM, Materials Park, OH, 1992, pp. 181–210. 26. S.M. Bruemmer, E.P. Simonen, P.M. Scott, P.L. Andresen, G.S. Was and J.L.

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Nelson, ‘Radiation Induced Material Changes and Susceptibility to Intergranular Failure of Light Water Reactor Core Internals’, J. Nucl. Mater., Vol. 274, 1999, pp 299–314. 27. G.S. Was and P.L. Andresen, ‘SCC Behavior of Alloys in Aggressive Nuclear Reactor Core Environments’, Corrosion, Vol. 63, No. 1, 2007, pp. 19–45. 28. G.S. Was and P.L. Andresen, ‘The Nature of SCC in Irradiated Stainless Steels and Nickel-base Alloys in LWR Environments’, 17th Int. Corrosion Congress, Las Vegas, NACE, Houston, TX, 2008. 29. ASME Boiler and Pressure Vessel Code, Sections III and XI, ASME, New York. 30. H. Hanninen and I. Aho-Mantila, ‘Environment-Sensitive Cracking of Reactor Internals’, Proc. Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, Traverse City, AIME, 1987, pp. 77–92. 31. R.L. Cowan and G.M Gordon, ‘Intergranular Stress Corrosion Cracking and Grain Boundary Composition of Fe-Ni-Cr Alloys’, Stress Corrosion Cracking and Hydrogen Embrittlement of Iron-Base Alloys, Firminy, France, June 1973, Ed., by R.W. Staehle, J. Hochmann, R.D. McCright and J.E. Slatern, NACE, Houston, TX, 1977, pp. 1063–1065. 32. J.S. Armijo, J.R. Low and U.E. Wolff, Nuclear Applications, Vol. 1, 1965, p. 462. 33. T.J. Pashos et al., ‘Failure Experience with Stainless Steel Clad Fuel Rods in VBWR’, Trans. Am. Nuclear Society, Vol. 7, No. 2, 1964, p. 416. 34. Y.J. Kim, L.W. Niedrach, M.E. Indig and P.L. Andresen, ‘Applications of Noble Metals in Coatings and Alloys for Light Water Reactors’, Journal of Metals, Vol. 44, No. 2, 1992, pp. 14–18. 35. P.L. Andresen, ‘Application of Noble Metal Technology for Mitigation of Stress Corrosion Cracking in BWRs’, Proc. Seventh International Symposium on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, NACE, 1995, pp. 563–578. 36. P.L. Andresen, Y.J. Kim, T.P. Diaz and S. Hettiarachchi, ‘Mitigation of SCC by Online NobleChem’, Proc. 13th Int. Symp. on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, Canadian Nuclear Society, 2007. 37 P.L. Andresen, Y.J. Kim, T.P. Diaz and S. Hettiarachchi, ‘Online Catalytic Mitigation of SCC at Parts Per Trillion Level’, Paper 1683, Corrosion/08, NACE, Houston, TX, 2008. 38. P.L. Andresen, K. Gott and J.L. Nelson, ‘Stress Corrosion Cracking of Sensitized Type 304 Stainless Steel in 288C Water: A Five Laboratory Round Robin’, Proc. Ninth Int. Symp. on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, AIME, 1999. 39. P.L. Andresen, P.W. Emigh, M.M. Morra and J. Hickling, ‘Effects of PWR Primary Water Chemistry and Deaerated Water on SCC’, Proc. 12th Int. Symp. on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, TMS, Snowbird, August 2005. 40. P.L. Andresen and J. Hickling, ‘Effects of B/Li/pH on PWSCC Growth Rates in Ni-base Alloys’, EPRI Final Report 1015008 (MRP-217), August 2007. 41. P.L. Andresen, ‘Mitigation of PWSCC in Nickel-base Alloys by Optimizing H2 in Primary Water’, Report to EPRI, Report 1016603 (MRP-252), December 2008. 42. P.L. Andresen, J. Hickling, K.S. Ahluwalia and J.A. Wilson, ‘Effects of Hydrogen on SCC Growth Rate of Ni Alloys in High Temperature Water’, Corrosion, Vol. 64, No. 9, 2008, p. 707.

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43. P.L. Andresen, J. Hickling, K.S. Ahluwalia and J.A. Wilson, ‘Effect of Dissolved H2 in Primary Water on the SCC Growth Rate of Ni Alloys’, Proc. Int. Conf. on Water Chemistry of Nuclear Reactor Systems, Berlin, VGB, 2008. 44. S.A. Attanasio and D.S. Morton, ‘Measurement of the Ni/NiO Transition in Ni-Cr-Fe Alloys and Updated Data and Correlation to Quantify the Effect of Aqueous Hydrogen on Primary Water SCC’, Proc. 11th Int. Symp. on Environmental Degradation of Materials in Nuclear Power Systems, ANS, 2003. 45. ‘Materials Reliability Program: Crack Growth Rates for Evaluating Primary Water Stress Corrosion Cracking (PWSCC) of Alloy 82, 182, and 132 Welds’, Report 1006696 (MRP-115), EPRI, Palo Alto, CA. 46. D. Morton, S. Attanasio, E. Richey, G. Young, ‘In Search of the True Temperature and Stress Intensity Factor Dependences for PWSCC’, 12th International Conference on Environmental Degradation of Materials in Nuclear Systems, 2005. 47. K. Arioka, T. Yamada, T. Terachi and G. Chiba, ‘Cold Work and Temperature Dependence of SCC Growth of Austenitic Stainless Steel in Hydrogenated and Oxygenated High Temperature Water’, Corrosion, Vol. 63, No. 12, 2007, pp. 1115–1123. 48. O.K. Chopra and D.J. Gavenda, ‘Effects of LWR Coolant Environments on Fatigue Lives of Austenitic Stainless Steels’, J. Pressure Vessel Technol., Vol. 120, 1998, pp. 116–121. 49. M. Higuchi and K. Iida, ‘Reduction in Low–Cycle Fatigue Life of Austenitic Stainless Steels in High–Temperature Water’, in Pressure Vessel and Piping Codes and Standards, PVP, Vol. 353, ASME, New York, 1997, pp. 79–86. 50. S. Ritter and H-P Seifert, ‘Corrosion Fatigue Crack Growth Behavior of Austenitic Stainless Steels under Simulated LWR Conditions’, Proc. 17th International Corrosion Congress, Las Vegas, NACE, Houston, TX, 2008. 51. K. Fukuya, K. Fujii, M. Nakano, N. Nakajima and M. Kodama, ‘Stress Corrosion Cracking on Cold-Worked 316 Stainless Steels Irradiated to High Fluence’, Proc. 10th Int. Symp. on Env. Deg. of Materials in Nuclear Power Systems – Water Reactors’, TMS, Snowbird, 2003. 52. K. Arioka, Y. Kanashima, T. Yamada and T. Terachi, ‘Influence of Boric Acid, Hydrogen Concentration and Grain Boundary Carbides on IGSCC Behaviors of SUS 316 Under PWR Primary Water’, Proc. 11th Int. Symp. on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, ANS, 2003. 53. S.A. Attanasio, D.S. Morton, M.A. Ando, N.F. Panayotou and C.D. Thompson, ‘Measurement of the Ni/NiO Phase Transition in High Temperature Hydrogenated Water Using the Contact Electrical Resistance (CER) Technique’, Proc. 10th Int. Symp. on Environmental Degradation of Materials in Nuclear Power Systems, NACE, 2001. 54. D.S. Morton, S.A. Attanasio and G.A. Young, ‘Primary Water SCC Understanding and Characterization Through Fundamental Understanding in the Vicinity of the Ni/NiO Phase Transition’, Proc. 10th Int. Symp. on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, NACE, 2001. 55. S. Attanasio, J. Mullen, J. Wuthrich, W. Wilkening, D. Morton, ‘SCC Growth Rates of Alloy 182 and 82 Welds’, NRC Conference on PWR Vessel Penetration Inspection, Cracking and Repair, Gaithersburg, MD, September 2003. 56. D. Morton, S. Attanasio, E. Richey, G. Young and R. Etien, ‘Updated Data and Correlation to Quantify the Effect of Aqueous Hydrogen and Low Temperature on the SCC Growth Rate of Nickel-base Alloys in Primary Water’, Proc. Alloy 600 Conference, Atlanta, June 2007, EPRI, Palo Alto, CA. © Woodhead Publishing Limited, 2010

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57. S.M. Bruemmer, J.S. Vetrano and M.B. Toloczko, ‘Microstructure and SCC Crack Growth of Nickel-Base Alloy 182 Weld Metal in Simulated PWR Primary Water’, Proc. 13th Int. Symp. on Env. Degradation of Materials in Nuclear Power Systems – Water Reactors, CNS, 2007. 58. C.L. Briant and P.L. Andresen, ‘Role of S, P and N Segregation on Intergranular Environmental Cracking of Iron and Nickel Base Alloys in High Temperature Water’, Proc. Third Int. Conf. Degradation of Materials in Nuclear Power Industry – Water Reactors, Traverse City, TMS-AIME, Warrendale, PA, 1989, pp. 371–382. 59. P.L. Andresen, ‘Effect of Noble Metal Coating and Alloying on the Stress Corrosion Crack Growth Rate of Stainless Steel in 288C Water’, Proc. Sixth Int. Symp. on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, AIME, 1994, pp. 245–253. 60. S.M. Bruemmer, B.W. Arey and L.A. Charlot, Proc. 6th Int Symp on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, AIME, 1994, pp. 277–285. 61. T.M. Angeliu, P.L. Andresen, J.A. Sutliff and R.M. Horn, ‘Intergranular Stress Corrosion Cracking of Unsensitized Stainless Steels in BWR Environments’, Proc. Ninth Int. Symp. on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, AIME, 1999, pp. 311–318. 62. T.M. Angeliu, P.L. Andresen, E. Hall, J.A. Sutliff, S. Sitzman, ‘Strain and Microstructure Characterization of Austenitic Stainless Steel Weld HAZs’, Corrosion/2000, Paper 00186, NACE, 2000. 63. P.L. Andresen, T.M. Angeliu, W.R. Catlin, L.M. Young and R.M. Horn, ‘Effect of Deformation on SCC of Unsensitized Stainless Steel’, Corrosion/2000, Paper 00203, NACE, 2000. 64. P.L. Andresen, T.M. Angeliu and L.M. Young, ‘Effect of Martensite and Hydrogen on SCC of Stainless Steels’, Corrosion/01, Paper #01228, NACE, 2001. 65. P.L. Andresen, L.M. Young, P.W. Emigh and R.M. Horn, ‘Stress Corrosion Crack Growth Rate Behavior of Ni Alloys 182 and 600 in High Temperature Water’, Corrosion/02, Paper 02510, NACE, 2002. 66. P.L. Andresen, ‘Similarity of Cold Work and Radiation Hardening in Enhancing Yield Strength and SCC Growth of Stainless Steel in Hot Water’, Corrosion/02, Paper 02509, NACE, 2002. 67. D.S. Morton, S.A. Attanasio, J.S. Fish, and M.K. Schurman, ‘Influence of Dissolved Hydrogen on Nickel Alloy SCC in High Temperature Water’, Corrosion/99, Paper 99447, NACE, 1999. 68. D.S. Morton, S.A. Attanasio, G.A. Young, P.L. Andresen and T.M. Angeliu, ‘The Influence of Dissolved Hydrogen on Nickel Alloy SCC: A Window to Fundamental Insight’, Corrosion 2001, Paper 01117, NACE, 2001. 69. A.J. Jacobs, ‘Hydrogen Buildup in Irradiated Type 304 Stainless Steel’, 13th Symp. Radiation Induced Changes in Microstructure, Ed. by F.A. Garner, N.H. Packan and A.S. Kumar, STP 956, Vol. II, ASTM, 1985, p. 239. 70. P.L. Andresen, ‘SCC Testing and Data Quality Considerations’, Ninth Int. Symp. on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, AIME, 1999. See also, P.L. Andresen, ‘Experimental Quality Guidelines for SCC Testing’, GE CRD, January 30, 1998. 71. P.L. Andresen and M.M. Morra, ‘Effects of Positive and Negative dK/da on SCC Growth Rates’, Proc. 12th Int. Symp. on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors’, TMS, Snowbird, August 2005.

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72. P.L. Andresen and M.M. Morra, ‘Effect of Rising and Falling K Profiles on SCC Growth Rates in High Temperature Water’, Journal of Pressure Vessel Technology, Vol. 129, No. 3, 2007, pp. 488–506. 73. P.L. Andresen and M.M. Morra, ‘Effects of Si on SCC of Irradiated and Unirradiated Stainless Steels and Nickel Alloys’, Proc. 12th Int. Symp. on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors’, TMS, Snowbird, August 2005. 74. P.L. Andresen, ‘Emerging Issues and Fundamental Processes in Environmental Cracking in Hot Water’, Corrosion, Vol. 64, No. 5, 2008, pp 439–464. 75. C.M. Brown and W.J. Mills, ‘Load Path Effects on the Fracture Toughness of Alloy 82H and 52 Welds in Low Temperature Water’, Proc. 12th Int. Symp. on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors’, TMS, Snowbird, August 2005. 76. P.L. Andresen, F.P. Ford, T.M. Angeliu and R.M. Horn, ‘Stress Corrosion Cracking Initiation in Austenitic Stainless Steel in High Temperature Water’, Proc. Ninth Int. Symp. on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, AIME, 1999. 77. T. Shoji, ‘Progress in the Mechanistic Understanding of BWR SCC and Its Implications to the Prediction of SCC Growth Behavior in Plants’, Proc. 11th Int. Symp. on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, ANS, 2003. 78. Y.C. Gao and K.C. Hwang, ‘Elastic Plastic Fields in Steady Crack Growth in a Strain-Hardening Material’, 5th Int. Conf. on Fracture, 1981, pp. 669–682. 79. L.M. Young, P.L. Andresen and T.M. Angeliu, ‘Crack Tip Strain Rate: Estimates Based on Continuum Theory and Experimental Measurement’, Corrosion/2001, Paper 01131, NACE, 2001. 80. H. P. Offer, R.M. Horn, A.Q. Chan and M.M. Morra, ‘Assessment of the Mitigation of SCC by Surface Stress and Material Improvements’, 13th Int. Conf. on Environmental Degradation of Materials in Nuclear Power Systems, Canadian Nuclear Society, 2007. 81. B.M. Gordon, R.L. Cowan, C.W. Jewett and A.E. Pickett, ‘Mitigation of Stress Corrosion Cracking through Suppression of Radiolytic Oxygen’, Proc. 1st Int. Symp. on Environmental Degradation of Materials in Nuclear Power System – Water Reactors, NACE, 1983, p. 893. 82. F.P. Ford, P.L. Andresen, T.M. Angeliu, H.D. Solomon, R.M. Horn, R.L. Cowan, ‘Prediction and Mitigation of Cracking in BWR Core Components’, Proc. Ninth Int. Symp. on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, AIME, 1999. 83. P.L. Andresen, M.M. Morra, J. Hickling, K.S. Ahluwalia and J.A. Wilson, ‘Effect of Deformation and Orientation on SCC of Alloy 690’, Corrosion/09, Paper 4840, NACE, Houston, TX, 2009. 84. D.J. Paraventi and W.C. Moshier, ‘Alloy 690 SCC Growth Rate Testing’, Workshop on Cold Work in Iron- and Nickel-Base Alloys, Ed. by R.W. Staehle and J. Gorman, June 2007, EPRI, Palo Alto. 85. D.J. Paraventi and W.C. Moshier, ‘Alloy 690 SCC Growth Rate Testing’, Proc. EPRI Alloy 690 Workshop, Atlanta, 31 October 2007.

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10

Void swelling and irradiation creep in light water reactor (LWR) environments

F. A. G a r n e r, Radiation Effects Consulting, USA

Abstract: Until 1993 it was assumed that swelling and irradiation creep were phenomena of little importance to light water cooled reactors. It is now recognized that swelling and irradiation creep are in progress in austenitic internals of pressurized water reactors (PWRs) especially, with boiling water reactors (BWRs) not being as vulnerable to these processes. Some manifestations of swelling and irradiation creep are already being observed in PWRs. Owing to the non-linear development of swelling with increasing neutron exposure, it is expected that consequences of swelling and irradiation creep will accelerate, especially as PWRs move beyond their original design lives of forty years. Key words: light water reactors, fast reactors, neutron irradiation, austenitic steels, pressure vessel internals, void swelling, irradiation creep, license extension, long-term operation, PWRs, BWRs.

10.1

Introduction to void swelling and irradiation creep

In the various national programs conducted on fast breeder reactors in the period 1970–90, it was universally found that the dominant life-limiting irradiation phenomenon for austenitic structural materials was a process called ‘void swelling’ with ‘irradiation creep’ following as a close second. Until 1993, however, it was generally assumed that the phenomena of irradiation creep and especially void swelling were not problems that would seriously impact the operation of light water cooled reactors (LWRs) [1, 2]. It is now known that LWRs may be prone to experiencing these phenomena, especially now that the licenses of currently operating nuclear plants are being renewed from 40 to 60 years and eventually as high as 80 years [3–5]. In order to cover this LWR-relevant subject efficiently in a report of reasonable length, it is necessary to limit the amount of background information that is presented and cited. The reader is therefore directed toward a comprehensive review article that develops this subject in much more detail but which focuses primarily on data produced in sodium-cooled fast reactors [6]. It should be noted that fast reactors generally operate at fast neutron fluxes that are one to two orders higher than fluxes experienced by LWR 308 © Woodhead Publishing Limited, 2010

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components. Thus, swelling and creep were first discovered in fast reactors where in-core structural components reach high lifetime exposures in only several years. At the lower neutron fluxes characteristic of LWRs, equivalent neutron exposures require decades to accumulate. Note that in fast reactors the fast flux is usually characterized in terms of neutrons with energy >0.1 MeV, while in LWRs this energy threshold traditionally has been >1.0 MeV, the difference reflecting primarily the fact that the energetic portion of the neutron spectrum in sodium-cooled fast reactors is somewhat less energetic or ‘softer’ than that of water-cooled reactors. What are the nature and origins of these phenomena in metals? The basic driving force arises from neutron collisions with atoms in a crystalline metal matrix. When exposed to displacive irradiation by energetic neutrons or charged particles, the atoms in a metal are sometimes displaced from their crystalline position. The displacements can be in the form of single displacements resulting from a low-energy neutron collision with a single atom. More frequently, however, the ‘primary knock-on’ collision involves a larger energy transfer and there occurs a ‘cascade’ of defects that result from subsequent atom to atom collisions. For structural components of various types of nuclear reactors, it is traditional to express the accumulated damage exposure in terms of the number of times, on 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 on the order of 100–200 dpa can be accumulated over the lifetimes of some reactor components in various reactor concepts. 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. In boiling water reactors (BWRs) the stainless steel shrouds constructed of AISI 304 are not positioned very close to the core and therefore accumulate a maximum of <5 dpa over their 40-year lifetime. This low dose is a consequence of a relatively large water gap between the core and the shroud. Pressurized water reactors (PWRs), however, utilize stainless steels in closer proximity to the core and can accumulate doses of ~100 dpa in some nearcore locations over a 40-year lifetime. Therefore PWRs are more prone to atomic displacement-induced problems than are BWRs. The displacement process produces two types of crystalline 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. Both defect types have the ability to recombine with the opposite type (annihilation) or to form agglomerations of various types and geometries. The developing ensemble of various defect types with increasing dose induces significant changes

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in physical and mechanical properties, as well as resulting in significant dimensional distortion. With one exception, mechanical property changes are not covered in this chapter and the reader is directed to ref. 6 for a fuller discussion.

10.1.1 Void swelling Interstitial agglomerations are generally one- or two-dimensional in nature, but vacancy agglomerations can also exist in three-dimensional forms such as stacking fault tetrahedra and cavities. This mismatch in dimensionality, especially for the case of the cavity, allows accumulation of significant amounts of ‘voidage’ that is accompanied by significant decreases in material density and concurrent increases in volume. This process is usually referred to as ‘void swelling’ or ‘radiation swelling’. In general most cavities are not spherical in shape but tend to develop crystallographically-faceted shapes defined by close-packed crystal planes having the lowest surface energy. The sharp corners where close-packed planes meet are frequently ‘truncated’ by the next most densely populated planes. In austenitic steels this results in voids which are truncated octrahedra defined by (111) faces and (110) corners. An example is shown in Fig. 10.1 for AISI 304, a PWR-relevant steel irradiated at PWR-relevant temperature and dpa rate [4]. Cavities are usually small, ranging from tens to thousands of nanometers in diameter, with both the mean size and concentration changing strongly with

50 nm

10.1 Voids and M23C6 precipitates observed in annealed AISI 304 irradiated in EBR-II fast reactor at 380 °C to 21.7 dpa [4]. Line dislocations or dislocation loops are not in contrast.

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irradiation temperature. Cavities can form as ‘voids’, which are essentially vacuum-filled holes or they may accumulate gases such as nitrogen, oxygen, helium and hydrogen. In some cases, especially involving helium and hydrogen, such cavities are often characterized as ‘bubbles’ rather than voids. Note in Fig. 10.2 that ‘line dislocations’ coexist in a strongly voided microstructure [7]. These dislocations move in response to either defect accumulation and/or applied stresses. It is this movement of line dislocations that results in dimensional changes of irradiated metals. Initially, however, radiation produces sessile two-dimensional clusters of interstitials called ‘Frank‘ dislocation loops, and these must evolve to form somewhat more mobile ‘perfect’ dislocation loops which in turn evolve to form very mobile line dislocations. In some crystal systems, especially simple body-centered cubic (bcc) metals, the void swelling process is self-limiting, usually saturating at some value below 5%. Such saturation is accompanied by a process referred to as ‘self-organization’ whereby voids arrange themselves in three-dimensional arrays that exhibit the same crystalline orientation as that of the crystal structure. Unfortunately for most face-centered cubic (fcc) metals, especially stainless steels, self-organization and saturation of void swelling do not

10.2 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 [7]. Average void size is ~70 nm. Line dislocation segments end either on void surfaces or on upper and lower surfaces of the thin microscopy specimen.

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operate under most reactor-relevant conditions, and as a result swelling in austenitic stainless steels is an inherently unsaturable process. Tens of percent swelling can be reached during many reactor-relevant irradiation histories, but values of 80–90% swelling without hint of impending saturation have been attained during neutron irradiation [6, 8]. An example of non-saturable void swelling in a PWR-relevant steel (AISI 316 stainless) is presented in Fig. 10.3 [8]. Note that the onset of swelling, defined by a ‘transient’ regime, is dependent on both irradiation temperature and dpa rate. The dpa rate dependence is not easily discerned in these data but each irradiation temperature is matched with a specific dpa rate, with the range of rates increasing ~65% from lowest to highest. It will be shown later that dpa rate is a very strong determinant of void swelling. The duration of the transient regime of swelling in austenitic and high-nickel steels is known to be exceptionally sensitive not only to these irradiation parameters but also to fine details of composition, heat treatment and processing. The transient is sensitive to a lesser extent on applied stresses. As also shown in Fig. 16.3, the maximum post-transient swelling rate of AISI 316 is typical of all austenitic stainless steel at ~1%/dpa, essentially independent of most irradiation or material variables [6, 8, 9].

510 °C 80

1%/dpa 538 °C 482 °C

60 Swelling %

593 °C

40 427 °C 20

0 0

650 °C 454 °C

400 °C

0.2%/dpa

10 20 30 ¥ 1022 Neutron fluence, n cm–2 (E > 0.1 MeV)

10.3 Swelling as a function of irradiation temperature and dose observed in 20% cold-worked AISI 316 irradiated in the EBR-II fast reactor [8].

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In the absence of applied or internally generated stresses, void swelling distributes the increased volume isotropically, but in the presence of a stress field some portion of the increased volume can be partitioned anisotropically. Figure 10.4 shows an often-shown example of isotropic swelling [6] and Fig. 10.5 presents examples of the sensitivity to void swelling in fuel pins to variations in temperature, dpa rate and minor element composition [10, 11].

10.1.2 Irradiation creep While void swelling is non-conservative of volume, it is often accompanied by a process called ‘irradiation creep’ which is fully volume-conservative in nature. As shown in Fig. 10.6, when stress is applied to a metal, irradiation creep occurs at rates orders of magnitude greater than that of thermal creep at most reactor-relevant temperatures [12]. In general the radiation-induced creep rate of austenitic steels is directly proportional to the dpa rate and the magnitude of the applied stress under most reactor-relevant conditions, but

Unirradiated fuel cladding tube

1 cm

10.4 Isotropic increase of ~10% in dimensions of 20% cold-worked 316 tube irradiated without constraints to 80 dpa at 510 °C in the EBR-II fast reactor [6]. Swelling was measured by density change to be ~33%.

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(a)

(b)

10.5 (a) Fuel assembly from the BN-600 fast reactor showing larger swelling-induced elongation of annealed EI-847 steel in pins with slightly lower silicon content [10]. (b) Fuel assembly from the FFTF fast reactor showing larger swelling-induced elongation of pins having slightly lower phosphorus content [11]. The gradual variations in height across the FFTF fuel assembly result from gradients in irradiation temperature and neutron flux.

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Void swelling and irradiation creep in LWR environments 138 MPa 454 °C

15 ¥ 10–4

Tensile strain

315

10 Irradiation creep

5

Thermally induced densification and creep

0 0

500

1000 Time, h

1500

2000

10.6 Comparison of creep rates observed in 20% cold-worked 316 stainless steel in uniaxial creep tests during thermal aging or neutron irradiation in the EBR-II fast reactor [12]. Precipitation of carbides at elevated temperatures leads to a small densification and shrinkage of the creep specimen as shown in the thermal creep behavior. A similar process occurs during irradiation but is overwhelmed by the creep strain.

irradiation creep in the absence of swelling is not particularly sensitive to temperature or alloy composition [6]. In the steady-state pre-swelling creep regime, the creep coefficient of austenitic steels is ~1 ¥ 10–6 MPa–1 dpa–1. Irradiation creep also frequently exhibits a transient regime at lower doses but the magnitude of the transient appears to be very sensitive to details of alloy preparation, especially those details that influence texture of the alloy and its relationship to the applied stress state [6]. In the presence of stress, irradiation creep precedes the onset of swelling but is strongly accelerated once swelling begins. When swelling is in progress the irradiation creep rate becomes almost completely proportional to the swelling rate. It is important to note that unlike thermal creep, irradiation creep is inherently a non-damaging process on the microstructural level, always working to reduce to very low levels any stress concentrations or stress gradients that might arise in the steel. It is the shear component of the stress state operating on the dislocation population that drives the partition of mass flow away from the fully isotropic distribution of mass that occurs in the absence of stress. Irradiation creep will be activated by any type of stress state, whether continuously imposed or preloaded. Preloaded stresses of springs or bolts,

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regardless of their magnitude, will be progressively relaxed by irradiation creep. A general rule of thumb for austenitic stainless steels is that >90% relaxation of any preload will occur by 10 dpa.

10.1.3 Distortion of structural components It is the combined and interactive nature of void swelling and irradiation creep that produces the often spectacular deformation produced in some reactor components. Void swelling, when constrained by another structural component or when producing a strong gradient in swelling rate, tends to generate stress fields. These fields activate irradiation creep which attempts to reduce the stress fields, thereby producing mass flows in the unconstrained direction, producing an anisotropic distribution of mass flow. Figure 10.7 shows an example of creep-induced distortion in fast reactor fuel pins that were experiencing significant levels of swelling [13]. In some cases, the most pronounced consequences of swelling are driven more by differential swelling than by swelling itself. Differential swelling can arise from the interaction of two components with greatly different swelling rates or by strong gradients in swelling across a single component in response to gradients in temperature and dpa rate. As will be discussed later, both cases are important in the determination of distortion and/or component failure in PWR internal components.

10.1.4 Consequences of swelling and irradiation creep In the fast reactor research community it was eventually recognized that ‘reasonable’ levels of swelling and creep could be incorporated into or allowed for in the design of most components without significant consequences on operation or safety. Various mitigation strategies could also be employed. For instance, stresses that drive creep deformation could be reduced by increasing wall thickness or by allowing larger gas plenums. Welds which

10.7 Swelling-creep interaction in fuel pin bundle clad with 20% CW 316 stainless steel following irradiation in the FFTF fast reactor. The fuel pins were spirally wrapped with spacing wire that swelled less than the cladding, creating a constraint that activates irradiation creep to deform the pins in a spiral manner [13].

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in general swell more than the base metal can also be placed far outside the active core zone. It was also found that replacement of swelling-prone materials with lesser prone materials would delay swelling and swelling-driven creep, thereby delaying the necessity to remove an affected structural component from the reactor. This approach was suitable for components that had relatively short in-core lifetimes. As swelling becomes increasingly prominent, however, there are additional consequences that need to be considered. Unfortunately these consequences are not so easily mitigated for PWR application where the swelling-vulnerable components are already in place and are intended to remain in service, without replacement, for 40 to perhaps 80 years. As swelling increases, almost all properties of engineering interest become dominated by swelling, including irradiation creep, mechanical integrity and physical properties [6]. Most importantly, several new but related forms of severe embrittlement emerge. In the first form of embrittlement, growing levels of void surface become progressively enriched in nickel by radiation-induced segregation and depleted in chromium, moving the alloy matrix between the voids toward a composition prone to martensite instability, especially at lower temperatures characteristic of reactor shutdown conditions. This failure mode is characterized by intense flow localization and a zero tearing modulus. This produces essentially zero deformation at failure, resulting from a propagating crack forming the very brittle alpha-martensite phase at its tip as it progresses [14]. Very little energy is required to extend a crack once it has been initiated and failure occurs very quickly. Even at higher temperatures characteristic of reactor operating conditions there is a pronounced tendency toward intensive flow localization and failure without significant deformation [15–17]. This process does not involve martensite instability, however, and involves large amounts of local plasticity on the eventual failure surface. This process might best be characterized as ‘quasi-embrittlement’. This failure mode involves a suppression of uniform elongation, as opposed to true embrittlement, which involves the complete suppression of the metal’s capability for plastic deformation. Two examples of such failure modes are shown in Figs 10.8 and 10.9 [18, 19] while Figs 10.10 and 10.11 show microscopic examples of flow localization both before and after failure [20, 21]. Such failure modes tend to change our perception of swelling from being a mere operational concern to a potential safety issue. It has been shown in a number of studies that these failure modes correlate directly with the swelling level, with ~10% swelling roughly defining the transition from a reasonably ductile state to a very brittle condition [14, 15, 17, 22]. For PWR applications it is thought to be more prudent to assume that swelling not be allowed to exceed ~5% in locations where brittle failure would be considered to be a safety issue, especially since by this level the swelling

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10.8 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. Swelling at the failure point was ~14% [18].

rate should be at or near the terminal rate of ~1%/dpa and therefore 10% swelling would be reached within another 5 dpa or slightly more.

10.2

Potential for swelling and irradiation creep in light water cooled reactors (LWRs)

10.2.1 Differences between BWRs and PWRs Unfortunately, AISI 304 austenitic stainless steel in the annealed condition is the most swelling-prone, commercially available steel identified to date [23], but this steel was used to construct the major internal components of BWRs and PWRs before swelling was discovered and before its potential consequences were realized. As mentioned previously, austenitic steels used in BWRs do not receive very large exposures over a 40–60 year lifetime due a relatively large water gap between the core and the shroud. Therefore void swelling is not a significant concern in BWRs, even though voids at low levels are now beginning to be observed in the shroud assemblies of BWRs [24]. An example is shown in Fig 10.12.

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319

By-92

53 dpa 27.8%

52 dpa 29.8%

U-796

34 dpa max 14% swelling

10.9 Severe embrittlement arising from void swelling of 12X18H9T fuel assembly wrappers in the BOR-60 fast reactor. Three assemblies were broken during refueling operations with the fuel pins, lower portions of the wrappers and wrapper debris left in the core [19]. Swelling not only embrittled the wrappers, but caused high withdrawal loads due to combined ‘fattening’ and bowing of the assemblies, both of which contributed to the failure. Maximum dpa levels and maximum swelling levels for each assembly are shown.

PWRs, on the other hand, have large amounts of 304 steel framing the core (baffle-former-barrel assembly) with only several millimeters of clearance from the fuel. The re-entrant corners of the baffle-former assembly can reach very large dpa levels in 40–60 years and experience higher local temperatures due to proximity to the fuel. For most PWRs it is not considered to be economically feasible to replace these components and the potential for swelling in such locations must be closely monitored, especially in the period beyond 40 years. Three Japanese utilities, however, have successfully replaced their PWR internals at ~30 years of operation in order to avoid potentially developing problems with irradiation-assisted stress corrosion cracking and void swelling [25]. Until 1993 it was assumed that PWRs were immune from swelling due to the lower operating coolant temperatures (290–340 °C) compared to those of most Western design fast reactors (370–600 °C) [1]. It was also assumed that the much lower dpa rates characteristic of PWRs would reduce the tendency to form voids. It is now known, however, that the temperature regime of void swelling shifts toward lower temperatures as the dpa rate decreases, moving the swelling regime into PWR-relevant conditions. Even more importantly, it has been shown in various fast reactor studies that lower dpa rates cause a significant reduction in the duration (in dpa, not time) of the transient

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10.10 Channel shearing of voids resulting from severe flow localization in annealed AISI 304 stainless steel after irradiation to 38 dpa at 387 °C and subsequent tensile testing at 370 °C [20, 21]. The sheared voids indicate local deformation levels of 100–200% while the surrounding matrix is undeformed.

10.11 Failure surface of specimen shown in Fig. 10.10 showing planar facets produced by intense flow localization [21].

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Loops

Cavities

20 nm

10.12 Fresnel image showing very small cavities and edge-on Frank dislocation loops observed in BWR shroud after 25 years, reaching 1.9 dpa at ~290 °C [24]. Significant under-focusing was required to image the cavities which were not visible at focus conditions.

regime of swelling in both model and commercial austenitic alloys [26–30]. Thus, swelling occurs both at lower dpa and lower temperature, becoming a previously unanticipated concern for PWR internal components. There are three other factors that are now known to potentially affect and possibly accentuate the tendency of void swelling to occur in PWRs compared to that in fast reactors [31]. These three factors arise primarily from the differences in neutron spectra between PWRs and fast reactors. First, as shown in Fig. 10.13 the high energy part of the neutron spectrum of various light water test reactors is more energetic than that of fast reactors, producing more dpa per fast neutron. Second, LWR spectra have a significant amount of very slow ‘thermal’ neutrons which produce very large levels of helium and hydrogen (arising from reactions with isotopes of nickel) compared to that produced in fast reactors [1, 32–34]. These two transmutant gases are known to stabilize void nuclei and often accelerate the onset of swelling. Compared to test reactors like HFIR or ORR the ratio of thermal to fast neutron ratio (T/F) in PWRs is smaller by a factor of 5–10 and the thermal peak of the neutron spectrum is not as pronounced, as shown in Fig. 10.14. The thermal neutron population also tends to peak just above and below the fuel bearing region as shown in Fig. 10.15. It peaks just outside the core in the radial direction as well, becoming more important in the baffle-former region just outside the core [31].

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Flux per unit lethargy

1015

HFIR 1014 ORR 1013 EBR II 1012

10–8 10–7 10–6 10–5 10–4 10–3 10–2 10–1 100 101 102 Neutron energy, MeV

10.13 Difference in neutron flux-spectra of two water-cooled test reactors (high-flux HFIR and lower-flux ORR) and one sodium-cooled fast reactor (EBR-II) [31]. The majority of the displacement damage occurs above ~0.03 MeV in both types of reactors. In LWRs thermal neutrons do not cause a significant fraction of the displacements. 1014 T/F ~0.15

Flux per unit lethargy

1013 Baffle bolt

Top of bolt head

1012

1011 Upper core plate 1010

109 10–8 10–7 10–6 10–5 10–4 10–3 10–2 10–1 100 101 102 Neutron energy, MeV

10.14 Typical neutron flux-spectra of PWR internal components, having a T/F ratio smaller than that of typical test reactors [31].

Third, high levels of thermal neutrons are accompanied by higher rates of gamma ray production, especially via absorption of thermal neutrons in the steel. The resultant ‘gamma heating’ arises from both fission events in

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Fast flux, E > 0.1 MeV

200 150

Distance in cm

100 50 0 –50 –100 –150 –200 –250 0.001

Thermal flux 0.01 0.1 1 Flux ¥ 1014 n/cm2 sec

10

10.15 Typical axial neutron flux profiles through a PWR core. The T/F ratio in the core is ~0.2. There is a ‘bump’ in the thermal neutron population just outside the core and a corresponding increase in T/F ratio [31].

the core and thermal neutron absorption in the plates, significantly raising the local temperature of thick baffle-former plates above that of the ambient coolant temperature. Increasing temperature is known to increase the swelling of AISI 304 in fast reactors [23]. Examples of the internal components of Western PWRs are shown in Fig. 10.16. A set of somewhat overly conservative estimates of dose and temperature maps of a ‘typical’ re-entrant corner of a mid-core former plate is presented in Fig. 10.17 and demonstrate that local near-core gradients in temperature and neutron flux can generate a situation that favors the development of highly localized swelling [35]. The estimate of the temperature distribution is based on possibly somewhat larger than realistic gamma heating rates than were available at the time these estimates were made. A prediction of swelling for this ‘hot corner’ will be presented later. It should be noted that pressure vessel internals of PWRs come in two basic types, those that are of welded construction and those that are bolted together. Bolts are usually, but not always, constructed from ‘harder’ steels such as cold-worked AISI 316 austenitic stainless steel. In some reactors AISI 347 or 304 have been used, sometimes in the annealed condition. The use of cold-worked 316 steel in most bolted designs is significant in that

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Core baffle structure

Reactor core

(a)

(b)

10.16 Schematic of components of the PWR vessel, core and baffleformer assembly [4, 5].

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Baffle

Bolt

Core barrel Former (c)

D

D

10.16 Continued

370 °C on 1022 n/cm2 (E > 0.1 MeV)

380 °C

390 °C

3

F

E

400 °C

410 °C

4

~100 dpa

Dose

53 dpa

G

8

10

A

12

B

14

A G

B

5 420 °C

Temperature

10.17 Estimated temperature distribution and 40-year dose profiles assuming 75% availability of typical PWR baffle-former junction at mid-thickness [35].

it swells later than does annealed 304 stainless steel, leading to significant consequences arising from differential swelling that will be discussed later. When we address PWRs it is appropriate to note that Russian WWERs (water-cooled, water-moderated energetic reactors or VVERS when using the Russian word for water) are also of the PWR type. The primary differences between internals of Russian and Western reactors are the use of a Russian analog of AISI 321, which is also a swelling-prone steel, and the use of a much thicker baffle ring instead of a baffle-former-barrel assembly. As

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shown in Fig. 10.18 this arrangement leads to much higher temperatures and a wider range of temperatures induced by gamma heating [36]. The displacement doses shown are calculated for a 30-year lifetime, the current regulatory limit for most WWERs.

10.2.2 Review of data supporting the potential for void swelling in PWRs From irradiation experience in fast reactors, as well as from charged particle simulation studies, it is known that AISI 304 swells earlier than AISI 316, and that cold-working further delays the onset of swelling as shown schematically in Fig. 10.19. While swelling of these steels has been identified as a potential issue for PWRs during long-term operation (e.g. 60+ years), the degree of concern for PWRs is still being debated. Therefore one must review the available data in order to make an assessment of potential swelling-related problems that might be encountered at higher exposures. The swelling data derived from PWRs and WWERs on 304, 321 or 316 steels are rather limited but have established the potential for swelling-induced distortion. The first clear example of void swelling was found in a coldworked 316 baffle bolt removed from the Tihange PWR reactor located in Belgium [37]. The bolt was removed in response to an ultrasonic indication of cracking under the bolt head.

420 °C 12 dpa 460 °C 16 dpa 1–5 dpa 400 °C 4 dpa

49 dpa 420 °C 24 dpa

10.18 Calculated 30-year exposure dose and irradiation temperature distributions in a Russian WWER-1000 power reactor [36].

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Typical swelling response for stainless steels at breeder reactor temperatures

% Swelling

SA304SS > SA316SS > CW316SS

SA 304SS

SA 316SS CW 316SS

Irradiation dose, dpa

10.19 Schematic showing relative swelling behavior of PWR-relevant austenitic steels always observed in either fast reactor or charged particle irradiations. Some data from HFIR irradiations also support this trend.

Although the bolt shown in Fig. 10.20 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, wellfacetted 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 reactors. The worrisome inference is that the 304 plate surrounding the bolt might be swelling at higher levels. Subsequently, voids were observed in other AISI 316 bolts from this same reactor by other researchers [38, 39], often at lower doses and temperatures, producing lesser but measurable amounts of swelling. An example is shown in Fig. 10.21, but it should be noted that there appear to be two populations of cavities, those 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. Voids have been sometimes but not always observed in bolts of various steels removed from US PWRs [40, 41]. These studies were conducted before the need for defocusing was recognized, however. Small voids or ‘cavities that could be either voids or bubbles’ have also been observed in thin-walled flux thimble tubes removed from various PWRs [42–45]. A special case of ‘sub-visible’ voids [45] in one of these flux thimble tubes will be covered later. Neustroev and co-workers also found voids in a thimble tube removed from a WWER operating in the Ukraine, noting that voids were observed at unexpectedly low temperatures and dpa levels [46].

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10 nm Avg. size = 8.6 nm Density = 0.61 ¥ 1022 m–3 Swelling = 0.20% 20 nm

10.20 Voids observed in Tihange baffle-former bolt designated 2K1R5 made with cold-worked 316 stainless steel after irradiation at ~345 °C to 12 dpa producing ~0.2% swelling [37].

The potential for void swelling at PWR-relevant dpa rates and temperatures is best demonstrated in more comprehensive studies conducted in four Soviet 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 [47–53]. A summary paper containing an overview of most 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 [47]. First, whenever the dpa rate was lower at any investigated temperature, swelling was observed at surprisingly very low dpa levels. An excellent example is shown in Fig. 10.22 where significant void swelling was observed at only 0.64 dpa at 350 °C [48]. Second, whenever a comparison could be made within one reactor at a given temperature, the transient duration decreased with lower dpa rate [49]. Most importantly, whenever temperatures approaching 280 °C could be reached, swelling was observed not only at these low temperatures

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50 nm (a)

50 nm (b)

10.21 (a) Voids at very low density (see arrows) and (b) an exceptionally high density of sub-visible cavities or ‘nano-bubbles’ observed in another Tihange baffle-former bolt designated 2K1R1 after 8.5 dpa at ~299 °C [38]. The smaller cavities can only be seen with significant under-focusing.

but at surprisingly low dpa levels [50–53]. Some examples are shown in Figs 10.23 and 10.24. It might be tempting to dismiss each of these individual fast reactor observations as being unrepresentative of PWR experience, but when the

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50 nm

50 nm

10.22 Voids observed in annealed 12X18H9T steel at 350 °C in the BR-10 fast reactor at only 0.64 dpa produced at 1.9 ¥ 10–9 dpa/s [48]. This steel is analogous to AISI 321. 50 nm



3 dpa

6.5 dpa

22 dpa

10.23 Microstructure of annealed 12X18H9T 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 [47, 49].

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0.65 dpa 281 °C



12.3 dpa, 363 °C

331

7.7 dpa, 285 °C

12.6 dpa, 380 °C

8.8 dpa, 430 °C

10.24 Void microstructure observed in a wrapper duct constructed from annealed 12X18H9T stainless steel and irradiated in the BN-350 reactor at various axial distances from the midplane [47, 53]. Lowest temperatures correspond to the bottom of duct.

uniformity of behavior is considered in the total database, it is reasonable to deduce that a similar effect must be occurring in PWRs. It might also be tempting to dismiss these Russian steels as being unrepresentative of AISI 304 or 316 steels, but we have already presented the observations of swelling of cold-worked 316 baffle bolts from Western PWRs. As will be shown in the next section, however, swelling of AISI 304 stainless steel has been investigated in the EBR-II fast reactor at dpa rates and temperatures representative of the ‘hot corners’ discussed earlier.

10.2.3 Dependence of swelling in AISI 304 on dpa rate in EBR-II Much of the data shown above strongly imply that swelling increases at lower dpa rates, possibly by reducing the duration of the transient regime of swelling. In order to conclusively demonstrate the effect of dpa rate on void swelling, Garner and co-workers conducted a more comprehensive experiment on annealed AISI 304 stainless steel [4, 54, 55]. This experiment isolated the effect of dpa rate by concentrating on a limited range of temperatures (373–444 °C, with the majority of the data at 373–410 °C), but a very large range of dpa rates (0.06–3.8 ¥ 10–7 dpa/sec), with no significant difference in helium/dpa ratio. These temperatures are characteristic only of the upper range of PWR interest, especially in the hot corner regions of former plates, and not the lower range experienced by the majority of PWR internals, but the dpa rates span the full range of PWR internal conditions.

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The experiment involved the examination 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. A total of 280 disks (20 mm diameter ¥ 1 mm thick) were cut from these assemblies. Each disk had a unique combination of average irradiation temperature, dpa and dpa rate, produced by both axial and radial variations in these parameters. Swelling was measured by immersion density for all 280 disks and was confirmed by microscopy on 40 specimens spanning the full range of the experiment. Voids were found in all examined specimens with swelling ranging as high as 2.8% [54, 55]. Examples of the void microstructure and its sensitivity to dpa rate are shown in Fig. 10.25. Figure 10.26 shows a subset containing about one half of the collected swelling data. All data in the temperature range 373–388 °C are plotted vs. dpa. Note that there are relatively small variations in dpa rate of the data from a given subassembly in this temperature subset, but the data clearly show that the transient regime of swelling is progressively shortened as the dpa rate decreases, such that only 10 dpa are required to reach 1% swelling in row 14. In previous publications it was shown that 30–50 dpa were required to exceed 1% swelling when data were collected at these temperatures from rows 2–4 inside the EBR-II core at higher dpa rates [23]. The complete data set from the 280 disks confirms the general validity of the effect of decreasing dpa rate to strongly increase swelling at all examined temperatures. This increase occurs by shortening the duration of the transient regime. At temperatures above 388 °C the impact of dpa rate was found to be even more pronounced than shown in Fig. 10.26. Based on the success of using EBR-II data to assist in PWR predictive swelling equations, a series of other old EBR-II experiments that were not completed or sufficiently analyzed are being examined to provide additional insight on the interaction of temperature and dpa rate on swelling. A summary report is contained in ref. 56. In general these experiments confirm the previously derived dependencies of swelling on temperature and dpa rate. These other experiments also confirm the generality of the 1%/dpa terminal swelling rate of AISI 304 when swelling exceeds ~4%.

10.3

Predictions of void swelling and associated uncertainties

10.3.1 Currently available predictive equations for deformation of annealed AISI 304 stainless steel Over the years a number of empirical swelling equations have been developed to describe the swelling of AISI 304 austenitic stainless steel. The one

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10 dpa 0.15 ¥ 10–7 dpa/sec 1.2% swelling

100 nm

14.3 dpa 1.8 ¥ 10–7 dpa/sec 0.42% swelling

100 nm

10.25 Void microstructures observed in annealed AISI 304 reflector ducts from EBR-II showing figure variation of swelling in response to differences in dpa rate at 379 °C [54–56]. Small dark features are M23C6 precipitates that form concurrently.

that was most used toward the end of the US fast reactor program was the Foster–Flinn equation which was developed from mostly in-core EBR-II data at higher than PWR-relevant dpa rates [57]. This equation was expressed only in terms of fast neutron fluence and temperature. It has no explicit dependence on neutron flux or dpa rate. Based on the data presented in the previous section a new equation was developed that explicitly includes dose rate as well as temperature and dose. Both this equation and the Foster–Flinn equation suffer from the same general deficiency in that no data are available below the EBR-II inlet temperature of ~370 °C. The majority of the PWR baffle-former assembly exists at

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3.0 U1603 Row 14 0.062 – 0.156 ¥ 10–7

2.5

U1603 U9009 U8972 U9807

U9009 Row 10 0.38 – 0.96 ¥ 10–7

Swelling (%)

2.0 U8972 Row 9 1.00 – 2.05 ¥ 10–7

1.5 1.0

U9807 Row 8 1.25 – 0.60 ¥ 10–7 dpa/sec

0.5 0.0

0

5

10

15

DPA

20

25

30

35

10.26 Swelling of annealed 304 stainless in the range 373–388 °C measured by density changes in the lower halves of four EBR-II reflector subassemblies, designated by identification numbers such as U9807 etc. [4, 56]. The ranges of dpa rates from bottom to center of the duct are shown for each assembly.

temperatures below 370 °C, however. Therefore the major uncertainty in the application of this equation lies in its extrapolation below 370 °C where swelling is most likely lower than in the hotter regions where the equation should be generally applicable. The new empirical equation preserves the temperature dependence of the Foster–Flinn equation while separating the dependence of dpa and dpa rate. The spectral differences between EBR-II and PWR have been accommodated in the equation by employing dpa rather than fast neutron fluence. The well-known curvature of the transient regime in AISI 304 [23] is captured in the dependence on the square of the accumulated dpa. However, once the swelling rate reaches 1%/dpa, this equation is no longer applicable and swelling proceeds at 1%/dpa thereafter.

S = (dpa)2 (dpa rate)–0.731 F(T)

where F(T) = EXP (22.106 – (18558/(T + 273.15))) with swelling S given in %, temperature T in °C and dpa rate is in units of 10–7 dpa/sec. While internally or externally applied stresses are known to accelerate the onset of swelling by shortening the duration of the transient regime, its influence is relatively small compared to that of temperature and dpa rate [6, 58]. Also, for easily swelling steels such as AISI 304 the effect is not very pronounced and the magnitude of stresses for most PWR applications is not

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very large [5]. Therefore, given the relative uncertainties in prediction of swelling for PWRs, there is no significant advantage to providing a stressdependent swelling equation at this time. Figure 10.27 presents a schematic representation of the effect of the two major and one minor variable on void swelling of AISI 304 in the temperature range of PWR interest. The equation for irradiation creep was easier to develop since most of its dependence on temperature, steel processing and composition are those associated with swelling. It is important that swelling and creep equations be directly coupled or very unrealistic stress levels will be predicted. The creep rate is defined as

e ¢ = B + DS ¢ o s

where e is the effective plastic strain, and the ‘prime’ indicates a derivative with respect to dose (swelling rate/dpa), s is the von Mises effective stress, Bo is creep modulus, D is the creep-swelling coupling coefficient and S¢ is the derivative of swelling with respect to dose. In this equation a transient creep term has not been included because it is relatively small in AISI 304 in the annealed condition and is relatively hard to define without data on the specific heat of steel and the relationship of texture to stress state. Any predictions of stress magnitude arising from this transient-free equation will be conservatively high compared to any transient-included equation. If strain and swelling are given in percent, and stress is given in units of MPa, then the recommended value for Bo is 1 ¥ 10–4 %/(MPa dpa), and the value for D is 0.6 ¥ 10–2 MPa–1. These values are considered to be reasonable for temperatures from 250 °C to 500 °C. Using the stress-free swelling equation above, a prediction was made of the spatial dependence of swelling after 40 years in the mid-plane of the former 1%/dpa is the maximum swelling rate

dpa rate decreasing Stress increasing

Swelling %

Temperature decreasing

dpa

10.27 Schematic representation of the parametric dependence of swelling of annealed AISI 304 in the temperature and dpa rate of PWR interest.

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plate at the hot corner described in Fig. 10.17. Note in Fig. 10.28 the large and very pronounced peak in local swelling in the hot corner. The reader should not accept this value as being fully realistic since the calculation depends on several conservative assumptions concerning the gamma heating rates, local coolant flow, etc. The reader should extract the conclusions, however, that high swelling levels are very localized and are not a generally widespread problem. More importantly, even if the prediction was too high by even a factor of five, the 10% swelling-induced embrittlement limit would still have been exceeded. Of course the swelling will continue to increase in a non-linear manner beyond 40 years. With respect to the uncertainty associated with such an equation in extrapolation to lower temperatures, there are other uncertainty factors to consider. In addition to very large differences in helium and hydrogen gas generation rates, the neutron spectral differences between EBR-II and PWR will produce some small variations in transmutants and transmutation rates, primarily in the progressive loss of manganese to form iron, and the production of vanadium from chromium [31, 59]. Most likely these small changes in composition may not affect swelling, but this assumption cannot be stated with certainty. Additionally, the inlet temperature of a PWR appears to coincide with

% swelling 5–15 15–25 25–35 35–45 45–55

10.28 Calculated mid-plane swelling distribution for baffle-former conditions shown in Fig. 10.17.

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the near-bottom limit of the temperature regime of swelling. Since void nucleation and void growth may respond differently to temperature, some additional uncertainty resides here. Such a consideration becomes more important when one considers the time dependence of temperature in PWR internals. While non-fueled structural components of fast reactors operate at a relatively constant temperature, the situation in the baffle-former assembly of PWRs is rather different. As discussed earlier, gamma heating in the baffle-former assembly arises from two primary sources. The first is fission-born gammas whose intensity is strongest at the core surface and decreases in intensity with penetration into the baffle-former assembly. The second source arises from absorption of thermal neutrons in the metal of the assembly and to a lesser extent from absorption in the water. In PWRs, boric acid is added to the water at the start of each irradiation cycle to serve as a burnable poison with 10B (20% of natural boron) being a strong thermal neutron absorber. In addition to the previously mentioned increase of thermal/fast ratio (T/F) with distance from the core boundary, the time-dependent burn-up of 10B leads to a progressive increase in T/F ratio and a concurrent increase in gamma heating during each reactor cycle, as shown in Fig. 10.29. At the beginning of each cycle the boron is replenished. Over successive cycles there is a saw-tooth variation of gamma heating rate in the baffle-former assembly and therefore in DT, with the latter reaching values as large as ±20 °C in the worst case. Since swelling is very sensitive to irradiation temperature, this induces an additional uncertainty in our ability

Centerline bolt temperature underneath the bolt head, °F

660

640

620

600

580

Time state point

10.29 Typical calculated temperature history of baffle bolt head, showing cycle-to-cycle variations arising from burn-up and periodic replacement of boric acid in the cooling water, as well as a mid-life change in fuel loading pattern [5].

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to predict swelling based on fast reactor data, which are generated without such temperature variations. Note that additional temperature complexity can arise from operational considerations. The most relevant example for PWRs is the mid-life introduction of low leakage fuel loading to reduce the neutron flux and displacement rate experienced by the reactor pressure vessel. Such practice also reduces both the neutron flux and the gamma heating rate experienced by the baffle-former assembly, adding significantly to the difficulty of determining both the timeaveraged neutron flux-spectra and the irradiation temperature experienced by a given reactor component. In summary, our knowledge of void swelling in AISI 304, as encapsulated in the current swelling equation, is sufficient to predict where problems will arise, but there is insufficient confidence that the magnitude of swelling or the timescale of its development can be accurately predicted. There is a similar level of uncertainty in the swelling of cold-worked 316 bolts. It can only be stated that at a given dpa rate and temperature that the annealed 304 steel will be swelling more and at a higher rate than the cold-worked 316 steel. The emphasis on higher swelling rate as opposed to higher swelling is deliberate since differential swelling can cause as many problems as does larger swelling.

10.4

Potential swelling/creep consequences

What types of creep- and swelling-related problems might arise in PWR internals? While it is obvious that large local levels of swelling can lead to brittleness, small levels of swelling operating over longer distances can also cause problems. 1. In response to relatively small levels of average swelling along the height of baffle plates between formers, the baffle plate may bow inward or outward between the constraints imposed by the former plates, possibly contacting the fuel or reducing cooling flow. 2. Transverse movement of the baffle plate relative to the former will put a lateral stress on the bolt head and it will distort via irradiation creep. In itself this is not an inherently damaging process, but the still-intact but deformed bolt will now be difficult and eventually impossible to remove. Such removal difficulties have already been observed but it is not yet certain that they were caused by swelling. 3. Movement of the baffle plate relative to the former will tend to turn the head of the bolt, raising one side above the surface of the plate, thereby presenting a ‘scratching’ hazard during fuel removal. Scratches possibly arising from such behavior have already been observed. 4. Differential swelling of two adjacent structural components may also open unanticipated or undesirable channels of coolant flow. © Woodhead Publishing Limited, 2010

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5. Relaxation of bolts by irradiation creep can also open up undesirable flow channels. Earlier it was found that if undesirable and/or unanticipated high-pressure flow channels were directed toward the fuel, a process called ‘baffle jetting’, the fuel integrity was imperiled. In some PWR cores the overall flow pattern of the coolant was reversed to direct such jetting away from the fuel. 6. Axially preloaded bolts will immediately begin to relax by irradiation creep and will be essentially unloaded by 10 dpa in the absence of swelling, but since most bolts are constructed from lesser-swelling steels, differential swelling will eventually reload the bolt. This reloading occurs in the 5–15 dpa interval and has both good and bad consequences [60]. This reloading has already been observed, mostly as a consequence of measuring unbolting torques and finding them higher than predicted by creep relaxation alone. 7. While swelling-induced reloading will return the bolt toward or even above its originally intended load, the maintenance of such loading will contribute to cracking and failure via irradiation-assisted stress corrosion cracking (IASCC). 8. The cracking of baffle-former bolts has been observed to accelerate at higher dose and has been identified as a concern for plant life extension, requiring that bolts be replaced periodically [61]. 9. Replacement into an already swelling plate of a cracked or broken bolt with an unirradiated bolt will subject the new bolt to an almost immediate increase in load arising from a relatively large difference in swelling rate, possibly leading to a much shorter failure time via IASCC for the replacement bolt [60]. 10. When bolts are subjected to both transverse and axial loading as a consequence of swelling, they will be even more sensitive to failure via IASCC. Ongoing analyses show that the most pronounced effect of such complex loading will occur at the formers which are one level removed from either the top and bottom former levels. In ref. 61 it was shown that these two former levels exhibited the earlier failures and highest bolt failure rates. 11. It is known that IASCC is driven partially by radiation-induced segregation at grain boundaries. Void swelling also involves segregation at void surfaces. It is therefore reasonable to anticipate that the cracking characteristics of austenitic steels may change as swelling advances. This subject has not received any attention to date. The US PWR industry is now designing and conducting a comprehensive surveillance program, not only on swelling-creep effects, but on the full range of material issues that impact all components and systems in a PWR plant [62]. The first publication addressing both prediction and surveillance

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of swelling-creep issues of a specific plant is being published as this book goes to press [63].

10.5

Second-order effects associated with or concurrent with void swelling

There are some easy to ignore second-order effects that may move to firstorder status as swelling increases during long-term plant operation. In general these second-order effects have not received much attention. The first of these is listed above as item 11 in the previous section but its existence is rather speculative in nature at this point in time.

10.5.1 Swelling-induced changes in physical properties In another less speculative example, void swelling decreases elastic moduli by ~2% per 1% swelling with possible consequences on mechanical properties and failure modes [64–66]. Void swelling also increases the electrical and thermal resistivities by ~1% per 1% swelling [67, 68], as shown in Fig. 10.30. These changes in resistivities and moduli have been proposed as a way to measure swelling in situ in PWRs [69, 70]. While changes in electrical properties are not very relevant to PWR operation, progressive changes in thermal properties require some evaluation 15 Resistance = 0.66 + 1.14*S r2 = 0.8 Shear = –0.40 – 17.45*S r2 = 0.98 Young’s = 0.17 – 2.00*S r2 = 0.98 Linear regression

Material property change, %

10

5 Resistivity 0

–5 Elastic moduli –10

–15

0

1

2

3 4 Void swelling, %

5

6

7

10.30 Void-induced property changes measured in a Russian stainless steel irradiated in the BN-350 fast reactor, showing dependence of physical properties on void swelling [67].

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in regions where large swelling is anticipated. To the first order, increasing thermal resistance will drive the steel toward higher temperature and therefore potentially higher swelling, producing a positive feedback loop to accelerate swelling. This tendency will be partially counteracted, however, by a swelling-induced reduction in the volumetric gamma heating rate and possibly by changes in dimension [68], but such a possibility has been and should continue to be examined for plant long-term operation scenarios.

10.5.2 Increasing martensite instability It appears that under some PWR-relevant irradiation conditions, low-nickel steels such as austenitic AISI 304 and 316 steels used in pressure vessel internals can develop a new form of deformation and perhaps failure when exposures greater than ~20 dpa are reached [71, 72]. Whereas current design equations for radiation-induced changes in mechanical properties assume that ductility will decrease initially and then saturate with increasing exposure, it appears that the ductility loss will be reversed at higher exposure for deformation temperatures characteristic of zero-power shutdown conditions. The enhanced ductility is a consequence of radiation-enhanced martensite formation during deformation that precludes necking and thereby produces a deformation wave that moves through the steel. Such behavior may make the steel less vulnerable to failure and may extend predicted lifetimes of components of baffle-former assemblies. However, this increasing tendency toward martensite instability may change other characteristics such as resistance to IASCC. More research is needed on this subject. While the development of void swelling is occurring concurrent with the observed onset of martensite stability, it cannot be confidently stated at this time that swelling participates directly in causing or assisting the observed instability.

10.5.3 Consequences of transmutation of nickel isotopes In nickel-containing alloys irradiated in thermalized neutron flux-spectra, the formation and reaction of nickel isotopes with thermal neutrons can lead to significant time-varying changes in dpa rate, gas formation and nuclear heating. Depending on the neutron spectrum and the nickel level of the alloy, these changes can range from insignificant to completely dominating [73]. For AISI 304 and 316 steels the effects may develop into significant considerations at higher dose levels over a 40–80-year period. Nickel has five naturally occurring stable isotopes with 58Ni comprising ~67.8% natural abundance, 60Ni comprising 26.2%, and ~6.1% total of 61 Ni, 62Ni and 64Ni. During irradiation in a highly thermalized neutron spectrum, all nickel isotopes are transmuted, primarily to the next higher

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isotopic number of nickel. Even before transmutation via thermal neutrons becomes important, however, nickel contributes to the majority of transmutant helium and hydrogen, primarily arising from reactions with neutrons above ~6 MeV. There is no natural 59Ni or 63Ni at the beginning of radiation. However, 59 Ni which has a half-life of 76,000 years is formed at significant levels from 58 Ni via thermal neutron absorption. The recoil of the 59Ni upon emission of the gamma ray produces about five displacements per event. The isotope 59Ni undergoes three strong reactions with thermal and resonance (~0.3 keV) neutrons. These reactions in order of higher cross-section are (n, g) to produce 60Ni and (n, p) and (n, a) to produce hydrogen and helium, respectively. Helium/dpa ratios on the order of 10–25 appm per dpa can be experienced along the length of a baffle bolt [1, 31] while comparable rates in fast reactors are on the order of 0.1–0.2 appm/dpa. In thermalized spectra the latter two reactions can quickly overwhelm the gas production produced at high neutron energies. Most importantly, these 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 5.10 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 event. However, the recoiling 56Fe carries 340 keV which is very large compared to most primary knock-on energies, and produces ~1701 displacements per event [73, 74]. The thermal (n, p) reaction of 59Ni produces about 1 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 [33]. Note that only 4.9 displaced atoms are created by each (n, g) recoil of 60Ni. Since 59Ni is progressively transmuted to 60Ni and 58Ni is continuously reduced in concentration, the 59Ni concentration rises to a peak level at 4  ¥  1022 n/cm2 where the 59/58 ratio peaks at ~4% and then declines, as shown in Fig. 10.31. Given the long half-life of 59Ni, its decay is not a factor and the increased damage rate is determined only to the accumulated thermal neutron fluence and the nickel content of the alloy. An extreme example of this increase in dose is shown for pure nickel in Fig. 10.32. Note that this calculated increase arises only from 59Ni (n, a) reaction. An additional but smaller increase will occur as a result of the 59Ni (n, p) reaction. At the peak 59Ni level the heating rates from the energetic (n, a) and (n, p) reactions are 0.377 and 0.023 watts per gram of nickel, significantly

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Ratio to initial value

1.6

60

Ni Natural nickel Ni-58 67.85% Ni-60 26.2%

1.2

58

0.8

0.4

0.0 21 10

343

59

Ni-61 Ni-62 Ni-64

Ni

6.1% total

Ni

1022 1023 1024 Thermal fluence, n cm–2

10.31 Transmutation-induced evolution of three nickel isotopes during irradiation in thermalized neutron spectra [31]. 100 Pure nickel in hfir-ptp Percentage increase

80

60 56

40 4

20

0

He

Fe

340 keV 1701 displacements

4.8 MeV 62 displacements

20 40 60 80 Displacements (dpa) neglecting

59

100 120 140 Ni (n, a) 56Fe reaction

160

10.32 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 T/F ratio is 2.0 [31]. The rate of increase will be increased another few percent if the 59Ni (n, p) reaction is taken into account.

larger than the neutron heating level of ~0.03 watts per gram of nickel. Thus an increase in nuclear heating of ~0.4 watts per gram of nickel must be added to the gamma heating rate at the peak 59Ni level. Fractions of the

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peak heating rates that are proportional to the 59Ni level should be added at non-peak conditions. Gamma heating is the primary cause of temperature increases in the interior of thick plates and temperature is a major variable that influences void swelling. Gamma heating is also a strong function of the thermal-to-fast neutron ratio and the neutron flux, being ~40 watts per gram in the center of the HFIR test reactor where the T/F ratio is ~2.0. In PWR near-core internals, however, the T/F ratios are lower by a factor of 2 to 10, depending on location, and the gamma heating rates in the baffle-former assembly are ~1–3 watts per gram. In this case an additional 0.4 watts per gram of nuclear heating can be a significant addition to total heating, especially for high nickel alloys. Previously the LMR and LWR communities have focused primarily only on the effect of 59Ni reactions on the gas generation rates, but it is now obvious that the displacement and heating effects must also be taken into account. Additionally, another concern may arise in that small nickel-rich phases such as gamma-prime, Ni-phosphides and G-phase may become less stable due to recoil dissolution as the 56Fe recoils originating in the precipitates, thereby altering the phase evolution in thermalized neutron spectra compared to non-thermalized spectra such as found in fast reactors. These precipitates are known to form as a direct result of irradiation and to contribute to hardening, swelling and irradiation creep processes [6]. The size of these precipitates at PWR-relevant temperatures is often comparable to or smaller than the ~80 nm range of the recoiling 56Fe atom.

10.5.4 Impact of creep

59

Ni effects on swelling and irradiation

As the dpa rate increases, the apparent swelling rate will also increase as a consequence. To the first order this effect might seem negligible for AISI 304 and 316 steels. However, the possibility should not be dismissed lightly. The potential for significant impact may be demonstrated by a recent publication on the impact of 59Ni on interpretation and extrapolation of data on irradiation creep. Since analyses of the 59Ni effect were not published until the early 1980s [74] and not widely appreciated and incorporated into dpa calculations until several years later, its unsuspected action may have influenced the results of earlier studies, as can be demonstrated by several examples. In a review on irradiation creep in 1971, Gilbert [75] showed an estimate of the temperature dependence of irradiation creep that could be derived from the very limited data available at that time. This result implicitly assumes that the composition of the alloy and the magnitude of the T/F ratio are unimportant. However, as the temperature decreased, the limited data in this set move not only toward higher nickel content but from fast reactor toward

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thermal reactor spectra, both of which would tend to underestimate the dpa level and artificially increase the creep coefficient as temperature decreased. This result has been cited several times as supporting an observed increase in the creep rate at relatively low temperatures [76, 77]. Another indication that changes in nickel content may give rise to differences in creep rate in thermalized neutron spectra is contained in a creep relaxation study by Causey and co-workers who produced the original NRU data on Inconel X-750 [76]. In the same study they also irradiated pure nickel and 304 stainless steel. As shown in Table 10.1 their derived creep coefficients at 70 and 200 °C decreased steadily with nickel content. Given the behavior observed in X-750, this result may at least partially be another expression of the 59Ni effect. Finally, in three papers by Foster and coauthors [78–80], it was shown that steady-state creep coefficients derived from thermal reactors were consistently larger than coefficients derived from fast reactors. In the first study four separate steels exhibited creep coefficients about 2.5 times larger than that observed in the same steels in fast reactors [78]. In the second study it was shown that the steady-state creep rate of Inconel X-750 increased from 1.1 to 1.9 ¥ 10–6 (MPa dpa)–1 when going from a fast reactor (EBR-II) to a thermal reactor (ETR) [79]. Although Foster’s first two papers were published in 1980 and 1988, the analysis of thermal data was based on publications from the early 1970s, before the publication of the 59Ni effect on helium production and dpa rate. Foster noted that the thermal reactor data were derived from lower flux reactors and that Lewthwaite and Mosedale had demonstrated that fast reactor (DFR) creep rates for austenitic steels increase with decreasing dose rate [81]. Thus, Foster’s analysis was consistent with their results. However, it was later shown by Garner and Toloczko that the data of Lewthwaite and Mosedale had been misinterpreted and that there was no dependence of the steady-state creep rate on displacement rate [82, 83]. Therefore it is likely that the enhancement of creep in thermal reactors observed by Foster might instead be attributed to the effect of 59Ni. Foster and coworkers have just recently presented a third paper on several austenitic steels (CW 316 SS, CW 316LN SS and SA 304L SS) and found essentially the same relative behavior where creep rates in thermal reactors are greater than in fast reactors Table 10.1 Creep coefficients (C) in units of 1030 (n/m2)–1 MPa–1, where neutrons are measured above 1.0 MeV Alloy

C (70 °C)

C (200 °C)

Ratio C(70)/C(200)

Pure nickel Inconel X-750 AISI 304

2.4 1.3 0.28

1.8 0.5 0.24

1.33 2.6 1.1

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[80]. In this third paper Foster cites the possibility that the 59Ni effect may account for the difference in creep rates between the two types of reactor spectra. However, there is still another possible reason why the creep rate may increase in thermalized spectra. As noted earlier, swelling increases strongly as swelling begins. A similar phenomenon occurs for gas bubbles produced by the 59Ni effect [83, 84], even for bubbles too small to easily image in a microscope. When the swelling rate arising either from voids or bubbles reaches only 0.01%/dpa the effective creep rate is doubled. Thus 59Ni has two powerful modes by which to accelerate the rate of irradiation creep, one involving an increase in dose rate and another by cavity-acceleration of creep.

10.5.5 Interaction between swelling and gas production and possible consequences on IASCC In a series of recent papers Garner and co-workers have shown, in apparent contradiction of Seivert’s Law, significant amounts of hydrogen can be stored in pure nickel and various austenitic alloys when irradiated in water-cooled reactors under conditions where large levels of helium and helium-nucleated cavities are formed, including in the Tihange PWR baffle bolt discussed earlier [85–87]. Initially it was thought that such storage might lead to accelerated swelling since both gases are known to facilitate and accelerate void nucleation. More recently, however, it appears that extensive co-production of helium and hydrogen in PWRs might have another unsuspected consequence. Connerman and co-workers [88] with support by Edwards and co-workers [45] have recently examined intergranular stress corrosion cracking (IGSCC) of cold-worked 316 stainless steel thimble tubes irradiated to very high dpa levels in a PWR. As the %IGSCC measured in post-irradiation slow strain rate tests climbed to ~100% with increasing dpa level, retained hydrogen was measured to climb at a correspondingly increasing rate. At 70 dpa and 330 °C, for example, helium was measured to be ~600 appm and hydrogen to be ~2500 appm. Most significantly, electron microscopy by Edwards revealed a very high density (1.6 ¥ 1023 m–3) of exceptionally small (<3 nm) cavities (probably bubbles) that not only populate the matrix but strongly coat the grain boundaries, as shown in Fig. 10.33. To image these very hard-to-see bubbles Edwards found it necessary to strongly under-focus the electron beam, an uncommon procedure. This suggests the possibility that such cavities may have been overlooked in earlier studies on lower exposure specimens. Recall from Figs. 10.12 and 10.21 that Thomas used a similar technique to image such tiny cavities in a BWR shroud and also in a PWR baffle bolt. These rather consistent observations of sub-visible cavities suggest that bubbles

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347

CW 316SS, Thimble Tube 70 dpa, 330°C

Bubbles on grain boundary

Matrix Bubbles 1.6 ¥ 1023 m–3

20 nm

–256 nm UF

20 nm

10.33 Nano-bubbles (<3 nm diameter) observed in cold-worked AISI 316 PWR thimble tube at 70 dpa and 330 °C [45]. Note significant under-focusing was used as shown in the right-hand micrograph. The specimen was measured to contain ~600 appm helium and ~2500 appm hydrogen.

filled with hydrogen and helium residing on the grain boundaries may be accelerating the development of boundary cracking. Connerman and co-workers compared their results on PWR-irradiated steels with comparable specimens irradiated in the BOR-60 fast reactor [88]. There was very little IGSCC in these specimens, very little helium and essentially no hydrogen, again suggesting a role of hydrogen in cracking. Of all the second-to-first-order effects discussed in this chapter, this may be the one that requires the greatest vigilance, surveillance and testing in the future.

10.5.6 Gamma-to-ferrite transformations as a consequence of void growth AISI 304 stainless steel irradiated in the EBR-II fast reactor has been shown to decompose at ~400 °C into coated voids which are surrounded by shells of austenite phase while the matrix between the coated voids has transformed to ferrite [89]. This transformation is the result of radiation-induced segregation of nickel to void surfaces and increased concentration of chromium in the spaces between the voids. Figure 10.34 presents micrographs demonstrating this phenomenon. While this phenomenon has not yet been observed in PWRs, the potential for such a transformation may arise in the hot corners of the former plates as swelling becomes increasingly larger. Upon such transformation there is

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100 nm (a)

100 nm (b)

100 nm (c)

10.34 Micrographs of decomposed 304 stainless steel after irradiation in EBR-II at ~399°C to 17.5 dpa; (a) bright-field, showing coated voids, (b) weak-beam dark-field (110-a), showing ferrite matrix between the voids, and (c) dark-field (111-g), showing austenite shells on the voids [89]. © Woodhead Publishing Limited, 2010

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a volume increase in the transformed austenite to ferrite regions on the order of ~4% which will also contribute to strains and stresses.

10.6

Conclusion

Whereas swelling and irradiation creep were previously thought not to be major concerns for light water reactors, it now appears that these processes must be considered, especially in the case of austenitic steel internals in PWRs and even more importantly as these reactors move beyond their originally designed operational lives to long-term operation of 60 and perhaps 80 years. The most swelling-vulnerable components are the baffle-former-barrel assembly that frames the PWR core and the current expectation is that these assemblies should serve without replacement for the entire extended life, although they have been replaced in three Japanese PWRs, demonstrating that replacement can be performed, albeit with an economic penalty associated both with replacing the assembly and in the safe storage of the old highly radioactive assembly indefinitely. Plant outage costs may also be significant, as well as procurement delays. Unfortunately, the internals of Western PWRs and Russian WWERs were constructed with two of the most swelling-prone steels that are commercially available. Thus void swelling and its companion irradiation creep have already shown their potential to impact the operation and perhaps the safety of PWRs during extended operation. In particular there is a concern that perhaps various second-order and easily overlooked effects may grow to first-order importance during extended life operation. The most prominent of these second-order effects may be the retention and continued growth of hydrogen-storing cavities. In this review a summary of what is known or surmised about the interactive effect of void swelling and irradiation creep has been presented. The general outline of potential consequences appears to be reasonably well founded, but there are significant uncertainties in the magnitude and time-frame on which the consequences will present themselves. Additional information will only become available as surveillance of plants in the long-term operation mode continues. It is thus important that the PWR industry continuously and conscientiously conducts such surveillance and monitoring activities.

10.7

References

1. F. A. Garner, L. R. Greenwood and D. L. Harrod, ‘Potential High Fluence Response of Pressure Vessel Internals Constructed from Austenitic Stainless Steels’, Proc. Sixth Intern. Symp. on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, San Diego, CA, August 1–5, 1993, pp. 783–790.

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2. F. A. Garner, ‘Materials Issues Involving Austenitic Pressure Vessel Internals Arising From Void Swelling and Irradiation Creep’, Trans. Am. Nucl. Soc. 71 (1994) 190. 3. F. A. Garner and M. B. Toloczko, ‘Irradiation Creep and Void Swelling of Austenitic Stainless Steels at Low Displacement Rates in Light Water Energy Systems’, J. Nucl. Mater. 251 (1997) 252–261. 4. F. A. Garner, D. J. Edwards, S. M. Bruemmer, S. I. Porollo, Yu. V. Konobeev, V. S. Neustroev V. K. Shamardin and A. V. Kozlov, ‘Recent Developments Concerning Potential Void Swelling of PWR Internals Constructed from Austenitic Stainless Steels’, Proc. Fontevraud 5, Contribution of Materials Investigation to the Resolution of Problems Encountered in Pressurized Water Reactors, September 23–27, 2002, on CD format, no page numbers. 5. H. T. Tang and J. D. Gilreath, ‘Aging Research Management of PWR Internals’, Proc. Fontevraud 5, Contribution of Materials Investigation to the Resolution of Problems Encountered in Pressurized Water Reactors, September 23–27, 2002, on CD format, no page numbers. 6. F. A. Garner, Chapter 6: ‘Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors’, Materials Science and Technology, Vol. 10A: A Comprehensive Treatment, VCH Publishers, Weinheim, 1994, pp. 419–543. 7. H. R. Brager, F. A. Garner, D. S. Gelles and M. L. Hamilton, ‘Develop­ment of Reduced Activation Alloys for Fusion Service’, J. Nucl. Mater. 133–134 (1985) 907–911. 8. F. A. Garner and D. S. Gelles, ‘Neutron-Induced Swelling of Commercial Alloys at Very High Exposures’, Proceedings of Symposium on Effects of Radiation on Materials: 14th International Symposium, ASTM STP 1046, N. H. Packan, R. E. Stoller and A. S. Kumar, eds., American Society for Testing and Materials, Philadelphia, PA, 1990, Vol. II, pp. 673–683. 9. F. A. Garner, ‘Recent Insights on the Swelling and Creep of Irradiated Austenitic Alloys’, J. Nucl. Mater. 122–123 (1984) 459–471. 10. S. I. Porollo, S. V. Shulepin, Yu. V. Konobeev and F. A. Garner, ‘Influence of Silicon on Swelling and Microstructure in Russian Austenitic Stainless Steels Irradiated to High Neutron Doses’, J. Nucl. Mater. 378 (2008) 17–24. 11. B. J. Makenas, S. A. Chastain and B. C. Gneiting, ‘Dimensional Changes in FFTF Austenitic Cladding and Ducts’, Westinghouse Hanford Company Report WHCSA-0933VA, Richland WA, 1990. 12. E. R. Gilbert, D. C. Kaulitz, J. J. Holmes and T. T. Claudsen, Proc. Conf. on Irradiation Embrittlement and Creep in Fuel Cladding and Core Components, British Nuclear Energy Society, London, 1972, pp. 239–251. 13. B. J. Makenas, S. A. Chastain and B. C. Gneiting, ‘Dimensional Changes in FFTF Austenitic Cladding and Ducts’, Proc. LMR: A Decade of LMR Progress and Promise, ANS, La Grange Park, IL, pp. 176–183. 14. M. L. Hamilton, F. H. Huang, W. J. S. Yang and F. A. Garner, ‘Mechani­cal Properties and Fracture Behavior of 20% Cold-Worked 316 Stainless Steel Irradiated to Very High Exposures’, Effects of Radiation on Materials: Thirteenth International Symposium (Part II) Influence of Radiation on Material Properties, ASTM STP 956, F. A. Garner, N. Igata and C. H. Henager, Jr., eds., ASTM, Philadelphia, PA, 1987, pp. 245–270. 15. V. S. Neustroev and V. K. Shamardin, ‘Relation between the Microstructure and the Nature of the Fracture of Cr18Ni10Ti Steel Irradiated with Neutrons to 70 dpa’ (in Russian), Atomnaya Energiya 70, 4 (1991) 345–348.

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16. V. S. Neustroev and F. A. Garner, ‘Very High Swelling and Embrittlement Observed in a Fe-18Cr-10Ni-Ti Hexagonal Fuel Wrapper Irradiated in the BOR-60 Fast Reactor’, J. Nucl. Mater. 378 (2008) 327–332. 17. V. S. Neustroev and F. A. Garner, ‘Severe Embrittlement of Neutron Irradiated Austenitic Steels Arising from High Void Swelling’, J. Nucl. Mater. 386–388 (2009) 157–160. 18. D. L. Porter and F. A. Garner, ‘Irradiation Creep and Embrittlement of AISI 316 at Very High Neutron Fluences’, J. Nucl. Mater. 159 (1988) 114–121. 19. V. S. Neustroev, Z. E. Ostrovsky, A. A, Teykovtsev, V. K. Shamardin and V. V. Yakolev, ‘Experimental Studies of the Failure of Irradiated Ducts in the BOR-60 Reactor’, (in Russian) Proc. 6th Russian Conference on Reactor Material Science, September 11–15, 2000, Dimitrovgrad, Russia. 20. Micrograph courtesy of J. J. Laidler, Hanford Engineering Development Laboratory, Richland WA. 21. R. l. Fish, J. L. Straalsund, C. W. Hunter and J. J. Holmes, ‘Swelling and Tensile Property Evaluations of High Fluence EBR-II Thimble’, Effects of Radiation on Substructure and Mechanical Properties of Metals and Alloys, ASTM STP 529, American Society for Testing and Materials, Philadelphia, PA, 1973, pp. 149– 164. 22. A. Fissolo, R. Cauvin, J.-P. Hugot and V. Levy, ‘Influence of Swelling on Irradiated CW Titanium Modified 316 Embrittlement’, Effects of Radiation on Materials: Fourteenth Intern. Symp. (Vol. II), ASTM STP 1046, N. H. Packan, R. E. Stoller and A. S. Kumar, eds., American Society for Testing and Materials, Philadelphia, PA, 1990, pp. 700–713. 23. F. A. Garner and D. L. Porter, ‘A Reassessment of the Swelling Behavior of AISI 304L Stainless Steel’, 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. 24. L. E. Thomas, D. Edwards, K. Asano, S. Ooki and S. Bruemmer, ‘Crack-Tip Characteristics in BWR Service Conditions, 10th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, 2001, issued on CD format, no page numbers. 25. S. Yaguchi and J. Uchiyama, ‘Core Internals Replacement Method in Japanese PWR’, Fontevraud-6 Symposium on Contribution of Materials Investigations to Improve the Safety and Performance of LWRs’, September 18–22, 2006, Fontevraud, France, paper A153-T02. 26. T. Okita, N. Sekimura, F. A. Garner, L. R. Greenwood, W. G. Wolfer and Y. Isobe, ‘Neutron-Induced Microstructural Evolution of Fe-15Cr-16Ni Alloys at ~400 °C during Neutron Irradiation in the FFTF Fast Reactor’, 10th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, 2001, issued on CD format, no page numbers. 27. T. Okita, N. Sekimura, T. Iwai and F. A. Garner, ‘Investigation of the Synergistic Influence of Irradiation Temperature and Atomic Displacement Rate on the Microstructural Evolution of Ion-Irradiated Model Austenitic Alloy Fe-15Cr-16Ni’, 10th International Conference on Environmental, Deregulation of Materials in Nuclear Power Systems – Water Reactors, 2001, issued on CD. 28. T. Okita, W. G. Wolfer, T. Sato, N. Sekimura and F. A. Garner, ‘Influence of Composition, Helium Generation Rate and dpa Rate on Neutron-Induced Swelling of Fe-15Cr-1Ni-0.25Ti Alloys in FFTF at ~400 ºC’, 11th International Conference

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29.

30.

31.

32.

33.

34.

35. 36.

37.

38. 39.

40.

41.

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42. J. P. Foster, D. L. Porter, D. L. Harrod, T. R. Mager and M. G. Burke, ‘316 Stainless Steel Cavity Swelling in a PWR’, J. Nucl. Mater. 224 (1995) 207–215. 43. K. Fujii, K. Fukuya, G. Furutani, T. Torimaru, A. Kohyama and Y. Katoh, ‘Swelling in 316 Stainless Steel Irradiated to 53 dpa in a PWR’, 10th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, 2001, issued on CD format, no page numbers. 44. K. Fukuya, K. Fujii, H. Nishioka and Y. Kitsunai, ‘Evolution of Microstructure and Microchemistry in Cold-worked Stainless Steels under PWR Irradiation’, J. Nucl. Sci. and Tech. 43, 2 (2006) 159–173. 45. D. J. Edwards, F. A. Garner, S. M. Bruemmer and Pal Efsing, ‘Nano-cavities observed in a 316SS PWR Flux Thimble Tube Irradiated to 33 and 70 dpa’, J. Nucl. Mater. 384 (2009) 249–255. 46. V. S. Neustroev, V. G. Dvoretzky, Z. E. Ostrovsky, V. K. Shamardin and G. A. Shimansky, ‘Investigation of Microstructure and Mechanical Properties of 18Cr10Ni-Ti Steel Irradiated in the Core of VVER-1000 Reactor’, Effects of Radiation on Materials, 21st Intern. Symp., ASTM STP 1447, M. L. Grossbeck, T. R. Allen, R. G. Lott and A. S. Kumar, eds., ASTM International, Philadelphia, PA, 2004, pp. 32–45. 47. F. A. Garner, S. I. Porollo and Yu. V. Konobeev, V. S. Neustroev and O. P. Maksimkin, ‘Void Swelling of Russian Austenitic Stainless Steels at PWR-relevant Displacement Rates and Temperatures’, Fontevraud-6 Symposium on Contribution of Materials Investigations to Improve the Safety and Performance of LWRs, September 18–22, 2006, Fontevraud, France, pp. 637–648. 48. S. I. Porollo, A. M. Dvoriashin, Yu. V. Konobeev, A. A. Ivanov, S. V. Shulepin and F. A. Garner, ‘Microstructure and Mechanical Properties of Austenitic Stainless Steel 12X18H9T after Neutron Irradiation in the Pressure Vessel of BR-10 Fast Reactor at Very Low Dose Rates’, J. Nucl. Mater. 359 (2006) 41–49. 49. V. S. Neustroev, V. K. Shamardin, Z. E. Ostrovsky, A. M. Pecherin and F. A. Garner, ‘Temperature-Shift of Void Swelling Observed at PWR-relevant Temperatures in Annealed Fe-18Cr-10Ni-Ti Stainless Steel Irradiated in the Reflector Region of BOR-60’, Effects of Radiation on Materials: 19th International Symposium, ASTM STP 1366, M. L. Hamilton, A. S. Kumar, S. T. Rosinski and M. L. Grossbeck, Eds., American Society for Testing and Materials, Philadelphia, PA, 2000, pp. 792–800. 50. S. I. Porollo, Yu. V. Konobeev, A. M. Dvoriashin, A. N. Vorobjev, V. M. Krigan and F. A. Garner, ‘Void Swelling at Low Displacement Rates in Annealed X18H10T Stainless Steel at 4 to 56 dpa and 280–332 ∞C’, J. Nucl. Mater. 307–311 (2002) 339–342. 51. S. I. Porollo, Yu. V. Konobeev, A. M. Dvoraishin, V. M. Krigan and F. A. Garner, ‘Determination of the Lower Temperature Limit of Void Swelling of Stainless Steels at PWR-relevant Displacement Rates’, 10th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, 2001, issued on CD format, no page numbers. 52. O. P. Maksimkin, K. V. Tsai, L. G. Turubarova, T. Doronina and F. A. Garner, ‘Characterization of 08Cr16Ni11Mo3 Stainless Steel Irradiated in the BN-350 Reactor’, J. Nucl. Mater. 329–333 (2004) 625–629. 53. O. P. Maksimkin, K. V. Tsai, L. G. Turubarova, T. A. Doronina and F. A. Garner, ‘Void Swelling of AISI 321 Analog Stainless Steel Irradiated at Low dpa Rates in the BN-350 Reactor’, J. Nucl. Mater. 367–370 (2007) 990–994.

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54. F. A. Garner, M. L. Hamilton, D. L. Porter, T. R. Allen, T. Tsutsui, M. Nakajima, T. Kido and T. Ishii, ‘The Influence of Displacement Rate on the Void Swelling of Annealed AISI 304 Stainless Steel in the EBR-II Fast Reactor’, J. Nucl. Mater., in preparation. 55. G. M. Bond, B. H. Sencer, F. A. Garner, M. L. Hamilton, T. R. Allen and D. L. Porter, ‘Void Swelling of Annealed 304 Stainless Steel at ~370–385 °C and PWR-relevant Displacement Rates’, 9th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, 1999, pp. 1045–1050. 56. F. A. Garner and B. J. Makenas, ‘Recent Experimental Results on Neutron-induced Void Swelling of AISI 304 Stainless Steel Concerning its Interactive Dependence on Temperature and Displacement Rate’, Fontevraud-6 Symposium on Contribution of Materials Investigations to Improve the Safety and Performance of LWRs, September 18–22, 2006, Fontevraud, France, pp. 625–636. 57. J. P. Foster and J. E. Flinn, ‘Residual Stress Behavior in Fast Neutron Irradiated SA AISI 304L Stainless Steel Cylindrical Tubing’, J. Nucl. Mater. 89 (1980) 99–112. 58. F. A. Garner, J. J. Laidler and G. L. Guthrie, ‘Development and Evalua­tion of a Stress-Free Swelling Correlation for 20% Cold-Worked 316 Stain­less Steel’, Irradiation Effect on the Microstruc­ture and Properties of Metals, ASTM STP 611, F R Shober, ed., ASTM, Philadelphia, PA, 1976, pp. 208–276. 59. J. F. Bates, F. A. Garner and F. M. Mann, ‘The Effects of Solid Trans­mutation Products on Swelling in AISI 316 Stainless Steel’, J. Nucl. Mater. 103–104 (1981), 999. 60. E. P. Simonen, F. A. Garner, N. A. Klymyshyn and M. B. Toloczko, ‘Response of PWR Baffle-Former Bolt Loading to Swelling, Irradiation Creep and Bolt Replacement as Revealed Using Finite Element Modeling’, Proc. 12th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, 2005, pp. 449–456. 61. P. M. Scott, M-C. Meunier, D. Deydier, S. Silvestre and A. Trenty, ‘An Analysis of Baffle/Former Bolt Cracking in French PWRs’, Environmentally Assisted Cracking: Predictive Methods for Risk Assessment and Evaluation of Materials, Equipment and Structures, ASTM STP 1401, R. D. Kane, ed., American Society for Testing and Materials, Philadelphia, PA, 2000. 62. H. T. Tang, ‘Aging Management of Reactor Internals and License Renewal of US PWR Plants’, Fontevraud-6 Symposium on Contribution of Materials Investigations to Improve the Safety and Performance of LWRs, September 18–22, 2006, Fontevraud, France, Paper A145–T02, issued on CD. 63. G. A. Gardner, A. Demma, F. J. Marx, T. Liszkai and J. Rashid, ‘Modeling and Simulation of Aging Effects in Irradiated PWR Reactor Internal Components’, Proc. 14th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, 2009, Virginia Beach, VA, in press. 64. M. Marlowe and W. K. Appleby, ‘Measurements of the Effects of Swelling on the Young’s Modulus of Stainless Steels’, Trans. ANS 16 (1973) 95–96. 65. J. L. Straalsund and C. K. Day, ‘Effect of Neutron Irradiation on the Elastic Constants of Type 304 Stainless Steel’, Nucl. Tech. 20 (1973) 27. 66. R. L. Trantow, ‘Ultrasonic Measurement of Elastic Properties in Irradiated 304 Stainless Steel’, Hanford Engineering Development Laboratory Report HEDL-TME 73-92, Richland, WA, 1973. 67. A. V. Kozlov, E. N. Shcherbakov, S. A. Averin and F. A. Garner, ‘The Effect of Void Swelling on Electrical Resistance and Elastic Moduli in Austenitic Steels’,

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68.

69.

70.

71.

72.

73.

74. 75. 76.

77. 78.

79.

80.

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82.

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Effects of Radiation on Materials, ASTM STP 1447, M. L. Grossbeck, T. R. Allen, R. G. Lott and A. S. Kumar, eds., ASTM International, West Conshohocken, PA, 2004, pp. 66–77. W. G. Wolfer and F. A. Garner, ‘Effective Thermophysical and Elastic Properties of Materials with Voids’, Damage Analysis and Fundamental Studies Quarterly Progress Report No. 25 (May 1984) DOE/ER-0046/17, p. 58 I. I. Balachov, E. N. Shcherbakov, A. V. Kozlov, I. A. Portnykh and F. A. Garner, ‘Influence of Irradiation-Induced Voids and Bubbles on Physical Properties of Austenitic Structural Alloys’, J. Nucl. Mater. 329–333 (2004) 617–620. I. I. Balachov, F. A. Garner, Y. Isobe, M. Sagisaka and H. T. Tang, ‘NDT Measurements of Irradiation Induced Void Swelling’, 11th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, 2003, pp. 640–646. M. N, Gusev, P. Maksimkin, I. S. Osipov and F. A Garner, ‘Anomalously Large Deformation of 12Cr18Ni10Ti Austenitic Steel Irradiated to 55 dpa at 310 °C in the BN-350 Reactor’, J. Nucl. Mater. 386–388 (2009) 273–276. M. N. Gusev, O. P. Maksimkin, I. S. Osipov, N. S. Silniagina, and F. A. Garner, ‘Unusual Enhancement of Ductility Observed during Evolution of a ‘Deformation Wave’ in 12Cr18Ni10Ti Stainless Steel Irradiated in BN-350’, J. ASTM Int., 6, 7 (2009) paper ID JAI102062. F. A. Garner, M. Griffiths, L. R. Greenwood and E. R. Gilbert, ‘Impact of Ni-59 (n, a) and (n, p) Reactions on dpa Rate, Heating Rate, Gas Generation and Stress Relaxation in LMR, LWR and CANDU Reactors’, Proc. 14th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors 2009, Virginia Beach, VA, in press. L. R. Greenwood, ‘A New Calculation of Thermal Neutron Damage and Helium Production in Nickel’, J. Nucl. Mater. 116 (1983) 137–142. E. R. Gilbert, ‘In-Reactor Creep of Reactor Materials’, Reactor Technol. 14 (1971) 258–285. A. R. Causey, G. J. C. Carpenter and S. R. Ewen, ‘In-Reactor Stress Relaxation of Selected Metals and Alloys at Low Temperatures’, J. Nucl. Mater. 90 (1980) 216–223. M. L. Grossbeck and L. K. Mansur, ‘Low Temperature Irradiation Creep of Fusion Reactor Structural Materials’, J. Nucl. Mater. 179–181 (1991) 130–134. J. P. Foster and A. Boltax, ‘Correlation of Irradiation Creep Data Obtained in Fast and Thermal Neutron Spectra with Displacement Cross Sections’, J. Nucl. Mater. 89 (1980) 331–337. J. P. Foster and C. M. Mildrum, ‘Correlation of Inconel X750 Stress Relaxation Data Obtained in Thermal and Fast Neutron Reactors’, J. Nucl. Mater. 151 (1988) 135–139. J. P. Foster and T. Karlsen, ‘Irradiation Creep and Irradiation Stress Relaxation of 316 and 304L Stainless Steel’, Proc. 14th International Conference on Environmental Degradation of Materials in Nuclear Power Systems –Water Reactors, 2009, Virginia Beach, VA, in press. G. W. Lewthwaite and D. Mosedale, ‘The Effects of Temperature and Dose Rate Variations on the Creep of Austenitic Stainless Steels in the Dounreay Fast Reactor’, J. Nucl. Mater. 90 (1980) 205–215. F. A. Garner and M. B. Toloczko, ‘Irradiation Creep and Void Swelling of Austenitic Stainless Steels at Low Displacement Rates in Light Water Energy Systems’, J. Nucl. Mater. 251 (1997) 252–261. © Woodhead Publishing Limited, 2010

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83. F. A. Garner, M. B. Toloczko and M. L. Grossbeck, ‘The Dependence of Irradiation Creep in Austenitic Alloys on Displacement Rate and Helium to dpa Ratio’, J. Nucl. Mater. 258–263 (1998) 1718–1724. 84. C. H. Woo and F. A. Garner, ‘Contribution to Irradiation Creep Arising from GasDriven Bubble Growth’, J. Nucl. Mater. 271–272 (1999) 78–83. 85. F. A. Garner, B. M. Oliver, L. R. Greenwood, D. J. Edwards, S. M. Bruemmer and M. L. Grossbeck, ‘Generation and Retention of Helium and Hydrogen in Austenitic Steels Irradiated in a Variety of LWR and Test Reactor Spectral Environments’, 10th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors, 2001, issued on CD format, no page numbers. 86. L. R. Greenwood, F. A. Garner, B. M. Oliver, M. L. Grossbeck and W. G. Wolfer, ‘Surprisingly Large Generation and Retention of Helium and Hydrogen in Pure Nickel Irradiated at High Temperatures and High Neutron Exposures’, Effects of Radiation on Materials, ASTM STP 1447, M. L. Grossbeck, T. R. Allen, R. G. Lott and A. S. Kumar, eds., ASTM International, West Conshohocken PA, 2004, pp. 529–539, J. ASTM Int. 1, 4 (2004) Paper ID JAI11365. 87. F. A. Garner, E. P. Simonen, B. M. Oliver, L. R. Greenwood, M. L. Grossbeck, W. G. Wolfer and P. M. Scott, ‘Retention of Hydrogen in FCC Metals Irradiated at Temperatures Leading to High Densities of Bubbles or Voids’, J. Nucl. Mater. 356 (2006) 122–135. 88. J. Connerman, R. Shogan, K. Fujimoto, T. Yonezawa and Y. Yamaguchi, ‘Irradiation Effects in a Highly Irradiated Cold Worked Stainless Steel Removed from a Commercial PWR’, Proc. 12th International Conference on Environmental Degradation of Materials in Nuclear Power System – Water Reactors, 2005, pp. 277–287. 89. D. L. Porter, F. A. Garner and G. M. Bond, ‘Interaction of Void-Induced Phase Instability and Subsequent Void Growth in AISI 304 Stainless Steel’, Effects of Radiation on Materials: 19th International Symposium, ASTM STP 1366, M. L. Hamilton, A. S. Kumar, S. T. Rosinski and M. L. Grossbeck, eds., American Society for Testing and Materials, Philadelphia, PA, 2000, pp. 884–893.

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11

Irradiation hardening and materials embrittlement in light water reactor (LWR) environments

M. B r u m o v s k y, Nuclear Research Institute Rez plc, Czech Republic

Abstract: This chapter deals with the explanation of the effect of neutron irradiation damage as one of the most important ageing mechanisms in ferritic reactor pressure vessel (RPV) materials. Irradiation conditions, as well as the initial mechanical properties of the RPV material, are the leading parameters that determine the value and character/nature of irradiation damage that can be characterized as irradiation hardening and irradiation embrittlement. While the first one is usually detected by the increase in the yield stress of materials, the latter is characterized by the shift to higher temperatures of the ductile-to-brittle transition temperature (DBTT) obtained from different types of tests – Charpy notch impact or static/dynamic/arrest fracture toughness tests. Key words: neutron irradiation on RPV, radiation hardening, radiation embrittlement, transition temperature shift, yield stress increase.

11.1

Introduction

Nuclear power plant (NPP) operating equipment is subjected to a variety of chemical, mechanical and physical conditions during service. Such stressors lead to changes, over time, in the materials, which are caused and driven, for example, by the effects of varying loads, flow conditions, corrosion, temperature and neutron irradiation. Time-dependent changes in the mechanical and physical properties of these components are referred to as ageing. The effects of ageing become evident as a reduction in design margins. During the operation of a NPP, the wall of the RPV is exposed to high energy (e.g. >0.1 MeV) neutron irradiation, which results in localized embrittlement of the steel and welds in the area of the reactor core (i.e. the RPV core beltline region). Ageing effects of the RPV have the potential to be life-limiting for a NPP as it is impossible, or economically unviable, to replace the RPV if its mechanical properties (toughness levels) degrade significantly. Accordingly, research on irradiation embrittlement of RPV steels has been the subject of significant international efforts. Over the past three decades, developments in fracture mechanics have led to a number of 357 © Woodhead Publishing Limited, 2010

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consensus standards and codes for determining essential fracture toughness parameters and associated uncertainties as derived from the available RPV embrittlement databases. This understanding has resulted in remarkable progress in developing a mechanistic and also physical understanding of neutron irradiation embrittlement.

11.2

Irradiation conditions

The active fuel core of light water reactors (LWRs) is the source of neutron irradiation, since fuel fission processes of 235uranium (235U) creates energetic neutrons, isotopes and gamma radiation, for example. This main damage in RPV materials is caused by neutrons, and they have different energies, ranging from thermal (approximately 0.025 eV) to fast neutrons, with energies up to 17 MeV. Furthermore, high energy gamma radiation (mostly between 0.5 to several MeV) is also present. The original neutron fission energy spectrum changes due to scattering with hydrogen atoms in the coolant water: this spectrum is thermalized, i.e. its energy is decreased and also its flux attenuates rapidly as it leaves the active core. The water reflector around the active core serves to return some neutrons back to the active core, and also for decreasing of neutron flux before it approaches the RPV wall. Neutron flux levels and hence fluence with time, differs between reactor types – generally, the neutron flux (hence fluence over time) on pressurized water reactors (PWRs) RPVs is higher than that found in boiling water reactors (BWRs) RPVs due to their different design and operation conditions. But, there are also large differences between individual PWR RPVs – older designs are characterized by higher fluences, and WWER (Russian design of PWR) RPVs are damaged by higher fluence due to the requirements for possible transportation of RPVs by land that resulted in a smaller RPV diameter, and thus also smaller thickness of the water reflector (i.e. resulting relatively smaller water gap between the active fuel core and the RPV inner wall). Comparison of different designs is given in Table 11.1. The neutron energy spectrum changes when it crosses from the active core to the RPV wall and then through the vessel wall. Irradiation damage in ferritic RPV steels is created by high energy neutrons, with a threshold between 0.1 and 1 MeV. Thus, for comparison and easier use, fluxes and fluences are determined for energies above some threshold: 1 MeV is applied for PWR designs while 0.5 MeV is used for WWER designs. The ratio between fluxes with different threshold energies varies and depends on reactor design and the location under consideration, but approximately it can be taken that fluxes with a threshold of 0.5 MeV are 1.6 times larger than those with 1 MeV. Sometimes, another flux description is used – dpa = displacement per atom, which represents the calculated number of atoms in the cascade made by one knocked-on atom after its collision with a fast

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Table 11.1 Operating maximum lifetime fluence for WWERs, PWRs and the BWR RPVs Reactor type

Flux, n.m–2 · sec–1 (E > 1 MeV)

Lifetime* fluence, n.m–2 (E > 1 MeV)

WWER-440 core weld WWER-440 beltline centre WWER-1000 PWR (Westinghouse)/USA PWR (B&W Babcock and   Wilcox/USA) PWR (KWU) (Kraftwerk   Union/German) BWR (typical for General   Electric GE/USA)

1.2 ¥ 1015 1.5 ¥ 1015 3–4 ¥ 1014 4 ¥ 1014 1.2 ¥ 1014

1.1 1.6 3.7 4¥ 1.2

2.7 ¥ 1014(old) – 3 ¥ 1013(new) 4 ¥ 1013

2.8 ¥ 1023(old) – 3 ¥ 1022 (new) 4 ¥ 1022

¥ 1024 ¥ 1024 ¥ 1023 1023 ¥ 1023

neutron. The ratio between neutron fluence and dpa also depends on design and location, but usually a fluence of 1023 m–2 is taken approximately to be equal to 0.015 dpa. Irradiation temperature is a second important parameter that determines the value of radiation damage, as damage is a thermally activated process. Most PWRs and BWR are operated at 288 °C (550 °F) with the exception of some older ones, and all WWER-440 types are operated at a temperature of 270 °C. Quite different irradiation temperatures were present in gas-cooled reactor (GCR) RPVs, being in the range between 150 °C and 350 °C in different parts of the RPV.

11.3

Nature of radiation damage

The change in a material’s strength and toughness due to irradiation is determined by the interaction of irradiation-induced defects with an existing dislocation system and the system that develops in the process of plastic deformation of crystals. The theory of irradiation hardening was first proposed by Seeger [1]. This theory accounts for the role of focusing of inter-atomic collisions in cascade processes, by virtue of which a zone of high vacancy concentration appears where the cascading areas of damage occur. It possesses a high concentration of vacancies and is referred to as a depleted zone. Along its periphery, the knocked-out/displaced atoms are located at interstitial sites. The complex defects of this kind are strong barriers for mobile dislocations, but they may be overcome by the application of stresses with the participation of thermal activation processes. For crystals with such zones, Seeger’s theory predicts the critical shear stress increase is proportional to F1/2 (F is neutron fluence). As already mentioned, irradiation effects in RPV steels involve many complex and interacting mechanisms. However, most of the experimental and

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simulation results are in agreement with the following simplified scenario, explaining the neutron irradiation-induced damage in these steels [2]: 1. Direct matrix damage due to neutron bombardment can be assumed to be given by a square-root dependence on fluence for a given material and a given temperature. 2. During matrix damage formation, impurity copper (Cu) (in solution), together with other elements, is known to lead to a precipitation mechanism of matrix-coherent nano-precipitates and this also causes matrix hardening and embrittlement. Such a mechanism is assumed to continue until saturation occurs at the prevailing conditions, depending on the available amount of precipitates that can be generated (i.e. the impurity Cu concentration). These defects are usually called copper-rich precipitates. 3. Other elements, like phosphorus (P), can segregate, in grains or at grain boundaries, also in combination with matrix damage or attracted into the Cu-rich precipitates. Since diffusion of segregates also plays a role, this mechanism becomes rather difficult to understand in detail. This effect is not common for US types of PWR while it is common for the first generation of WWER-440/V-230 type RPVs with P content up to 0.055 mass %. These mechanisms are summarized in Fig. 11.1 and Table 11.2, and the effect of these mechanisms on the total level of irradiation embrittlement of RPV steel is shown in Fig. 11.2 [2]. In high-nickel (Ni) RPV steels (with Ni content more than approx. 1.2 mass %), another damage mechanism can be observed, as previously predicted Vacancies, intersticials, transmutations, etc.

Cu, P, Ni, etc.

Ultra-fine precipitates Cu, etc. PP Segregation at grain boundary P, Mn, etc.

Ultra-fine Precipitates Cu, etc.

ff Di

PPP

u

sio

n

P Ni

P P P P

Segregation at grain boundary P, Mn, etc.

Vacancies, intersticials, transmutations, etc.

11.1 Schematic embrittlement process for RPV materials.

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Table 11.2 Embrittlement mechanisms considered Embrittlement mechanism

Origin of the effect

Direct matrix damage Precipitation hardening the matrix Segregation

Due to neutron bombardment Cu is the leading element P is recognized as a segregating element

200 180 160

DTshift, °C

140

Total

120 100 80

Precipitation (Cu lead)

60

Segregation (P lead)

40

Direct matrix damage

20 0 1.00E+18 5.10E+19

1.01E+20 1.51E+20 Fluence, n cm–2

2.01E+20

11.2 Schematic diagram showing the effect of three damage mechanisms on irradiation embrittlement of RPV steel.

based on thermodynamic considerations and modelling. This ‘late blooming effect’ can cause an additional increase in yield stress and higher transition temperature shifts at larger neutron fluences, where a saturation trend usually takes place (see Fig. 11.3). Key issues are the combined effects of irradiation temperature, content of Ni, Mn and Cu and neutron flux/fluence. This effect is not properly studied as it is mostly observed for neutron fluences over 6 ¥ 1023 m–2 which are larger than end-of-life fluences for 40 years of PWR operation. Thus, most of these mechanisms are classified as ‘hardening’ ones (matrix damage, copper-rich precipitates, partially also segregations when they are inside grains), but P segregation on grain boundaries is classified as ‘non-hardening’ (non-hardening embrittlement, since not detectable with conventional hardness tests). The last type of embrittlement can manifest itself as intergranular (grain boundary) fracture, rather than the usual transgranular cleavage fracture. Thus, their effect on radiation hardening and embrittlement can be quite different and thus can affect, for example, the ratio between the tensile yield stress increase and the Charpy ductile-to-brittle-transition temperature (DBTT) shift.

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Low Ni – 1.2%

160 140

DBTTshift, °C

120 100 80 60 40 20 0

0

20

40

60

80 100 120 Fluence, 1018 n/cm2

140

160

180

200

11.3 ‘Late blooming effect’ in high-nickel WWER-1000 weld metals.

11.4

Irradiation hardening and embrittlement

Regarding their effects on material properties, the ultrafine (nanometre) microstructural features mentioned above act as effective dislocation obstacles and thus increased applied stress is required to move dislocations through and around them. As radiation exposure increases, the number of ultrafine obstacles increases and higher stresses are required to create dislocation motion, with a resulting increase in the yield stress of the material. The yield stress increase results in higher temperatures required to keep the yield stress below the cleavage fracture strength, especially near the tip of a crack where large stress and strain concentrations exist. Thus, the fracture toughness transition temperature is increased and is the measure used to describe the radiation induced embrittlement. The effect of the yield strength increase on the DBTT shift is shown in Fig. 11.4. Irradiation damage in RPV steels is characterized by the following changes: ∑ ∑

irradiation-induced hardening: – yield stress and ultimate tensile strength increase, – increase in hardness, irradiation embrittlement: – decrease in ductility – decrease in elongation and reduction of area (typically measured on tensile specimens) – decrease in toughness – shift of DBTT of Charpy impact test and fracture toughness tests to higher temperatures relative to unirradiated material.

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363

Flow stress irradiated

Stress

Fracture stress

Irradiation strengthening

DT shift

Flow stress unirradiated T1

T2

Temperature

Engineering stress

11.4 Schematic diagram showing how the irradiation-induced strength increase results in an upward shift in the DBTT.

Irradiated Unirradiated

Ferritic steel Engineering strain

11.5 Schematic diagram showing the effect of increasing neutron fluence on the tensile stress–strain diagram for typical ferritic RPV steel.

Typical irradiation hardening, resulting in the change in tensile properties (stress–strain diagram) for RPV steel is shown in Fig. 11.5: yield stress and ultimate tensile strength are increased with increase in neutron fluence. As the yield strength (Rp0.2) increase is faster than for that observed for the ultimate tensile strength (Rm), the ratio Rp0.2/Rm also increases and approaches unity. Due to the losts of ductility, uniform elongation decreases, and approaches zero for very high fluences. Neutron irradiation embrittlement is usually characterized by a shift to higher temperatures in the DBTT obtained from Charpy impact tests (Fig. 11.6). It is the dominant test in RPV surveillance programmes, but is also the most common test used in test reactor experiments, even though all RPV integrity assessments are based on fracture mechanics approaches with the main parameter being the fracture toughness. The reason for the common use

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Neutron embrittlement

200 180

KV, J

160 140 120 100 80 60 40

65 J 41 J

20 0 –150 –100

–50

0

50 100 150 Temperature, °C

200

250

300

350

11.6 Schematic diagramme showing the effect of increasing neutron fluence on the temperature dependence of Charpy impact energy for typical ferritic RPV steel.

of Charpy impact tests is connected with more than 100 years of experience in testing and characterization of RPV steels, as a huge database of these results is available. Thus, Charpy impact tests, due to their simplicity, serve as a main test in material qualification and monitoring, and also in components acceptance tests. Fracture toughness tests are a more modern approach, but their realization is more complicated, requiring more sophisticated testing equipment, and are much more expensive to be accepted by manufacturers as a part of acceptance tests. These days, fracture toughness tests are becoming more widely used not only in material qualification tests, but mainly in RPV surveillance specimen programmes due to current procedures for RPV integrity assessment. The Charpy test DBTT shift is not the only part of irradiation embrittlement detection: the so-called upper shelf energy is also seen to decrease in steels with a high content of impurities (Cu and P), its value can decrease in a substantial way, even below 70 J (50 ft-lb). The experimentally determined correlation between yield strength increase and DBTT shift is usually described as [3]:

DRp0.2 (MPa) ≈ 0.7 DDBTT (°C)

(11.1)

but the coefficient may vary between 0.4 and 0.9. The relationship between changes in yield strength increase and DBTT and fracture toughness transition temperature increases is shown in Fig. 11.7.

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Fracture toughness

Irradiation embrittlement

Irradiated

DJlc DT0

Irradiated

Temperature Produces

DYS

Unirradiated

Unirradiated

Duse

CVN Energy

Yield strength

Irradiation hardening

Unirradiated

Temperature

DT30

Irradiated

Temperature

11.7 Schematic relationship between changes in temperature dependences in yield stress, DBTT from Charpy impact and fracture toughness.

11.5

Main factors

Factors affecting the irradiation damage value or irradiation hardening and embrittlement, can be divided into several groups [4]: ∑ neutron field, like neutron flux, fluence and neutron energy spectrum, ∑ irradiation temperature, ∑ metallurgical variables, like content of impurities (Cu, P) as well as of alloying elements (e.g. Ni, Mn), ∑ annealing and re-embrittlement processes and mechanisms.

11.5.1 Neutron field The most important parameter is the neutron fluence; increase in the neutron fluence results in a substantial increase of radiation hardening and embrittlement; a fluence of 1022 m–2 (E > 1 MeV) is usually taken as a

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threshold for such effect. The exponent in the neutron fluence dependency changes with the neutron fluence The tendency for saturation can be observed for some steels when the neutron fluence is larger than 1024 m–2 (E > 1 MeV), but some steels (e.g., 15Kh2MFAA of Cr-Mo-V type for WWER-44% RPVs) does not show any tendency for saturation, even for fluences exceeding one order of magnitude higher. Knowledge about the effect of neutron flux is necessary for the application of irradiation test results obtained in experimental reactors to real RPVs as the lead factor (ratio between neutron flux in irradiated specimens and in the RPV) can lie in the range between 10 and 1000. Based on many experiments, it seems that some flux-dependent effects can be observed for high-Cu materials, whilst in other materials this effect has not been fully proved, even though correction factors exist in some predictive formulae. Standards usually require that the lead factor in surveillance specimens may not be larger than 3 (or now 5), which seems fully conservative. The neutron energy spectrum differs between individual locations even in one reactor and slightly even between the location of surveillance specimens and the RPV. Accordingly, the attenuation effect on the neutron flux through a RPV wall will substantially change this spectrum, too. In addition, differences in spectra also exist in experimental reactors – the effect of neutron spectrum was demonstrated when material was irradiated in heavy water reactors (HWRs) and then compared with LWR conditions.

11.5.2 Irradiation temperature Irradiation damage is a thermally activated process – its ab initio structure is affected by diffusion, recombination, annihilation and segregation of defects and the final situation is strongly dependent on the irradiation temperature. Generally, neutron irradiation damage in RPV materials decreases with increasing irradiation temperature, and a significant decrease in damage is seen in the temperature region of 150–400 °C, practically with the same temperature dependence and independent of irradiation temperature for all RPV materials even with different impurity contents. Levels of irradiation embrittlement can decrease by a factor of 10 in this temperature region. Odette and Lucas observed an average effect of about 1 °C/°C on DT41J for western steels [5].

11.5.3 Metallurgical variables The main effect on irradiation embrittlement (and also on hardening) is caused by impurities in the alloy, such as copper and phosphorus. Analysis of databases of experimental results showed that a copper content lower than approx. 0.08 mass % and phosphorus content of approx. 0.008 mass %

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have only small, or no, effect on irradiation embrittlement. Higher content of both impurities strongly affects the embrittlement, and the effect of copper is usually saturated at lower fluences than the effect of phosphorus. In reality, synergism of both impurities takes place, mostly in connection with other elements. In some experiments it was also shown that some other impurities like tin, arsenic and antimony can increase irradiation embrittlement in a similar way to phosphorus (mainly in WWER RPV types of steel), thus requirements for the purity of these steels also includes some limits. Generally, for the older types of RPV materials, PWR RPVs were manufactured with a higher, non-homogenously distributed content of copper in welds (as a result of the weld wire copper coating) up to 0.40 mass %, and WWER RPVs welds contained a higher content of phosphorus, up to 0.055 mass %. Current RPVs are manufactured more carefully with respect to impurity content, and their levels are substantially lower. The most recent RPVs have copper and phosphorus contents below their threshold for radiation embrittlement. The effect of alloying elements like nickel and manganese seems to be synergetic and it is not yet fully quantitatively determined if their contents are higher than approx. 1.2 mass %, even though their effect could be substantial and can lead to the ‘late blooming’ type effect.

11.5.4 Annealing and re-embrittlement Since irradiation damage is a thermally activated process, any increase in temperature over the irradiation one, especially for a longer time, can decrease the level of irradiation damage. This process is discussed in detail in Chapter 12.

11.6

Predictive formulae

The assessment of integrity and lifetime of RPVs requires a proper knowledge about the trends of irradiation embrittlement as a function of operation time, i.e. neutron fluence. Direct determination of this trend, based on testing surveillance specimens, is not always possible or it is not allowed by codes, thus predictive formulae were introduced practically in every code. Such formulae should have to be based on analysis of a wide database of results from surveillance specimen testing when requirements for lead factors are satisfied. Only in special cases can results from irradiations in experimental reactors be used in such analysis, but with restrictions. Two types of these formulae are included in the codes – either trend for mean values with given standard error values, or ‘upper boundaries’ that should cover all data (with at least 95% probability).

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These formulae are based mostly on national databases only, thus they can differ in different national codes, even for one type of steel. This can be explained either by a limited number of data points but also by the fact that metallurgical processes can differ at different manufacturers. The simplest formula has the following format:

DT = CF . Fn

11.2

where DT is the shift of Charpy DBTT and CF is a ‘chemical factor’ that depends on content of some chosen elements (mostly P and Cu) and n is an exponent that is mostly constant. This formula is still part of the Russian codes for WWER [6], FIS and FIM in France [7], KTA in Germany [8], JETE in Japan [9] or in former US NRC RG.199 Rev. 1 and 2 in the USA [10]. A summary of the principles of individual formulae is presented in Table 11.3 New predictions are based on physical models, like the new ASTM or regulatory guide (RG), where the formula has a format:

DT = SMD + CRP + bias

11.3

where SMD is stable matrix damage (including effect of P), CRP is copperrich precipitates (including effect of Ni and Cu), and bias is the effect of irradiation time.

11.7

Detection and measurement of irradiation hardening and embrittlement

Irradiation hardening and embrittlement cause changes in the mechanical properties of RPV materials. Irradiation hardening is represented by the increase in tensile yield stress DRp0.2 or DRe (yield strength Rp or physical yield stress Re). This property shows a larger change in comparison with the ultimate tensile strength and thus the yield stress is the most important for RPV integrity assessment. But this increase in yield stress is not taken into account in most RPV codes as the most conservative approach must be used. This hardening is connected with the loss of ductility in tensile tests, the most important being the reduction of area that is in the code (PNAEG in Russia) and is used for fatigue calculations. Irradiation embrittlement is mostly represented by the shift of temperature dependence of material toughness, using some index temperature. For Charpy impact tests, two different index temperatures are applied. PWR type codes define the so-called the reference temperature nil-ductility temperature (RTNDT ) [11] that in the initial unirradiated condition is based on the drop weight test and Charpy impact tests with a fixed energy value KV, and an additional requirement for lateral expansion of the Charpy specimens at the fracture surface, while the shift of this temperature is determined for a fixed

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Elements considered

Country

Fluence power exponent

Remarks

Cu, Cu, Cu, Cu, Cu, Cu, Cu,

USA USA Germany Russia France Japan Japan

0.5 0.28–0.10 log F Not given (graph) 0.33 0.35 0.29–0.04 log F 0.29–0.04 log F

No cross factors – Thresholds Cross factor Ni-Cu – No thresholds No cross factors – Thresholds No cross factors – No thresholds Cross factor Ni-Cu – Thresholds Cross factor Ni-Cu – No thresholds Cross factor Ni-Cu – No thresholds

P Ni P P P, Ni P, Ni P, Ni, Si

Reg.Guide 1.99 Rev.1 Reg. Guide 1.99 Rev.2 KTA PNAEG(x) FIS, FIM JEPE BASE JEPE WELD

(x) Formulae represent ‘upper boundary’ but are based only on data from irradiation in experimental reactors – new formulae are now being developed based on results from actual R

Irradiation hardening and materials embrittlement

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Table 11.3 Summary of national principles in construction of predictive formulae

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energy value KV equal to 41 J. The WWER code defines a so-called ‘critical temperature of brittleness’ Tk0 [6] that in the initial unirradiated condition is based on Charpy impact tests only, but with a notch toughness value KCV dependent on the yield stress of the material and an additional requirement on ductile fracture appearance. The shift produced by irradiation is also determined for the similarly chosen KCV value. Integrity of RPVs is based on a fracture mechanics approach that requires knowledge of the (mainly static) fracture toughness value of materials. This value can be obtained by quasi-static testing of fracture toughness specimens, either by three (four)-point bending or by excentric tension of compact type (CT) specimens, all having fatigue pre-cracking to create a sharp tipped crack, necessary for fracture mechanics testing. The index temperature for this property was defined at the 100 MPa.m0.5 fracture toughness level. In recent years, a new approach, i.e. ‘Master curve’, has been developed and widely applied – the same criterion for a reference temperature T0 is used whereby a temperature-toughness curve is obtained by re-calculating test results obtained with different size specimens to generate values that would be the same as those obtained with specimens having the standard thickness of 25.4 mm. Similarly, shifts in such reference temperatures can be obtained also for dynamic fracture toughness testing or for arrest fracture toughness. Charpy impact tests are technological dynamic tests where the impact notch toughness KCV describes the total energy necessary for the initiation of the crack from the notch, propagation and potential arrest of such a crack, while static fracture toughness KJC represents only the energy necessary for initiation of a fracture from an already existing crack. Nevertheless, experiments show that shifts of fracture toughness based reference temperature T0 are mostly larger than those of Charpy test DBTT (RTNDT or Tk0), and this relation can be expressed as [12]:

DT0 ≈ 1.1 DDBTT

11.4

The determination of irradiation hardening and embrittlement in RPV materials is realized on test specimens irradiated under defined and controlled irradiation conditions. Irradiation in experimental reactors are mostly carried out in specially heated irradiation rigs with high neutron flux lead factors, thus the use of such test results is limited and must be carefully analysed. The most representative source of material comes from the irradiation of RPV surveillance test specimens that are located in special water-tight containers directly in RPVs, mostly in their beltline region and close to the inner RPV wall to satisfy conditions for a low lead factor. To obtain representative test results applicable and useable for the integrity for a lifetime assessment of the RPV, the following conditions must be fulfilled (these requirements are similar in all codes and standards):

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∑

∑

371

specimens’ irradiation temperature must not be higher than +10 °C compared with the inner RPV wall temperature in the beltline region; this temperature is measured either by thermocouples in experimental reactors, or by temperature monitors of melting type in RPVs, neutron fluence in individual specimens of one group (used for determination of one curve, e.g. temperature dependence of Charpy impact toughness or static fracture toughness) must not differ by more than 10–15%; neutron fluences are measured by neutron monitors of activated fission type with a good knowledge of neutron energy spectrum.

11.8

Conclusions

Irradiation hardening and embrittlement are the most important damage mechanisms in reactor pressure vessels. Their levels depend on many factors, mainly: ∑ neutron field, ∑ irradiation temperature, ∑ RPV materials and their chemical composition – content of impurities and some alloying elements. Characteristic parameters of irradiation damage in RPV materials include the increase of yield stress and the shift to higher temperatures of a DBTT obtained either from Charpy notch impact tests or from static/dynamic/arrest fracture toughness tests. It must be mentioned that shifts of different DBTTs are not identical, and the reference temperature determined from static tests seems to be larger than those from dynamic ones. Integrity and lifetime assessment of RPVs is based on a good knowledge of irradiation embrittlement and predictive formulae are usually used. For a proper and reliable construction of such formulae, several formats are used, but they must be based on a wide and properly created and analysed database of results from surveillance specimen test results.

11.9

Sources of further information and advice

IAEA, Application of Surveillance Programme Results to Reactor Pressure Vessel Integrity Assessment, IAEA-TECDOC-1435, IAEA, Vienna (2005). IAEA, Effects of Nickel on Irradiation Embrittlement of Light Water Reactor Pressure Vessel Steels, IAEA-TECDOC-1441, IAEA, Vienna (2005). IAEA, Guidelines for Prediction of Radiation Embrittlement of Operating WWER-440 Reactor Pressure Vessels, IAEA-TECDOC-1442, IAEA, Vienna (2005). IAEA, Guidelines for Application of the Master Curve Approach to Reactor

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Pressure Vessel Integrity in Nuclear Power Plants, Technical Report Series, TRS-429, IAEA, Vienna (2005). IAEA, Integrity of Reactor Pressure Vessels in Nuclear Power Plants: Assessment of Irradiation Embrittlement Effects in Reactor Pressure Vessel Steels, Nuclear Energy Series No. NP-T-3.11, IAEA, Vienna (2009). IAEA, Master Curve Approach to Monitor Fracture Toughness of Reactor Pressure Vessels in Nuclear Power Plants, IAEA-TECDOC-1631, IAEA, Vienna (2009). IAEA, Pressurised Thermal Shock in Nuclear Power Plants: Good Practices for Assessment, IAEA-TECDOC-1627, IAEA, Vienna (2009). IAEA, Assessment and Management of Ageing of Major Nuclear Power Plant Components Important to Safety: BWR Pressure Vessels, IAEATECDOC-1470, IAEA, Vienna (2005). IAEA, Assessment and Management of Ageing of Major Nuclear Power Plant Components Important to Safety: BWR Pressure Vessel Internals, IAEA-TECDOC-1471, IAEA, Vienna (2005). IAEA, Assessment and Management of Ageing of Major Nuclear Power: PWR Pressure Vessels, IAEA-TECDOC-1556, IAEA, Vienna (2007). IAEA, Assessment and Management of Ageing of Major Nuclear Power Plant Components Important to Safety: PWR Vessel Internals, IAEATECDOC-1557, IAEA, Vienna (2007).

11.10 References [1] Seeger, A., in Proc. Second UN Int. Conference on Peaceful Uses of Atomic Energy (Geneva 1958). [2] Debarberis, L., Kryukov, A., Gillemot, F., Acosta, B., Sevini, F., ‘Semi-mechanistic model for radiation embrittlement and re-embrittlement data analysis’, Int. J. Pressure Vessel and Piping 82 (2005) 195–200. [3] Sokolov, M.A., Nanstad, R.K., ‘Comparison of irradiation-induced shifts of KJc and Charpy impact toughness for reactor pressure vessel steels’, Effects of Radiation on Materials (18th Int. Symp.), ASTM STPn1325 (1999), pp. 167–190. [4] IAEA, Integrity of Reactor Pressure Vessels in Nuclear Power Plants: Assessment of Irradiation Embrittlement Effects in Reactor Pressure Vessel Steels, Nuclear Energy Series No. NP-T-3.11, IAEA, Vienna (2009). [5] Odette, G.R., Lucas, G.E., ‘Irradiation embrittlement of reactor pressure vessel steels: Mechanisms, models, and data correlations’, Radiation Embrittlement of Nuclear Reactor Pressure Vessel Steels – An International Review, ASTM STP 909 (1986), pp. 206–241. [6] Code for Strength Calculations of Components of Reactors, Steam Generators and Piping of NPPs, Test and Research Reactors and Stations. Metallurgia, Moscow (1973). [7] Association française pour les règles de conception et de construction des materiels des chaudières électronucléaires, Règles de conception et de construction des materiels mécaniques des ilots nucleates PWR. RCC-M edition June 1993 + addenda June 1995, AFCEN, Paris (1995).

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[8] Safety Standards of the Nuclear Safety Standards Commission (KTA); KTA 3201.2 Components of the Reactor Coolant Pressure Boundary of Light Water Reactors; Part 2: Design and Analysis, Edition 06/1996. [9] Japanese Industrial Technical Standards: The Reactor Vessel Material Surveillance Test Methods, JEAC 4201-2000, Japan Electric Association, 2000. [10] Nuclear Regulatory Commission, Radiation Embrittlement of Reactor Vessel Materials, Office of Nuclear Regulatory Research Regulatory Guide 1.99, Revision 2, USNRC, Washington, DC (1988). [11] American Society of Mechanical Engineers, Asme Boiler and Pressure Vessel Code, Section III, ‘Nuclear Power Plant Components’, Appendix G, ‘Protection Against Non-ductile Failure’, ASME, New York (2004). [12] IAEA, Guidelines for Prediction of Irradiation Embrittlement of Operating Wwer440 Reactor Pressure Vessels, Iaea-Tecdoc-1442 IAEA, Vienna (2005).

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12

Reactor pressure vessel (RPV) annealing and mitigation in nuclear power plants

M. B r u m o v s k y, Nuclear Research Institute Rez plc, Czech Republic

Abstract: Radiation embrittlement is a key factor that determines the operational lifetime of any reactor pressure vessel (RPV). Several mitigation procedures can be applied for improving the mechanical state of the vessels: decreasing the neutron flux on the RPV wall (low-leakage core by inserting dummy elements), but the most effective is thermal annealing of the RPV, principally of the most critically affected area, i.e. weldments in the reactor active core beltline. Key words: reactor pressure vessel, thermal annealing, radiation embrittlement, radiation re-embrittlement, annealing device.

12.1

Introduction

Neutron irradiation embrittlement in ferritic reactor pressure vessels (RPV) is evident in two effects: firstly, it narrows the ‘pressure-temperature’ operation window for normal operating conditions, and secondly, it limits RPV lifetime as the transition temperature of RPV materials cannot be higher than that determined from pressurized thermal shock (PTS) calculations. Several mitigation measures can be applied to decrease radiation embrittlement of RPV beltline materials: ∑

use of a ‘low-leakage core’ that could decrease the neutron flux on the RPV wall by 30–40%, ∑ use of ‘dummy elements’ in reactor core periphery/corners that could decrease the original peak flux by a factor of 4.5 and the ‘new’ peak flux by a factor of about 2.5 (absolute values depend on the real core configuration), ∑ recovery annealing as the most effective measure, as it could practically restore initial mechanical properties of RPV core beltline (welds) and base materials. All three measures have been applied in different types of reactors: ∑

‘low-leakage core’ is used in practice in all reactors throughout the world as it is the cheapest measure, even though it has a limited efficiency,

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∑ ∑

375

dummy elements were inserted mostly only to WWER-440/V-230 type reactors where a substantial decrease of neutron flux was required; in most cases this insertion was connected with RPV annealing, recovery annealing was applied in many WWER-440/V-230 type reactors as the most effective mitigation measure.

12.2

Structures and materials affected

The most radiation embrittled zone of RPVs is the beltine region, i.e. the cylindrical part that is adjacent to the reactor active fuel core. Maximum fluence is usually found near the reactor core axial centre with a cosine distribution at both ends, sometimes deformed by the function of control rods and with only small fluxes at the core boundaries. The current design of RPVs requires exclusion of welding joints from the active core region, but in older vessels there are not only circumferential welds but also many axial welds in the beltline region since vessels were manufactured from plates. There is also a circumferential distribution of neutron fluxes in this direction depending on active core design (square for PWR and hexagonal for WWER ones). Even though this distribution can be substantial, it is taken into account only in probabilistic assessment of RPV failure probability, in deterministic evaluation no such distribution is applied as the most severe conditions are already put into calculations. Neutron flux distribution and resulting radiation embrittlement must be taken into account when mitigation measures are to be applied – either for low-leakage core or for potential thermal annealing.

12.3

Main mitigation measures

The radiation embrittlement can be mitigated by either flux reductions (operational methods aimed at managing the mechanism) or by thermal annealing of the RPV (maintenance method aimed at managing ageing effects). Flux reductions can be achieved by either fuel management (‘inside-out’ configuration) or direct shielding of the RPV from neutron exposure.

12.3.1 Fuel management The neutron flux (hence fluence with time) can be reduced by initiating a fuel management programme preferably early in the life of a given plant. Such fuel management is carried out by implementing a low neutron leakage core (LLC). A LLC is a core that utilizes either spent fuel elements or dummy (stainless steel) fuel elements on the periphery of the core which reflect neutrons back into the core or absorb them rather than allowing them to bombard the RPV wall. LLCs can result in a reduction in power and/or increase in

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cost to the NPP owner. Most of the western PWRs as well as the WWER plants have implemented LLC management programmes using spent fuel elements on the periphery of the core, but generally only after some period of operation. LLCs have been effective in reducing the re-embrittlement of the WWER-440 RPVs after thermal annealing. A more drastic reduction of neutron flux can be achieved by inserting shielding dummy elements into the periphery of an active core, for example into the corners of the WWER active core hexagons. Dummy elements were inserted into most of the WWER- 440/V-230 reactors in the mid-1980s. Dummy fuel elements were also used in some of the WWER-440/V-213 plants with RPVs with relatively high impurity (phosphorus) content (e.g. Loviisa, Rovno). Up to 32 dummy elements are usually inserted into the core periphery. They cause not only a significant flux reduction but also a shifting of the maximum neutron flux by an angle of about 15° relative to both sides of the hexagon corners. Thus 12 new peak values of neutron flux are created on the pressure vessel wall. The original peak flux is decreased by a factor of 4.5 and the ‘new’ peak flux is decreased by a factor of close to 2.5 (see Fig. 12.1). Thus, the cumulative effect of flux reduction must be calculated for both locations. Again, this method is most effective when applied during the first years of operation or just after thermal annealing. The use of dummy elements usually results in a significantly different neutron balance in the core. The radial gradient is increased and thus the power distribution is disturbed in such a way that the peak power may exceed certain limits. Thus, a reduction in the fuel cycle length or a reduction of the reactor output is often necessary.

12.3.2 RPV shielding Flux (hence fluence with time) can also be reduced by further shielding the RPV wall from neutron bombardment. The reactor internals, the core barrel and thermal shield provide design basis shielding of the RPV. However, if it is judged that the design basis neutron exposure will result in significant radiation damage such that limitations are placed on the heating up and cooling down of the plant and/or accident/unusual/upset/transient conditions, such as PTS, becomes a potential safety issue, additional shielding is required. Shielding of the RPV wall from neutron exposure can be accomplished by increasing the thickness of the thermal pads that exist on the thermal shield at locations where the fluence is high or by placing shielding directly on the RPV wall.

12.3.3 Thermal annealing Once a RPV is degraded by radiation embrittlement (e.g. significant increase in the Charpy ductile-brittle transition temperature (DBTT) or reduction © Woodhead Publishing Limited, 2010

Neutron flux without dummy elemens

36° 48°

0°

60 °

1. 0. 0 9 0. 0. 8 0. 7 0. 6 0. 5 0. 4 0. 3 2 0. 1

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v

Re

l.

Je

dn

.

Neutron flux with dummy elements

RPV

Dummy elements

12.1 WWER flux distributions in low leakage core with dummy elements [1].

Reactor pressure vessel (RPV) annealing and mitigation

24° 12°

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of fracture toughness), thermal annealing of the RPV is the only way to recover the RPV material toughness properties. Thermal annealing is a method by which the RPV (with all internals removed) is heated up to some temperature by use of an external heat source (electrical heaters, hot air), held for a given period and slowly cooled. The restoration of material toughness through post-irradiation thermal annealing treatment of RPVs has received considerable attention recently, due to the fact that a number of operating plants will be approaching the PTS screening criteria during their license renewal period. The first RPV annealings were carried out using primary coolant and nuclear heat (US Army SM-1A) [2] or primary pump heat (Belgian BR-3) [3]. The annealing temperature in the former case was 293–300 °C (72–79 °C above the service temperature). The degree of recovery in this case was about 70%. In the BR-3 reactor the service temperature was 260 °C and the vessel was annealed at 343 °C. The recovery was estimated to be at least 50%. The planned annealing of the Yankee Rowe vessel at 343 °C (83 °C above the service temperature) was estimated to give a 45–55% recovery. The ‘wet’ annealing technique is easy to implement because usually only the fuel is needed to be removed from the RPV, but unfortunately it can be utilized only in reactors (RPVs) which have operated at a low service temperature. RPVs are not designed to withstand the pressure of water at higher temperatures and the critical point of water is reached already at 374 °C (pcrit = 22 MPa). Due to very limited recovery, wet annealing with water is not a practical solution for power reactors and in any case it needs to be repeated frequently. Following the publication of the Westinghouse conceptual procedure for dry thermal annealing an embrittled RPV, the Russians (and recently, the Czechs) undertook the thermal annealing of several highly irradiated WWER-440 RPVs [1]. To date, at least 15 vessel thermal annealings have been realized (and others are under consideration). The WWER experience, along with the results of relevant laboratory scale research with western RPV material irradiated in materials test reactors and material removed from commercial RPV surveillance programmes, are consistent and indicate that an annealing temperature at least 150 °C more than the irradiation temperature is required for at least 100 to 168 hours to obtain a significant benefit. A good recovery of all of the mechanical properties was observed when the thermal annealing temperature was about 450 °C for about 168 hours (1 week). Moreover, The re-embrittlement rates upon subsequent re-irradiation were similar to the embrittlement rates observed prior to the thermal anneal. The dominant factors which influence the degree of recovery of the properties of the irradiated RPV steels are the annealing temperature relative to the irradiation (service) temperature, the time at the annealing temperature, the impurity and alloying element levels, and the type of product (plate, forging, weldment, etc.)

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Mitigation mechanisms including microstructure changes

The only mitigation measure that directly affects material properties and its microstructure is thermal annealing. The main purpose of such annealing is to restore initial mechanical properties, mainly the limiting material’s (most embrittled) DBTT transition temperature and its toughness, as much as possible. The efficiency of recovery annealing and, consequently, the lifetime of operating after annealing the RPV are defined by two factors: firstly, by the degree of transition temperature, Tk shift recovery or residual irradiation embrittlement value and, secondly, by the rate of irradiation embrittlement during re-irradiation. Hence it is vital to understand irradiation embrittlement before and after RPV annealing, and even after repeated annealing actions. The principal scheme of the whole process of RPV material embrittlement before and after annealing is shown in Fig. 12.2. In this diagram, C indicates ‘conservative shift’ where re-embrittlement rate is equal to the initial one; L indicates ‘lateral (or horizontal) shift’ where re-embrittlement rate was shifted horizontally to the right from zero fluence, and V indicates ‘vertical shift’ when re-embrittlement rate is equal to the initial rate for fluences larger than those before annealing. Most of the experimental data lie between the lateral and vertical shifts, thus the ‘lateral shift’ could be taken as the most conservative one. In this diagram, two important parameters are included: first, residual value of transition temperature after annealing, DTTannealing, and re-embrittlement rate after annealing during further operation. The residual value, DTTannealing, in practice represents the efficiency of

Re-irradiation

1

DTTF

C Annealing

Tk shift

Initial irradiation

L

V

DTTresidual Fequivalent

2 FI

Fluence

FR

12.2 Scheme of embrittlement of reactor pressure vessel under reirradiation of sequentially irradiated and annealed materials [1].

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the annealing process. This value depends on the annealing temperature and annealing time (holding at annealing temperature). Most of the experimental results showed that annealing temperatures between 430 and 475 °C are sufficient for substantial and efficient recovery of the initial mechanical properties. Annealing time for efficient recovery was determined as a minimum of 100 hours, which is usually the minimum applied time. The effect of both annealing parameters – time and temperature – is summarized in Fig. 12.3 obtained for steel of A 533-B type but a similar diagram was also constructed for WWER-440 type of RPV steels. As WWER/V-230 type weld metals are characterized by relatively high phosphorus (up to 0.055 wt%), necessary research was performed which produced the results shown in Fig. 12.4. It is seen that the residual value of transition temperature after annealing is increasing with higher phosphorus content: a conservative value of +30–40 °C was recommended for further RPV lifetime assessment. Regarding neutron fluence value at the annealing time, no substantial effects were found, and residual embrittlement, as well as further re-embrittlement rate, are practically independent of this neutron fluence. The lateral (horizontal) shift approach was also approved in the IAEA Round Robin Exercise on Radiation Embrittlement of WWER-440 Weld Metal. The microstructure investigation results demonstrate that impure copper clusters that are formed in the material under primary irradiation are not

Rel. recovery of I-induced hardness (%)

100 90 465 °C

80 70

500 °C

480 °C

440 °C

60 50 40 30 20 10 0

Fluence = 2.23 ¥ 1019/cm2, E > 1MeV 1

10

100 Time (min)

1000

10000

12.3 Summary of isothermal annealing on irradiated A 533-B type steel [4].

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150 460 °C 420 °C 340 °C

DTres °C

100

50

0 0.00

0.01

0.02

0.03 P, %

0.04

0.05

0.06

12.4 Dependence of residual embrittlement on phosphorus content in WWER-440 materials for different annealing temperatures [5].

recovered during annealing (i.e. not taken back into solid solution, as would be the case for the non-irradiated but heat-treated RPV prior to its going into service). Annealing of the materials at 455–470 °C causes changes in the morphology and distribution of copper-rich clusters, namely their growth and a decrease in their number density. The low density of (now larger) copper precipitates in thermally annealed RPVs has little influence on the mechanical properties. Thus, if annealing leads to a low density of nearly pure copper precipitates [5] and low matrix copper content, further neutron irradiation of this neutron irradiated and annealed material should not produce less transition temperature shifts as under primary irradiation. As follows from atom probe test results [5] the phosphorus content in the matrix after annealing is recovered approximately to the level of unirradiated material. It means the phosphorus influence for the material embrittlement under reirradiation is substantial.

12.5

Application of research and operational experience to the practical solution of problems

Design of the annealing device must be closely connected with the real RPV – its dimensions and location and size of the radiation embrittled zone. There is a principal difference between old RPVs of PWR and WWER design: while old PWR RPVs were usually manufactured from plates and thus also

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contain axial welds, all WWER RPVs were manufactured only with rings, i.e. only with circumferential welds. Thus, the critical zones of WWER RPVs are relatively narrow areas around one critical circumferential weld, while the critical zones of PWR RPVs are much wider as they contain not only circumferential but also axial welds and thus the heated region can reach also the RPV nozzle region. At the same time, the support area of the RPVs must also be taken into account. If the RPV is supported in the lower dome (as in the PWR), then RPV is extended by thermal expansion and some bending of primary piping may occur. The WWER RPVs are supported below the nozzle ring and thus such a problem is not important. Thus, annealing of the WWER RPVs can be limited only to a zone around a critical weld (in the lower part of the active core), while ‘old’ PWR RPVs, with axial welds, cannot be annealed by this method – heating of the whole vessel is preferred. Generally, two different designs for dry annealing have been applied up to now: ∑ ∑

electric heating furnace inserted into the RPV and annealing of only the critical circumferential weld and its surroundings – this design was used for annealing all WWER RPVs, indirect gas-fired ‘can’ process that was used for a demonstration project in Marble Hill RPV.

12.5.1 Electric furnace annealing Electric furnaces were used for annealing of WWER-440 RPVs by both Russian and Czech companies. In both cases, the annealing equipment is a ring-shaped furnace with heating elements on its external surface. Annealing equipment basic parameters are a maximum diameter of 4.27 m, a height of about 10 m and a total weight of more than 60 tonnes. Installed power output of heating elements is more than 500 kW, while approximately only 200–400 kW is sufficient for the annealing. Heating elements are connected to five adjustable heating sections. The equipment also consists of control boxes, a transformer, a power supply cable network, and a control system. Power supply is drawn from the main circulation pump feed system. The control system works in a semi-automatic mode where surface temperatures are determined in individual heating sections and these are automatically maintained by the control system. The same is applicable for heating and cooling rates. Control correction can also be made manually at any time. The annealing temperature was increased from 450 °C (which was applied for the first several units) to 475 °C+25°C –0°C while the initial holding time 168 hours (1 week) was decreased to 100 hours. The heating rate is 20 °C per hour, and the cooling rate is between 20 and 30 °C per hour. The main reason for such slow rates is to achieve minimum residual stresses after annealing. © Woodhead Publishing Limited, 2010

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The SKODA RPV annealing device is shown in Fig. 12.5 – two additional heating zones above and below the circumferential weld should decrease the thermal axial gradient, and thus residual stresses, after annealing.

v.

IV.

19° viii.

22°

i.

22°

iii.

vi.

25° ii.

vii.

12.5 Scheme of SKODA RPV annealing device [1].

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12.5.2 Indirect gas fired ‘can’ process A project was conducted at the cancelled Marble Hill nuclear power plant as a demonstration of the engineering feasibility of performing a thermal annealing treatment on a US-designed RPV. The Marble Hill plant was partially completed with the vessel in place and provided a unique opportunity to test the logistics of performing a dry anneal on a large commercial vessel. The Marble Hill demonstration was completed in 1997. The Marble Hill demonstration results are documented in the proceedings of an EPRI Reactor Pressure Vessel Thermal Annealing Demonstration Workshop held in Santa Fe, New Mexico, USA, in 1998 and in EPRI TR108316. The Marble Hill RPV was a Westinghouse design four-loop pressurized water reactor (PWR) with nozzle supports, similar to the US-Palisades RPV. The heating and cooling arrangement used an indirect gas-fired method through a heat exchanger (‘can’) as illustrated in Fig. 12.6. The heat exchanger was designed for potential re-use and easy clean-up after the annealing procedure. The results were successful in showing that annealing could be performed

Reactor containment

Exhaust

Burners Blowers

Gas control system

Propane storage Reactor vessel Heat exchanger

12.6 Marble Hill RPV heating system [1].

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at a nominal temperature of 454 ± 14 °C at the inside surface of the RPV for a time period of one week. Analytical models of the Marble Hill RPV and the reactor coolant system were shown to be correct based on measured temperatures and strains in the actual vessel during the annealing process. Documentation of critical vessel dimensions both before and after the annealing procedure confirmed that all vessel interfaces and dimensions were maintained within acceptable tolerances.

12.6

Conclusions

1. An effective RPV annealing can be realized by ‘dry’ annealing only, the most effective parameters lie between 450 and 475 °C for 100–168 hours. 2. Since the RPV is life-determining, annealing is a cost-effective way to ensure operation for times even in excess of the original design. 3. The operational flexibility (pressure-temperature window) of the RPV is widened after annealing recovery has been achieved. 4. RPV-PTS issues are less critically restrictive after annealing due to the recovery of the toughness levels (i.e. low ductile-to-brittle transition temperature). 5. It is essential to follow re-embrittlement rates after annealing and further irradiation using focused surveillance programmes.

12.7

Sources of further information

Four survey documents can be recommended for further study of the annealing effects on RPV material conditions, integrity and lifetime assessment: Brumovsky, M. et al., Annealing and re-embrittlement of reactor pressure vessel materials, State-of-the-art report, ATHENA WP-4, AMES Report No. 19. JRC 46534. EUR 23449 EN (2008). IAEA, Assessment and Management of Ageing of Major Nuclear Power Plant Components Important to Safety: PWR Pressure Vessels. 2007 Update. IAEA-TECDOC-1556 (2007). Pelli R., Torronen, K., State of the art review of thermal annealing. AMES Report No. 2. EUR 16278 EN (1995). Planman, T., Pelli, R., Torronen, K., Irradiation embrittlement mitigation. AMES Report No. 1. EUR 16072 EN (1994).

12.8

References

[1] Brumovsky, M. et al., Annealing and re-embrittlement of reactor pressure vessel materials, State-of-the-art report, ATHENA WP-4, AMES Report No. 19. JRC 46534. EUR 23449 EN (2008). © Woodhead Publishing Limited, 2010

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[2] Potapovs, U., Hawthorne, J.R., Serpan, C.Z. Jr., ‘Notch ductility properties of SM1A reactor pressure vessel following the in-place annealing operation’, Nucl. Appl. 5 (6) (1968), 389–409. [3] Motter, F., ‘Low-temperature annealing of the BR-3 reactor vessel’, NUREG/CP0058, Vol. 4 (1985), pp. 144–175. [4] Nanstad, R., Tipping, Ph., Waeber, W., Kallehof, R.D., ‘Effects of irradiation and post-annealing re-irradiation of reactor pressure vessel steel heat JRQ’, Proceeding of the IAEA Specialists’ Meeting, Gloucester, May 2001, TWG-LMNPP-01/2, pp. 42–58. [5] Pareige, P., Stoller, R.E., Russell, K.F., Miller, M.K., ‘Atom probe characterization of the microstructure of nuclear pressure vessel surveillance materials after neutron irradiation and after annealing treatment, J. Nucl. Mater. 249 (1997), 165–174.

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Part III Analysis of nuclear power plant materials, and application of advanced systems, structures and components (SSC)

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13

Characterization techniques for assessing irradiated and ageing materials in nuclear power plant systems, structures and components (SSC)

S. L o z a n o - P e r e z, University of Oxford, UK

Abstract: In recent years, several techniques have been able to provide information on the microstructure or chemical composition of materials with a lateral resolution better than 100 nm. Some methods can even consistently reach atomic resolution both in 2D and 3D. This chapter will review several of these high-spatial resolution techniques and how they have been successfully applied to the understanding of degradation of materials from nuclear reactors. Recent advances and future prospects will also be discussed. Key words: characterization techniques, microanalysis, (scanning) transmission electron microscopy (S)TEM, X-ray tomography, atom-probe, secondary ion mass spectroscopy (SIMS), focused ion beam (FIB), small angle neutron scattering (SANS), positron annihilation spectroscopy (PAS), scanning auger microscopy (SAM), irradiation damage.

13.1

Introduction

The control of microstructure, segregation and precipitation is often crucial in producing serviceable components in metal alloys and ceramics. It is therefore very important to be able to obtain an accurate description of the material in question through quantitative measurements and imaging techniques. Various techniques have traditionally been used for this purpose, providing data at different scales and accuracy. Only a few of them, however, can provide information at the nanoscale, revealing features smaller than 100 nm. These methods/techniques will be the main topic of this chapter. Experimental methods can be divided into direct and indirect ones, depending on whether the information they provide can be directly interpreted or requires previous knowledge and fitting to a model. Direct techniques can reveal information on the microstructure or on the composition, in two or three dimensions. In Fig. 13.1, a comparison of the direct techniques that are covered in this chapter is provided. Indirect techniques, on the other hand, can provide information on particular properties or features of the sample, relying on the goodness of fit to a model or certain assumptions. Two of 389 © Woodhead Publishing Limited, 2010

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Understanding and mitigating ageing in nuclear power plants SEM (S)TEM

X-ray tomography (3D) FIB slicing (3D)

100% (S)TEM: EDX and EELS mapping, EFTEM, electron tomography (3D)

10%

Detection range (at%)

1% 0.1%

Scanning auger + depth profiling (3D)

100 ppm 10 ppm Atom probe tomography (3D)

1 ppm

NanoSIMS + depth profiling (3D)

100 ppb 10 ppb

1 Å

10 Å

100 Å

1000 Å 1 mm 10 mm 100 mm 1 mm Scale covered

1 cm

13.1 Comparison of different imaging (top) and analytical techniques detection range vs. scale covered.

the most widely used for the characterization of nuclear materials, namely positron annihilation spectroscopy (PAS) and small angle neutron scattering (SANS) will be described. It should be noted that the classification between direct and indirect techniques is, not surprisingly, subject to discussion. For most techniques, it is not easy to decide where real data finishes and when the interpretation starts, so the previously stated distinction between indirect and direct techniques can be very subtle. As an example, a simple image acquisition of any kind using a charge-coupled device (CCD) requires several assumptions and corrections until the information provided can be used for any quantitative analysis (Moldovan, Li et al. 2008).

13.2

Non-destructive techniques

Non-destructive techniques can be divided into either volume or surface techniques, depending on whether they extract the information threedimensionally or in a ‘classical’ two-dimensional way. Secondary ion mass spectroscopy (SIMS) has been included in this section since, although removing several nanometres from the sample surface for analysis, the feature or region of interest can normally be analysed again by a different technique (e.g (scanning) transmission electron microscopy (S)TEM or atom-probe tomography).

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13.2.1 Volume techniques X-ray tomography X-ray tomography is a characterization technique which can offer different types of 3D information. It can analyse the microstructure, the defects and the crystallography of most types of materials with sub-micron resolution (Withers 2007). Two methods have been used recently for the characterization of stress corrosion cracking in steels from nuclear reactors. The first one is based on the variations in absorption coefficients along the path of the X-ray beam. Since the absorption coefficient is linked to the density and atomic number of the different materials which the beam encounters as it passes through the sample, the visualization of defects or second phases in the bulk is relatively easy. As in any tomography experiment, 3D information is gathered by acquiring a series of 2D images while rotating the sample (typically between 0 and 180º). An extra advantage is that experiments can be performed in situ, providing crucial information on, e.g., localized corrosion or transition from pitting to crack growth (Connolly, Horner et al. 2006). When the 3D characterization of the microstructure or crystallography of the sample is the target, X-ray diffraction tomography is the preferred choice (King, Johnson et al. 2008, Poulsen 2004). This technique facilitates the visualization of plastically non-deformed, polycrystalline materials, resolving the 3D grain shapes and crystallographic orientation. By simultaneously acquiring the absorption contrast, the information acquired can be combined with the microstructure. It also has been applied to in-situ experiments, enabling the behaviour of individual grains and grain boundaries to be characterized during straining experiments in corrosive environments. Examples of application of this technique can be found in the recent work from the Materials Performance Centre in Manchester (UK) which uses X-ray tomography to get a better insight into the dynamics and morphology of intergranular stress corrosion cracking in austenitic stainless steels in simulated light water environments. More importantly, by using in situ, three-dimensional X-ray tomographic images of intergranular stress corrosion crack nucleation and growth in sensitized austenitic stainless steel, evidence was provided for the development of crack bridging ligaments, caused by the resistance of non-sensitized special grain boundaries (Babout, Marrow et al. 2006) (see Fig. 13.2). Diffraction contrast tomography (DCT) also proved to be very useful to map the crystallographic orientation of the different grains in a sample tested in situ (see Fig. 13.3). A stress corrosion crack was grown through a volume of sensitized austenitic stainless steel mapped with DCT and observed in situ by synchrotron tomography. Boundaries which had shown an exceptional resistance to cracking were identified, revealing that they were not the twin variant type usually maximized during grain boundary engineering (King, Johnson et al. 2008). © Woodhead Publishing Limited, 2010

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Defect

PCCs

Bridge

~25 m m z

(a)

x

z

(b)

y

x

y

3

~25 m m Bridge z

(c)

x

y

13.2 Tomography data from in-situ stress corrosion cracking experiments: (a) reconstructed slice highlighting thin secondary cracks; (b) 3D isosurface; (c) combination of (a) and (b) showing that there is no phase contrast at crack bridging ligament. (From Babout, Marrow et al., 2006; courtesy of Maney Publishing. http://www. ingentaconnect.com/content/maney/mst/2006/00000022/00000009/ art00009).

SANS SANS is part of a family of techniques that rely on the diffraction of a transmitted beam, in a similar way to X-ray (SAXS) or light. SANS instruments can rely on a nuclear reactor for the production of monochromated neutron beams or a pulsed neutron source combined with a time of flight instrument. These techniques are useful because they can provide information on the size,

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Tensile axis

Crack path y x Bridge 1

z Bridge x (S1) z

(a)

(b)

13.3 Combined use of DCT and CT data to identify crack-bridging grain boundary structure: (a) cracks obtained from CT data are shown in black, at the final step before sample failure, and compared with DCT data of 3D grain shapes; (b) 2D section of the grain boundaries, identified by DCT, compared with the crack path identified by CT (from King, Johnson et al., 2008; reprinted with permission from AAAS).

shape or orientation of some secondary phases in a bulk. In a typical SANS experiment, a neutron beam a few mm in diameter is directed at a sample, which it will penetrate several mm. As an example, the SANS instrument at the ISIS spallation neutron source (UK), called LOQ, uses ‘cold’ neutrons with wavelengths between 0.2 nm (17 meV) and 1.0 nm (0.8 meV), allowing scales of between 0.4 and 80 nm to be probed (King 2000). The objective of any SANS experiment is the determination of the differential scattering cross-section, which contains the information on the shape, size and interactions between the scattering centres. For that purpose, the volume of material to be examined is bombarded with neutrons, which scattering angle and momenta after interacting with the sample can be measured by dedicated detectors. SANS data, however, has to be corrected, reduced and fitted to a model before an interpretation of the data can be made. Of particular relevance to nuclear materials is the characterization of precipitates in alloys. Most metals are hardened by the formation of nanometre-sized precipitates, which can then impede the movement of dislocations within the matrix. Precipitation can be induced either by thermal ageing or by irradiation, such as in the case of ferritic steels and welds from nuclear reactor pressure vessels (RPVs). Welds are more critical due to the presence of copper impurities in the welding rods, especially in first generation reactors. Copper can precipitate by thermal ageing or neutron irradiation, causing an undesired embrittlement of the RPV as the dislocations are blocked by the high density

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of Cu-rich precipitates (Barashev, Golubov et al. 2004). SANS has proven very successful to characterize the precipitate evolution with temperature or irradiation dose in base plates, welds and model alloys (Miller, Wirth et al. 2003, Carter, Soneda et al. 2001). Another important area of application of SANS has been the characterization of nano-structured ferritic alloys, which are promising candidates for advanced fission and fusion reactors. In-situ heating experiments can provide information on the dependence of precipitate size and number density with time or temperature (Miao, Odette et al. 2008) (see Fig. 13.4). Positron annihilation (PA) Positrons were discovered by Anderson in 1932, and identified as antiparticles to electrons. As they enter a metal and annihilate when interacting with electrons, they emit photons with energies, momenta and time of emission which can be accurately measured. Positron annihilation (PA) was established in the 1970s as a reliable technique to characterize vacancies in thermal equilibrium and other radiation-induced defects (Gauster 1976, Gil, De Lima et al. 1989). Sample preparation is relatively simple, and plates of ~10 ¥ 10 20 TEM measurement/3000 h

18

Predicted/3000 h SANS measurement

Diameter of particles (nm)

16

SANS/as-extruded

14

Extrapolated

12 10 243 h

8

480 h

6

480 h 4 2 0



As-extruded 900 1000 1100 Temperature (°C)

1200

1300

13.4 Size variation of nm-scale solute cluster-oxide features, (NFs) with temperature for the as-extruded and the aged MA957. The filled circles: NF sizes measured by TEM, open diamonds: NF sizes measured by SANS, and open circles: NF sizes predicted for 1150 °C/3000 h, 1200 °C/3000 h and 1250 °C/3000 h ageing (from Miao, Odette et al., 2008; reprinted with permission from Elsevier).

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¥ 0.5 mm are the only requirement. Since PA techniques are sensitive to the identification of vacancy-type defects in metals, they are very effective to characterize microstructure evolution during irradiation. When the coincidence Doppler broadening (CDB) technique is used, the momentum distributions of the core electrons specific to each element around the vacancies can be measured, so information on the elements around the annihilation sites can be extracted. This has been proven very useful to characterize fine precipitation of Cu in ferritic steels depending on the ageing time (Nagai, Hasegawa et al. 2000) or irradiation dose (Fujii, Fukuya et al. 2005). Figure 13.5 shows the CDB spectra for the same Fe-1 wt% Cu alloy as quenched and after 2 h ageing at 550 ºC, together with those for pure (bulk) Fe and Cu. It can be easily observed that for the as-quenched sample, the CDB spectrum is identical to the bulk Fe, so the positrons are not trapped by the isolated Cu atoms still in solution. The spectrum for the 2 h ageing is almost identical to that for the pure Cu, indicating that all positrons annihilate with the electrons of Cu. This information was used to estimate that the number density of Cu precipitates was in the order of 1018/cm3 assuming that 10% of Cu atoms had precipitated from the matrix solid solution (see Fig. 13.5). In Fig. 13.6, Coincidence-Doppler broadening spectra Pure Fe Pure Cu

106

Fe-1.0wt%Cu (as quenched) Fe-1.0wt%Cu (2h-aging at 550°C)

Counts

105

104

103

102

0

10

20 PL (10–3 m0c)

30

13.5 CBD spectra for Fe-1.0 wt% Cu as quenched and after 2 h ageing at 550 °C, compared with those for pure (bulk) Fe and Cu. Each spectrum is normalized to the same total count. Reprinted with permission from Nagai, Hasegawa et al. (2000). Copyright ©2000 by the American Physical Society.

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Ratio to pure Fe

unirrad. 0.1 mdpa 1 mdpa 10 mdpa 22 mdpa

Pure Cu

1.5

1.0

0

0

10

20 PL (10–3m0c)

30

40

13.6 CDB ratio curves for a A533B steel after different irradiation doses at 290 °C (from Fujii, Fukuya et al., 2005; reprinted with permission from Elsevier).

CDB ratio curves are shown for irradiated and unirradiated A533B steel samples, together with the curve for a well-annealed pure Cu sample as a reference. The broad peaks at around 24 ¥ 10–3 m0c are characteristic of Cu-3d electrons, indicating clustering of Cu atoms.

13.2.2 Surface techniques Scanning auger microscopy (SAM) Auger electron spectroscopy (AES) was developed in the 1960s as a surface technique which takes advantage of the characteristics of the low energy electrons (100 eV to a few keV) emitted during an auger process. It is used for elemental analysis of surfaces, achieving high sensitivity for most elements and good quantitative results. The technique has been used for obtaining quantitative depth profiles and 2D elemental maps as scanning auger microscopy (SAM). SAM works in a similar way to a scanning electron microscope (SEM), with the advantage that the spatial resolution is improved with respect to X-ray mapping due to the much smaller interaction volumes. A typical value for the spatial resolution would be ~20 nm, as opposed to one micron with energy dispersive X-ray (EDX) mapping. SAM has a history of successful applications in the area of nuclear materials characterization. It is a powerful technique to characterize grain boundary segregation (Allen, Tan et al. 2007, Nettleship, Wild 1990) or

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to chemically map fracture surfaces (Terachi, Fujii et al. 2005) . In Fig. 13.7, the grain boundary chromium concentration was measured with high spatial resolution for several austenitic alloys to reveal its dependence with irradiation temperature, indicating that the segregation reaches a maximum at around 400 ºC. The same experiment also revealed that grain boundary segregation also depended on irradiation dose, peaking at 0.2 dpa. In Fig. 13.8, a cross-sectional view is used to map the fracture surface of an austenitic stainless steel tested under pressurized water reactor (PWR) primary water conditions revealing the formation of a dual oxide layer (inner Cr-rich and outer Fe-rich). SIMS Secondary ion mass spectroscopy (SIMS) is a surface analysis technique in which a primary beam (Ga+, Cs+ or O–) is used to sputter material from the surface in a controlled way so that it can be analysed by a mass spectrometer. The emitted particles (secondary beam) are a combination of electrons, neutral species, atoms, molecules and clusters of ions. The bombardment of the surface by the primary beam will cause erosion and it is normally referred to as surface sputtering. The sputtering is controlled by the beam size, energy and current and it is normally restricted to a few monolayers. During the sputtering process, not all atoms will be ionized with the same probability.

Grain boundary Cr concentration (at. %)

16 15 Ni-18Cr-9Fe 14 13 12 11 10 9 150

Ni-18Cr 0.5 dpa AES measurements 200

250

300 350 400 Temperature (°C)

450

500

550

13.7 Grain boundary chromium concentration for Ni-18Cr and Ni18Cr-9Fe irradiated to 0.5 dpa from 200 to 500 °C (from Allen, Tan et al., 2007; reprinted with permission from Elsevier).

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1.0 mm

O

Cr

Ni

Fe

13.8 Secondary electron image (top left) and SANS elemental maps of a fracture surface from a cross-sectional view revealing the dual oxide layer formation in an austenitic stainless steel tested under PWR primary water conditions (from terachi, Fujii et al., 2005; reprinted with permission from the Atomic Energy Society of Japan).

The ionic efficiency is known as ion yield, and represents the fraction of atoms that can be ionized after sputtering. A reliable quantitative analysis will rely on a good previous knowledge of the different ionic yields for each element. Mass resolution is another important parameter, determining how easily two ions of similar masses can be separated. Although some elements and molecular fragments can have nominally the same mass (e.g. S and O2), differences in binding energies will result in different mass deficits which will be enough in most cases to separate adjacent peaks in the mass spectrum. If the beam is scanned over the surface and the spectra (or selected channels) acquired for each pixel, 2D maps can be obtained. Only with the development and use of the Cameca NanoSIMS (Conty 2001), can spatial resolutions of less than 100 nm be achieved routinely when mapping while keeping a high mass resolution. The technique has been recently applied (with excellent results) to the characterization of stress corrosion cracks in austenitic stainless steels (Alloy 304SS) from PWRs (Lozano-Perez, Schröder et al. 2008, Lozano-Perez, Kilburn et al. 2008). In Fig. 13.9, it can be observed how spatial resolution was sufficient to separate the two oxide layers (Cr-rich and Fe-rich) on the crack flanks of stainless steel tested under PWR primary water conditions. The high sensitivity of the technique allowed the detection of boron segregated to grain boundaries, as revealed in Fig. 13.10.

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Characterization techniques for irradiated and ageing materials 16

O–

10 mm

52

Cr16O–

10 mm

56

399

Fe16O–

10 mm

SE

10 mm

13.9 NanoSIMS maps and secondary electron image (SE) from a secondary crack in a 20% cold worked stainless steel tested under PWR primary water. Preferential oxidation along deformation bands is clearly visible (from Lozano-Perez, Kilburn et al., 2008; reprinted with permission from Elsevier).

13.3

Destructive techniques

This category of techniques, as suggested by its name, involves sample preparation which requires the destruction of the original sample or the extraction of the feature of interest from the bulk with the dimensions required by the technique to be used. The prepared specimen can be destroyed during the examination process, preventing any further examination, such as in focused ion beam (FIB) 3D slicing or atom probe tomography; or it can be re-used, as is the case with (S)TEM specimens which can be re-examined by SIMS or electron back scattered diffraction (EBSD), for instance. FIB 3D slicing The FIB has become a common tool for preparing TEM samples as well as to obtain quick cross-sectional views from selected regions of interest. In the last few years, it has proven once again a very versatile instrument when its capability to perform tomography by finely slicing the volume of interest

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Fe16O–

2 mm (a) 11 16

B O2

–

2 mm (b)

Normalized intensity

1.2 Cro S O BO2 FeO

1 0.8 0.6 0.4 0.2 0

3

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13.10 Dominant crack tip region from a 20%CW stainless steel sample: (a) NanoSIMS 56Fe16O– map showing the position of the line profile; (b) NanoSIMS 11B16O2– map from the same region; and (c) NanoSIMS line profiles (normalized) (from Lozano-Perez, Kilburn et al., 2008; reprinted with permission from Elsevier).

was demonstrated (Inkson, Steer et al. 2001, Kotula, Keenan et al. 2004, Claves, Bandar et al. 2004). Figure 13.11 shows a 3D reconstruction of a real stress corrosion cracking (SCC) crack tip from a Inconel 600 sample. The sample was tested in an autoclave under simulated PWR primary water conditions in order to induce SCC (Lozano-Perez, Yamada et al. 2008a). The 3D model was used to measure real crack openings, oxide widths and orientations between strain direction and grain boundary planes.

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1 mm

13.11 3D FIB slicing reconstruction from a stress corrosion crack in a Inconel 600 sample tested under PWR primary water conditions (Lozano-Perez, unpublished data).

(S)TEM (S)TEM is a well-established characterization technique which has the unique ability of providing both microstructural and chemical information with a resolution better than 0.1 nm in most cases. Microstructural information is obtained via diffraction experiments but also by understanding how the electron interacts with the sample. Most of the elastic and inelastic scattering theory was develop decades ago (Hirsch 1977), although some important developments have occurred over the years. Good examples are the development of the weak-beam diffraction technique (Cockayne 1973), which was proven very useful for the characterization of small defects, or the use of highly scattered electrons to obtain Z-contrast images of single atoms in the STEM mode (Crewe, Wall et al. 1970). Advances in hardware, computing power and electronics have allowed the acquisition of spectrum images, where an EELS and/or EDX spectrum is acquired serially for each pixel of an image in scanning mode. Although the effects of the neutron irradiation can manifest themselves macroscopically, they are always the result of events occurring at the atomic scale. If these effects are to be understood, a good knowledge of the microstructure evolution during irradiation is needed. Irradiation might induce both structural and compositional changes. The structure might be modified by the creation of point defects which usually cluster. These might include

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dislocation loops of vacancy or interstitial nature, stacking fault tetrahedral, gas bubbles and voids. In alloys, precipitation of second phases might occur, together with radiation induced atom segregation to grain boundaries or dislocations. The above-mentioned capability of imaging small defects has made TEM an ideal technique to characterize radiation damage by directly imaging the microstructure. It should be noted that indirect techniques, such as positron annihilation or SANS, rely on previous knowledge which was normally acquired via electron microscopy, since they lack the possibility of directly imaging the features of interest. A review of all the TEM techniques available for this purpose can be found in Jenkins (2001). Reliable characterization of the radiation damage requires a sample preparation methodology that ensures that no extra defects have been introduced during the process. Several approaches have successfully overcome this problem. The radiation damage group in Oxford Materials (UK) have developed a method which combines a selective mechanical grinding with electropolishing to produce good quality specimens that contained radiationinduced defects (Yao, Xu et al. 2008). This can be appreciated in Fig. 13.12, where different surface quality in the same irradiated specimen has proven crucial in imaging point defects. In-situ irradiation experiments, where a beam line is connected to a TEM and the evolution of the microstructure can be characterized in ‘real-time’ is

50 nm

13.12 Ion-irradiated pure Fe specimens polished under the same conditions: (a) shows a well-polished specimen with a high-quality surface, and (b) shows a poorly polished specimen with surface oxide (from Yao, Xu et al., 2008; reprinted with permission from Oxford University Press).

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one of the key techniques to understand the series of events under, for example, cascade irradiation conditions. These experiments can be performed under controlled temperature, so valuable information on the temperature or dose dependence can easily be extracted. Any theoretical modelling of irradiation damage should eventually be validated by experimental observations, and in-situ experiments provide an excellent source of input data. In Fig. 13.13 the different dislocation loop populations are characterized for particular temperature and irradiation conditions, in order to understand how and when they form and how they can contribute to the overall point-defect mobility (Meslin, Barbu et al. 2008). In the (S)TEM, EDX and/or electron energy loss spectroscopy (EELS) are typically used to obtain chemical information. A major advantage in using EDX over EELS is the ability to detect and quantify elements over most of the periodic table. When a material contains several alloying elements and has the potential for containing traces of many other impure elements, then EDX is especially effective. However, EELS can provide a different type of information. Not only can the chemical composition be measured from a spectrum, but also the fine structure at the ionization edges (ELNES) contain information about the electron density of states that can be directly related to bonding through comparisons with theoretical models of interfacial structure (Egerton 1986). Besides, the spatial resolution of EELS measurements is generally superior to the corresponding EDX experiment because the EDX data are affected by beam-broadening. Modern (S)TEMs can be used to 010

g = 011

001

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100 nm

13.13 Dislocation loops formed after irradiation at 400 °C up to 0.5 dpa. Two loop populations are clearly visible. The diffracting condition is g = <011> and z = [011] (from Meslin, Barbu et al., 2008; reprinted with permission from Elsevier).

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obtain quantitative information in the nanometre (or even sub-nm) range through EELS or EDX analysis. Detection limits for most elements can be in the order of the fraction of an atomic monolayer or 0.1 at%. (Watanabe, Williams 2005; Williams, Goldstein et al. 1995). Many examples of the application of (S)TEM microanalysis to the characterization of radiation effects in nuclear materials can be found in the literature (Carter, Soneda et al. 2001, Fukuya, Fujii et al. 2006). EDX spectrum imaging (EDX SI) was used to quantify the composition of precipitates in neutron-irradiated low-alloy steels (Burke, Watanabe et al. 2006) achieving great spatial resolution (see Fig. 13.14). Energy-filtered TEM was found equally useful to characterize small precipitates and obtain quantitative data (Lozano-Perez, Titchmarsh et al. 2006). In Fig. 13.15, the location of Cu precipitates smaller than 5 nm diameter in a thermally aged ferritic steel are revealed. The characterization of SCC-related phenomena has greatly benefited from the advances in sample preparation for (S)TEM. Locating a crack tip in an electron-transparent region proved a very challenging task which was finally overcome in the late 1990s. The use of a comprehensive methodology for mounting an ion beam and thinning samples containing crack tips allowed Thomas and Bruemmer to initiate a series of (S)TEM characterization

80 (a) Fe

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13.14 Compositional maps of the neutron-irradiated low-alloy steel from a EDX SI acquisition reconstructed using multivariate statistical analysis (from Burke, Watanabe et al., 2006; with kind permission from Springer Science+Business Media).

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2

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13.15 Fe-L23 elemental map showing a distribution of copper precipitates with diameter <5 nm (from Lozano-Perez, Titchmarsh et al. 2006; with kind permission from Springer Science+Business Media).

experiments that provided valuable data on crack tip chemistry and microstructure (Thomas, Charlot et al. 1996). The use of the FIB for preparation of similar samples meant that a specific crack could be selected from TEM observation and, more importantly, from regions that had previously been characterized by EBSD, SIMS, SAM, etc. (Lozano-Perez 2008a) (see Fig. 13.16). Specimens prepared by FIB with the method shown in the previous figure can be thinned further until the desired thickness is achieved. This allows the most demanding techniques, such as EELS spectrum imaging (EELS SI), to be used (Lozano-Perez, Yamada et al. 2008b) as shown in Fig. 13.17. Atom probe tomography Atom probe tomography (APT) has the unique ability to identify and quantify individual chemical species in three dimensions. The technique is a natural complement to other major microscopy techniques such as (S)TEM and SIMS, but APT provides the highest available spatial resolution for chemical analysis. The key to expanding the materials science problems accessible to APT has been, and will continue to be, the ability to craft specimens from a wide spectrum of materials in their as-available state. The technique

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Box 2 5 mm

Pt welding

Micromanipulator

13.16 FIB SE image showing how a volume containing the stress corrosion crack tip is cut from the bulk sample and welded to a micromanipulator so it can be lifted and later welded to a dedicated Cu grid (from Lozano-Perez 2008a; reprinted with permission from Elsevier).

O

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13.17 EELS SI elemental maps from a dominant crack tip extracted from a stainless steel sample tested under PWR primary water conditions (Lozano-Perez, unpublished data).

has difficulties associated with the sample preparation and with the correct interpretation of the evaporation phenomena. The use of FIB milling for the preparation of atom probe needles has enabled the preparation of specimens from materials that are difficult to electropolish and, more importantly, to select a specific orientation in a specimen or microstructural features such as coarse or low volume fraction precipitates or grain boundaries (Cerezo, Clifton et al. 2007b, Miller, Russell 2007a, Takahashi, Kawakami et al. 2007). Modern atom probes are capable of achieving a mass resolution of up to 1000 (full width at half maximum (FWHM)), a field of view up to 200 nm diameter and a typical volume analysis of 200 000 nm3 in 1 hour. APT has a successful history of contributions to the understanding of neutron irradiation embrittlement of RPVs (Miller, Russell 2007b). It was used to demonstrate the relevance of the post-weld stress relief treatment in reducing the matrix impurity copper content in high copper-contaminated ferritic alloys and welds, the formation of small precipitates and their subsequent

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coarsening during post-irradiation heat treatments. It also demonstrated the formation of nanoprecipitates rich in nickel, silicon and manganese in low copper and copper-free ferritic RPV alloys after neutron irradiation, and has been systematically used to quantify solute segregation and precipitation to dislocations and grain boundaries (see Fig. 13.18). It is also proving an excellent technique for characterizing the effects of radiation in oxidedispersion strengthened (ODS) Fe-Cr alloys, revealing that the precipitates chemistry is far more complex than anticipated (Marquis 2008) (see Fig. 13.19).

13.4

Recent advances, future trends and new techniques

Most limitations found in current techniques can be grouped into two categories. On the one hand is the maximum resolution or elemental detectability that can be achieved by current instruments. These numbers are always evolving and continuously being improved as new generations of instruments are developed. On the other hand, there are limitations imposed by the materials or features to be analysed, for which there might not be a suitable sample preparation technique currently available. X-ray tomography is expected to benefit from new generation synchrotron sources, such as diamond in the UK (McClarence 2008, Smith 2007). Such

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13.18 Atom maps from a KX-01 weld that was neutron irradiated to a fluence of 0.8 ¥ 1023n m–2 (E > 1 MeV) at a temperature of 288 °C. A high number density (~3 ¥ 1024 m–3) of Cu-, Mn-, Ni-, Si- and P-enriched precipitates with an average radius of 3 nm is evident (from Miller, Russell 2007b; reprinted with permission from Elsevier).

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(a)

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13.19 Slices through 3D reconstructions showing the evaporation structures of larger oxide NFs and profiles in (a) MA957, (b) ODS Fe12Cr alloy, and (c) ODS Eurofer 97 alloy. All samples were analysed in laser pulsing mode (reprinted with permission from Marquis 2008; copyright © 2008, American Institute of Physics).

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sources will be capable of producing brighter and more coherent X-ray beams that will facilitate experiments with faster acquisition times and closer to ‘real-time’ imaging. This will be of special interest when designing in-situ experiments, such as those involving oxidation, straining or heating. The development of new instrumentation for SAM has not been as rapid as expected, with many commercial instruments still relying on traditional cylindrical mirror or concentric hemispherical analysers. More importantly, they still make use of serial spectral acquisition, making spectrum imaging unrealistic. Fortunately, recent advances, such as the introduction of hyperbolic field analysers, have enabled spectrum images to be made in reasonable times (Jacka 2001). SIMS is another technique in constant evolution towards better spatial and mass resolution. Recent advances in ion sources and probe forming lenses encourage the hope that time-of-flight SIMS (ToF SIMS) will soon join the ‘select’ group of techniques that can provide information in the sub 100 nm spatial resolution region. This would mean that full mass spectrum maps could be acquired adding spectrum imaging capabilities to high-resolution SIMS. FIBs will remain one of the most versatile instruments available in modern laboratories. Its associated techniques are in constant evolution and only limited by the imagination of the operator. Recent advances include the successful 3D reconstruction of real crack tips and, combined with an EBSD detector, the generation of 3D EBSD maps. The characterization of real stress corrosion crack tips reached a key milestone when it was finally possible to prepare samples containing oxidized tips for atom probe tomography (Cerezo, Clifton et al. 2007a) (see Fig. 13.20). (S)TEM is a technique in continuous development. The recent addition of spherical aberration correctors (Cs-correctors) to the column means that CrO2

5 nm

13.20 Study of cracks in a type-304 stainless steel after stress corrosion cracking. APT maps of Cr (left) and CrO2 (right) species from a volume taken from the vicinity of a crack tip, showing O diffusion and Cr-rich oxide formation along a shear band (from Cerezo, Clifton et al., 2007a; reprinted with permission from Elsevier).

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the attainable resolution in both TEM and STEM modes can be considerably improved upon. More importantly, in the STEM mode, Cs-corrected probes can provide higher currents for microanalysis with improved beam sizes down to 0.1 nm. Other advances in instrumentation include more stable high tension power supplies, optimized environments (lower noise levels, more controlled temperature, no stray fields, etc.) and detectors with improved collection efficiency (Klie, Johnson et al. 2008, Shenkenberg 2007). As an example of the type of routine analysis that can be expected, an EELS line profile across an oxidized crack flank in a 304SS, which was tested in an autoclave simulating PWR primary water conditions, is shown in Fig. 13.21 (Lozano-Perez, Yamada et al. 2008a). A 0.1 nm probe was used and data sampled every 0.3 nm. The internal Cr-rich and the external Fe-rich oxides are clearly visible. Note that the Fe-rich oxide has been imaged with atomic resolution (see Fig. 13.21). New techniques to characterize small defects also look very promising. A good example is the use of diffuse scattering to characterize small loops and its ability to discern its interstitial or vacancy nature (Zhou, Dudarev et al. 2007, Kirk, Jenkins et al. 2006). In Fig. 13.22, experimental and simulated diffuse scattering patterns are compared for a Frank loop of known Burgers vector. Multivariate statistical analysis (MSA), which can be considered a group of processing techniques designed to analyse the information contained in large multidimensional datasets, has started to be used regularly to process experimental data. In the last two decades it has been successfully applied to the area of analytical electron microscopy, in particular to electron energy loss (Trebbia, Bonnet 1990, Titchmarsh 1999) and energy dispersive X-ray

Cr-rich oxide

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13.21 EELS line profile across crack flank in 304SS tested under PWR primary water conditions to induce SCC (Lozano-Perez, unpublished data).

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13.22 Images (a, b and c) of Frank loop based defects in g¼200 in bright-field kinematical diffraction condition and associated diffuse scattering (d, e and f, respectively) around the 400 Bragg peak with 800 excited. Simulated intensity contours (g, h and i) for symmetric Huang electron scattering for three distinct Frank loop orientations and Burger’s vectors in the same [011] orientation, and adjusted to agree with experimentally defined directions (from Kirk, Jenkins et al., 2006; reprinted by permission of the publisher (Taylor & Francis Group, http://www.informaworld.com)).

spectra (Burke, Watanabe et al. 2006, Kotula, Keenan et al. 2003) both in scanning and transmission electron microscopy (SEM and TEM). Recent advances in hardware and software are allowing the automatic acquisition of EELS or EDX datasets containing more than 100 million data points. Traditional methods for the extraction of chemical information rely on background subtraction and edge or peak signal integration. However, only a relatively small fraction of the available information is actually used. With the advent of modern affordable computers and their high computing power, MSA can finally be applied to any dataset. This way, the whole dataset can be analysed in a purely mathematical and unbiased way, extracting the main

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sources of information, which can be used to reconstruct the original data in a ‘noise-free’ way (Lozano-Perez 2008b). Tomographic data acquisition in the (S)TEM has also become a common technique for materials characterization in the last decade. Data is usually acquired by using a high-angle annular dark field (HAADF) detector in STEM mode (Z-contrast mode), or by energy-filtered TEM (Midgley, Ward et al. 2007, Midgley, Weyland et al. 2006). However, this technique has rarely been applied to the characterization of materials from nuclear reactors. Its recent application to the characterization of stress corrosion crack tips indicates just how relevant it might become in the future. With its ability to visualize the interaction of the crack with the microstructure in 3D and the possibility of measuring real angles, crack openings, etc. (Lozano-Perez, Yamada et al. 2008a), it opens the door to a new type of information not previously available. However, the next revolution in sample characterization is expected to come from a new multi-technique approach. Recently developed sample preparation methods and techniques have finally allowed the combined use of a range of microstructural techniques to enable a comprehensive characterization of crystallographic and compositional features over a range of scale lengths from millimetres to sub-nanometre. One example of this powerful holistic approach is the characterization of stress corrosion cracks. SCC has been traditionally investigated using indirect methods, either because the available techniques did not have high enough resolution or because the region of interest (crack tip) was not accessible for higher resolution techniques. It has been shown that the same crack tip can be characterized by techniques as diverse as optical microscopy, auger scanning microscopy, scanning electron microscopy (SEM), transmission electron microscopy (TEM), nanoscale secondary ion mass spectroscopy (NanoSIMS) and atom probe tomography. The complexity of the data acquisition was underpinned by the development of measurement-specific scripts that control data acquisition, thereby ensuring reproducibility of experimental conditions (Lozano-Perez 2008c).

13.5

References

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Burke, M.G., Watanabe, M., Williams, D.B. and Hyde, J.M., 2006. Quantitative characterization of nanoprecipitates in irradiated low-alloy steels: advances in the application of FEG-STEM quantitative microanalysis to real materials. Journal of Materials Science, 41(14), 4512–4522. Carter, R.G., Soneda, N., Dohi, K., Hyde, J.M., English, C.A. and Server, W.L., 2001. Microstructural characterization of irradiation-induced Cu-enriched clusters in reactor pressure vessel steels. Journal of Nuclear Materials, 298(3), 211–224. Cerezo, A., Clifton, P., Galtrey, M.J., Humphreys, C.J., Kelly, T.F., Larson, D.J., LozanoPerez, S., Marquis, E.A., Oliver, R.A., Sha, G., Thompson, K. and Zandbergen, M., 2007a. Review: atom probe tomography today. Materials Today, 10(12), 36–42. Cerezo, A., Clifton, P.H., Lozano-Perez, S., Panayi, P., Sha, G. and Smith, G.D.W., 2007b. Overview: recent progress in the 3-dimensional atom probe instruments and applications. Microscopy and Microanalysis, 13(6), 408–417. Claves, S.R., Bandar, A.R., Misiolek, W.Z. and Michael, J.R., 2004. Three-dimensional (3D) reconstruction of AlFeSi intermetallic particles in 6xxx aluminum alloys using the Focused Ion Beam (FIB). Microscopy and Microanalysis, 10(2), 1138–1139. Cockayne, D.J.H., 1973. The principles and practice of the weak beam method of electron microscopy. Journal of Microscopy, 98(2), 116–134. Connolly, B.J., Horner, D.A., Fox, S.J., Davenport, A.J., Padovani, C., Zhou, S., Turnbull, A., Preuss, M., Stevens, N.P., Marrow, T.J., Buffiere, J.Y., Bolller, E., Groso, A. and Stampanoni, M., 2006. X-ray microtomography studies of localised corrosion and transitions to stress corrosion cracking. Materials Science and Technology, 22(9), 1076–1085. Conty, C., 2001. Today’s and tomorrow’s instruments. Microscopy and Microanalysis, 7(2), 142–149. Crewe, A.V., Wall, J. and Lanomore, J., 1970. Visibility of single atoms. Science, 168(3937), 1338–1340. Egerton, R.F., 1986. Electron energy-loss spectroscopy in the electron microscope. New York: Plenum. Fujii, K., Fukuya, K., Nakata, N., Hono, K., Nagai, Y. and Hasegawa, M., 2005. Hardening and microstructural evolution in A533B steels under high-dose electron irradiation. Journal of Nuclear Materials, 340(2–3), 247–258. Fukuya, K., Fujii, K., Nishioka, Y. and Kitsunai, Y., 2006. Evolution of microstructure and microchemistry in cold-worked 316 stainless steels under PWR irradiation. Journal of Nuclear Science and Technology, 43(2), 159–173. Gauster, W.B., 1976. Positron annihilation as a non-destructive monitor of radiation damage in reactor pressure vessel steels. Journal of Nuclear Materials, 62(1), 118–120. Gil, C.L., De Lima, A.P., De Campos, N.A., Fernandes, J.V., Kögel, G., Sperr, P., Triftshäuser, W. and Pachur, D., 1989. Neutron-irradiated reactor pressure vessel steels investigated by positron annihilation. Journal of Nuclear Materials, 161(1), 1–12. Hirsch, P.B., 1977. Electron microscopy of thin crystals. Malabar, FA: Krieger. Inkson, B.J., Steer, T., Möbus, G. and Wagner, T., 2001. Subsurface nanoindentation deformation of Cu-Al multilayers mapped in 3D by focused ion beam microscopy. Journal of Microscopy, 201(2), 256–269. Jacka, M., 2001. Scanning Auger microscopy: Recent progress in data analysis and instrumentation. Journal of Electron Spectroscopy and Related Phenomena, 114–116, 277–282. Jenkins, M.L., 2001. Characterization of radiation damage by transmission electron microscopy. Bristol: Institute of Physics.

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King, A., Johnson, G., Engelberg, D., Ludwig, W. and Marrow, J., 2008. Observations of intergranular stress corrosion cracking in a grain-mapped polycrystal. Science, 321(5887), 382–385. King, S.M., 2000. Using SANS to study adsorbed layers in colloidal dispersions. In: B.J. Gabrys, ed., Applications of neutron scattering to soft condensed matter, 1st edn. Amsterdam: Gordon and Breach. Kirk, M.A., Jenkins, M.L., Zhou, Z., Twesten, R.D., Sutton, A.P., Dudarev, S.L. and Davidson, R.S., 2006. Diffuse elastic scattering of electrons by individual nanometersized dislocation loops. Philosophical Magazine, 86(29–31), 4797–4808. Klie, R.F., Johnson, C. and Zhu, Y., 2008. Atomic-resolution STEM in the aberrationcorrected JEOL JEM2200FS. Microscopy and Microanalysis, 14(1), 104–112. Kotula, P.G., Keenan, M.R. and Michael, J.R., 2003. Automated analysis of SEM X-ray spectral images: a powerful new microanalysis tool. Microscopy and Microanalysis, 9(1), 1–17. Kotula, P.G., Keenan, M.R. and Michael, J.R., 2004. Tomographic spectral imaging with a dual-beam FIB/SEM: 3D microanalysis. Microscopy and Microanalysis, 10(2), 1132–1133. Lozano-perez, S., 2008a. A guide on FIB preparation of samples containing stress corrosion crack tips for TEM and atom-probe analysis. Micron, 39(3), 320–328. Lozano-Perez, S., 2008b. Improving EFTEM analysis using multivariate statistical analysis. Journal of Physics: Conference Series, 126, 12040. Lozano-Perez, S., 2008c. Novel characterization of stress corrosion cracks. Journal of Physics: Conference Series, 126, 12078. Lozano-Perez, S., Kilburn, M.R., Yamada, T., Terachi, T., English, C.A. and Grovenor, C.R.M., 2008. High-resolution imaging of complex crack chemistry in reactor steels by NanoSIMS. Journal of Nuclear Materials, 374, 61–68. Lozano-Perez, S., Schröder, M., Yamada, T., Terachi, T., English, C.A. and Grovenor, C.R.M., 2008. Using NanoSIMS to map trace elements in stainless steels from nuclear reactors. Applied Surface Science, 255(4), 1541. Lozano-Perez, S., Titchmarsh, J.M. and Jenkins, M.L., 2006. Quantitative EFTEM measurement of the composition of embedded particles. Journal of Material Science, 41(14), 4394–4404. Lozano-Perez, S., Yamada, T. and Terachi, T., 2008a. 3-D Characterization of Crack Tips, in G. Ilevbare, M. Costello and R.W. Staehle, eds., Proceedings of the Detection, Avoidance, Mechanisms, Modeling, and Prediction of SCC Initiation in Water-Cooled Nuclear Plants workshop, 8–12 September, 2008, EPRI. Lozano-Perez, S., Yamada, T. and Terachi, T., 2008b. New ways of characterizing stress corrosion cracking, in B.L. Eyre and I. Kimura, eds., Proceedings of the International Symposium on Research for Aging Management of Light Water Reactors and its Future Trend, 22–23 October 2007 2008b, INSS 255. Marquis, E.A., 2008. Core/shell structures of oxygen-rich nanofeatures in oxide-dispersion strengthened Fe-Cr alloys. Applied Physics Letters, 93, 181904. Mcclarence, E., 2008. Bright lights – synthetic diamond plays its role in the new diamond synchrotron. Industrial Diamond Review, 68(1), 41–44. Meslin, E., Barbu, A., Boulanger, L., Radiguet, B., Pareige, P., Arakawa, K. and Fu, C.C., 2008. Cluster-dynamics modelling of defects in a-iron under cascade damage conditions. Journal of Nuclear Materials, 382(2–3), 190–196. Miao, P., Odette, G.R., Yamamoto, T., Alinger, M. and Klingensmith, D., 2008. Thermal stability of nano-structured ferritic alloy. Journal of Nuclear Materials, 377(1), 59–64. © Woodhead Publishing Limited, 2010

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Midgley, P.A., Ward, E.P.W., Hungría, A.B. and Thomas, J.M., 2007. Nanotomography in the chemical, biological and materials sciences. Chemical Society Reviews, 36(9), 1477–1494. Midgley, P.A., Weyland, M., Yates, T.J.V., Arslan, I., Dunin-Borkowski, R.E. and Thomas, J.M., 2006. Nanoscale scanning transmission electron tomography. Journal of Microscopy, 223(3), 185–190. Miller, M.K. and Russell, K.F., 2007a. Atom probe specimen preparation with a dual beam SEM/FIB miller. Ultramicroscopy, 107(9), 761–766. Miller, M.K. and Russell, K.F., 2007b. Embrittlement of RPV steels: an atom probe tomography perspective. Journal of Nuclear Materials, 371(1–3), 145–160. Miller, M.K., Wirth, B.D. and Odette, G.R., 2003. Precipitation in neutron-irradiated Fe-Cu and Fe-Cu-Mn model alloys: a comparison of APT and SANS data. Materials Science and Engineering A – Structural Materials Properties Microstructure and Processing, 353(1–2), 133–139. Moldovan, G., Li, X., Wilshaw, P. and Kirkland, A.I., 2008. Counting electrons in transmission electron microscopes. Microscopy and Microanalysis, 14(2), 912–913. Nagai, Y., Hasegawa, M., Tang, Z., Hempel, A., Yubuta, K., Shimamura, T., Kawazoe, Y., Kawai, A. and Kano, F., 2000. Positron confinement in ultrafine embedded particles: quantum-dot-like state in an Fe-Cu alloy. Physical Review B – Condensed Matter and Materials Physics, 61(10), 6574–6578. Nettleship, D.J. and Wild, R.K., 1990. Segregation to grain boundaries in nimonic PE16 superalloy. Surface and Interface Analysis, 16(1–12), pp. 552–558. Poulsen, H.F., 2004. Three-dimensional X-ray diffraction microscopy: mapping polycrystals and their dynamics. Berlin: Springer. Shenkenberg, D.L., 2007. Team develops electron microscope with 0.5-Å resolution. Photonics Spectra, 41(11), 108. Smith, D., 2007. Diamond synchrotron prepares to light up x-ray optics in the UK. Physics Education, 42(1), 106–109. Takahashi, J., Kawakami, K., Yamaguchi, Y. and Sugiyama, M., 2007. Development of atom probe specimen preparation techniques for specific regions in steel materials. Ultramicroscopy, 107(9), 744–749. Terachi, T., Fujii, K. and Arioka, K., 2005. Microstructural characterization of SCC crack tip and oxide film for SUS 316 stainless steel in simulated PWR primary water at 320 degrees C. Journal of Nuclear Science and Technology, 42(2), 225–232. Thomas, L.E., Charlot, L.A. and Bruemmer, S.M., 1996. High-resolution analytical electron microscopy of intergranular stress corrosion cracks. New Techniques for Characterizing Corrosion and Stress Corrosion. Proceedings. TMS – Miner. Metals & Mater. Soc, Warrendale, PA, USA. Titchmarsh, J.M., 1999. Detection of electron energy-loss edge shifts and fine structure variations at grain boundaries and interfaces. Ultramicroscopy, 78(1–4), 241–250. Trebbia, P. and Bonnet, N., 1990. EELS elemental mapping with unconventional methods. I. Theoretical basis: image analysis with multivariate statistics and entropy concepts. Ultramicroscopy, 34(3), 165–178. Watanabe, M. and Williams, D., 2005. X-ray analysis in the AEM with angstrom-level spatial resolution and single-atom detection. Microscopy and Microanalysis, 11(2), 1362–1363. Williams, D.B., Goldstein, J. and Newbury, D., 1995. X-Ray Spectrometry in Electron Beam Instruments. New York: Plenum Press. Withers, P.J., 2007. X-ray nanotomography. Materials Today, 10(12), 26–34.

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Yao, Z., Xu, S., Jenkins, M.L. and Kirk, M.A., 2008. Preparation of TEM samples of ferritic alloys. Journal of Electron Microscopy, 57(3), 91–94. Zhou, Z., Dudarev, S.L., Jenkins, M.L., Sutton, A.P. and Kirk, M.A., 2007. Diffraction imaging and diffuse scattering by small dislocation loops. Journal of Nuclear Materials, 367–370 A, 305–310.

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14

On-line and real-time corrosion monitoring techniques of metals and alloys in nuclear power plants and laboratories

L. Y a n g and K. T. C h i a n g, Southwest Research Institute, USA

Abstract: Corrosion monitoring plays an important role in corrosion control and mitigation. This chapter discusses the techniques that are capable of on-line and real-time measurements of corrosion under industrial plant or laboratory conditions. It presents the state-of-the art, science and technology for monitoring both general corrosion and localized corrosion. It provides discussions on the advantages and limitations of the different methods. This chapter also includes a section on the measurement of electrochemical potential (ECP), which is one of the most important parameters related to corrosion in high temperature and high pressure water systems, such as in nuclear power plants. Key words: corrosion monitoring, corrosion sensors, pitting corrosion, crevice corrosion, localized corrosion, general corrosion, linear polarization resistance (LPR), galvanic sensor, electrical resistance sensor, differential flow through cell, radio tracer, non-destructive evaluation (NDE), electrochemical noise, ultrasonic testing, multielectrode sensor, multielectrode array (CMAS), electrochemical potential (ECP), reference electrode.

14.1

Introduction

Corrosion monitoring is the practice of acquiring information on the progress of corrosion damage to a material on a frequent and regular basis. Corrosion monitoring plays an important role in corrosion control and mitigation. In nuclear power plants, various parameters are strictly controlled and chemical agents are added to the coolant to minimize corrosion of system components. For example, amines such as hydrazine (N2H4) are added to the secondary side of the coolant/heat transfer systems in many pressurized water reactor plants. The hydrazine is for scavenging oxygen with the ultimate goal of controlling the extent and rate of corrosion. At the present time, the additions of the chemicals are usually performed manually, based on the measured parameters such as the concentrations of the chemicals, oxygen, conductivity and pH of the water. While these parameters are important to corrosion, they do not supply information on corrosion rate. Certain values of these 417 © Woodhead Publishing Limited, 2010

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parameters may be important for corrosion control under ideal laboratory conditions, but may not be adequate or may be difficult to maintain under actual plant operating conditions. The corrosion rate of system components should be used as the basis to adjust the addition of the corrosion inhibiting chemical agents. Furthermore, system components are subject to corrosion during chemical cleaning and decontamination tasks since usually relatively aggressive chemicals are present in cleaning agents. Highly sensitive and robust realtime corrosion sensors should be used to monitor the cleaning agent and associated process damage to the system components. The corrosion test coupon method has been a simple and long-established method for evaluating corrosion or corrosion monitoring. The general corrosion rate is usually obtained from the weight loss or weight gain measured before and after the exposure to the environment of interest and the duration of the exposure.1 The coupon method is also widely used to evaluate localized corrosion such as pitting corrosion2 and crevice corrosion.3 If properly implemented, coupon methods are the most reliable method for corrosion monitoring. However, this method is slow; it usually requires an exposure time of three months to one year. The evaluation of coupons is also labour intensive and the coupons must be taken out of the monitoring environment before the evaluation can be performed. In contrast, corrosion sensors can provide day-to-day and even minute-to-minute real-time corrosion rate information for metal components in a system, without the need to retrieve the sensors within a closed system. This chapter focuses on the techniques that have been used, or have the potential to be used, as on-line and realtime tools for corrosion monitoring of metals and alloys in nuclear power plants and laboratories.

14.2

General corrosion monitoring

General corrosion is characterized by the corrosive attack that extends over the whole exposed surface or at least over a large area. The term ‘general corrosion’ is often used synonymously with uniform corrosion. However, purely uniform corrosion is a rare occurrence. The morphology of the corrosion surface produced by general corrosion always exhibits irregularities and roughness to some degree.4 Since general corrosion takes place over large areas of the metal surface, monitoring of general corrosion is relatively easy. The following section discusses the methods that may be used for monitoring general corrosion under high temperature and high pressure nuclear power plant conditions.

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14.2.1 Electrical resistance (ER) probes Principle The electrical resistance (ER) technique is probably the most widely used method for quantitative monitoring of metal cross-sectional loss caused by general corrosion. This technique measures the change in electrical resistance of a metallic element exposed to a corrosive environment and converts the electrical resistance, R, into metal loss due to corrosion according to:

R = r l/A

14.1

where r is the resistivity, l is the length, and A is the cross-sectional area of the sensing element, respectively. For a given shape of the sensing element, r and l are fixed. A, which can be converted to thickness change as the sensing element is corroded, can be obtained from the measurement of R. When R is plotted as a function of time, the corrosion rate (CR) in terms of millimeter per year (mm/yr) or mil per year (mpy) can be obtained from the slope of the R versus time plot (Fig. 14.1).5 Since the electrical resistance of a metal or alloy changes with temperature, the effect of temperature must be compensated for in ER probes. Figure 14.2 shows a typical circuit that measures the resistance of the sensing element

Connection probe units

CR = < 0.025 mm/yr (< 1 mpy)

CR = 0.88 mm/yr (34 mpy)

CR = 0.025 mm/yr (1 mpy)

Lost baffles

Days

14.1 Typical response of an electrical resistance probe in a flow control baffle system.5

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Sealed (reference) element

Resistance meter

Probe body

14.2 Typical measurement circuit for an ER probe. Note: The resistance from the sealed element provides a compensation reference for temperature effect. Cylindrical sensing element

Electrical connection Fitting for installation to process streams

14.3 Typical commercial probe for corrosion monitoring in pressurized systems. Courtesy of Metal Samples Company, Munford, AL, USA.

and the resistance of the reference element, which is a sealed element that senses the same temperature of the sensing element, but does not corrode. The simultaneous measurement of the resistance of the reference element allows cancellation of the temperature effect on the sensing element. Applications ER probes are available commercially in various designs. Figure 14.3 shows a typical commercial ER probe for application in a pressurized system. The sensing element in Fig. 14.3 has a cylindrical design. During service, the outside surface of the cylindrical element is exposed to the monitoring environment. The resistance of the sensing element is measured between the two ends of the cylindrical element. This cylindrical design allows convenient connection of the electrical wires to the bottom end of the sensing element. There are many other shapes of the sensing elements for the ER probes to suit different application needs. Figure 14.4 shows some of the typical shapes of the sensing elements in commercial probes. The flush-mounted type is suitable for erosion-corrosion monitoring and the wire loop design offers higher sensitivity. © Woodhead Publishing Limited, 2010



Spiral loop probe

Cylindrical probe

Tube loop probe

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Large flush probe

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Atmospheric probe

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14.4 Various forms of the sensing elements in commercial ER probes. Courtesy of Metal Samples Company, Munford, AL, USA.

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Wire loop probe

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The inside wall of a tubular sensing element can also be used as the sensing surface of the ER probe.6 In the design shown in Fig. 14.5(a), the inside diameter of the tubular sensing element is small (1 mm) in order to achieve a high fluid velocity in a regular laboratory-scale test loop to study flow accelerated corrosion. Figure 14.5(b) shows typical measurement results obtained with the tubular ER probe under a simulated piping condition in the primary coolant (heat transport system) of a Canadian Deuterium (CANDU) reactor at 310 °C, 10.1 MPa, and 2.1 m/s fluid velocity. The response of the wall thickness reduction to the addition of acid can clearly be seen. The pH of the fluid measured at room temperature was about 9 and 3 prior to and after the addition of sulphuric acid (H2SO4), respectively. Inductance probes A variation of the ER method is the inductance approach.7,8 With the inductance approach, the metal loss is derived by the measurement of the Fitting

Fitting

Coolant

Tubular sensing element (a)

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–5 175 195

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14.5 (a) An ER probe with tubular sensing element which allows high velocity fluid to flow inside for flow assisted corrosion studies and (b) the response of the resistance and wall thickness of the ER probe to the addition of acid in a LiOH solution at 310 °C.6

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inductive resistance (or the magnetic permeability of the sensing element). When the sensing element is corroded, the inductive resistance increases. It was reported that this approach is significantly more sensitive than the conventional electrical resistance approach and the response time can be reduced by a factor of 100 to 2500. For example, at corrosion rates of 10 to 100 mm/yr (0.4 to 4 mpy), the response times for typical ER probes are between 48 and 200 hours, depending on the life of the probe (total thickness of the sensing element).5 In principle, the response time of the inductive resistance method can be as short as 10 minutes and this method holds great promise for on-line and real-time corrosion monitoring. Some commercially available brands of probes, such as Microcor® by Rohrback Cosasco (USA) and CEION® by Cormon (UK), are supposed to be based on the inductive resistance concept; but no documentation is available in the public domain for verification. Since the inductive resistance method relies on the measurement of magnetic permeability of the sensing element, the application of this method is limited to highly magnetic materials such as carbon steel. This method cannot be applied to non-magnetic materials such as aluminum and copper, or to weakly magnetic materials such as some nickel-base alloys.5 Advantages and limitations The advantages of ER probes include: (1) the ability to measure corrosion in almost any environment – aqueous, non-aqueous; (2) the measured signal is directly related to the metal loss – no effect of side reactions. The limitations of the ER probes include: (1) the effect of temperature on measured signal; (2) the change of electrical resistance due to corrosion is extremely small. Therefore, the response of the ER probes is slow (weeks to hours). For better sensitivities, the probe sensing element must be extremely thin and the life of the probe is limited. In addition, ER probes are not sensitive to localized corrosion, such as pitting.

14.2.2 Linear polarization resistance (LPR) probes Principle The LPR method is probably the most commonly used fast response method for quantitative monitoring of corrosion in aqueous systems. Figure 14.6 shows a typical potential-current plot for metal electrodes in activationcontrolled systems.9 The potential-current relationship is essentially linear near the corrosion potential and the corrosion current, Icorr, may be calculated using the following equation:

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(–)

Slope = Rp

(–)

Current density (+)

14.6 Hypothetical linear polarization resistance plot.9



ba · bc ˆ Ê I corr = 1 Á RP Ë 2.303 (b a + b c )˜¯

14.2

where Rp is the slope as shown in Fig. 14.6 (called polarization resistance), and ba and bc are anodic and cathodic Tafel constants, respectively. Equation 14.2 is commonly known as the Stern–Geary equation, named after the first researchers to introduce the explicit form of the equation 50 years ago.10 The concept of Equation 14.2 may be traced back to 1938 when Wagner and Traud reported the current-potential relationship that can be used to determine the corrosion rate of metals in activation-controlled systems.11 Under certain conditions, the ba and bc in Equation 14.2 may be treated as constants and combined:

B=

ba · bc 2.303 (b a + b c )

14.3

where B is the Stern–Geary coefficient (or B-value). Therefore, the corrosion current is related to the polarization resistance by the following simple equation:

I corr = B Rp

14.4



Therefore, the corrosion current can be determined by simply measuring the slope near the corrosion potential. As shown in Figure 14.6, the potential-

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current relationship is linear only near the corrosion potential. The measurement should be conducted at ±30 mV, or typically ±10 mV within the corrosion potential.12,9 In practice, the B-value in a given system is often treated as a constant. Some of the B-values in selected systems can be found from the literature.9 According to Faraday’s Law, CR can be calculated as: ÊI ˆ CR = K1 Á corr ˜ EW Ë dA ¯

14.5 where CR is given in mm/yr, Icorr in mA, K1 is 3.27 ¥ 10–3 [mm g/mA cm yr], A is the surface area in cm2, d is the density in g/cm3, and EW is the equivalent weight which is defined as the mass in grams that will be oxidized by the passage of one Faraday [96 489 C mol] of electric charge. The EW for pure elements is given as: 14.6 EW = W n where W is the atomic mass of the element and n is the number of electrons involved in the oxidation of an atom of the element in the corrosion process, i.e., the valence of the element. To calculate the EW of an alloy, the following formula may be used: 1 14.7 S nWi fi i where ni is the valence of the ith element of the alloy; wi is the atomic mass of the ith element of the alloy, and fi is the mass fraction of the ith element of the alloy. It should be noted that valence assignments for elements that exhibit multiple valences under the testing conditions involve uncertainties. It is best if an independent technique is used to establish the proper valence for each alloying element. The EW values for selected metals may be found in ASTM G102.13 EW =

Applications The guidance for using the LPR methods for corrosion testing and corrosion monitoring may be found in ASTM G0314 and ASTM G59.12 Since the potential value as shown in Fig. 14.6 does not need to be absolute, it can be measured against a pseudo reference electrode. As a matter of fact, nearly all commercial LPR probes for field applications use pseudo reference electrodes made of the same metal as the sensing electrode so that the LPR probe can be installed in harsh environments such as pressurized systems and with no need for special maintenance or care of the reference electrode. Therefore,

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many of the industrial LPR probes have three electrodes (sensing electrode, counter electrode, and reference electrode) made of the same metal. Figure 14.7 shows the schematic diagram for a three-electrode commercial LPR probe and its installations in the piping system in the field.15 The electrodes of commercial LPR probes are usually made of small cylinders (3–5 mm in diameter and 20–40 mm in length) with a thread connection at one end so that these electrodes can be easily removed or attached to the LPR probe body for easy replacement (see Fig. 14.7). In cases where some errors in corrosion rate are tolerable or inevitable, the counter electrode and the reference electrode may be combined. The LPR probes for these applications have only two electrodes. Figure 14.8 is a typical commercial LPR probe with two electrodes. LPR probes are widely used in industry, especially for corrosion monitoring in cooling and waste water systems. Figure 14.9 shows the typical results from an LPR probe installed in a once-through cooling water system.16 The corrosion rate was high initially, probably due to the freshly polished surface of the probe sensing electrode. When inhibitor was added, the corrosion rate decreased over a period of several days. It should be mentioned that a localized corrosion monitor (differential flow cell technique, see Section 14.3.3) was also used during the measurement. The results showed that the Removable electrodes

(a)

(b)

(c)

(d)

14.7 Schematic diagram of a typical three-electrode polarization resistance probe (a) for determining corrosion rate in the field and installation of such a probe in pipe fitting (b), in weld line (c), and in pipe tee (d).9

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14.8 Typical commercial linear polarization probe with two electrodes for corrosion monitoring in pressurized systems. Courtesy of Metal Samples Company, Munford, AL, USA. 700

Corrosion rate (mm/year)

600 500 Feed 25 ppm inhibitor

400 300 200 100 0 0

5

10

15 Time (day)

20

25

30

14.9 Typical results from an LPR probe installed in a once-through cooling water system in a nuclear power station. Modified from B. Yang.16

response from the localized corrosion monitor for localized corrosion was much clearer and faster than the response from the LPR probe for general corrosion. Advantages and limitations The LPR method is based on the Stern–Geary equation, which is derived on the assumption that the corrosion process is activation controlled. It applies

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to cases where corrosion is controlled by activation processes. In most general corrosion cases, the corrosion process is usually under activation control. Therefore, the LPR method is an excellent method for following general corrosion. Oldham and Mansfeld showed that the LPR method also applies in cases where one of the half-reactions is totally under diffusion control (metal dissolution in passive state or reduction reaction under oxidant diffusion control) where ba or bc in Equation 14.2 is infinity.17 Compared with the ER probe, it has a much shorter response time (six minutes for a potential scan cycle as recommended by the ASTM standard14) and offers near instantaneous rate measurement. It is a well developed technology and has many decades of application history for real-time corrosion monitoring. The LPR method is currently the most widely used instantaneous method to monitor general corrosion in industrial applications. Owing to the assumption of activation control in the Stern–Geary equation, strictly speaking, the LPR method is not applicable to corrosion processes that are controlled by both diffusion and activation processes. The use of the LPR method in these systems should be carefully verified with other methods. There are also other restrictions for the LPR methods.17–19 These include: (1) the corrosion potential does not lie close to the reversible potentials of the metal/metal ion or oxidizing agent/reduction product couples, and (2) no thick film of corrosion products covers the sensing electrode. In addition, the B-value for a given metal varies with the corrosion environments, and may change with time in the same environment. The B-values in many commercial LPR probe instruments are assumed to be constant. Therefore, the accuracy of the measured corrosion rates by these instruments may be limited in many systems, especially if the corrosion process is not purely controlled by activation.

14.2.3 Other methods for monitoring general corrosion Electrochemical noise (EN) sensors Electrochemical noise method measures the fluctuations in potential and current that occur on a corroding metal electrode.20 Figure 14.10 shows a schematic diagram for electrochemical noise measurement using three electrodes. The fluctuation of currents is measured between two identical sensing electrodes (Sensing electrode #1 and Sensing electrode #2). The voltage fluctuation is measured between the two sensing electrodes and a reference electrode. The electrochemical noise method has been used to measure the general corrosion rate based on noise resistance21,22 which is defined as23:

Rn = sV/sI

14.8

where sV and sI are the standard deviations of voltage and current values, respectively, measured during a given time period as defined by: © Woodhead Publishing Limited, 2010

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Sensing electrode #1

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V

Sensing electrode #2

Reference electrode

14.10 Electrochemical noise measurement using three electrodes.



sV2 = ∑ (Vj – Vm)2/(n – 1)



sI2 = ∑ (Ij – Im)2/(n – 1)

14.9 14.10

In Equations 14.9 and 14.10, Vj is the voltage value measured at the jth time interval, Vm is the mean voltage in the given time period, Ij is the current value measured at the jth time interval, Im is the mean current in the given time period, and n is the number of time intervals. In deriving the corrosion rate, the noise resistance, Rn, is treated as the polarization resistance Rp (see Section 14.2.2) and Equations 14.2 or 14.4 are used to calculate the corrosion current.20 Electrochemical noise method is also used for detection of localized corrosion (see Section 14.3.1 for more information). Galvanic sensors In a galvanic sensor, a corroding metal of interest is electrically coupled to another metal which is more corrosion-resistant (copper, stainless steel, or gold foil if the metal of interest is carbon steel) to raise the electrochemical potential of the metal of interest.24 At such raised potential, the corrosion of the metal is accelerated in a corrosive environment. Therefore, it is a reliable way to detect the corrosivity of the metal in the environment. Since the method only requires two metal electrodes, coupled by an ammeter (usually zero resistance ammeter), a galvanic probe is simple and low cost. Galvanic probes have been widely used in the industry, especially for atmospheric corrosion monitoring.25 In these environments, the galvanic sensors effectively measure the formation of water on the metal, especially if the water contains corrosive salts (such as NaCl) or acids. It should be noted that galvanic

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sensors cannot be used to give the quantitative corrosion rate because it operates under accelerated conditions (raised potential). Ultrasonic testing (UT) The ultrasonic technique has been widely used to detect flaws or wall thinning as a non-destructive inspection tool. Increasingly more work has been reported on the use of the ultrasonic technique to monitor wall thinning caused by corrosion as a monitoring tool.26 Ultrasonic thickness measurement is based on the time required for the ultrasonic pulse to travel from the front surface of the component to the back and return, based on the known ultrasonic velocity in the component. This method has been used in laboratories27 and commercial nuclear power plants28 to measure the pipe wall thinning caused by flow assisted corrosion at high temperatures. Figure 14.11 shows an online ultrasonic thickness measurement system and a high temperature transducer installed on a nuclear reactor pipe. This system has been installed in the high temperature (310 °C) piping system of several commercial reactors in Canada. According to the manufacturer’s product brochure, the transducer has a resolution of better than 0.1 mm and is capable of continuously operating at temperatures up to 500 °C. Radioactive tracer method The radioactive tracer methods have been used to measure metal loss caused by a wide variety of mechanisms including corrosion, erosion and mechanical wear.29 With this method, specimens of interest are made radioactive by exposing them to a thermal neutron flux (bulk or neutron activation) near the core of a nuclear reactor. The high energy neutrons penetrate the nucleus of a small number of atoms within the test specimen. This results in an increase in the mass of the nucleus, transmuting the atom into a heavier isotope, which is often radioactive. The now radioactive specimen of interest is then installed in a test loop with a gamma spectroscopy system (Fig. 14.12). As corrosion takes place on the specimen, corrosion products from the specimen are transported into the test solution and cause the build-up of small amounts of radioactivity. The build-up of the radioactivity is measured by the gamma ray detector and used to derive the corrosion rate. Although this method may not be used in the actual plant because the background level of radiation is probably higher than the source radiation in the coupon, it is an excellent laboratory tool for corrosion studies under high temperature and high pressure conditions. This is because the specimens can be machined into almost any shape (to study the geometry effect) and easily placed anywhere in the pressurized test loop without the need to penetrate the pressure boundary. Compared with the non-intrusive ultrasonic method, the

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(a)

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(b)

14.11 An on-line ultrasonic thickness measurement system (a) and a high temperature transducer installed on a nuclear reactor pipe (b). Courtesy of Research and Product Council, New Brunswick, Canada.

radio tracer method is more sensitive and can be used to measure corrosion rates as low as nanometres per year.29

14.3

Localized corrosion monitoring

Localized corrosion is characterized by the corrosive attack that is localized over an isolated surface. Examples of localized corrosion include pitting corrosion and crevice corrosion. Localized corrosion for metal components is usually the major concern in chemical plants or nuclear power plants because localized corrosion is difficult to detect at an early enough stage to mitigate it. Evidence of pitting or crevice corrosion, for example, may be hidden under gaskets, or between steam generator tubes and the tube sheet. In addition, once initiated, localized corrosion can propagate rapidly and result in either component failures (such as through wall penetrations) or trigger other modes of failures, such as stress-corrosion cracking, even though the

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14.12 A high temperature flow loop for corrosion measurement using the radioactive tracer method. Note: The cylinder-shaped equipment on the right is a gamma detector. Courtesy of Southwest Research Institute®, San Antonio, TX, USA.

majority of the surface area is not affected. Therefore, localized corrosion monitoring is vitally important in corrosion control in many industries where small and through-wall penetration would cause equipment failures. Unlike general corrosion monitoring, which started more than half a century ago as discussed in Section 14.2, the attempts for localized corrosion measurements started less than 30 years ago when the electrochemical noise method first became available.20 The following sections discuss three methods that have been used to detect or quantitatively measure localized corrosion.

14.3.1 Electrochemical noise (EN) As discussed in Section 14.2, the electrochemical noise method has been used to measure the general corrosion rate. Since the LPR method was available much earlier than the electrochemical noise method, and it has been widely accepted as an industrial method for real-time general corrosion monitoring, the unique value of the electrochemical noise method for corrosion monitoring is really its ability to detect localized corrosion, as reported by many investigators.30,31 The detection of localized corrosion in the EN method is usually based on an empirical parameter, pit index (PI) also called localization index (LI),31 which is calculated according to:

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PI = sI/Irms

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14.11

where sI is the standard deviation of current values measured during a given time period as defined by Equation 14.10, and Irms is the root mean square as defined by:

Irms2 = S Ij2/(n – 1)

14.12

where Ij and n are as defined in Equation 14.10. By definition, 0 ≤ PI ≤ 1. According to the classification proposed by Eden31 if the PI approaches 1, the corrosion process is unstable and, therefore, more likely to be stochastic; conversely, if the PI approaches the order of 0.001, the corrosion is mainly uniform. Good correlations between the pitting index and localized corrosivity have been reported in the literature.30,32 However, it was also shown that the pitting index is not a reliable pitting indicator by other investigators.23, 33,34 The use of pitting index as pitting indicator for a given system should be carefully evaluated in laboratories prior to employing the probes in the field. Electrochemical noise method was selected as the tool for monitoring pitting and stress corrosion cracking of the steel tanks containing nuclear waste liquid in Hanford, Washington (USA). The first EN probe was installed in one of the tanks in 1996 and five more were deployed between 1996 and 2005.35 Large amounts of data were collected during the monitoring period and the pitting risk was considered low based on the current and potential measurements, except for a few weeks when one of the probes showed a large degree of fluctuation in current (up to 500 nA from a mean of 130 nA) and in potential (up to 50 mV from a mean value of 40 mV). When the current fluctuation was high, the risk of pitting corrosion was considered to be high. Localized corrosion monitoring using the electrochemical noise method is limited to qualitative indication for the probability or risk assessment for localized corrosion. The electrochemical noise method does not provide the quantitative rate for localized corrosion.

14.3.2 Coupled multielectrode array sensor (CMAS) Coupled multielectrode array sensors (CMAS) are a recently emerged technology for corrosion monitoring, especially for localized corrosion monitoring.36 The coupled multielectrode array was initially introduced by Fei et al., in 1996, for studying the spatial and temporal electrochemical behaviours of an iron metal in solutions.37 Shortly after, Tan reported the use of the coupled multielectrode array for corrosion measurements in 1997 and 1998.38,39

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Principle When a metal undergoes non-uniform corrosion, particularly localized corrosion such as pitting corrosion or crevice corrosion in an electrolyte, electrons are released from the anodic sites where the metal corrodes and travel to the cathodic sites where the metal corrodes less or does not corrode to support the cathodic reaction and maintain the balance of charge (Fig. 14.13).40,41 If the metal is separated into small areas, some of the local electrodes have properties that are close to the anodic sites and others have properties that are close to the cathodic sites of the corroding metal. When these small local electrodes are coupled electrically by connecting each of them to a common contact through an external circuit, the electrodes that exhibit anodic properties simulate the anodic areas, and the electrodes that exhibit the cathodic properties Electrolyte (liquid, wet gas, bio-film…) Cathodic sites: O2 + 4e– + 2H2O = 4OH– e– e–

e–

e–

e– e–

e–

e–

e–

Anodic sites: M – ne– + nH2O = M(OH)n + nH+ (Electrons flow internally from anodic sites to cathodic sites.)

Metal

Anodic and cathodic sites are separated but coupled externally. Electrons are forced to flow externally. e–

e–

e– e e–

–

e

–

e–

e–

e– e–

e– e– e– e–

e–

Insulators

Instrument measures electrons from and to individual electrodes.

Most corroded electrode simulates maximum penetration

14.13 Schematic diagram showing the principle of coupled multielectrode array sensors for localized corrosion monitoring.40 „ NACE International 2006.

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simulate the cathodic areas of the corroding metal (Fig. 14.13).40 The electrons released from the anodic electrodes are forced to travel through the external circuit to the cathodic electrodes. Thus there are anodic currents flowing into the more corroding electrodes and cathodic current flowing out of the less corroding or non-corroding electrodes. The resulting electrical currents can be measured and the quantitative localized or non-uniform corrosion rates can be determined by the CMAS instrument.23, 41–45 The CMAS probes can be made in many difference configurations and sizes, depending on the application. Figure 14.14 shows typical commercial CMAS probes and a system for corrosion monitoring under high temperature and high pressure conditions. Figure 14.15 shows the principle of a CMAS probe, assuming that one electrode on the probe is anodic and all the other electrodes are cathodic.36, 46 Since localized corrosion often involves small areas of corroded anodic sites accompanied by large areas of cathodic sites, such an assumption is often reasonable in many environments. The thin solid curves represent the dissolution and reduction polarization behaviours on the anodic electrode, respectively. The thick solid curves represent the combined dissolution and reduction polarization behaviours, respectively, on the rest of the electrodes (the cathodic electrodes) if these cathodic electrodes are coupled as a single electrode. The dashed lines represent the reduction curve for all electrodes, or the dissolution current for all electrodes on the CMAS probe, respectively. For a passive metal, in the cathodic area (or the cathodic electrodes in a CMAS probe) where no localized corrosion has been initiated, the anodic current is usually extremely low due to the protective layer of the oxide formed on the metal and the corrosion potential for the cathodic electrodes, Eccorr, is high (or noble). For the anodic electrode where localized corrosion has been initiated and the protective layer has been compromised, however, the anodic current is usually high and the corrosion potential for the anodic electrode, Eacorr, is low (or active). Note in Fig. 14.15, the cathodic current on the combined cathodic electrodes is significantly higher than that on the anodic electrode. This is because it is assumed that the surface area on the anodic electrode is significantly smaller than that of the cathodic electrodes (one anodic electrode versus many cathodic electrodes). In addition, the cathodic reactions deep in an anodic pit on the anodic electrode require more effort for the reactants (O2 or H+) to overcome the mass transfer barriers. When the anodic electrode and the combined cathodic electrodes are coupled, the potential changes to a new value, Ecoup (or Ecorr for all coupled electrodes), and the total anodic dissolution currents equal the total cathodic reduction currents (see the dashed lines in Fig. 14.15):

Icorr + Icin = Iain + Ic

14.13

where Icorr is the corrosion current (total dissolution current) on the anodic

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(a)

(b)

14.14 (a) Typical CMAS probes and (b) systems for corrosion monitoring under high temperature and high pressure conditions. Courtesy of Corr Instruments, LLC, San Antonio, TX, USA for (a) and Southwest Research Institute, San Antonio, TX, USA for (b).

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437

Dissolution of anodic electrode Dissolution of all electrodes

E ccorr SIDissolution = SIReduction Icorr + I cin = I ain + I c

Ecoup

Reduction on cathodic electrodes Reduction on all electrodes

E acorr Reduction on anodic electrode logI cin

logI ain

logIex

logI c logIcorr

Icorr = I ain + Iex I ain << Iex Icorr ~ Iex

14.15 Schematic diagram for the polarization curves on one anodic electrode and several cathodic electrodes that are connected together on a coupled multielectrode sensor.36

electrode, Icin is the internal dissolution (anodic) current on all the cathodic electrodes (anodic current that flows within all the cathodes), Iain is the internal reduction current on the anode (the cathodic current that flows within the anode), and Ic the cathodic current on the combined cathodes. On the anodic electrode, the corrosion current (total dissolution current), Icorr, is equal to the sum of the externally flowing anodic currents, Iex, and the internally flowing anodic currents which is equal to the internally flowing cathodic currents (or the internal reduction current), Iain. Therefore,

Icorr = Iex + Iain

14.14

Since the Iain for the anodic electrode, especially when the anodic electrode is the most anodic electrode of the CMAS probe, is often much smaller than its Iex at the coupling potential in a localized corrosion environment, the externally flowing current from such an anodic electrode of the probe can often be used directly to estimate the localized corrosion current:

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Icorr ≈ Iex

14.15

In some cases of general corrosion, however, there would be less separation between the anodic electrodes and the cathodic electrodes. The behaviour of even the most anodic electrode may be similar to the other electrodes in the CMAS probe. In these cases, the Iain for the anodic electrode would be close to Iex, and Icin would be close to the Ic. On the most anodic electrode, there may still be significant internal electrons flowing from the anodic sites to the cathodic sites within the same electrode. Therefore, Iain in Equation (14.14) cannot be ignored in the calculation of the corrosion current in these cases. In a corrosion management programme for engineering structures, field facilities, or plant equipment, the most important parameter is the remaining life (often the remaining wall thickness) of the systems. If localized corrosion is of concern, the remaining wall thickness in the most corroded area or site is often used to evaluate the remaining life. Therefore, the maximum corrosion depth (the corrosion-induced wall thinning at the most corroded area) for localized corrosion is often the most important parameter in an operator’s mind. Since the corrosion depth is a parameter that takes a long time (often many years) to accumulate, the corresponding parameter that is important to the day-to-day operation would be the corrosion rate or the maximum localized corrosion rate. Since the anodic electrodes in a CMAS probe simulate the anodic sites on a metal surface, the maximum anodic current (the current from the most anodic electrode) may be considered as the corrosion current from the most corroding site on the metal. Therefore, the maximum anodic current should be used as one of the most important parameters for the CMAS probes.41,42 Accordingly, the maximum localized corrosion rate (or maximum localized corrosion penetration rate) may be derived from the maximum anodic current (assuming no internal current effect):41,42

CRmax = Imax EW/(F r A)

14.16

where CRmax is the calculated maximum penetration rate (mm/yr), Imax is the maximum anodic current, or the most anodic current, F is the Faraday constant (96 485 C/mol), A is the surface area of the electrode (mm2), r is the density of the alloy or electrode (kg/m3), EW is the equivalent weight (kg/mol) (see Equation 14.7). Equation 14.16 assumes that corrosion on the most corroded electrode is uniform over the entire surface. Since the electrode surface area is usually between 1 and 0.03 mm2, which is approximately 2 to 4 orders of magnitude less than that of a typical linear polarization resistance (LPR) probe or a typical electrochemical noise (EN) probe, the prediction of penetration rate or localized corrosion rate by assuming uniform corrosion on the small electrode is realistic in most applications.41,42

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The corrosion depth or penetration is related to the total damage accumulated in a given time period. The corrosion depth of the ith electrode may be derived from the cumulative charge that can be obtained by integrating the corrosion current through the electrode from time zero to time t:

Qi =

Ú

I i (t ) dt



14.17

where Qi is the cumulative charge of the ith electrode. Similar to the maximum localized corrosion rate, the following equation has been used to calculate the maximum cumulative localized corrosion depth or penetration (mm):

CDmax = Qmax  EW/(FrA)

14.18

where Qmax is the maximum of the cumulative charges (coulomb) from all the electrodes. The cumulative charge of each electrode is calculated individually using Equation 14.17. Applications Figure 14.16 shows typical responses of the maximum localized corrosion rate of a low carbon steel CMAS probe.47 The maximum localized corrosion rate in air was close to the instrument theoretical detection limit (10 nm/yr). The initial maximum localized corrosion rate was 10 mm/yr in distilled water. The maximum localized corrosion rate in simulated seawater was approximately 1 mm/yr. When 10 mM H2O2 was added to the simulated seawater, the corrosion rate was 10 mm/yr. It was later verified that the corrosion of the electrode in the distilled water and in the simulated seawater were mainly in the form of non-uniform corrosion.40 Therefore, the maximum corrosion rate in Fig. 14.16 represents the maximum non-uniform corrosion rate. A large number of corrosion monitoring case studies at low temperatures have been reported for laboratory and field applications.46 The application areas include cooling water, simulated seawater, salt-saturated aqueous solutions, concentrated chloride solutions, concrete, soil, low-conductivity drinking water, process streams of chemical plants, H2S systems, and oil/ water mixtures. CMAS probes are also used for monitoring corrosion under coatings, deposits of sulphate-reducing bacteria, deposits of salt in air, and for cathodically protected systems. Crevice effect for high temperature applications Similar to all electrochemical methods in corrosion rate measurements, the electrode surface area must be well defined to apply Faraday’s law to calculate the corrosion rate. For highly corrosion-resistant alloys or for corrosion in less corrosive environments, CMAS probes are sensitive to the

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100000.00

3937 3%wt Sea salt+ 10 mMH2O2

Corrosion rate (micrometre/year)

10000.00

3%wt Sea salt

394

3%wt Sea salt

1000.00

39.4 Distilled water

100.00

3.94

Distilled water

10.00

0.394

1.00

0.0394

0.10 0.01

0.00394

Dry air

Dry air

Corrosion rate (mil/year)

440

0.000394 0.0

1.0

2.0

3.0 Time (day)

4.0

5.0

6.0

14.16 Typical response of the maximum localized corrosion rate from a 16-electrode carbon steel probe to the changes in solution chemistry.47 „ NACE International 2004.

presence of crevices that may be formed between each sensing electrode and the surrounding insulator. This is because the CMAS probes have a much smaller sensing surface area than the probes based on other electrochemical techniques. In addition, the crevice between the sensing electrode and the surrounding insulator, if present, may also promote localized corrosion inside the crevice. If the objective is to measure localized corrosion, including crevice corrosion, such as in the case with steam generator tube corrosion under the tube sheet, the measurement with a probe that has a crevice would not be a concern. But for cases where no crevice exists, the measurement with a probe that has a crevice would give false results. It should be mentioned that the concern only exists for highly corrosion-resistant alloys or for corrosion in less corrosive environments. If the solution is highly corrosive or the metal is highly active, such as in the case with carbon steel in seawater, most of the corrosion reactions would take place on the baldly exposed surfaces rather than deep in the crevice because of the limitation by mass transfer through the relatively thin crevice.48 If the solution is not corrosive, or the metal is highly corrosion-resistant, however, mass transfer would not be a limiting factor and corrosion reactions may take place anywhere the metal is in direct contact with the solution, even if the location is deep inside the crevice. In this case, the effective corrosion area is the total area of the metal in contact with the solution.

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Coating

Epoxy

Alloy 22

Coated (a)

Localized corrosion rate (mm/yr)

1000.00 Uncoated wire 100.00

10.00

Average = 30.0 mm/yr Coated wire

1.00

0.10

0.01

Average = 0.43 mm/yr

–1

0

1

2 3 4 Time (day)

5

6

7

(b)

14.17 (a) Scanning electron microscope image of the interface between the DLC coating and the metal substrate of a DLC-coated crevice-free sensing electrode after exposure to brine at 150 °C and (b) typical localized corrosion rates measured from Alloy 22 CMAS probes with and without DLC coating in brine at 150 °C.50,51 „ NACE International 2007 and 2008.

For applications at low temperatures (<80 °C), epoxy coating may be used on the sensing electrodes to successfully avoid the formation of crevices. However, for applications at high temperatures (>100 °C), most epoxy coatings are not effective in preventing the formation of a crevice. Diamond-like carbon (DLC) is mechanically rugged, thermally stable, chemically inert, and electrically insulating. It has been demonstrated that if applied properly,

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DLC can form an excellent coating on a variety of metal electrodes for the CMAS probes and give a well-defined exposed surface area for the sensing electrode.49 Figure 14.17(a) shows the interface between the DLC coating and the metal substrate of an Alloy 22 (UNS N06022) electrode of a CMAS probe after a one-week test in a highly concentrated NaCl-NaNO3-KNO3 brine at 150 °C.49,50 No crevice is visible between the DLC coating and the electrode substrate after the exposure of the electrode at 150 °C. Figure 14.17(b) shows typical measurements from the CMAS probes made of electrodes with or without the DLC coating in a NaCl-NaNO3-KNO3 brine at 150 °C.51 The probe with DLC coating exhibited a reasonable corrosion rate throughout the measurement, while the probe with the electrode that was not coated with the DLC (only supported by an epoxy) gave much higher and exaggerated corrosion rate. The scanning electron microscope (SEM) picture revealed that there was a visible crevice between the supporting epoxy and the electrode for the uncoated CMAS probe after the test at 150 °C. Advantages The CMAS probe is a highly sensitive method (down to several tens of nanometres per year) for monitoring the quantitative rate of non-uniform corrosion, especially localized corrosion. CMAS probes gave continuous real-time readings (every 10–30 seconds) and may be used for process control. CMAS probes are not polarized and there is no disturbance to the electrode. Since the sensing electrodes can be closely packed, the method may be used to measure corrosion under deposits and in humid air. In addition, the sensing electrodes are vertically embedded in an insulator and the life of a CMAS probe can be virtually unlimited. It can be repeatedly polished and reused. Limitations The CMAS method is an electrochemical method. It only works if the corrosion is electrochemical in nature (i.e., in the presence of electrolyte, a liquid phase or a thin film formed by condensation, adsorption or salt deliquescence). CMAS probes cannot be used for monitoring corrosion caused by purely chemical reactions, such as dry air oxidation, and mechanical actions, such as erosion. CMAS probes measure only non-uniform corrosion and cannot be used if the corrosion is dominantly uniform. The CMAS probes respond to the non-uniform portion of the corrosion if there is a mix of uniform and non-uniform corrosion. In addition, CMAS probes are sensitive to crevice corrosion and chemically stable and robust coatings are required for the sensing electrode when the CMAS probe is used for quantitative localized corrosion monitoring in certain environments.

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14.3.3 Differential flow through cell method The galvanically coupled differential flow through cell52 employs essentially a slow-moving (relative to solution) electrode and a fast-moving electrode, both being made of identical corrodible metals and coupled together through a zero resistance ammeter (ZRA). Due to the differences between the electrolyte flow pattern on the surfaces, the slow-moving electrode acts as an anode and the fast-moving electrode acts as a cathode in a localized corrosion environment. Thus, electrons flow from the anodic electrode to the cathodic electrode through the ammeter. The localized corrosion current from the small anodic electrode can be derived from the measurement of the electron flow by the ZRA. The total corrosion current is the sum of the uniform corrosion current and the localized corrosion current. If a separate LPR probe (see Section 14.2.2) is used, the total corrosion current is the sum of the corrosion current from the flow through cell probe and the corrosion current from the LPR probe. This method has been tested for localized corrosion monitoring in cooling water systems.16 The measurement conducted in a cooling water system of an integrated plastics plant shows that the sensor responded to upset or non-ideal operating conditions and the dosing of inhibitors. It was shown to be a useful tool in optimizing treatment performance and identifying upset conditions at the plant as well as providing field verification of the effectiveness of a new treatment programme in reducing the localized corrosion rate of carbon steel.

14.4

Electrochemical potential (ECP) monitoring

Electrochemical potential is one of the most important parameters that relate to corrosion.53 ECP is probably the most commonly measured electrochemical parameter under high temperature and high pressure conditions in the nuclear industry. There is a large body of literature on the measurement of ECP both in laboratory studies27,54 and in operating nuclear power plants for corrosion control.55 The measurement of ECP under reactor or steam generator conditions requires reference electrodes that are chemically stable and mechanically robust at high temperature and high pressures at or near reactor or steam generator operating conditions. There are two types of reference electrodes that can be used under high temperature and high pressure conditions: (1) internal reference electrode operating in the high temperature environment, and (2) external reference electrode in which the electroactive element is maintained at room temperature but is connected to the high temperature environment via a non-isothermal electrolyte bridge.56–58

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14.4.1 Internal reference electrode With an internal reference probe 56 in which the reference electrode (electroactive element) is inside the pressurized vessel and exposed to the same system temperature, the electrode potential can be easily converted to the standard hydrogen electrode (SHE) scale based on thermodynamic data. For the silver-silver chloride/potassium chloride (Ag/AgCl-KCl) system, the electrode potential versus the SHE scale can be expressed as:

o DEAgCl,T = DEAgCl,T + (RT /F ) ln (aCl– )



14.19

o where DEAgCl,T and DEAgCl,T are the electrode potential and the standard electrode potential of the Ag/AgCl system against SHE at temperature T, respectively, and aCl– is the activity of the Cl– at temperature T. Values of DEAgCl,T for selected KCl concentrations (e.g., 0.1 M) are available in the literature.59,60 If the potential of a working electrode measured against the internal reference electrode at temperature T is DEmeas, the potential of the working electrode on the SHE scale at temperature T is:



DEW,SHE,T = DEmeas + DEAgCl,T

14.20

where DEW,SHE,T is the potential of the working electrode against the SHE at temperature T.

14.4.2 External reference electrode Most of the electroactive elements (such as AgCl) are not stable at elevated temperatures. The external pressure-balanced reference electrode (EPBRE) has been the type of choice for applications at high temperatures (above 200 to 250 °C).56 Figure 14.18 shows the schematic of an EPBRE probe.56 The reference electrode is located outside of a pressurized vessel (external) and maintained at room temperature. The liquid junction is maintained by the balancing gas pressure. Since the reference electrode is at room temperature, the chemical stability of the electroactive elements (usually Ag/AgCl) is not affected by the elevated temperature and the probe can be used in almost any temperature that is tolerable by the probe packing materials and pressure boundary tubes and fittings (see Fig. 14.18). With the EPBRE, where the electroactive species is at an ambient temperature, the thermocell potential difference must be accounted for when the EPBRE is used to report the potential of a metal component (working electrode) at system temperature T. The thermocell potential difference can be expressed as:

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Pressure seal Balancing gas pressure Reference electrode Liquid junction solution

Pressure boundary for mounting to pressure vessel

Porous packing (e.g., KCl + glass)

14.18 Schematic of an external pressure-balanced reference electrode.56



DEth = DEAgCl,T – DEAgCl,To + DETLJP



14.21

where DEth is the thermocell potential difference (the experimentally measured potential difference between the internal reference electrode at temperature T and the external reference electrode at temperature TO (usually 25 °C)), o DEAgCl,T and DEAgCl,T are the electrode potentials of the Ag/AgCl system against SHE at temperature T and T o, respectively, and DETLJP is the thermal liquid junction potential. When the potential of a working electrode is measured against the potential of the EPBRE at T o, it can be corrected to the potential of the Ag/AgCl at the system temperature, T, using the following equation:

DEW,AgCl,T = DE¢meas – DETh

14.22

where DEW,AgCl,T is the potential of the working electrode against the Ag/ AgCl electrode at temperature T (or the potential of the working electrode against an internal Ag/AgCl reference electrode at temperature T), and DE¢meas is the potential of the working electrode measured against the EPBRE at T o. Figure 14.19 shows typical values of DEth for different KCl concentrations as a function of the temperature difference (T – T o).60,61 The DEth values increase with the decrease in KCl concentration. For KCl concentration less than 0.1 M, the DEth increases initially with the increase of temperature,

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Understanding and mitigating ageing in nuclear power plants 0.100 m (molal) 0.005 0.080

0.01 0.025 0.051

DEth (V)

0.060

0.102 0.040 0.252 0.505

0.020

0.0

0

50

100

150 DT (K)

200

250

14.19 Dependence of thermocell potential difference for the Ag/ AgCl-KCl system on temperature difference (T–To) at different KCl concentrations (T ° = 25 °C).60

reaching a maximum at the temperature difference of approximately 170 K. The maximum DEth values are 40 mV for 0.1 M KCl, and 22 mV for 0.505 M KCl, respectively. When the concentration of KCl increases further, the DEth value should be less than 22 mV. Therefore, for some applications (such as corrosion potential or redox potential measurements) where 10–20 mV of difference is tolerable, the correction of the thermocell potential effect may not be needed if a concentrated KCl solution (>0.5 M) is used in the salt bridge. The following equation can be used to correct the potential of the working electrode measured against the EPBRE at T o to the SHE scale at temperature T:

DEW,SHE,T = DE¢meas – DEth + DEAgCl,T



= DE¢meas – DETLJP + DEAgCl,T°

14.23

The DEAgCl,T or DEAgCl,T° can be calculated and DEth can be measured (Fig. 14.19). The regression polynomial for DETLJP is available in the literature. After substituting the regression polynomial for DETLJP in a 0.1 M KCl

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system, the potential against the SHE scale may be obtained by using the following Equation 14.24:61

DEW,SHE,T = ΔE¢meas + DEoAgCl,T – [–3 ¥ 10–6  DT 3



+ 0.0024DT 2 + 0.7485DT]

14.24

where DEW,SHE,T and DE¢meas are in mV and DT is the temperature difference (T – T o) in K. When T o = 298 K, DE AgCl,T° = 288 mV. Therefore the following polynomial may be used to convert the potential measured against a 0.1 M KCl EPBRE at 298 K to the SHE scale:

DEW,SHE,T = DE¢meas  (T o = 25 °C) + 288



– [–3 ¥ 10–6 DT 3 + 0.0024 DT 2 + 0.7485 DT] 14.25

Figure 14.20 shows a typical pressure balanced reference electrode probe for applications under high temperature and high pressure conditions. The probe is an external type (EPBRE) if the length of the probe is long (>500 mm) and the electroactive species (AgCl) is near the top of the probe in which case the AgCl is maintained at room temperature. The probe is an internal type if the electroactive species (AgCl) is near the bottom of the probe in which case the AgCl is maintained at the system temperature.

14.4.3 Zirconia membrane pseudo-reference electrode Yttrium stabilized zirconia membrane electrodes filled with electroactive materials such as Cu/Cu2O, Fe/Fe3O3 have been used for pH measurements at elevated temperatures.62,63 These types of electrodes were also used as pseudo-reference electrodes in systems where the pH is well defined.64–66 Many investigators believed that these types of reference electrodes with Cu/Cu2O or Fe/Fe3O3 internal filling materials are the most reliable and least prone to bias and errors at elevated temperatures.66 This is because the electroactive materials (such as Cu/Cu2O or Fe/Fe3O3 ) are inside the zirconia membrane, which has excellent chemical stability in solutions at

14.20 Typical pressure-balanced reference electrode for application under high temperature and high pressure conditions. Note: If the probe is an external type (EPBRE), the length of the probe is long (>500 mm) and the electroactive species is near the top (the left-hand side is the top and the right-hand side is the bottom of the probe) and maintained at ambient temperature. If the probe is an internal type, the electroactive species is at the bottom and maintained at system temperature. Courtesy of Corr Instruments, LLC, San Antonio, TX, USA.

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elevated temperatures and prevents the internal electroactive materials from being contacted by the solutions. The drawback of this type of sensor is that they respond to the changes in solution pH. In addition, the performance of these types of sensor cannot be verified at ambient temperatures and they cannot be used at temperatures below 90 °C because of the extremely low conductivity of the zirconia membrane at temperatures below 90 °C.

14.4.4 Bare metal reference electrode Bare metal electrodes have also been used to provide reference potentials under high temperature and high pressure conditions. An example is the two platinum electrode system67 in which one platinum (2nd platinum electrode) is used as a counter electrode and the other electrode (1st platinum electrode) is used to obtain the cyclic voltammogram using a potentiostat. The decomposition potential of water to form either hydrogen or oxygen was used as the reference potential. During the measurement, the electrical connection for the reference electrode on the potentiostat is connected to the working electrode or the system component whose potential is to be measured and the electrical connection for the working electrode on the potentiostat is connected to the 1st platinum electrode. Since the electroactive species is the water itself in the pressurized system, this type of reference electrode is highly robust under reactor operating conditions (temperature, pressure and radiation). Since the decomposition potential of water is affected by pH, this type of reference system should be considered as a pseudo reference electrode.

14.4.5 Applications It is well recognized that ECP is directly related to the intergranular stress corrosion cracking (IGSCC) of nuclear reactor austenitic stainless steel components. IGSCC only occurs when the alloy’s ECP is above a certain level in the reactor coolant system. To prevent IGSCC of stainless steel components in reactors, hydrogen water chemistry (HWC) has been adopted in the nuclear industry.68–71 Under HWC, hydrogen is added to the reactor water system to lower the electrochemical potential of system components. The measurement of ECP is an indicator of the effectiveness of the HWC. Therefore, it is one of the most important parameters in reactor operations.55,71

14.5

Conclusion

14.5.1 Corrosion modes Corrosion is categorized as general corrosion and localized corrosion. General corrosion is characterized by corrosive attack that extends over the whole © Woodhead Publishing Limited, 2010

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exposed surface or at least over a large area. However, general corrosion may not be uniform, and purely uniform corrosion is rare. The morphology of the corrosion surface produced by general corrosion always exhibits some sort of irregularity and roughness. Localized corrosion is characterized by corrosive attack that is localized over an isolated surface. Examples of localized corrosion include pitting corrosion, crevice corrosion and stress corrosion cracking. Localized corrosion for metal components is usually the major concern in chemical plants and nuclear power plants because localized corrosion is difficult to detect at an early enough stage to mitigate. Once initiated, localized corrosion can propagate rapidly and result in either component failures (such as through wall penetrations) or trigger other modes of failure, such as stress-corrosion cracking, even though the majority surface area is not affected. Therefore, localized corrosion monitoring is vitally important in corrosion control in many industries where small and through wall penetration would cause equipment failures.

14.5.2 Corrosion monitoring methods Coupon techniques for the measurement of the wall thickness of actual system components are the most reliable methods for corrosion monitoring for both uniform and localized corrosion. However, the coupon method is slow and the results are not available until the coupons are retrieved after an exposure (2 to 12 months). In addition, the evaluation of the coupons is labour intensive and time consuming. The real-time monitoring tools can provide on-line and continuous corrosion data and can be used to guide the day-to-day operations. But real-time monitoring tools rely on electronics and involve the use of assumptions or assumed parameters to derive the corrosion rate. Calibration for the real-time tools must be performed frequently. Coupon methods or thickness measurement of the system components should be used as the basis for calibration/validation of the real-time corrosion monitoring tools. Corrosion monitoring techniques for uniform corrosion There are a number of techniques available for monitoring uniform corrosion. The electrical resistance and the ultrasonic techniques are effective methods for monitoring uniform corrosion. The ultrasonic technique is a non-intrusive method and has been used in the field at reactor temperatures. The electrical resistance technique is usually intrusive and the installation of such probes in high pressure systems requires the penetration of the probe through the pressure boundary. Both the ultrasonic and the electrical resistance techniques are based on cumulative metal loss and the response time is long. They cannot be used to guide the day-to-day operations.

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The electrochemical linear polarization resistance technique and the electrochemical noise technique are near instantaneous methods (2–10 minutes of response time) for uniform corrosion rate measurements. They can be used to guide the day-to-day operations. The galvanic technique is a low cost method for the detection of the corrosivity of the corrosion environments. The galvanic technique does not measure the corrosion rate and cannot be used for quantitative corrosion monitoring. Corrosion monitoring techniques for non-uniform and localized corrosion Coupled multielectrode array sensor (CMAS) is a newly emerged technology. It measures the quantitative rate of non-uniform corrosion, including localized corrosion. It provides an instantaneous corrosion rate (down to 30 seconds of response time) and may be integrated with inhibitor addition controllers for dosing control. For quantitative corrosion rate measurements, CMAS probes are more sensitive to the effect of crevices formed between the sensing electrode and the insulator than the other types of electrochemical probes because the CMAS probe uses small electrodes. For low temperature (<80–100 °C) applications, the crevice may be minimized or eliminated by using a proper sealing insulator such as epoxy. For high temperature (>100 °C) applications, the elimination of the crevice has been a challenge. Diamond-like carbon coating has been used to form crevice-free CMAS electrodes for high temperature applications. However, the cost of the probes with diamond-like carbon coating is relatively high at the present time.

14.5.3 Electrochemical potential monitoring Electrochemical potential is the most commonly measured electrochemical parameter under high temperature and high pressure conditions, even under reactor operating conditions. Electrochemical potential probes have been used to maintain the reducing condition in reactors to avoid intergranular stress-corrosion cracking (IGSCC) of reactor components.

14.6

Acknowledgements

A great amount of information for this chapter was taken from Techniques for Corrosion Monitoring (L. Yang, ed., Woodhead Publishing, Cambridge, 2008). The authors of this chapter owe a debt of gratitude to the many contributing authors of the book. The authors also acknowledge the reviews of Drs T. Mintz and D. Ferrill and the assistance of C. Patton in preparing the manuscript.

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14.7

451

References

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16. B. Yang, ‘Corrosion Monitoring in Cooling Water Systems Using Differential Flow Cell Techniques’, in Techniques for Corrosion Monitoring, L. Yang, ed., Woodhead Publishing, Cambridge (2008), Chapter 23. 17. K.B. Oldham and F. Mansfeld, ‘Corrosion Rates from Polarization Curves: A New Method’, Corrosion Science, 13, 813–819 (1973). 18. F. Mansfeld and K.B. Oldham, ‘A Modification of the Stern–Geary Linear Polarization Equation’, Corrosion Science, 11, 787–796 (1971). 19. J.R. Scully, ‘Polarization Resistance Method for Determination of Instantaneous Corrosion Rates’, Corrosion, 56 (2), 199–218 (2000). 20. R.A. Cottis, ‘Electrochemical Noise for Corrosion Monitoring’, in Techniques for Corrosion Monitoring, L. Yang, ed., Woodhead Publishing, Cambridge (2008), Chapter 4. 21. D.A. Eden, K. Hladhy, D.C. John and J.L. Dawson, ‘Electrochemical Noise Simultaneous Monitoring of Potential and Current Noise Signals from Corroding Electrodes’, CORROSION/1986, Paper 274 (Houston, TX: NACE International, 1986). 22. Y.J. Tan, S. Bailey and B. Kinsella, ‘The Monitoring of the Formation and Destruction of Corrosion Inhibitor Films Using Electrochemical Noise Analysis (ENA)’, Corrosion Science, 38, 1681–1695 (1996). 23. L. Yang, N. Sridhar, C.S. Brossia and D.S. Dunn, ‘Evaluation of the Coupled Multielectrode Array Sensor as a Real Time Corrosion Monitor’, Corrosion Science, 47, 1794–1809 (2005). 24. R.D. Klassen and P.R. Roberge, ‘Zero Resistance Ammetry And Galvanic Sensors’, in Techniques for Corrosion Monitoring, L. Yang, ed., Woodhead Publishing, Cambridge (2008), Chapter 5. 25. V.S. Agarwala and S. Ahmad, ‘Corrosion Detection and Monitoring – A Review’, CORROSION/2000, Paper 271 (Houston, TX: NACE International, 2000). 26. G. Light, ‘Nondestructive Evaluation Technologies for Monitoring Corrosion’, in Techniques for Corrosion Monitoring, L. Yang, ed., Woodhead Publishing, Cambridge (2008), Chapter 12. 27. N.Y. Lee, S.G. Lee, K.H, Ryu and I.S. Hwang ‘On-line Monitoring System Development for Single-phase Flow Accelerated Corrosion’, Nuclear Engineering and Design, 237, 761–767 (2007). 28. P. Kielczynski and J. Goszczynski, ‘Recent Developments in Ultrasonic Devices for Monitoring Critical Parameters in Canadian Nuclear Reactors’, 1998 IEEE Ultrasonics Symposium Proceedings, 1, 793–802 (1998). 29. D.C. Earble, ‘Radioactive Tracer Methods’, in Techniques for Corrosion Monitoring, L. Yang, ed., Woodhead Publishing, Cambridge (2008), Chapter 10. 30. N. Rothwell, D.A. Eden and G. Row, ‘Electrochemical Noise Techniques for Determining Corrosion Rates and Mechanisms’, CORROSION/1992, Paper 223 (Houston, TX: NACE International, 1992). 31. D.A. Eden, ‘Electrochemical Noise’, in Uhlig’s Corrosion Handbook, 2nd edn, R.W. Revie, ed., John Wiley and Sons, New York, 2000 (Chapter 69). 32. E. Garcia-Ochoa, R. Ramirez, V. Torres, F.J. Rodriguez and J. Genesca, ‘Comparison of Electrochemical Noise and Wire-on-Screw Technique in Simulated Marine Atmospheres’, Corrosion, 58, 756–760 (2002). 33. R.G. Kelly, M.E. Inman and J.L. Hudson, ‘Analysis of Electrochemical Noise for Type 410 Stainless Steel in Chloride Solutions’, in Electrochemical Noise Measurement for Corrosion Applications, J.R. Kearns, J.R. Scully, P.R. Roberge, D.L. Reichert,

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J.L. Dawson (eds.), ASTM Special Technical Publication 1277, Conshohocken, PA, 1996, pp. 101–113. 34. S.T. Pride, J.R. Scully and J.L. Hudson, ‘Analysis of Electrochemical Noise from Metastable Pitting in Aluminum, Aged Al–2%Cu, and AA 2024-T3; in Electrochemical Noise Measurement for Corrosion Applications, J.R. Kearns, J.R. Scully, P.R. Roberge, D.L. Reichert, J.L. Dawson, eds., ASTM Special Technical Publication 1277, Conshohocken, PA, 1996, p. 307. 35. G.L. Edgemon, ‘Electrochemical Noise Corrosion Monitoring in Radioactive Liquid Waste Storage Tanks’, Materials Performance, 44, (2) 52–55 (2005). 36. L. Yang, ‘Multielectrode Systems’, in Techniques for Corrosion Monitoring, L. Yang, ed., Woodhead Publishing, Cambridge (2008), Chapter 8. 37. Z. Fei, R.G. Kelly and J.L. Hudson, ‘Spatiotemporal Patterns on Electrode Arrays’, J. Phys. Chem., 100, 18986–18991 (1996). 38. Y.J. Tan, ‘Wire Beam Electrode: A New Tool for Localized Corrosion Studies’, Proceedings of Australasian Corrosion Association Corrosion & Prevention 97, Paper No. 52, Australasian Corrosion Association, Australia., Nov. 9–12 (1997). 39. Y.J. Tan, ‘Monitoring Localized Corrosion Processes and Estimating Localized Corrosion Rates Using a Wire-beam Electrode’, Corrosion, 54 (5), 403–413 (1998). 40. X. Sun and L. Yang, ‘Real-Time Monitoring of Localized and General Corrosion Rates in Drinking Water Utilizing Coupled Multielectrode Array Sensors’, CORROSION/2006, Paper 06094 (Houston, TX: NACE, 2006). 41. L. Yang and N. Sridhar, ‘Coupled Multielectrode Online Corrosion Sensor’, Materials Performance, 42 (9), 48–52 (2003). 42. L. Yang, N. Sridhar, O. Pensado and D. Dunn, ‘An In-situ Galvanically Coupled Multi-Electrode Array Sensor for Localized Corrosion’, Corrosion, 58, 1004 (2002). 43. L. Yang and D. Dunn, ‘Evaluation of Corrosion Inhibitors in Cooling Water Systems Using a Coupled Multielectrode Array Sensor’, CORROSION/2002, Paper 02004 (Houston, TX: NACE International, 2002). 44. L. Yang, N. Sridhar and G. Cragnolino, ‘Comparison of Localized Corrosion of FeNi-Cr-Mo Alloys in Concentrated Brine Solutions Using a Coupled Multielectrode Array Sensor’, CORROSION/2002, Paper 02545 (Houston, TX: NACE International, 2002). 45. L. Yang and N. Sridhar, ‘Monitoring of Localized Corrosion’, in ASM Handbook, Volume 13A, Corrosion: Fundamentals, Testing, and Protection, Crammer and B.S. Covino, Jr., eds, ASM International, Materials Park, OH, 519–524 (2003). 46. L. Yang and K.T. Chiang, ‘A Review of Coupled Multielectrode Array Sensors for Corrosion Monitoring and a Study on the Behaviors of the Anodic and Cathodic Electrodes’, Journal of ASTM Vol. 6, No. 3, Paper ID JA101253, (2009). 47. X. Sun, ‘Online Monitoring of Corrosion under Cathodic Protection Conditions Utilizing Coupled Multielectrode Sensors’, CORROSION/2004, Paper 04094 (Houston, TX: NACE International, 2004). 48. X. Sun and L. Yang, ‘Real-time Measurement of Crevice Corrosion with Coupled Multielectrode Array Sensors’, CORROSION/2006, Paper 06679 (Houston, TX: NACE International, 2006). 49. K.T. Chiang, L. Yang, R. Wei and K. Coulter, ‘Development of Diamond-like Carbon-coated Electrodes for Corrosion Sensor Applications at High Temperatures’, Thin Solid Films, 517, 1120–1124 (2008).

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50. K.T. Chiang and L. Yang, ‘Development of Crevice-Free Multielectrode Sensors for Elevated Temperature Applications’, CORROSION/2007, Paper 07376 (Houston, TX: NACE International, 2007). 51. K.T. Chiang and L. Yang, ‘Development of Crevice-Free Electrodes for Multielectrode Array Sensors for Applications at High Temperatures’, Corrosion, 64, 805–812 (2008). 52. B. Yang, ‘Differential Flow through Cell Technique’, in Techniques for Corrosion Monitoring, L. Yang, ed., Woodhead Publishing, Cambridge (2008), Chapter 6. 53. S. Uchida, ‘Corrosion of Structural Materials and Electrochemistry in High Temperature Water of Nuclear Power Systems’, PowerPlant Chemistry, 10 (11), 1–19 (2008). 54. D.H. Lister, L. Liu, A.D. Feicht, M. Khatibi, W.G. Cook, K. Fujiwara, E. Kadoi, T. Ohira, H. Takiguchi and S. Uchida, ‘A Fundamental Study of Flow-Accelerated Corrosion in Feedwater Systems’, PowerPlant Chemistry, 10 (11) (2008). 55. S. Hettiarachchi, D.A. Hale, R. Burrill, L. Gorrochategui, R. Coello, S. Suzuki and M. Sambongi, ‘First Lower Plenum ECP Measurement in an Operating BWR’, in Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors IX, F.P. Ford, S.M. Bruemmer and G.S. Was, eds, TSM, Warrendale, PA, pp. 435–442 (1999). 56. D.D. Macdonald, ‘Reference Electrodes for High Temperature Aqueous Systems – A Review and Assessment’, Corrosion, 34, 75–84 (1978). 57. M.J. Danielson, ‘A Long-Lived External Ag/AgCl Reference Electrode for Use in High Temperature/Pressure Environments’, Corrosion, 39, 202–203 (1983). 58. C.M. Menendez, ‘Reference Electrodes for High Pressure and High Temperature Electrochemical Testing’, CORROSION/2001, Paper 01305 (Houston, TX: NACE International, 2001). 59. D.D. Macdonald, A.C. Scott and P. Wentrcek, ‘External Reference Electrodes for Use in High Temperature Aqueous Systems’, J. Electrochemical Society, 126, 908 (1979). 60. D. Macdonald, A.C. Scott and P. Wentrcek, ‘Silver-Silver Chloride Thermocells and Thermal Liquid Junction Potentials for Potassium Chloride Solutions at Elevated Temperatures’, J. Electrochemical Society, 126, 1618 (1979). 61. R.W. Bosch, W.F. Bogaerts and J.H. Zheng, ‘Simple and Robust External Reference Electrodes for High-Temperature Electrochemical Measurements’, Corrosion, 59, 162–171 (2003). 62. L.W. Niedrach, ‘A New Membrane Type pH Sensor for Use in High TemperatureHigh Pressure Water’, J. Electrochemical, Society., 127, 2122–2130 (1980). 63. T. Tsuruto and D.D. Macdonald, ‘Stabilized Ceramic Membrane Electrodes for the Measurement of pH at Elevated Temperatures’, J. Electrochemical Society, 129, 1221–1225 (1982). 64. L.W. Niedrach, ‘Use of a High Temperature pH Sensor as a “Pseudo-Reference Electrode” in the Monitoring of Corrosion and Redox Potentials at 285 °C’, J. Electrochemical, Society., 127, 1445–1449 (1982). 65. L.W. Niedrach and W.H. Stoddard, ‘Monitoring pH and Corrosion Potentials in High Temperature Aqueous Environments’, Corrosion, 41, 45–51 (1985). 66. Y.-J. Kim and P.L. Andresen, ‘Data Quality, Issues, and Guidelines for Electrochemical Corrosion Potential Measurements in High-Temperature Water’, Corrosion, 59, 584–596 (2003). 67. L. Yang, ‘A Bare Metal Reference Electrode for Application in the High-Temperature and High-Pressure Coolants of Nuclear Reactors’, CORROSION/1999, Paper 460 (Houston, TX: NACE International, 1999). © Woodhead Publishing Limited, 2010

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68. M.E. Indig, ‘Investigation of the Protection Potential Against IASCC’, CORROSION/1992, Paper 71 (Houston, TX: NACE International, 1992). 69. M.E. Indig, ‘Technology Transfer: Aqueous Electrochemical Measurements Room Temperature to 290 °C’, Corrosion, 46, 680–686 (1990). 70. M.E. Indig and J.L. Nelson, ‘Electrochemical Measurements and Modeling Predictions in Boiling Water Reactors under Various Operating Conditions’, Corrosion, 47, 202–209 (1991). 71. W. R. Kassen and D. Cubicciotti, ‘Proposed Guidelines for Implementing ECP Measurements in Boiling Water Reactors’, CORROSION/1990, Paper 485 (Houston, TX: NACE International, 1990).

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15

Multi-scale modelling of irradiation effects in nuclear power plant materials

L. M a l e r b a, SCK∑CEN, Belgium

Abstract: This chapter surveys the computer-based multi-scale modelling approaches currently being used to develop physical models of the effects of radiation on nuclear materials. The focus is on the problem of radiationinduced hardening (and embrittlement) in steels, limited to the scales ranging from the nucleus to the single crystal. First, the multi-scale nature of radiation effects is illustrated, including examples of microstructural and mechanical property changes observed in steels used in nuclear reactors. Then the chapter discusses the fundamental ideas upon which the multi-scale modelling approach is based. Next, an overview of the techniques of use in a multi-scale modelling framework is given, with an example of how these can be integrated. A discussion of the state-of-the-art and other general remarks conclude the chapter. Key words: multi-scale modelling, radiation effects.

15.1

Introduction

Since the early 1990s, the development of physical models that describe radiation damage effects in solids based on the extensive use of numerical techniques has received a tremendous boost, on the wave of the exponential performance growth of computers. In a matter of a few years, the application of numerical tools in ‘brute force’ computer simulations resulted in a leap forward in our understanding of how pure metals (e.g., face-centred-cubic (fcc) Cu, Al and Ni; body-centred-cubic (bcc) Fe; and hexagonal-close-packed (hcp) Zr; see Almazouzi et  al., 2000) behave under irradiation. The rapid success of this approach, commonly known as ‘multi-scale modelling’ or ‘multi-scale simulation’ (Díaz de la Rubia and Bulatov, 2001; Lu and Kaxiras, 2004), in providing not only qualitatively but also sometimes quantitatively valid predictions, created much expectation. So much so that, from 2000, even large utilities and industries became interested in such models, as a promising way to develop physics-based tools for predicting the in-service behaviour of nuclear components (Jumel et  al., 2000a; 2000b; Malerba et al., 2002; Massoud et al., 2006). Multi-scale modelling tools hold the promise to support and complement the pressure vessel surveillance programmes of current and future nuclear power plants, as well as the related research programmes in test reactors, 456 © Woodhead Publishing Limited, 2010

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whose aim is to guarantee safe operation until the end of service life and possibly to extend it (Massoud et  al., 2006). Such tools are also expected to allow improved prediction of the time-to-failure of in-core components, which are exposed to the risk of irradiation-assisted stress-corrosion cracking (IASCC), thereby reducing the number of expensive component replacement operations needed to avoid potential failure (Massoud et al., 2006). At the moment, most investigations into the behaviour of materials used in nuclear power plants are conducted by irradiating specimens, either in surveillance capsules, or in test reactors, and subsequently performing mechanical tests and (to a lesser extent) microstructural characterisations in hot cell facilities, where the samples are handled remotely and safely. Pressure vessel surveillance capsules are positioned nearer to the core than the vessel wall and thus receive a slightly higher flux. The specimens therefore reach higher fluence more quickly than the vessel wall, and their periodic testing enables the evolution of embrittlement in the wall to be predicted (Kirk et al., 2003). In materials test reactors, specimens are irradiated at much higher fluxes, in conditions otherwise similar to those experienced in service, so as to reach the same fluences expected at the end of the service life and beyond in a much shorter amount of time. However, the number of capsules available for evaluating the performance of the pressure vessel materials beyond the originally envisaged service life is limited. Furthermore, this procedure is expensive, which limits the number of tests that can be conducted and different irradiation conditions that can be explored. Moreover, in the last couple of decades the number of operating test reactors and hot cell laboratories has been decreasing steadily worldwide, stricter safety requirements have made managing such facilities more and more expensive and even the relevant expertise is slowly disappearing. Thus, the amount of data available for longterm predictions is reaching its limit and alternative means for evaluating the behaviour of materials under irradiation are needed. Multi-scale simulation is potentially a valuable alternative. At present, long-term predictions of the lifetimes of nuclear components are made using semi-empirical correlations that provide trend curves (US NRC, 1988, 2007; Eason et al., 1998; ASTM E900-02, 2007). These correlations are based on large reference databases, which include surveillance data and also data from materials test reactors. Their reliability will benefit from a deeper understanding of radiation damage mechanisms, which can be achieved using multi-scale modelling approaches and tools. Modelling becomes even more important in connection with the design of future nuclear options: Generation IV reactors (GenIV, 2002); spallation neutron sources, such as those in accelerator-driven systems (ADS), for partitioning and transmutation of isotopes (OECD/NEA, 2002); and fusion systems (see e.g. http://www.ofes.fusion.doe.gov; Bloom, 1998). In this case, not only economics, safety, reliability and efficiency, but also feasibility will

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ultimately depend on the ability of the structural materials to maintain their dimensional and mechanical integrity, under extremely hostile conditions (Bloom et al., 2004; Mansur et al., 2004; Zinkle, 2008). Furthermore, the radioactivity induced in the materials during their service life is expected to be an important issue at decommissioning; thus, the development of low, or reduced, activation materials is a priority, especially for fusion applications. A certain degree of success has already been achieved in this direction (Bloom, 1998; Bloom et al., 2004). In these innovative reactor concepts, operating temperatures are expected to be significantly higher than in current nuclear power plants (600–1000 °C versus ~300 °C). Moreover, the final neutron fluences are expected to be tens to hundreds of times larger: up to 200 displacements per atom (dpa) (Norgett et  al., 1975; ASTM E693-01, 2007), compared with fractions of dpa in present generation reactor pressure vessels and a few tens of dpa in internals. Finally, the materials will be in contact with potentially aggressive coolants, such as liquid metals or molten salt, depending on the type of reactor (Bloom, 1998; Bloom et al., 2004; Mansur et al., 2004; Zinkle, 2008). There is currently no direct experience of the performance of materials in such a combination of extreme conditions. Demonstrating the feasibility and safe design of future reactor concepts will thus depend largely on the possibility of testing materials in the laboratory, subjecting them to conditions that mimic those expected in operation. However, these conditions cannot be fully reproduced in any existing irradiation facility, and facilities formerly used to reach high fluences are no longer operational (Bloom et al., 2004). New, expensive and dedicated facilities must be built for this purpose and costly irradiation campaigns carried out. Even so, these facilities will only partially simulate the operating conditions, and in some cases, there will only be space for miniaturised specimens (e.g. IFMIF, see Noda et  al., 1998, or Möslang, 2008). Thus, these experiments will be in practice only experimental simulations, providing scattered data that will then need to be properly combined and extrapolated to real conditions. A guide to choosing the conditions to be explored, as well as allowing safe data extrapolation, will be necessary. Proper physical models, based on a precise and quantitative understanding of the fundamental mechanisms of the onset and evolution of radiation damage, are essential for the correct interpretation, rationalisation and extrapolation of data obtained in these planned experimental irradiation facilities (Ishino, 1996; Stoller et al., 2004; Zinkle, 2005; 2008). In this chapter, a short overview of the multi-scale modelling approach is given, as applied to describing the behaviour of structural materials subjected to irradiation. In Section 15.2, the multi-scale nature of radiationinduced degradation processes in materials is illustrated, along with a short summary of the main microstructural features observed in steels used for

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nuclear applications, and their mechanical behaviour under irradiation. In Section 15.3, the essential features of the multi-scale modelling approach are summarised and compared with other approaches, while in Sections 15.4–15.7 the main multi-scale modelling tools specifically used for nuclear materials studies are briefly presented. In Section 15.8, an example of how this approach can be applied is provided, showing the problems that must still be overcome. A short discussion of the state-of-the-art is given in Section 15.9, and concluding remarks in Section 15.10. Although it is the belief of the author that modelling-oriented experiments are also an integral part of the multi-scale modelling approach, for reasons of space and compactness this aspect will only be mentioned briefly here. In addition, although fuel modelling is equally important, especially in connection with innovative nuclear systems based on fission (see e.g. F-Bridge project, at http://www.fbridge.eu/index.php/Project-Description/Objectives.html), the examples cited here are limited to a few specific structural materials and phenomena. Finally, although electron and ion irradiation facilities are often used as a practical means for studying radiation effects in solids, the type of radiation referred to in this chapter will be neutrons. Thus, only changes in the mechanical properties (hardening and embrittlement) of metallic crystalline materials (specifically Fe alloys as models for steels) subjected to neutron irradiation will be discussed.

15.2

An overview of radiation effects

15.2.1 Radiation effects as a multi-scale problem The macroscopic behaviour of materials is necessarily always the result of atomic-level processes. However, in many instances, the discrete atomic nature of materials, and the fact that their chemical composition may change locally, is neglected for modelling purposes and continuum approaches can be in practice very effectively used. The thermomechanical behaviour of plant components, also for nuclear applications, is typically modelled using finite element (FE) techniques. With these techniques, the continuum equations governing elastic and plastic behaviour, coupled if needed to heat or mass transport equations, are solved with appropriate boundary conditions. The core of the methodology in this case is constituted of phenomenological constitutive laws, which provide the relationship between, for example, stress and strain for each phase, or between temperature gradient and heat flux, or between concentration gradient and mass flux. The effective physical parameters that appear in these laws, such as elastic moduli, thermal conductivity or diffusivity, must be known. Within this continuum approach, the main concern is that the phenomenological constitutive laws and the parameters that appear in them

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should be representative of the actual properties of the real material, which is not a continuum. If proper constitutive laws are given, several methodologies are traditionally used to compute, for example, the macroscopic stress–strain response of the material and therefore its effective plastic behaviour. Refined calculations can be performed using numerical homogenisation methods with FE calculation of microstructures (Barbe, 2001a, 2001b). In parallel, analytical homogenisation techniques alternative to FE methods are well developed (Bornert et  al., 2001), recent work based on fast Fourier transformation being especially interesting (Moulinec and Suquet, 1998; Lebensohn, 2001). Based on work at these scales, the real component scale can eventually be addressed, even though it is computationally impossible to treat the real component scale down to the detail of single grains. In the case of radiation effects in solids, however, no physically-grounded model can completely ignore the atomic nature of materials and the presence of different chemical species, because most processes of importance are strictly atomic (or even nuclear) in nature. The development of continuum physical models is possible, so long as sufficient information translating the effect of atomic, nano- and microstructure is brought into the constitutive laws. As will be seen, the bridge from atomistic to continuum models can in principle be built using mesoscale discrete models. Radiation effects originate in the interaction of energetic (usually >1 MeV) neutrons entering the material and colliding with the atoms composing it. The interaction is nuclear in nature, takes only a fraction of a femtosecond (10–15 s) and has long been known to lead to three main phenomena: activation, transmutation and atomic displacement (Seitz, 1952; Greenwood, 1994). Activation is the process whereby, upon the absorption of a neutron (inelastic reaction), a previously stable nucleus becomes radioactive. Transmutation is the production of chemical elements initially absent in the material, due to either the absorption of a neutron in a nucleus, which induces a change in atomic number, or the emission of a particle from an activated nucleus (a proton becoming a hydrogen (H) atom or an a-particle becoming a helium (He) atom). Finally, atomic displacements occur mainly when the neutron is not absorbed, but bounces off the target nucleus (elastic scattering), making the latter recoil and, if the transferred energy is high enough, also causing it to be ejected from its initial position. All three effects are pernicious and the extent of their negative consequences depends largely on the energy spectrum of the impinging neutrons, their rate of arrival (flux), and the duration of the exposure (fluence). Activation may pose radiological safety problems during plant operation and decommissioning. In the case of high-energy neutrons (e.g. 14 MeV neutrons from fusion reactions), activation is expected to be significant, hence the need to develop low, or reduced, activation structural materials (Conn et al., 1984; Bloom, 1998; Bloom et al., 2004; Zinkle, 2005). Transmutation

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is expected to have an impact only in the case of high-energy neutron spectra and after prolonged exposure to radiation, when the production of He and H becomes significant. It is therefore of concern mainly for future fusion and spallation sources, as well as, to a lesser extent, for Gen IV reactors (Mansur et al., 2004; Smith, 2004; Haight, 2008). Atomic displacements, on the other hand, affect the materials’ properties from the very early stages of exposure to irradiation. Of concern is the fact that, after prolonged exposure to highenergy neutrons, the synergistic effect of transmutation (especially He and H) and atomic displacement may exacerbate the degradation (Trinkaus and Singh, 2003; Henry et al., 2003; Tanaka et al., 2004). However, despite the radiological safety concerns stemming from activation and the problems posed by transmutation, the focus here will be mainly on the effect of atomic displacements and on the effort made to model their consequences at different length and timescales. Recoiling atoms lose energy by inducing electronic excitation in the host material but, provided that the energy transferred from the neutron to the atom after the collision is higher than a threshold energy for displacement, Ed, energy is also lost in elastic and inelastic collision events with other atoms (Robinson, 1994; Greenwood, 1994; Averback and Díaz de la Rubia, 1998). The energy that is not lost in electronic excitation is called damage energy (ED) because, if high enough, it causes the displaced atom to induce other atomic displacements, thereby damaging the crystal lattice. The atom hit by the neutron is customarily called the primary knock-on atom (PKA). When there are also many secondary displaced atoms, a branching atomic displacement sequence called a displacement cascade is produced (Brinkman, 1954, 1956; Seitz and Koehler, 1956; Seeger, 1958). Displacement cascades are complex phenomena in which different overlapping phases can be distinguished: the ballistic phase, thermal spike, and recombination (or cooling) phase. This was already understood in the 1950s (Brinkman, 1954, 1956; Seitz and Koehler, 1956; Seeger, 1958) and has been confirmed more recently by advanced atomic-level studies (Calder and Bacon, 1993; Averback and de la Rubia, 1998; Terentyev et al., 2006; Malerba, 2006). However, the overall lifetime of a displacement cascade is only a few picoseconds (10–12 s) and the region affected by it has a characteristic length of only a few nanometers (10–9 m). If the cascade energy is high (e.g. above ~20 keV in Fe), many subcascades are produced, rather than one single cascade (Stoller and Greenwood, 1999; Terentyev et  al., 2006). At the end of the displacement cascade process, a number of pointdefects are left in the affected region, i.e. vacancies (empty lattice sites) and self-interstitial atoms (SIA, atoms occupying off-lattice positions). These can either be isolated or form clusters. The distribution of defects at this point, reached in a few tens of picoseconds, defines the so-called primary state of damage, or cascade debris (Averback and Díaz de la Rubia, 1998).

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Figure 15.1 provides a pictorial representation of the different phases of a displacement cascade process, specifying the relevant timeframe. Cascades are produced continuously within the material. However, in most cases, the cascades in a given volume of material are produced so far from, or so much later than, previous cascades that they do not have any direct interaction with each other. This can be estimated easily by orders of magnitude. If a cascade that produces about 100 displacements is conservatively assumed to affect a region of about 1 million atoms (i.e. each cascade brings 10–4 dpa) and if the cascade debris lifetime is, also conservatively, assumed to be 1 s (Malerba et al., 2005), the dpa rate required for cascade overlap is 10–4 dpa/s. This rate should be compared with, for example, ~10–10 dpa/s experienced by a reactor vessel, or ~10–7 dpa/s typical of material test reactors. Thus, it is mainly the further evolution of the defects forming the cascade debris, in interplay with the chemical elements composing the material (microstructural evolution), that determines the material’s property changes observed at the macroscopic level (Bullough and Wood, 1986; Eyre and Matthews, 1993; Phythian and English, 1993; Mansur, 1994; English et  al., 1997; Singh et al., 1997b; Singh, 1998; Trinkaus et al., 2000; Gan et al., 2001). These changes occur at the pace of the diffusion properties of the defects, which ranges from micro- and milliseconds to seconds (Mansur, 1994; Singh et al., 1997b). Self-interstitial-type defects generally migrate faster than vacancytype defects (Mansur, 1994; Singh et al., 1997b). Defects migrate until they are absorbed at so-called sinks, therefore the size and density of the sinks, together with the defect migration mechanism, determine the mean distance covered by migrating defects and their lifetime (Bullough and Wood, 1986; Trinkaus et al., 2000, 2002; Barashev et al., 2001). A sink is any microstructural feature capable of absorbing a specific defect. For example, if a self-interstitial and a vacancy meet, they annihilate each other (recombination). Therefore, they can be regarded as sinks for each other. A cluster of point-defects is a sink for single point-defects: upon absorption of single point-defects (a process that is energetically favoured), the cluster grows, becoming for example a three-dimensional cavity (void) or a platelet (dislocation loop). In turn, these clusters may migrate (Trinkaus et al., 1992; Osetsky et al., 2003; Terentyev et al., 2007a; Fu et al., 2005; Djurabekova et  al., 2007a). Dislocations, grain boundaries and free surfaces, typically present in materials, are extended sinks for both single point-defects and migrating clusters (Bullough and Wood, 1986; Barashev et al., 2001). While migrating to sinks, defects cause a redistribution of chemical elements by diffusion. Radiation-generated defects, therefore, assist generally lengthy diffusion processes (Sizmann, 1978), such as precipitation (Sklad and Mitchell, 1974; Barbu and Martin, 1977; Odette, 1983; Mathon et al., 2003, 2005) or segregation (Kameda and Bevolo, 1989; Allen and Was, 1998). Irradiation makes these processes possible even at temperatures at which

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15.1 Different phases of a 10 keV displacement cascade in iron as simulated by molecular dynamics (only selfinterstitials are visualised). The ballistic phase lasts only a fraction of a picosecond. The thermal spike, at which in some metals a local melt can be produced, lasts only a few picoseconds. The primary damage state is stabilised after tens of picoseconds. At that point, much slower diffusion phenomena determine further recombination and migration of defects away from the cascade region. (The cascade snapshots are a courtesy of D. Terentyev.)

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they would not take place in normal conditions (radiation-enhanced, e.g. Sklad and Mitchell, 1974; Odette, 1983), or even induces them outside the temperature and concentration ranges in which they are thermodynamically expected (radiation-induced, e.g. Barbu and Martin, 1977). In turn, the kinetics of defect cluster formation and defect recombination is influenced by the interaction of the defects with the chemical species, both impurities and solute atoms, that influence their mobility (Cottrell et al., 2004; Terentyev et al., 2005, 2007b; Fu et al., 2008). These processes, depending on temperature and other factors, develop over timescales ranging from seconds to years, and produce features at the nano-to-micrometre scales that generally only a combination of advanced and often sophisticated experimental techniques can unambiguously detect (Eyre and Matthews, 1993; Phythian and English, 1993; English et al., 1997; Carter et al., 2001). Thus, atomic-level diffusion processes lead to the formation of new microstructural and microchemical features in the material subjected to the continuous production of defects in displacement cascades. A few important examples are described below.

15.2.2 Examples of microstructural features observed in steels for nuclear applications Reactor pressure vessels (RPV) are typically made from copper-containing bainitic steels (bcc structure). In service, these vessels are irradiated up to relatively low doses (~0.1 dpa) at ~300 °C (service temperature). After neutron irradiation, small-angle neutron scattering (SANS) and tomographic atom probe (TAP) techniques reveal the appearance in these steels of relatively dilute copper-rich precipitates (copper is almost insoluble in iron), which contain also sometimes phosphorus and typically nickel, manganese and silicon (English et al., 1992a; Pareige et al., 1997; Auger et al., 2000; Carter et  al., 2001; Miller et  al., 2003; Miller and Russel, 2007, and references therein). The latter elements are also found to form copper-free precipitates, at sufficiently high dose (in excess of ~0.1 dpa), in low-copper or copperfree alloys, especially those with high nickel content (Miller and Russel, 2007, and references therein; Meslin, 2007; Meslin et  al., 2009, 2010). Furthermore, phosphorus is found to distribute heterogenously and to segregate at dislocations and grain boundaries (Kameda and Bevolo, 1989; Miller and Russel, 2007, and references therein). Electron microscopy reveals the existence of interstitial-type dislocation loops (typically with either ½·111Ò or ·100Ò Burgers vectors) in low-dose, neutron-irradiated pure iron (Eyre and Bartlett, 1965; Nicol et  al., 2001; Zinkle and Singh, 2006, and references therein) and model alloys for RPV steels, i.e. iron alloys containing only some of the elements found in the steels, selected to investigate their effects separately (Hoelzer and Ebrahimi, 1995; Hernández-Mayoral and Gómez-Briceño, 2010). In pure iron also

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small voids (i.e. vacancy clusters grown into cavities) are visible to the microscope, while positron annihilation experiments reveal the existence of invisible vacancy clusters in the model alloys (Lambrecht, 2009; Lambrecht, et al., 2010). However, neither voids, nor loops are observed in actual RPV steels irradiated under customary conditions (Phythian and English, 1993; Meslin et al., 2010; Lambrecht, 2009; Lambrecht et al., 2010), most likely because they are too small to be detectable. ·100Ò and ½·111Ò loops are on the contrary the most frequently observed features in other ferritic alloys irradiated up to doses in excess of ~0.5 dpa, for example high-chromium ferritic/martensitic steels, and model alloys of similar chromium content (Matijasevic et al., 2008; Matijasevic and Almazouzi, 2008; Yao et  al., 2008; Hernández-Mayoral et al., 2008). Figure 15.2 shows how ½·111Ò and ·100Ò loops appear from an atomic-level perspective. Both molecular dynamics simulations (see Section 15.5.1) and experiments show that ½·111Ò loops migrate easily in one dimension, along their Burgers vector direction. Technically, they are said to be glissile. (Computer studies of loop mobility are found, e.g. in Osetsky et  al., 2003, and references therein, or Terentyev et  al., 2007a, and references therein; in Arakawa et al., 2007, as well as in Yao et al., 2008, and Hernández-Mayoral et al., 2008, experimental evidence of the simulation prediction is given.) Both simulation and experiments also show that the presence of defects different from loops, e.g. single vacancies, as well as the presence of impurities or solute elements, may drastically reduce the effective loop mobility (Puigvi et al., 2004; Cottrell et al., 2004; Terentyev et al., 2005, 2007b; Tapasa et al., 2007; Arakawa et  al., 2007). The mobility of ·100Ò loops is, on the other hand, still debated, but it is certainly much lower than that of ½·111Ò loops (Osetsky et al., 2003; Yao et al., 2008). As will be repeatedly stressed later (e.g. Sections 15.5.1 and 15.6.1), the diffusion properties of SIA clusters have important consequences for the microstructure evolution in metals under irradiation. Voids become clearly visible in ferritic model alloys at temperatures above 350 °C and doses of a few dpa (Porollo et al., 1998; Konobeev et al., 2006), although in ferritic/martensitic steels of similar chromium (Cr) content they remain undetected up to 10 dpa (Gelles, 2004). However, at very high dpa (at least tens of dpa), voids appear in ferritic steels as well (Kohno et al., 1992), when loops evolve into dislocation networks (Katoh et al., 1995; Gelles, 1995). Moreover, when the chromium content exceeds ~7% in these steels, there are strong indications suggesting the radiation-enhanced formation of coherent chromium-rich a¢ precipitates (Mathon et al., 2003, 2005; Gelles, 2004). In austenitic stainless steels (whose crystallographic structure is fcc), the microstructural features observed under conditions relevant to light water reactors change with both temperature and dose (Was and Andersen, 2007,

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15.2 Perfect ½·111Ò loop (left-hand side) and ·100Ò loop (right-hand side) in bcc Fe. Upper panel: atomic-level 3D view (only atoms displaced from the perfect lattice position are visualised): for crystal symmetry reasons, a perfect ½ ·111Ò loop has hexagonal shape, while a perfect ·100Ò loop has square shape. Lower panel: cross-section on crystallographic planes (110) (left) and (100) (right), showing how the regular distribution of atoms on lattice sites is distorted by the presence of a loop (different symbols represent atoms of different parallel atomic planes). The thick, upward arrows represent one Burgers vector (the Burgers circuit is also indicated) and also denote the direction of glide of the loop, if mobile. (Courtesy of D. Terentyev.)

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and references therein). Point defect clusters (‘black dot damage’) form at very low doses. Dislocation loops appear with increasing dose over several dpa, finally evolving into dislocation networks. At high dose and temperature, voids and He bubbles may also form and grow. In more detail, the microstructure is dominated by small clusters and dislocation loops at temperatures below 300 °C. Near 300 °C, the microstructure contains larger loops, as well as, at higher doses, dislocation networks and cavities. At temperatures above 300 °C, the formation of voids, bubbles and a dislocation network is observed even at lower doses, enhanced by higher defect mobility. Overall, the dominant microstructural features in austenitic steels are by far the so-called Frank dislocation loops. These are interstitial-type dislocation loops (such as those mentioned above for ferritic alloys), but they are faulted, i.e. the SIAs are oriented in such a way that they cannot migrate as ½·111Ò loops do in ferritic alloys. In order to migrate, they need to unfault first (Hull and Bacon, 2001). Technically, they are therefore said to be sessile (as opposite to glissile). Figure 15.3 shows how a perfect loop and a Frank (faulted) loop look from an atomic-level perspective. The unfaulting of sessile Frank loops is responsible for the formation of dislocation networks at doses in excess of a few dpa. At that point, voids appear as well, their formation being promoted by the reduction of sink strength that accompanies loop unfaulting. Finally, the so-called stacking-fault tetrahedra (SFT) are a type of defect typical of pure fcc metals (Singh and Zinkle, 1993; Kiritani, 1997, and references therein), which are also observed in austenitic steels under certain conditions (Dai et al., 2001; Schäublin et al., 2005; Li and Almazouzi, 2009). Figure 15.4 shows an atomic-level view of an SFT (other figures and text explaining the formation of SFTs from the atomic-level perspective can be found at http:// iron.nuc.berkeley.edu/~bdwirth/Public/WRG/documents/SFT_form.pdf).

15.2.3 Correlation between microstructure and mechanical property changes The formation of microstructural features (see previous section) changes the macroscopic mechanical properties of the material significantly. Cavities, dislocation loops, dislocation networks and precipitates formed due to irradiation act as additional obstacles to dislocation motion. These obstacles can generally be regarded as distributed points where the dislocation line is pinned, so a higher stress is required to set the dislocations into motion and keep them moving (dispersed barrier model, Bement, 1970). In reality, defects are often found to decorate dislocations (e.g. Miller and Russel, 2007, and references therein) and this fact may introduce a different mechanism to obstruct dislocation motion, as briefly discussed later (Singh, 1998; Singh et al., 2002, and references therein). In either case, however, the macroscopic effect is higher yield strength, which is the technical definition of hardening

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15.3 Perfect ·110Ò loop (lefthand side) and faulted Frank loop (right-hand side) in fcc Cu. Upper panel: atomiclevel 3D view (only atoms displaced from the perfect lattice position are visualised): for crystal symmetry reasons, the perfect ·100Ò loop has square shape, while the Frank loop, lying on a {111} plane, is hexagonal. Lower panel: crosssection on crystallographic planes (1 1 1) (left) and (1 1  2) (right), showing how the regular distribution of atoms on lattice sites is distorted by the presence of a loop (different symbols represent atoms of different parallel atomic planes). The thick, upward arrows represent one Burgers vector. The Burgers circuit is indicated only for the perfect loop (for which it also denotes the direction of glide of the loop, if mobile). For the faulted Frank loop no proper circuit can be built, because of the presence of the stacking fault at the loop, as indicated. (Courtesy of D. Terentyev.)

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15.4 Atomic-level appearance of a stacking fault tetrahedron: the atoms shown are those displaced from their perfect lattice positions (courtesy of D. Terentyev).

(Olander, 1976). Figure 15.5 depicts qualitatively the changes observed in the engineering stress–strain curve for tensile tests in ferritic and austenitic steels under increasing neutron dose, showing, among other things, the hardening process. The difficulty of setting dislocations into motion also affects the ability of the material to resist crack propagation, thereby becoming more brittle. The continuous production of dislocations from the crack tip as the crack propagates is accepted to be the energy dissipation mechanism that causes the work of fracture to exceed that expected from surface creation alone (e.g. Tanaka et al., 2008, and references therein), thereby hindering or delaying crack propagation and allowing ductile fracture. Thus, the presence of obstacles to dislocation movement will limit the efficiency of this mechanism of fracture energy dissipation, rendering the material more brittle. However, ductility or brittleness depend strongly on temperature. Although embrittlement should be measured physically in terms of deformation before fracture (Olander, 1976), technically it is more often the irradiation-induced shift of the measured ductile-brittle transition temperature (DBTT) that defines embrittlement. The difficulty of setting dislocations into motion translates into reduced elongation before fracture and higher DBTT, which may approach the service temperature and thus put the integrity of the structural component at risk. The fact that the fundamental mechanism in hardening and embrittlement is largely the same explains why the two (yield strength increase and DBTT increase) are often phenomenologically related, this being especially true in the case of RPV steels (e.g. Sokolov and Nanstad, 1999).

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15.5 Illustration of radiation-induced changes of mechanical properties in steels: qualitative examples of engineering stress–strain curves obtained from tensile tests after irradiation for increasing neutron dose are given. Left-hand side: ferritic steels; right-hand side: austenitic steels. In the lower panel, a pictorial representation is also provided of how radiation-induced defects pin a dislocation line, thereby impeding its glide.

However, there are cases where embrittlement is not accompanied by hardening and cannot be detected with tensile tests. One example is when radiation-induced segregation of specific chemical elements occurs at grain boundaries, decreasing their cohesion (for example, phosphorus in the case of RPV steels, see, e.g., Nishiyama et al., 2007, and references therein) and promoting intergranular fracture. Crack initiation and intergranular fracture may also be enhanced by void formation, especially when voids are formed close to grain boundaries and stabilised by helium (Klueh and Alexander, 1995; Trinkaus and Singh, 2003; Henry et al., 2003; Klueh et al., 2008). This can happen even at elevated temperatures, when embrittlement is a priori not expected to be an issue (Klueh et al., 2008). The latter is an example of the much-feared synergy between the consequences of atomic displacement and transmutation.

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Once the dislocations are set in motion, the presence of obstacles may or may not have significant consequences for plastic flow behaviour. Two regimes can be distinguished in structural alloys, depending on neutron fluence. At low fluence, the yield strength increases relatively fast with dpa, but this increase is accompanied only by a minor reduction in elongation, without loss of work-hardening. At high fluence, on the other hand, the yield strength increase is large, although it grows more slowly, and dramatic loss of work-hardening is observed (Farrell et al., 2004, and references therein) – see Fig. 15.5 (left-hand side) for a qualitative illustration of this effect. In RPV steels, due to the relatively low fluences usually attained, the only significant effect is an increase in yield strength. The stress–strain curve describing the plastic flow behaviour is thus almost rigidly shifted upward, with only a slight reduction of work-hardening and elongation (Farrell and Byun, 2003; Farrell et  al., 2004), as exemplified by the first curve after irradiation on the left-hand side of the graph in Fig. 15.5. However, at high enough fluence, in both bainitic (RPV) and ferritic/martensitic (high-Cr) steels, work-softening and loss of elongation (plastic instability) is observed, together with radiation-induced hardening (Byun and Farrell, 2004a, 2004b), as exemplified by the higher curves after irradiation on the left-hand side of the graph in Fig. 15.5. The actual fluence at which this loss of elongation appears will depend on temperature, as well as on the type of alloy. As a rule of thumb, above 0.3–0.4Tm (Tm = melting point) radiation hardening and embrittlement cease to be a problem (Singh, 1998), e.g. above 425– 450 °C in ferritic alloys (Klueh and Nelson, 2007). Below this temperature, however, plastic instability is invariably observed at high fluence, generally accompanied by the appearance of channels denuded of defects (clear bands) in the microstructure of the deformed materials (Byun and Farrell, 2004a, 2004b). The detailed mechanism of formation of these channels remains largely unclear and is specific for the type of obstacles removed; it is even debatable whether the plastic instability should always be associated with their appearance (Byun and Farrell, 2004a, 2004b). However, the general idea is that radiation-produced defects lock dislocation sources, so that higher stress is required to activate them (yield strength increase). Once a stress capable of activating them is attained, the sources emit avalanches of dislocations that sweep through the obstacles ahead of them, thereby clearing channels and creating sorts of highways for plastic deformation. So, once the channels are formed, dislocations are locally free to glide, less force is required to maintain the imposed strain rate, and the stress drops (Farrell et al., 2004, and references therein). This can be partially rationalised under the cascade-induced source-hardening model, which is based on the idea of dislocation decoration, as an alternative to the dispersed barrier model mentioned above (Singh, 1998; Singh et al., 2002, and references therein). A similar behaviour pattern is observed in austenitic steels. However,

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in these steels the onset of plastic instability appears at significantly higher fluence, due to the existence of different active modes of deformation (Farrell et al., 2004; Byun and Farrell, 2004a, 2004b). This somewhat less dramatic effect is illustrated on the right-hand side of Fig. 15.5. Above 0.3–0.4 Tm, radiation hardening and embrittlement cease to be a problem, but other phenomena, such as swelling, appear. This corresponds to isotropic increases in the dimensions of the irradiated material and is associated with the presence of voids and enhanced by the presence of gaseous transmutants, i.e. He and H (Olander, 1976). Swelling is the main limiting factor to the long-term use of austenitic steels in internal reactor components (Klueh and Nelson, 2007). Qualitatively, swelling appears after an incubation fluence, the value of which can greatly change depending on steel composition and irradiation conditions, especially dpa rate. After incubation, a material-dependent characteristic swelling rate is observed: 1%/ dpa in austenitic steels and 0.2%/dpa in ferritic/martensitic steels (Garner et al., 2000). Recent evidence suggests that the combined presence of helium and hydrogen may spectacularly increase swelling, even in ferritic steels (Tanaka et al., 2004). These and other macroscopic effects, all of them originating from the sudden production of defects in nanometric displacement cascades, but taking place over times that may reach the order of years, may seriously compromise the ability of a material to maintain its integrity in operation, with observable and measurable consequences at the macroscopic level. Thus, the radiation effects described above and summarised in the schematic diagram given in Fig. 15.6 illustrate that these are inherently a multi-scale problem. It is also sometimes said that radiation effects are a multiphysics problem (Odette et al., 2001). This expression means that different branches of physics (and chemistry), and the corresponding experimental techniques, must be combined in order to fully understand such effects. These branches include: nuclear, atomic, and solid state physics; thermodynamics; diffusion and dislocation theory; elasticity, plasticity and fracture mechanics; and more. A broad theoretical framework to qualitatively understand and explain radiation effects exists and behaviour patterns common to different materials and radiation sources may be found (Seeger, 1958; Stoller et al., 2004). Nonetheless, as illustrated by the above examples, the actual effects expected in a given material, their relative importance and the moment of their appearance during operation are largely specific to the type of material, its history, its texture, its crystallographic structure and its composition, including impurities present only in very low concentration. Environmental conditions, such as type of radiation, flux, fluence, temperature and chemical environment, also play an important role, of course. These facts and the multi-scale and multiphysics nature of radiation effects are what make the quantitative prediction of such effects a challenge. The complexity of the

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15.6 Schematic diagram describing radiation effects as inherently multi-scale phenomena. (Vacs and SIAs stand for vacancies and self-interstitial atoms, respectively; p-def = point defects; dislo’s = dislocations; GBs = grain boundaries; ppts = precipitates; SFTs = stacking fault tetrahedra; elong. = elongation.)

problem, and the many variables involved, inherently call for the support of numerical tools. At the same time, resorting to computers adds the need to master (advanced) numerical algorithms and address computer science problems to the multiple types of expertise required to develop models.

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Multi-scale modelling

15.3.1 What is multi-scale modelling Defining multi-scale modelling is a difficult task that borders on a philosophical exercise. The activity of modelling corresponds first and foremost to understanding the fundamental mechanisms governing the behaviour of a physical system, and to identifying and quantifying the variables and parameters influencing and determining them, so as to be able to translate them into a suitable physical model. The process of identifying and quantifying mechanisms goes through different phases and, accordingly, different modelling approaches can be distinguished. Broadly, the possible approaches can be classified as empirical, mechanistic and physical. A purely empirical approach consists of interpolating laws from an experimental database of macroscopic properties as functions of variables of practical interest that are expected to influence those properties (e.g. increase in yield strength versus neutron fluence). This often allows simplified analytical expressions to be written that describe the correlation, even without any specific knowledge of the fundamental mechanisms involved (Lucas, 1994). If this approach is coupled to a series of microstructural studies, so as to identify the possible origin of a certain macroscopic behaviour, depending on what is observed at the microscopic level, one can talk of a mechanistic approach (Phythian and English, 1993; English et al., 1997). The resulting models, while still oversimplified, allow, for example, a certain contribution to be included or excluded depending on the chemical composition, or other features, of the material. This approach is de facto currently applied in the nuclear industry, and broadly accepted by the regulatory authorities. For example, in the case of RPV steels, empirical correlations explicitly including terms associated with, e.g., copper-rich precipitate formation (including a dependence on Ni content), matrix damage accumulation (voids, loops, etc.) and phosphorus segregation at grain boundaries, have been fitted and are part of the standards used for integrity evaluation (USNRC, 1988, 2007; Eason et al., 1998; ASTM E90002, 2007). In a physical approach, the goal is to explain how the microstructural features observed are produced, and how they influence the mechanical behaviour of the material, based on precise physical mechanisms. Proper physical models, ideally endowed with predictive capability, can only be developed when all the important mechanisms have been identified, together with the variables influencing them and their values. Computer simulations are, in this context, nothing else than virtual experiments, aimed at identifying and quantifying mechanisms. As such, computer simulations stand on an equal footing with real experiments, so long as the latter are also aimed at identifying and quantifying mechanisms (modelling-oriented experiments),

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rather than just providing values of engineering use. This chapter focuses on physical modelling. It will become clear in the following sections that computer experiments (physical models) are used mainly for two purposes: (i) to assess quantities, or study phenomena, that are barely accessible, or totally inaccessible, to real experiments; (ii) to test possible physical mechanisms, in order to verify if they can explain the results of real experiments. The first type of computer experiment includes atomic-level models, such as density functional theory calculations (addressed in Section 15.4.2) and energy minimisation techniques, molecular dynamics (MD) simulations and, partly, Monte Carlo (MC) models (addressed in Section 15.5). The latter are generally based on the use of interatomic potentials (also addressed in Section 15.4.2) to describe interatomic forces. Atomic-level modelling presents problems vis-à-vis the so-called experimental validation. In most cases, atomic-level modelling results cannot be directly validated by experiment, and there are phenomena whose existence and features are mainly or only known through computer simulations (for example displacement cascades). In this case, the implicit standpoint is that the model is not developed to be compared with experiments. Instead, it is based on well-established physics (e.g. quantum or classical mechanics), and used in order to gain insight into phenomena that are known, or supposed to occur in reality, but which cannot be directly observed in experiments. The experimental validation of these models has supposedly been carried out already, because they are based on established physics. Nonetheless, in practice it remains true that direct or indirect ways to validate these models must be found, although performing this exercise is not always straightforward. The second class of models includes mesoscopic models, such as kinetic Monte Carlo, rate theory and dislocation dynamics. Here, the ability of computers to handle a large number of variables is exploited in order to discriminate between mechanisms and parameters that are supposed to be driving the evolution of a system, thereby identifying the important ones, by comparing the results of the models with experimental results. In a subsequent stage, when the important mechanisms and their parameters have been identified, these models can be used to explore situations for which experimental data are lacking, and to derive laws or correlations. If computer simulation techniques capable of modelling radiation effects at all different scales were properly developed, parameterised and coupled, this physical approach might make it possible to describe reliably the process by which radiation damage is produced, evolves and consequently causes materials to degrade, over all relevant time and length scales. It must be emphasised, however, that applying a physical approach exploiting multi-scale simulation techniques is of use also independently, regardless of whether the goal of a fully integrated multi-scale model is attained. A multi-scale,

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computer-based, physical approach provides knowledge on radiation effects which, even without being integrated on a common computer platform, can be used to: (i) interpret experimental results and guide further experimental work; (ii) refine existing empirical correlations (e.g. by introducing terms that explicitly account for a mechanism hitherto neglected); and (iii) support the selection or rejection of a certain material and suggest possible ways to improve the performance of nuclear materials. The multi-scale modelling approach can be defined, summarising in sentence a widespread view, as the ‘construction of a physical model based on the harmonised application of different, dedicated computer simulation tools, each describing phenomena occurring at a definite space and time scale, according to known physical mechanisms’. However, it is actually more than just this, it is a way of thinking and looking at problems that seeks the origin of a certain phenomenon at the proper scale and in the right place. It is hence a powerful tool to be used in combination with other approaches, and with experiments, to assist in the quantitative explanation, and possible prediction, of the behaviour of materials (under irradiation in the specific case of interest here).

15.3.2 Multi-scale modelling techniques for radiation effects Different theoretical frameworks and computer simulation tools exist that can describe the physical mechanisms governing radiation effects and simulate the production and evolution phases of radiation damage, from the atomic to the meso/macroscopic level. The most important of these are briefly reviewed below. Broadly speaking, in multi-scale modelling approaches to radiation effects in metals, ab initio data are transferred to interatomic potentials, which are in turn used for large-scale molecular dynamics and/or Monte Carlo simulations of the production and short-term evolution of radiation damage (Díaz de la Rubia et al., 1990). The information thereby obtained can be transferred to long-term mesoscopic models of microstructural and microchemical evolution, such as those based on kinetic Monte Carlo or rate theory approaches, and then to dislocation dynamics, to describe the plastic behaviour of the material at single-crystal level. The combination of the predictions of microstructural/ microchemical evolution models and dislocation dynamics models is expected to provide the basis to formulate physics-based constitutive laws that are the core of continuum models, thereby building the link to the macroscale. Following this broad scheme (summarised in Fig. 15.7), and taking into account Fig. 15.6, multi-scale modelling techniques are presented here under five headings: nuclear interactions, atomic-level interactions, atomic-level modelling, microstructure evolution modelling and mechanical property

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Ab Initio

Interatomic potentials

Reference values

Energy & Forces

477

Molecular dynamics

Migration energies Atomistic kinetic Monte Carlo

Diffusion parameters Migration & binding energies

Object kinetic Monte Carlo/Rate theory

Nanostructure

Migration energies

Diffusion parameters & Mechanisms

Nanostructure

Crystal plasticity

Physics based constitutive laws

Dislocation dynamics

Constitutive/Homogenisation laws

Dislo mobility & local rules (dislo/defect interaction) Mechanics/Macroscopic continuum models

15.7 Flowchart of the multi-scale modelling approach applied to radiation effects in metals: ab initio data are transferred into interatomic potentials for large-scale molecular dynamics and/ or Monte Carlo simulations of radiation damage production and short-term damage evolution. The information thereby obtained is transferred to long-term microstructural and microchemical evolution mesoscopic models (kinetic Monte Carlo or rate theory), and then to dislocation dynamics, for the description of the material plastic behaviour at single-crystal level. The combination of the predictions of the two latter mesoscopic tools is expected to provide the input to formulate physics-based constitutive laws that are the core of continuum models, thereby building the link to the macroscale. This link, however, in practice has not yet been clearly established.

modelling. This is only one of many possible ways of classifying the techniques, all of them equally valid, but also debatable at the same time. Inevitably, this overview will not be exhaustive and the reader is referred to more extensive and specialised articles, books and theses on the subject for further details. In addition, no attempt is made to review crystal plasticity models, which are often based on using finite element techniques to solve continuum equations. As a matter of fact, multi-scale modelling approaches to crystal plasticity do exist (e.g. McDowell, 2000; Van Houtte et al., 2006; Šiška et al., 2007; Zeman and Šejnoha, 2007). However, the interpretation of the word multi-scale itself is different in plasticity and in the present context. For plasticity applications, the microstructure is related to the size, shape and

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phase distribution of grains. This is also the lowest scale considered in plasticity models. Applying a multi-scale approach in plasticity means, therefore, that the model takes grain information into explicit consideration. The different scales involved in plasticity are thus the single crystal, the aggregate (group of grains), the representative volume element (RVE), and finally the actual component. (The RVE is the smallest statistically representative volume of material containing all microstructural information at the origin of its mechanical behaviour, in terms of heterogeneities, and radiation damage if present.) In radiation damage problems, on the other hand, the fundamental scale is atomic, and what is referred to as microstructure here corresponds to what is referred to as substructure in plasticity, i.e. what is found inside the grains in terms of defects or precipitates, the effects of which are generally not explicitly taken into account in crystal plasticity models. The effort put in developing physical multi-scale models applied to radiation damage has thus far focused on describing the microstructure evolution in a reference volume where grain boundaries act mainly as sinks. The mechanical behaviour modelled in this framework is limited to a single crystal or to a single grain. This fact can be easily deduced by inspecting a number of popular and recent reviews devoted to the field (Odette et al., 2001; Wirth et al., 2004; Becquart, 2005; Nguyen-Manh and Dudarev, 2006; Victoria et  al., 2007; Dudarev et al., 2009). Thus up to now, the multi-scale modelling approach applied to radiation damage has essentially stopped where it begins in crystal plasticity. This would not be a problem if a link between the two approaches was established, and this is of course the declared ultimate goal of all multi-scale modelling programmes on radiation damage. However, fully integrating physical information about radiation damage from lower scale models into the constitutive equations used in crystal plasticity models is yet to be achieved, to the best of the author’s knowledge. The only attempt at doing so is probably the example cited in Section 15.8. Thus, this chapter will stop at the single crystal scale. The only aspect concerning larger scales addressed here will concern how the bridge to crystal plasticity can be built in principle. This point will be succinctly addressed in Section 15.7.3.

15.4

Nuclear- and atomic-level interactions

15.4.1 Nuclear-level interactions As mentioned, radiation effects originate when neutrons interact with the nuclei of the atoms composing a material, resulting in activation, transmutation and atomic displacements. Part of the modelling effort has therefore been devoted to predicting which nuclear reactions will take place, which new nuclides will be formed (activation and transmutation), at which rate, and which recoil energy spectrum (number of PKAs per unit volume and per

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unit time of a certain energy) is to be expected. This is done for a material of given composition, subjected to a given flux of neutrons, distributed in energy according to a known spectrum, up to a certain fluence. The results concerning activation are particularly useful for assessing the radiological risks associated with the selection of a certain material, and have been used, for example, to guide the selection of low and reduced activation materials (Conn et  al., 1984). An estimate of the PKA spectrum, on the other hand, is needed to assess exposure parameters such as displacements per atom, customarily used to characterise the effect of irradiation in terms of damage. The main advantage of the dpa as exposure parameter is that it allows a comparison of results obtained under different neutron spectra and even different radiation sources (Greenwood, 1994; ASTM E693-01, 2007). In addition, the PKA spectrum, the dose rate (in dpa/s) and the total dose (dpa) are part of the information needed for parameterising microstructure evolution models (Section 15.6), to assign correctly the so-called ‘sourceterm’ (number of defects produced per unit time), and to know when to stop the simulation (when the total number of dpa is reached). Given a particular neutron spectrum, the activation and transmutation rates are calculated by integrating the appropriate neutron cross-sections (Robinson, 1994; Greenwood, 1994). The value of these must be given, and they are available in the form of Evaluated Nuclear Data Files (ENDF, National Nuclear Data Centre, Brookhaven National Laboratory). The reliability of the calculation will depend largely on the accuracy of these cross-sections, which has improved over the years. The calculation of PKA spectra in pure elements is traditionally performed using widespread codes such as SPECTER (Greenwood and Smither, 1985) or NJOY (Macfarlane et  al., 1984). For compounds and alloys, the calculations may be more involved. The SPECOMP package has been developed to address this problem (Greenwood, 1989). The results of neutronics calculations are relatively easy to validate by direct measurements, and code benchmarking exercises have been conducted in the past, (e.g. OECD/NEA, 2000). In addition, neutron spectra can be both accurately calculated and measured (Greenwood, 1994). Thus, incorporating nuclear interactions in a multi-scale modelling framework does not pose especially harsh problems, and this part of the calculation is certainly not the weakest link.

15.4.2 Atomic-level interactions The necessary prerequisite for atomic-level modelling is a correct description of the interaction between the atoms composing the material of interest. Moreover, in order to study radiation effects, particular types of atomic-level information must be known and transferred to larger scale models. Examples

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include the interaction energies between atomic species and defects in the material. In this section, the methods generally used to obtain reliably this type of information in radiation effect modelling, namely ab initio (or more precisely density functional theory, DFT) calculations, and the development of semi-empirical, many-body interatomic potentials (henceforth simply potentials), are briefly described. DFT calculations Ideally, the goal of ab initio calculations should be to obtain the value of a physical quantity by solving fundamental equations of physics, where no empirically fitted parameters (except universal constants) appear. The closest one can get to achieving this goal in the case of a material is by solving the Schrödinger equation for the system of interest, as composed of electrons and nuclei, where the only input parameter is the atomic number of the chemical elements present. In practice, however, this problem is generally unsolvable and must be simplified. The most effective technique, nowadays routinely used for ab initio calculations in materials science, is the application of density functional theory. DFT is an exact one-body reformulation of the many-body quantum mechanical problem governed by the Schrödinger equation, which can be used effectively (after introducing a number of approximations) to determine the ground-state energy of a system of interacting particles. An especially clear and succinct description of the theory and approximations made, including all key references, can be found, for example, in Martin (2004) or Hasnip (2005). Due to the existence of approximations, DFT techniques are not, strictly speaking, ab initio. Nonetheless, they are currently the most reliable tool at our disposal for describing the interaction between atoms, without a priori any restriction on the number and type of chemical elements that can be treated. The main problem is that their application remains computationally very demanding, since the algorithms scale as the cube of the number of atoms, so in practice only systems of about 1000 atoms can be studied effectively (for metals this is an upper limit, while insulators have been simulated using several thousand atoms). In addition, although methods for performing dynamic simulations do exist (Car and Parrinello, 1985; Payne et al., 1992), their application is extremely expensive in terms of computer time and power. Thus, in most cases DFT calculations are performed in a static manner, i.e. the atomic positions may be relaxed to accommodate the strain due to defects or impurities (using energy minimisation techniques, such as conjugate gradient methods or others, see Press et al., 1992), but, aside from this, they do not move during the simulation. Even so, the quantity of information that can be extracted from DFT calculations is enormous. Defects in pure elements or interacting with

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impurities and solute atoms can be studied in terms of characteristic energies (formation energies, binding energies, migration energies) and even migration mechanisms can be explored (see Malerba et al., 2010a, for a review of examples). These quantities are either extremely difficult or totally impossible to measure experimentally. By just knowing these energies for all the important elements of an alloy it is often possible to build directly a qualitative picture to interpret experimental results, not only in the case of modelling-oriented experiments on model alloys, but also of design-oriented experiments on materials of technological interest (Malerba et  al., 2008; Lambrecht et al., 2008; Van den Bosch, 2008). The main limitation of DFT remains the fact that, due to the restricted size of the system, only clusters of a few point defects can be studied, and either complicated solutions or oversimplified geometries must be adopted when studying extended defects, such as dislocations and grain boundaries (Ventelon and Willaime, 2007; Wachowicz and Kiejna, 2008, and references therein). In addition, difficulties arise when handling concentrated random alloys, especially with more than two chemical elements, because with a limited amount of atoms it is difficult to reproduce a realistic random distribution of species (Klaver et al., 2007). Interatomic potentials Another very important use for the results of DFT calculations is to fit potentials. Potentials are mathematical functions of the relative positions of atoms (distances and, in some cases, angles as well) that describe, after properly fitting the parameters that appear in them, the acting force field, i.e. the potential energy landscape of the system. The use of potentials still is the only practical way to simulate the dynamic behaviour and evolution of systems containing a large number (millions) of atoms with some degree of realism, by means of molecular dynamics and Monte Carlo tools. The main limitation is that reliable potentials are already relatively difficult to produce for pure elements and, to date, at most only potentials for binary and a few ternary alloys have been published, to the best of the author’s knowledge. Thus, only model alloys can be studied by means of interatomic potentials. By simply deriving the potential (U) with respect to the atomic positions, the force (F) acting on each atom can be obtained (F = –—U). In principle, if each of the N atoms interacts with all others (and if angles are not considered), N2 interactions must be calculated. In practice, this problem can often be avoided, certainly in the case of metals, as atoms can generally be assumed to interact only with close neighbours, within a cut-off distance, and neighbour lists can be built to avoid continuously checking the neighbours of a given atom. Therefore, the computing time required for force calculations scales

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generally only as N, which makes an enormous difference compared to DFT calculations. Knowing the forces, the Newtonian equations of motion can be written and solved for all the atoms, so their evolution starting from given initial conditions can be predicted. At the same time, given a potential, characteristic energies can be statically calculated by means of energy minimisation techniques (Press et al., 1992), in exactly the same way as with ab initio calculations, but with much less stringent limitations as to the size of the system. This enables the dynamic study of large point-defect clusters and also extended defects, such as dislocations and grain boundaries. Potentials must fulfil two main requirements. Firstly, they must be able to predict, in an acceptably correct way, the largest amount of physical properties possible for the material of interest. As a matter of fact, the reliability of the potential decides whether it is the behaviour of Fe or Cu being simulated, rather than that of a virtual, non-existent element. Secondly, they must require the minimum computing time possible, while fulfilling the previous condition. Many differently classified mathematical formalisms exist in the literature, each of them often more suitable for a specific class of materials than for others (Carlsson, 1990; Robinson, 1993; Adams et al., 1994). These formalisms are generally derived either as approximations of first principles expressions, or based on heuristic considerations of some physical significance. In the case of metals, many-body potentials of the embedded atom method (EAM) type have been the state-of-the-art for the last 25 years and are still widely used. (Here, EAM-type is used as a generic denomination covering the mathematically equivalent ‘glue’ model (Ercolessi et  al., 1986), the Finnis–Sinclair model, or second moment tight-binding approximation (Finnis and Sinclair, 1984), the effective medium method (Jacobsen et  al., 1987), and of course the proper EAM formalisms (Daw and Baskes, 1983, 1984).) EAM-type potentials offer the advantage of being computationally cheap, while providing high flexibility and often surprisingly good results. Their overall performance has lately been boosted by the availability of large ab initio databases, used as a reference for fitting. This has allowed the recent development of extremely accurate pure element potentials (for Fe see Malerba et al., 2010b; for other bcc metals see Derlet et al., 2007; for Al and Ni, see Mishin et al., 1999; for Cu see Mishin et al., 2001). Moreover, DFT-based potentials capable of describing concentration-dependent physical and thermodynamic properties of binary and even ternary alloys reasonably well have started to appear, based on modifications or extensions of the EAM scheme (Olsson et al., 2005; Caro et al., 2005; Malerba et al., 2008, 2010a, and references therein; Bonny et al., 2009a, 2009b, 2009d, 2009e). The challenge for further improving the performance of potentials is to devise new formalisms, with deeper physical foundations (especially for alloys and magnetic materials), offering better accuracy and reliability, with

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similar or slightly lower computational efficiency to EAMs (Ackland, 2006; Nguyen-Manh et al., 2007; Ma et al., 2008).

15.5

Atomic-level modelling

This heading includes all modelling techniques aimed at describing the evolution of a material taking into account explicitly the existence of atoms. In practice, there are only two families of techniques of this type: molecular dynamics (MD) and Monte Carlo (MC). The former is deterministic (at least in principle) and dynamic; the latter is probabilistic and quasi-static. Classical textbooks exist on these techniques, and the interested reader is referred to them for further information (Allen and Tildesley, 1987; Frenkel and Smit, 2001). Both techniques are often presented as numerical statistical-mechanics tools used to sample possible microstates within a certain macrostate of equilibrium, in order to calculate the expectation value of some thermodynamic quantity (e.g. internal energy, enthalpy, etc.). This is done by averaging over the sampled microstates, weighted with their respective Boltzmann probabilities (exp(-E/kBT), where E is the energy of the state, kB Boltzmann’s constant and T the absolute temperature), thereby spontaneously including any entropic effect at the given temperature. However, in radiation effect studies, both techniques are mainly used to study phenomena out of equilibrium and to determine the evolution of the system towards a dynamic steady state (if any), which may in fact be far from thermodynamic equilibrium. In addition, the MC techniques in this field are generally considered as a ‘prolongation’ of MD in time and length scales.

15.5.1 Molecular dynamics Molecular dynamics is a numerical method to trace the time evolution of a physical system by solving the classical Newtonian equations of motion of the set of N interacting atoms composing it (Fi = mai; i = 1,…, N; ai = d2ri/ dt2), given the interaction potential (on page 481) and starting from assigned initial conditions. If needed, convenient constraints are used to control thermodynamic variables, such as temperature and pressure. Molecular dynamics (MD) is a powerful, highly flexible technique used very widely to study innumerable physical problems in current materials science. It is irreplaceable whenever the knowledge of atomic-level detail is required in order to understand the phenomena of interest. An especially clear and succinct description of the technique can be found in Ercolessi (1997); the technique is explained in full detail, including examples of applications, in classical books such as the already cited Allen and Tildesley (1987), and Frenkel and Smit (2001). In contrast to Monte Carlo methods, MD is a deterministic technique:

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given an initial set of positions and velocities, the subsequent evolution is, in principle, completely determined (in practice, round-off numerical errors lead to a loss of memory of the initial conditions, but this is not a shortcoming for statistical mechanics studies). However, in order for the equations of motion to be solved in a convergent way, using finite difference methods, the time elapsed between the initial set of positions and velocities and the subsequent one (time step), has to be much smaller than the typical period of oscillation of atoms in condensed matter. This means an order of magnitude of one femtosecond (10–15 s). Hence, after one million time steps in a standard MD run, a simulated timespan of ‘only’ one nanosecond (10–9 s) has been covered. In most cases, the limitation on the timespan that can be simulated is therefore the main shortcoming of MD, because many phenomena of practical interest need times much longer than nanoseconds to develop. Nevertheless, for studies at the atomic scale, MD is a tool of tremendous effectiveness. What makes MD superior to other numerical techniques is its inherent capability of dealing with systems too complex to be modelled using any analytical approach without the need of simplifying hypotheses or approximations (except in the cohesion model – interatomic potential – used to determine the interatomic forces). Using MD as a statistical mechanics tool, it is possible to calculate all the usual variables that characterise thermodynamically a physical system, starting from their statistical mechanics definition (the easiest examples are temperature and pressure). It is possible to assess equilibrium properties that are potentially experimentally verifiable (e.g. pair correlation functions, velocity correlation functions, diffusion coefficients of chemical species), but it is also possible to determine properties that are experimentally inaccessible, or extremely difficult to measure, with high accuracy (e.g. diffusion coefficients of point-defect clusters). More importantly, MD allows atomic-level mechanisms to be identified and quantified, so that it is irreplaceable for radiation effect studies. In addition, it naturally allows the study of both equilibrium and non-equilibrium conditions, thereby embracing systems in a stable phase or in a phase transition, ordered and disordered, in the presence of complex defects, and so on. The only real limit to its capabilities is the computational cost of the simulation. With current computers the evolution of systems of up to 107 atoms, for times up to tens of nanoseconds, is accessible. The fact that MD algorithms are easily parallelised is helpful here. Larger sizes or longer times can be considered, but not simultaneously, i.e. the extent of the method is a trade-off between size and time. The physical reliability of the results then relies totally on the accuracy and acceptability of the interatomic potentials used. In radiation effect studies, the main applications of MD can be summarised as follows:

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∑

Simulation of displacement cascades, for which MD is the technique par excellence (English et al., 1992b, and references therein; Averback and Díaz de la Rubia, 1998, and references therein; Bacon et  al., 2000; Crocombette, 2005; Malerba, 2006). For example, the snapshots of displacement cascade phases in Fig. 15.1 have been produced by MD. ∑ Simulation of point-defects and point-defect clusters, in order to establish their stable configurations (formation energies, binding energies, etc.), how they interact with each other (reaction mechanisms), and how they diffuse (migration energies and mechanisms), also in interplay with alloying elements (e.g. Marian et al., 2001, 2002; Osetsky et al., 2003, and references therein; Puigvi et al., 2004, and references therein; Terentyev et al., 2005, 2007a, 2007b, 2008a; Kulikov et al., 2006). ∑ Simulation of extended defects, such as dislocations and grain boundaries, and their interaction with point-defects, point-defect clusters and different alloying chemical species (Bacon and Osetsky, 2005, and references therein; Domain and Monnet, 2005; Gao et al., 2009). MD is the most natural, reliable and complete technique for studying displacement cascades. Since these phenomena cannot be directly observed in experiments, all the available information about cascades comes essentially from MD simulations. The literature on the subject is vast. For early and recent reviews, the reader can consult English et  al. (1992b), Averback and Díaz de la Rubia (1998), Bacon et al. (2000), Crocombette (2005) and Malerba (2006). An important result of MD simulations of cascades is the now wellestablished fact that the number of atomic displacements per cascade, n, calculated using the Norgett–Robinson–Torrens (NRT) formula (Norgett et al., 1975):

n = 0.8

ED 2Ed

15.1

is actually an overestimation. As a consequence of intra-cascade recombination, the number of surviving defects at the end of the cascade process is about one third smaller, depending on the actual material, cascade energy and te