Corrosion Issues in Light Water Reactors: Stress Corrosion Cracking (EFC51)

i Corrosion issues in light water reactors ii iii European Federation of Corrosion Publications NUMBER 51 Corros...

Author: D Fron

146 downloads 1832 Views 8MB Size Report

This content was uploaded by our users and we assume good faith they have the permission to share this book. If you own the copyright to this book and it is wrongfully on our website, we offer a simple DMCA procedure to remove your content from our site. Start by pressing the button below!
Report copyright / DMCA form

DOWNLOAD PDF

i

Corrosion issues in light water reactors

ii

iii

European Federation of Corrosion Publications NUMBER 51

Corrosion issues in light water reactors Stress corrosion cracking Edited by D. Féron and J.-M. Olive

Published for the European Federation of Corrosion by Woodhead Publishing and Maney Publishing on behalf of The Institute of Materials, Minerals & Mining

CRC Press Boca Raton Boston New York Washington, DC

WOODHEAD

PUBLISHING LIMITED

Cambridge England

iv Woodhead Publishing Limited and Maney Publishing Limited on behalf of The Institute of Materials, Minerals & Mining Woodhead Publishing Limited, Abington Hall, Abington Cambridge CB21 6AH, England www.woodheadpublishing.com Published in North America by CRC Press LLC, 6000 Broken Sound Parkway, NW, Suite 300, Boca Raton, FL 33487, USA First published 2007 by Woodhead Publishing Limited and CRC Press LLC © 2007, Institute of Materials, Minerals & Mining 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 publishers cannot assume responsibility for the validity of all materials. Neither the authors nor the publishers, 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. Library of Congress Cataloging in Publication Data A catalog record for this book is available from the Library of Congress. Woodhead Publishing ISBN-13: 978-1-84569-242-1 (book) Woodhead Publishing ISBN-13: 978-1-84569-346-6 (e-book) CRC Press ISBN-1: 978-1-4200-6001-0 CRC Press order number: WP6001 ISSN 1354-5116 The publishers’ 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 elementary chlorine-free practices. Furthermore, the publishers ensure that the text paper and cover board used have met acceptable environmental accreditation standards. Typeset by Replika Press Pvt Ltd, India Printed by TJ International Limited, Padstow, Cornwall, England

Contents

Contributor contact details

xiii

Series introduction

xix

Volumes in the EFC series

xxi

Preface PART I 1

xxvii

Overviews An overview of materials degradation by stress corrosion in PWRs

3

P. M. SCOTT, Framatome ANP, France

1.1 1.2 1.3 1.4 1.5 1.6 1.7 2

Introduction Nickel base alloys in PWR primary water Nickel base alloys on the secondary side of PWR steam generators Stainless steels in PWR primary circuits Low alloy steels Concluding remarks References

3 5 11 15 19 21 22

Corrosion potential monitoring in nuclear power environments

25

A. MOLANDER, Studsvik Nuclear AB, Sweden

2.1 2.2 2.3 2.4 2.5 2.6 2.7

Introduction Measurements in BWRs PWR primary system PWR secondary systems Summary and conclusions Acknowledgements References

25 26 34 36 41 42 43 v

vi

Contents

3

Kinetics of passivation of a nickel-base alloy in high temperature water

44

A. MACHET, A. GALTAYRIES and P. MARCUS, Laboratoire de Physico-Chimie des Surfaces, France and P. JOLIVET, M. FOUCAULT, P. COMBRADE and P. SCOTT, Framatome ANP, France

3.1 3.2 3.3 3.4 3.5 3.6 Part II 4

Introduction Experimental procedure Results Discussion Conclusion References

44 44 46 53 54 55

Stress corrosion cracking: susceptibility and initiation IASCC susceptibility under BWR conditions of welded 304 and 347 stainless steels

59

M.L. CASTAÑO, CIEMAT, Spain, B. VAN DER SCHAAF, NRG, Holland, A. ROTH, Framatome ANP, Germany, C. OHMS, JRC-IE, Holland, D. GAVILLET, PSI, Switzerland and S. VAN DYCK, SCK·CEN, Belgium

4.1 4.2 4.3 4.4 4.5

Introduction Experimental procedure Results and discussion Conclusions References

59 60 63 68 69

5

The effect of lead on resistance of low alloy steel to SCC in high temperature water environments

70

K. MATOCHA and G. ROžNOVSKÁ, VÍTKOVICE, Czech Republic and V. HANUS, NPP Czech Republic

5.1 5.2 5.3 5.4 5.5 5.6 5.7

Introduction Testing material Experimental procedure Results and discussion Conclusions Acknowledgement References

70 70 71 71 74 75 75

6

Effect of cold work hardening on stress corrosion cracking of stainless steels in primary water of pressurized water reactors

76

O. RAQUET and E. HERMS, CEA/Saclay, France and F. VAILLANT, T. COUVANT and J. M. BOURSIER, EDF/Les Renardières, France

6.1

Introduction

76

Contents

vii

6.2 6.3 6.4 6.5 6.6

Experimental procedure Results and discussion Conclusions Acknowledgement References

77 81 85 85 85

7

Effect of strain-path on stress corrosion cracking of AISI 304L stainless steel in PWR primary environment at 360 ∞C

87

T. COUVANT, F. VAILLANT and J.M. BOURSIER, EDF R&D - MMC, France and D. DELAFOSSE, Ecole des Mines de St-Etienne, France

7.1 7.2 7.3 7.4 7.5 7.6

Introduction Experimental procedure Results Discussion Conclusions References

87 88 91 98 101 101

8

Dynamic strain ageing of deformed nitrogen-alloyed AISI 316 stainless steels

103

U. EHRNSTÉN and A. TOIVONEN, VTT Technical Research Centre of Finland, Finland and M. IVANCHENKO, V. NEVDACHA, Y. YAGOZINSKYY and H. HÄNNINEN, Helsinki University of Technology, Finland

8.1 8.2 8.3 8.4 8.5 8.6 8.7

Introduction Experimental procedure Results Discussion of results Conclusions Acknowledgements References

103 104 106 114 117 117 117

9

Laboratory results of stress corrosion cracking of steam generator tubes in a ‘complex’ environment – an update

119

O. HORNER, E.-M. PAVAGEAU and F. VAILLANT, EDF R&D, France and O. DE BOUVIER, EDF Nuclear Engineering Division, France

9.1 9.2 9.3 9.4 9.5 9.6

Introduction Experimental procedure Results Discussion Conclusions References

119 120 120 127 128 129

viii

Contents

10

The effect of sulphate and chloride transients on the environmentally-assisted cracking behaviour of low-alloy RPV steels under simulated BWR conditions

130

S. RITTER and H.P. SEIFERT, Paul Scherrer Institute (PSI), Switzerland

10.1 10.2 10.3 10.4 10.5 10.6

Introduction Experimental procedure Results and discussion Summary and conclusions Acknowledgements References

130 132 136 146 147 147

11

Transgranular stress-corrosion cracking in austenitic stainless steels at high temperatures

149

A. BROZOVA, Nuclear Research Institute, Czech Republic and S LYNCH, Monash University, Australia

11.1 11.2 11.3 11.4 11.5 Part III 12

Introduction Experimental procedure Results Discussion References

149 151 152 153 160

Stress corrosion cracking: propagation Crack growth behaviour of low-alloy steels for pressure boundary components under transient light water reactor operating conditions – CASTOC, Part 1: BWR/NWC conditions

165

S. RITTER and H.P. SEIFERT, Paul Scherrer Institute (PSI), Switzerland, B. DEVRIENT and A. ROTH, Framatome ANP GmbH, Germany, U. EHRNSTÉN, VTT Industrial Systems, Finland, M. ERNESTOVÁ and M. ŽAMBOCH, Nuclear Research Institute (NRI), Czech Republic, J. FÖHL and T. WEISSENBERG, Staatliche Materialprüfungsanstalt (MPA), Germany and D. GOMÉZ-BRICEÑO and J. LAPEÑA, Centro de Investigaciones Energéticas Medioambientales y Tecnológicas (CIEMAT), Spain

12.1 12.2 12.3 12.4 12.5 12.6

Introduction Experimental procedures Results and discussion Summary and conclusions Acknowledgements References

165 166 169 183 184 184

Contents

13

Crack growth behaviour of low-alloy steels for pressure boundary components under transient light water reactor operating conditions – CASTOC, Part 2: VVER conditions

ix

186

M. ERNESTOVÁ and M. ŽAMBOCH, Nuclear Research Institute (NRI), Czech Republic, B. DEVRIENT and A. ROTH, Framatome ANP GmbH, Germany, U. EHRNSTÉN, VTT Industrial Systems, Finland, J. FÖHL and T. WEISSENBERG, Staatliche Materialprüfungsanstalt (MPA), Germany, D. GOMÉZ-BRICEÑO and J. LAPEÑA, Centro de Investigaciones Energéticas Medioambientales y Tecnológicas (CIEMAT), Spain and S. RITTER and H.P. SEIFERT, Paul Scherrer Institute (PSI), Switzerland

13.1 13.2 13.3 13.4 13.5 13.6


14

Effect of yield strength on stress corrosion crack propagation under PWR and BWR environments of hardened stainless steels M.L. CASTAÑO, M.S. GARCÍA, G. CIEMAT, Spain

186 187 190 196 198 198

DE

200

DIEGO and D. GOMÉZ-BRICEÑO,

14.1 14.2 14.3 14.4 14.5


200 201 203 208 209

15

Corrosion fatigue crack growth behaviour of low-alloy RPV steels at different temperatures and loading frequencies under BWR/NWC environment

211

S. RITTER and H.P. SEIFERT, Paul Scherrer Institute (PSI), Switzerland

15.1 15.2 15.3 15.4 15.5 15.6


211 212 217 228 229 229

x

Contents

16

Effect of cyclic loadings on the stress corrosion crack growth rate in alloy 600 in PWR primary water

231

C. GUERRE, O. RAQUET and L. DUISABEAU, CEA, France and G. TURLUER, IRSN, France

16.1 16.2 16.3 16.4 16.5 16.6 16.7 16.8

Introduction Materials and specimen Experimental procedure Results Discussion Conclusions Acknowledgments References

231 231 234 236 241 244 244 244

17

Pattern recognition model to estimate intergranular stress corrosion cracking (IGSCC) at crevices and pit sites of 304 SS in BWRs environments

245

M. URQUIDI-MACDONALD, Penn State University, USA

17.1 17.2 17.3 17.4 17.5 17.6 17.7 17.8 17.9 17.10 17.11 17.12 17.13

Introduction Objective and procedure Effect of pH Effect of fluid velocity Effect of electrochemical corrosion potential (ECP) Effect of conductivity Effect of sensitization (EPR) Effect of stress intensity Data collection Non-deterministic approach: ANN Results Conclusions References

245 246 246 247 247 248 248 249 250 250 252 258 258

18

Fatigue crack growth in austenitic steel AISI 304L in PWR primary water at room temperature and elevated temperature

260

I. NEDBAL, J. KUNZ and J. SIEGL Czech Technical University, Czech Republic

18.1 18.2 18.3 18.4 18.5

Introduction Fatigue experiments Macroscopic crack growth rate Fractographic analysis Conclusions

260 260 261 263 267

Contents

18.6 18.7

Acknowledgements References

xi

268 268

Part IV Practical experience 19

Corrosion damage to 18Cr-9Ni-Ti steel after 25 years of operation in steam-water environments of the VK-50 reactor

273

G.V. FILYAKIN, V.K. SHAMARDIN, YU.D. GONCHARENKO and V.A. KAZAKOV, FSUE ‘SSC RIAR’, Russia

19.1 19.2 19.3 19.4 19.5 19.6

Introduction Material – operation conditions Experimental results Discussion Conclusions References

273 274 275 281 287 288

20

Comprehensive investigation of the corrosion state of the heat exchanger tubes of steam generators

289

K. VARGA, Z. NÉMETH, A. SZABÓ, K. RADÓ, D. ORAVETZ and K. É. MAKÓ, University of Veszprém, Hungary, Z. HOMONNAY, E. KUZMANN and S. STICHLEUTNER, Eötvös Loránd University, Hungary and P. TILKY, J. SCHUNK and G. PATEK, Paks Nuclear Power Plant Ltd., Hungary

20.1 20.2 20.3 20.4 20.5 20.6

Introduction Experimental procedure Results and discussion Conclusions Acknowledgements References

289 290 292 300 304 304

21

Stress corrosion cracking of a Kori 1 retired steam generator tube

306

H. P. KIM, S. S. HWANG, D. J. KIM, J. S. KIM, Y. S. LIM and M. K. JOUNG, Korea Atomic Energy Research Institute, Korea

21.1 21.2 21.3 21.4 21.5 21.6

Introduction Experimental method Results and discussion Summary Acknowledgement References

306 306 307 314 314 314

xii

Contents

22

A systematic study of the corrosion effects of the FRAMATOME CORD-UV technology

316

K. RADÓ, K. VARGA, Z. NÉMETH, I. VARGA, J. SOMLAI, D. ORAVETZ and K. É. MAKó, University of Veszprém, Hungary, Z. HOMONNAY and E. KUZMANN, Eötvös Loránd University, Hungary, J. BORSZÉKI and P. HALMOS University of Veszprém, Hungary and P. TILKY and J. SCHUNK, Paks Nuclear Power Plant Ltd., Hungary

22.1 22.2 22.3 22.4 22.5


316 318 319 326 327

Index

328


(* = main contact)

Editors

Chapter 2

D. Féron CEA-Saclay DPC/SCCME, bât. 458, P.C. 50 91191 Gif-sur-Yvette Cedex France

A. Molander Studsvik Nuclear AB SE-611 82 Nyköping Sweden E-mail: [email protected]

E-mail: [email protected]

J.-M. Olive HYDROGENIUS-AIST-Kyushu University 744 Moto-oka, Nishi-ku 819-0395 Fukuoka Japan E-mail: [email protected]

Chapter 1 P. M. Scott Framatome ANP Tour AREVA 92084 Paris La Défense Cedex France E-mail: [email protected]

Chapter 3 A. Machet, P. Jolivet and P. Scott Framatome ANP Tour AREVA F-92084 Paris-la-Défense France A. Galtayries* and P. Marcus Laboratoire de Physico-Chimie des Surfaces Ecole Nationale Supérieure de Chimie de Paris 11 rue P. et M. Curie F-75005 Paris France E-mail: [email protected]

M. Foucault and P. Combrade Framatome ANP Centre Technique F-71205 Le Creusot France xiii

xiv


Chapter 4

Chapter 7

M.L. Castaño*, B. van der Schaaf, A. Roth, C. Ohms, D. Gavillet and S. van Dyck CIEMAT Complutense 22 28040, Madrid Spain

T. Couvant*, F. Vaillant and J. M. Boursier EDF R&D - MMC Avenue des Renardières - Ecuelles 77818 Moret-sur-Loing Cedex France


Chapter 5 Karel Matocha*, Václav Hanus and Gabriela Rožnovská VÍTKOVICE – Research & Development, Ltd V. Hanus NPP Temelin Czech Republic E-mail: [email protected]

Chapter 6 O. Raquet and E. Herms CEA/Saclay DEN/DPC – 91191 Gif sur Yvette Cedex France E-mail: [email protected] [email protected]

T. Couvant*, F. Vaillant and J. M. Boursier EDF R&D - MMC Avenue des Renardières - Ecuelles 77818 Moret-sur-Loing Cedex France E-mail: [email protected] [email protected] [email protected]

E-mail: [email protected] [email protected] [email protected]

D. Delafosse Ecole des Mines de St-Etienne 157 Cours Fauriel 42023 St-Etienne cedex 2 France

Chapter 8 U. Ehrnstén* and A. Toivonen VTT Technical Research Centre of Finland Industrial Systems Kemistintie 3 P.O. Box 1704 FIN-02044 VTT Finland E-mail: [email protected]

M. Ivanchenko, V. Nevdacha, Y. Yagozinskyy and H. Hänninen Helsinki University of Technology Department of Mechanical Engineering Puumiehenkuja 3 P.O. Box 4200 FIN-02015 HUT Finland


Chapter 9 O. Horner*, E-M. Pavageau, F. Vaillant, EDF R&D - MMC Avenue des Renardières - Ecuelles 77818 Moret-sur-Loing Cedex France E-mail: [email protected] [email protected] [email protected]

O. de Bouvier EDF Nuclear Engineering Division Centre d’Expertise et d’Inspection dans les Domaines de la Réalisation et de l’Exploitation 93206 Saint Denis France E-mail: [email protected]

Chapter 10 S. Ritter* and H.P. Seifert Paul Scherrer Institute (PSI) Nuclear Energy and Safety Research Department Laboratory for Materials Behaviour CH-5232 Villigen PSI Switzerland E-mail: [email protected]

xv

S. Lynch School of Physics and Materials Engineering Monash University Victoria 3800 Australia


B. Devrient and A. Roth Framatome ANP GmbH Erlangen Germany U. Ehrnstén VTT Technical Research Centre of Finland Industrial Systems Kemistintie 3 P.O. Box 1704 FIN-02044 VTT Finland

Chapter 11


A. Brozova* Nuclear Research Institute Rez, plc. 25068 Rez Czech Republic

M. Ernestová and M. Žamboch Nuclear Research Institute Rez, plc. 25068 Rez Czech Republic

E-mail: [email protected] [email protected]


xvi


J. Föhl and T. Weissenberg Staatliche Materialprüfungsanstalt (MPA) Stuttgart Germany D. Goméz-Briceño and J. Lapeña Centro de Investigaciones Energéticas Medioambientales y Tecnológicas (CIEMAT) Madrid Spain

Chapter 13 M. Ernestová* and M. Žamboch Nuclear Research Institute Rez, plc. 25068 Rez Czech Republic E-mail: [email protected]

B. Devrient and A. Roth Framatome ANP GmbH Erlangen Germany U. Ehrnstén VTT Technical Research Centre of Finland Industrial Systems Kemistintie 3 P.O. Box 1704 FIN-02044 VTT Finland E-mail: [email protected]

J. Föhl and T. Weissenberg Staatliche Materialprüfungsanstalt (MPA) Stuttgart Germany

S. Ritter* and H.P. Seifert Paul Scherrer Institute (PSI) Nuclear Energy and Safety Research Department Laboratory for Materials Behaviour CH-5232 Villigen PSI Switzerland E-mail: [email protected]

D. Goméz-Briceño and J. Lapeña CIEMAT Nuclear Fission Department Structural Materials Project Avda. Complutense 22 28040 Madrid Spain

Chapter 14 M.L. Castaño, M. S. García*, G. de Diego, D. Goméz-Briceño CIEMAT Nuclear Fission Department Structural Materials Project Avda. Complutense 22 28040 Madrid Spain E-mail: [email protected]



Chapter 16

Chapter 18

C. Guerre* and O. Raquet CEA DEN/DPC/SCCME/LECA bât.458 91191 Gif-sur-Yvette Cedex France

I. Nedbal*, J. Kunz and J. Siegl CVUT - FJFI - KMAT Trojanova 13 PRAHA 2 CZ 120 00 Czech Republic

E-mail: [email protected] [email protected]

E-mail: [email protected] [email protected] [email protected] [email protected]

L. Duisabeau CEA DEN/DMN/SEMI/LCMI bât.625 91191 Gif-sur-Yvette Cedex France E-mail: [email protected]

G. Turluer IRSN DSR/SAMS BP17 92262 Fontenay-aux-roses Cedex France E-mail: [email protected]

Chapter 17 M. Urquidi-Macdonald Penn State University 203 Earth-Engineering Science Building University Park PA 16801 USA E-mail: [email protected]

xvii

Chapter 19 G. V. Filyakin, V. K. Shamardin*, Y. D. Goncharenko and V. A. Kazakov FSUE ‘SSC RIAR’ Dimitrovgrad -10 Ulyanovsk region 433510 Russia. E-mail: [email protected]

Chapter 20 K. Varga*, Z. Németh, A. Szabó and K. Radó University of Veszprém Department of Radiochemistry H-8201 Veszprém P.O. Box 158 Hungary E-mail: [email protected] [email protected]

D. Oravetz and K. É. Makó University of Veszprém Department of Silicate Chemistry and Materials Engineering H-8201 Veszprém P.O. Box 158 Hungary

xviii


Z. Homonnay, E. Kuzmann and S. Stichleutner Eötvös Loránd University Department of Nuclear Chemistry H-1518 Budapest P.O. Box 32 Hungary P. Tilky, J. Schunk and G. Patek Paks Nuclear Power Plant Ltd H-7031 Paks P.O.Box 71 Hungary

Chapter 21 H. P. Kim, S. S. Hwang, D. J. Kim, J. S. Kim, Y. S. Lim*, M. K. Joung Korea Atomic Energy Research Institute P.O. Box 105 Yusong Taejon 305-600 Korea E-mail: [email protected]

Chapter 22 Krisztián Radó, K. Varga, Z. Németh, I. Varga* and J. Somlai University of Veszprém Department of Radiochemistry H-8201 Veszprém P.O. Box 158 Hungary E-mail: [email protected] [email protected]

D. Oravetz and K. É. Makó University of Veszprém Department of Silicate Chemistry and Materials Engineering H-8201 Veszprém P.O. Box 158 Hungary Z. Homonnay and E. Kuzmann Eötvös Loránd University Department of Nuclear Chemistry H-1518 Budapest P.O. Box 32 Hungary J. Borszéki and P. Halmos University of Veszprém Department of Analytical Chemistry H-8201 Veszprém P.O. Box 158 Hungary P. Tilky and J. Schunk Paks Nuclear Power Plant Ltd H-7031 Paks P.O. Box 71 Hungary

European Federation of Corrosion (EFC) publications: Series introduction

The EFC, incorporated in Belgium, was founded in 1955 with the purpose of promoting European co-operation in the fields of research into corrosion and corrosion prevention. Membership of the EFC is based upon participation by corrosion societies and committees in technical Working Parties. Member societies appoint delegates to Working Parties, whose membership is expanded by personal corresponding membership. The activities of the Working Parties cover corrosion topics associated with inhibition, education, reinforcement in concrete, microbial effects, hot gases and combustion products, environment sensitive fracture, marine environments, refineries, surface science, physico-chemical methods of measurement, the nuclear industry, the automotive industry, computer based information systems, coatings, tribo-corrosion and the oil and gas industry. Working Parties and Task Forces on other topics are established as required. The Working Parties function in various ways, e.g. by preparing reports, organising symposia, conducting intensive courses and producing instructional material, including films. The activities of Working Parties are co-ordinated, through a Science and Technology Advisory Committee, by the Scientific Secretary. The administration of the EFC is handled by three Secretariats: DECHEMA e.V. in Germany, the Société de Chimie Industrielle in France, and The Institute of Materials, Minerals and Mining in the United Kingdom. These three Secretariats meet at the Board of Administrators of the EFC. There is an annual General Assembly at which delegates from all member societies meet to determine and approve EFC policy. News of EFC activities, forthcoming conferences, courses, etc. is published in a range of accredited corrosion and certain journals throughout Europe. More detailed descriptions of activities are given in a Newsletter prepared by the Scientific Secretary. The output of the EFC takes various forms. Papers on particular topics, for example reviews or results of experimental work, may be published in scientific and technical journals in one or more countries in Europe. Conference proceedings are often published by the organisation responsible for the conference. xix

xx

Series introduction

In 1987 the, then, Institute of Metals was appointed as the official EFC publisher. Although the arrangement is non-exclusive and other routes for publication are still available, it is expected that the Working Parties of the EFC will use The Institute of Materials, Minerals and Mining for publication of reports, proceedings, etc. wherever possible. The name of The Institute of Metals was changed to The Institute of Materials on 1 January 1992 and to The Institute of Materials, Minerals and Mining with effect from 26 June 2002. The series is now published by Woodhead Publishing and Maney Publishing on behalf of The Institute of Materials, Minerals and Mining. P. McIntyre EFC Series Editor The Institute of Materials, Minerals and Mining, London, SW1Y 5DB UK EFC Secretariats are located at: Dr B A Rickinson European Federation of Corrosion, The Institute of Materials, Minerals and Mining, 1 Carlton House Terrace, London, SW1Y 5DB, UK Dr J P Berge Fédération Européenne de la Corrosion, Société de Chimie Industrielle, 28 rue Saint-Dominique, F-75007 Paris, FRANCE Professor Dr G Kreysa Europäische Föderation Korrosion, DECHEMA e.V., Theodor-Heuss-Allee 25, D-60486 Frankfurt, GERMANY


1 Corrosion in the nuclear industry Prepared by the Working Party on Nuclear Corrosion 2 Practical corrosion principles Prepared by the Working Party on Corrosion Education (out of print) 3 General guidelines for corrosion testing of materials for marine applications Prepared by the Working Party on Marine Corrosion 4 Guidelines on electrochemical corrosion measurements Prepared by the Working Party on Physico-Chemical Methods of Corrosion Testing 5 Illustrated case histories of marine corrosion Prepared by the Working Party on Marine Corrosion 6 Corrosion education manual Prepared by the Working Party on Corrosion Education 7 Corrosion problems related to nuclear waste disposal Prepared by the Working Party on Nuclear Corrosion 8 Microbial corrosion Prepared by the Working Party on Microbial Corrosion 9 Microbiological degradation of materials – and methods of protection Prepared by the Working Party on Microbial Corrosion 10 Marine corrosion of stainless steels: chlorination and microbial effects Prepared by the Working Party on Marine Corrosion 11 Corrosion inhibitors Prepared by the Working Party on Inhibitors (out of print) xxi

xxii


12 Modifications of passive films Prepared by the Working Party on Surface Science and Mechanisms of Corrosion and Protection 13 Predicting CO2 corrosion in the oil and gas industry Prepared by the Working Party on Corrosion in Oil and Gas Production (out of Print) 14 Guidelines for methods of testing and research in high temperature corrosion Prepared by the Working Party on Corrosion by Hot Gases and Combustion Products 15 Microbial corrosion (Proc. 3rd Int. EFC Workshop) Prepared by the Working Party on Microbial Corrosion 16 Guidelines on materials requirements for carbon and low alloy steels for H2S-containing environments in oil and gas production Prepared by the Working Party on Corrosion in Oil and Gas Production 17 Corrosion resistant alloys for oil and gas production: guidance on general requirements and test methods for H2S Service Prepared by the Working Party on Corrosion in Oil and Gas Production 18 Stainless steel in concrete: state of the art report Prepared by the Working Party on Corrosion of Reinforcement in Concrete 19 Sea water corrosion of stainless steels – mechanisms and experiences Prepared by the Working Parties on Marine Corrosion and Microbial Corrosion 20 Organic and inorganic coatings for corrosion prevention – research and experiences Papers from EUROCORR ’96 21 Corrosion – deformation interactions CDI ’96 in conjunction with EUROCORR ’96 22 Aspects on microbially induced corrosion Papers from EUROCORR ’96 and the EFC Working Party on Microbial Corrosion 23 CO2 corrosion control in oil and gas production – design considerations Prepared by the Working Party on Corrosion in Oil and Gas Production


xxiii

24 Electrochemical rehabilitation methods for reinforced concrete structures – a state of the art report Prepared by the Working Party on Corrosion of Reinforcement in Concrete 25 Corrosion of reinforcement in concrete – monitoring, prevention and rehabilitation Papers from EUROCORR ’97 26 Advances in corrosion control and materials in oil and gas production Papers from EUROCORR ’97 and EUROCORR ’98 27 Cyclic oxidation of high temperature materials Proceedings of an EFC Workshop, Frankfurt/Main, 1999 28 Electrochemical approach to selected corrosion and corrosion control Papers from 50th ISE Meeting, Pavia, 1999 29 Microbial corrosion (Proc. 4th Int. EFC workshop) Prepared by the Working Party on Microbial Corrosion 30 Survey of literature on crevice corrosion (1979–1998): mechanisms, test methods and results, practical experience, protective measures and monitoring Prepared by F. P. Ijsseling and the Working Party on Marine Corrosion 31 Corrosion of reinforcement in concrete: corrosion mechanisms and corrosion protection Papers from EUROCORR ’99 and the Working Party on Corrosion of Reinforcement in Concrete 32 Guidelines for the compilation of corrosion cost data and for the calculation of the life cycle cost of corrosion – a working party report Prepared by the Working Party on Corrosion in Oil and Gas Production 33 Marine corrosion of stainless steels: testing, selection, experience, protection and monitoring Edited by D. Féron on behalf of Working Party 9 on Marine Corrosion 34 Lifetime modelling of high temperature corrosion processes Proceedings of an EFC Workshop 2001. Edited by M. Schütze, W. J. Quadakkers and J. R. Nicholls 35 Corrosion inhibitors for steel in concrete Prepared by B. Elsener with support from a Task Group of Working Party 11 on Corrosion of Reinforcement in Concrete

xxiv


36 Prediction of long term corrosion behaviour in nuclear waste systems Edited by D. Féron and Digby D. Macdonald on behalf of Working Party 4 on Nuclear Corrosion 37 Test methods for assessing the susceptibility of prestressing steels to hydrogen induced stress corrosion cracking Prepared by B. Isecke on behalf of Working Party 11 on Corrosion of Steel in Concrete 38 Corrosion of reinforcement in concrete: mechanisms, monitoring, inhibitors and rehabilitation techniques Edited by M. Raupach, B. Elsener, R. Polder and J. Mietz on behalf of Working Party 11 on Corrosion of Steel in Concrete 39 The use of corrosion inhibitors in oil and gas production Edited by J. W. Palmer, W. Hedges and J. L. Dawson 40 Control of corrosion in cooling waters Edited by J. D. Harston and F. Ropital 41 Metal dusting, carburisation and nitridation Edited by H. Grabke and M. Schütze 42 Corrosion in refineries Edited by J. Harston 43 The electrochemistry and characteristics of embeddable reference electrodes for concrete Prepared by R. Myrdal on behalf of Working Party 11 on Corrosion of Steel in Concrete 44 The use of electrochemical scanning tunnelling microscopy (ECSTM) in corrosion analysis: reference material and procedural guidelines Prepared by R. Lindström, V. Maurice, L. Klein and P. Marcus on behalf of Working Party 6 on Surface Science 45 Local probe techniques for corrosion research Edited by R. Oltra on behalf of Working Party 8 on Physico-Chemical Methods of Corrosion Testing 46 Amine unit corrosion in refineries Prepared by J. D. Harston and F. Ropital on behalf of Working Party 15 on Corrosion in the Refinery Industry 47 Novel approaches to the improvement of high temperature corrosion resistance Edited by M. Schütze and W. Quadakkers on behalf of Working Party 3 on Corrosion in Hot Gases and Combustion Products


xxv

48 Corrosion of metallic heritage artefacts: investigation, conservation and prediction of long term behaviour Edited by P. Dillmann, G. Béranger, P. Piccardo and H. Matthiessen on behalf of Working Party 4 on Nuclear Corrosion 49 Electrochemistry in light water reactors: reference electrodes, measurements, corrosion and tribocorrosion issues Edited by R.-W. Bosch, D. Féron and J.-P. Celis on behalf of Working Party 4 on Nuclear Corrosion 50 Corrosion behaviour and protection of copper and aluminium alloys in seawater Edited by D. Féron on behalf of Working Party 9 on Marine Corrosion 51 Corrosion issues in light water reactors: stress corrosion cracking Edited by D. Féron and J.-M. Olive on behalf of Working Party 4 on Nuclear Corrosion 52 (to come) 53 Standardisation of thermal cycling exposure testing Edited by M. Schütze and M. Malessa

xxvi

Preface

Stress corrosion cracking is one of the major localised corrosion issues in light water nuclear reactors. The various structural materials used in nuclear power plants including low alloy steels, stainless steels, nickel base alloys, Zirconium base alloys, had shown cracks developed under the combined action of a mechanical stress and more or less polluted water. Both pressurised water reactors (PWRs) and boiling water reactors (BWRs) had to face stress corrosion cracking phenomena even if the water chemistry and the materials are different. In both cases, stress corrosion cracking is under control either by changing the material (use of Alloy 690 in PWRs conditions) or the chemistry (hydrogenated/noble water chemistry in BWRs). These are direct feedback of the research and development programmes where European teams played a significant role. Moreover, nuclear power plants were designed to operate for 30–40 years, and the extension of their life time to 60 years is now being envisaged which is a longer life time than in many other industries: the materials used for components and circuit pipes which typically rely on their passivity in the aqueous environment for corrosion protection, are being or are intended to be used for significantly longer periods than initially planned. Further research and developments are then needed to predict their behaviour, to prepare remedial and repair actions. The obtained data will be helpful for the industry to define a fitness-for-service strategy. The objective of this EFC book No. 51 is to give an overview of recent developments on stress corrosion cracking performed mainly by European teams, from laboratory investigations to field applications. The book has been divided in four main parts: (i) overviews, (ii) stress corrosion cracking: susceptibility and initiation, (iii) stress corrosion cracking: propagation and (iv) practical experience. The current state-of-the-art is described not only for stress corrosion cracking, but also for two main related subjects: corrosion potential monitoring and passivation. The book also covers topics ranging from initiation and susceptibility to propagation. It includes low alloy steels, stainless steels and nickel base alloys, boiling water and pressurised water reactor conditions. The editors would like to thank the authors who presented and wrote xxvii

xxviii

Preface

chapters of outstanding scientific and technical content and who responded enthusiastically to the questions and comments raised by the reviewers. They would also like to thank the members of the Working Party ‘Nuclear Corrosion’ (EFC WP4) and of the Working Party 5 ‘Environment Sensitive Fracture’ (EFC WP5) of the European Federation of Corrosion who reviewed these chapters. They would like also to thank Mylene Belgome, Secretary of the EFC WP4, who helped the editors in reviewing, correction and secretarial procedures. The editors hope that this book will be useful to scientists and engineers in the development of understanding of and resolution of stress corrosion cracking phenomena that they have to face in light water reactors. Damien Féron Chairman of the EFC WP4 and Jean-Marc Olive Chairman of the EFC WP5

Part I Overviews

1

2


1 An overview of materials degradation by stress corrosion in PWRs P. M. S C O T T, Framatome ANP, France

1.1

Introduction

Most of the world’s nuclear steam supply systems for generating electricity are based on water cooled and moderated systems of which the most widespread designs are the Pressurized Water Reactor (PWR) and the Boiling Water Reactor (BWR). Such power production systems are initially designed to operate for up to 40 years and extension to 60 years is now being envisaged in many cases. It is perhaps often overlooked that the materials of construction that typically rely on their passivity in the aqueous environment for corrosion protection are being, or are intended to be used for significantly longer periods than in most other industries. Thus, long-term operating experience is really only now being gained as many such nuclear power plants have reached 20 to 30 years old. It is not surprising, therefore, that as plants have aged, some serious corrosion problems have been encountered and remedied or repaired, of which one of the most serious is stress corrosion cracking. The main difference between a PWR and a BWR is that in the former, sub-cooled primary water cools the nuclear fuel and exchanges its heat via steam generators to create steam to drive a turbine and alternator in a secondary circuit. In the latter, water is boiled directly by the nuclear fuel and the steam is then separated and dried before passing directly to the turbine. Operating temperatures range between about 280 and 320 ∞C except for the PWR primary circuit pressurizer which operates at 343 ∞C. The fundamentals of water reactor chemistry treatment and control are described in reference [1] and a recent overview of PWR water chemistry operating experience in reference [2]. From a corrosion perspective, the operating environments in PWRs and BWRs are radically different as illustrated in Fig. 1.1 on a Pourbaix diagram for nickel and iron at 300 ∞C. (This figure also indicates the corrosion conditions for some common stress corrosion phenomena in both PWRs and BWRs that will be described later.) Thus, in PWRs, the water of the primary and secondary circuits are alkali treated and essentially oxygen-free to ensure minimum 3

4

Corrosion issues in light water reactors V(SHE) 1.50 O2 1.00

H2O

H+ 0.50

0.00

NiO2 1 ppb H2 1 atm 8.2 ppm

Acid SO4 cracking

++

Ni Ni

BWR NWC

NiO PWR Secondary side Primary side

++

Fe Fe

– 0.50

Ni3O4

BWR HWC

Fe3O4 Fe

Caustic cracking

PWSCC

– 1.00

Ni (OH)3 Caustic IGA

– 1.50

– 2.00 0

2

4

6

8

10

12

14

pH

1.1 Simplified Pourbaix diagram for nickel and iron at 300 ∞C showing the principal pH-potential combinations for PWR primary and secondary water, BWR Normal Water Chemistry (NWC) and BWR Hydrogen Water Chemistry (HWC) and the modes of stress corrosion cracking of Alloy 600.

general corrosion and corrosion product release rates of the structural materials. PWR primary water also contains about 3 ppm of dissolved hydrogen to suppress water radiolysis and, as a consequence, primary circuit corrosion potentials are about 200 mV lower compared to the secondary side, in both cases being close to the H2/H+ redox potential for virtually all structural materials. In direct cycle BWRs by contrast, extremely pure water is used to ensure the lowest possible general corrosion rates. For those BWR plants on Normal Water Chemistry (NWC), radiolytic decomposition of water in combination with removal of non-condensable gases at the turbine condenser

An overview of materials degradation

5

establishes electrochemically significant concentrations of dissolved oxygen and hydrogen peroxide in the recirculating water and consequently corrosion potentials are around 500 mV more positive than in PWR primary coolant circuits. In the case of the Hydrogen Water Chemistry (HWC) variant for BWRs, hydrogen at about 10% of the concentration typical of PWR primary circuits is used to depress corrosion potentials to values intermediate between those of BWR NWC and PWR primary circuits, specifically with the intention of protecting sensitized and cold worked stainless steels from Intergranular Stress Corrosion Cracking (IGSCC), as described briefly later. The main emphasis of the examples of corrosion related material failures described hereafter come mainly from the author’s experience of interpreting and modelling stress corrosion of structural materials in PWR systems. However, some examples of BWR experience are also provided for comparison as well as to illustrate the significantly different experience in many cases between the two water cooled nuclear reactor systems.

1.2

Nickel base alloys in PWR primary water

The most severe stress corrosion problem to affect PWRs is IGSCC of Alloy 600 in PWR primary water (sometimes called PWSCC for Pressurized Water Stress Corrosion Cracking, as in Fig. 1.1). It has become a generic issue rivalling that of IGSCC of sensitized and/or cold worked stainless steels in BWRs in terms of unanticipated outages and cost of repairs. In addition, high strength nickel base alloy fasteners and springs fabricated from precipitation hardened Alloys X750 or 718 are used extensively in PWR primary circuits and some service failures of these items have also occurred. Alloy 600, a nickel base alloy containing 14–17% Cr and 6–10% Fe plus various minor elements was initially adopted for use in PWRs for steam generator tubes because of its excellent resistance to chloride cracking (from the secondary side) compared to stainless steel. It was also attractive for primary circuit components because of the close similarity of its coefficient of thermal expansion to that of the low alloy steel used to fabricate the reactor pressure vessel, pressurizer and steam generator shells. The susceptibility of Alloy 600 to IGSCC in operational service in PWR primary water was first revealed in steam generator tubing in the early 1970s in tight U-bends and in rolled, cold-worked transitions in diameter within or just above the tube sheet [3]. This then became a major cause of steam generator tube cracking in the 1980s, and later, premature steam generator retirement and replacement. IGSCC of pressurizer nozzles and Control Rod Drive Mechanism (CRDM) penetrations in the upper heads of PWR reactor pressure vessels followed in the late 1980s and has continued for over a decade [4, 5]. Apparently interdendritic, but in fact intergranular, stress corrosion cracking of the compatible weld metals Alloys 182 and 82, the

6


former having a composition similar to that of Alloy 600, has also been observed more recently in major primary circuit welds of several PWR plants, often after very long periods in service ranging between 17 and 27 years [5]. To these can be added the experience of extensive IGSCC in the g ¢ strengthened analogue of Alloy 600, Alloy X750, which is used for split pins attaching the CRDM guide tubes to the upper core plate. Even Alloy 718, a high strength nickel base alloy containing 17–21% Cr, which is normally considered a very reliable high strength material in PWR primary water use, has occasionally exhibited IGSCC [4]. A common feature of service failures of Alloy 600 and its compatible weld metals is the presence of very high residual stresses exceeding the nominal yield strength, usually coupled with a roughly machined or heavily ground surface finish. High residual stresses may be induced by rolling operations as with steam generator tube expansion into the steam generator tube sheet mentioned above or by nearby welding operations as in the case of CRDM nozzles. If thermal or mechanical plastic straining results in a plastic compression/tension hysteresis cycle, then very high tensile stresses easily up to 1000 MPa can be generated. By contrast, stress relief (in practice of attached low alloy steel components) has a very favourable effect on IGSCC resistance and no failures of Alloy 600 components so stress relieved have occurred in service. The other major factors influencing IGSCC susceptibility are the material microstructure and the temperature, an activation energy of 44 kcal/mole being generally admitted for crack initiation. Much research into the metallurgical parameters affecting IGSCC of Alloy 600 and similar materials in PWR primary water has shown that chromium carbides precipitated on the grain boundaries improve resistance while intragranular carbides have the opposite effect. Thus material procurement specifications were developed to ensure that products were delivered with the carbon precipitated as far as possible as carbides on grain boundaries. Even ‘sensitized’ materials, that is those with grain boundary carbides but an adjacent narrow zone of chromium depletion have improved IGSCC resistance in PWR primary water, in sharp contrast to their very poor resistance in oxygenated BWR NWC. The generic mechanism IGSCC of the nickel base Alloy 600 and its high strength analogue, Alloy X750, in PWR primary water has been extensively studied. Despite considerable experimental efforts, no consensus exists as to the nature of the cracking mechanism [1] and both life modelling and remedial measures have relied on empirical, phenomenological correlations. In addition to the major influencing parameters of stress, cold work, temperature and carbide morphology mentioned above, a profound influence of hydrogen partial pressure (or corrosion potential) has been identified with a worst case centred on corrosion potentials near the Ni/NiO equilibrium (Fig. 1.1). The mechanism of cracking also does not apparently change between 300 ∞C sub-cooled water and 400 ∞C superheated steam.


7

It is interesting to note that despite the intense debate concerning the mechanism of IGSCC of Alloy 600 in PWR primary water, the most recent models incorporate the idea that solid state grain boundary diffusion is rate controlling [6]. This is independent of whether the mechanistic model considers that cracks advance by an oxidation process at the crack tip or due to embrittlement caused by hydrogen discharged by the matching cathodic reaction. Such models provide physically based support for the high value of the apparent activation energy, which is typical of solid state grain boundary diffusion in nickel. Physical support for a fourth power dependency of IGSCC on applied stress comes mainly from studies of grain boundary sliding (itself dependent on grain boundary diffusion) observed during primary creep in Alloy 600 at temperatures between 325 and 360 ∞C. Grain boundary sliding rates are also observed to depend on grain boundary carbide coverage, greater coverage being associated with slower grain boundary sliding rates and higher resistance to IGSCC. Various empirical models have been developed to predict IGSCC of Alloy 600 and similar materials in PWR primary circuits until, as sometimes is the case, replacement becomes unavoidable. The only presently perceived sure remedy for susceptible Alloy 600 components is replacement, usually by Alloy 690 (28–31% Cr and 7–11% Fe) and its compatible weld metals, Alloys 152 and 52, which have proved to be resistant to IGSCC in PWR primary water both in severe laboratory tests and, to date, after up to 15 years in service. Predictive equations for IGSCC in Alloy 600 were first developed for steam generator tubes and later extended to pressurizer nozzles and upper head CRDM penetrations [7, 8]. Both deterministic and probabilistic methods have been developed. Modelling of Alloy 600 component life is often based on the following empirical equation:

( )

–4 t f = C s exp E Im RT

1.1

where: tf is the failure time (hours), C is a constant, s is the applied stress (MPa), Im is a material susceptibility index (e.g. Table 1.1), E is the apparent activation energy (44 kcal/mole), R is the universal gas constant (1.987 cal/mole/∞K), T is the absolute temperature (∞K). Establishing the stress including residual fabrication stress on a given component is not trivial, but well tried and proven approaches based on finite element stress analysis or experimental techniques applied to mock-

8


ups are available. Dealing with material variability in susceptibility to IGSCC is not so straightforward, however, and in the case of classification of the susceptibility of CRDM nozzle cracking in US PWRs, has been ignored. One method to account for variability in material resistance to IGSCC has been based on a system of material indices, Im, in equation (1.1) [9]. At its simplest, with no direct information about IGSCC susceptibility of individual heats, the guidelines given in Table 1.1 were adopted. They were based on observations of minimum times to failure of plant components or, in cases where no service failures have been observed, of laboratory specimens in accelerated tests of representative plant materials. The constant C in equation (1.1) was adjusted so that an index of unity corresponds to a minimum failure time of 10,000 hours at a temperature of 325 ∞C and an applied stress of 450 MPa, as observed in practice in plant and in laboratory tests. In addition, temperature and stress indices were defined relative to the reference conditions of 325 ∞C and 450 MPa consistent with equation (1.1) as follows:

(

)

È –Eˆ 1 ˘ Iq = exp Í Ê – 1 ˙ R T Ë ¯ 598 Î ˚

Thus:

tf =

10000 I m ◊ Iq ◊ I s

Is =

s ( 450 )

4

1.2 1.3

In this way, the minimum time to cracking of each generic Alloy 600 primary circuit component was assessed after determining its operating temperature and stress. The results for different generic components of PWR primary circuits are shown in Table 1.1. Appropriate surveillance strategies were then established. The quantification of variability of Alloy 600 heat susceptibility to IGSCC has been developed further to assess cracking encountered in the upper head CRDM nozzles of French PWRs and extended to other large Alloy 600 primary circuit components [8]. Three main types of microstructure were recognized and related to the carbon content, thermal treatment, especially the temperature at the end of forging or rolling operations, and yield strength after hot-working: ∑ ∑ ∑

class A with mainly intergranular carbide precipitates; class B re-crystallized with carbides mainly on a prior grain boundary network; class C re-crystallized with randomized intragranular carbides as well as carbides on prior grain boundaries.

These classes were then linked to their IGSCC resistance (i.e. material susceptibility index) as determined from operating experience or in accelerated laboratory tests of archive materials mainly at 360 ∞C. Inevitably, such an approach to assessing IGSCC susceptibility reveals

Table 1.1 Minimum failure times for IGSCC of Alloy 600 components in PWR primary circuits [9] Alloy 600 parts

Material index

Stress index

Temperature index

Overall index

Time (hours)

Observation *

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18

Hydraulic expansion Divider plate Hard rolling on cold leg (Ringhals 2) Pressurizer nozzle (San Onofre 3) Nozzle (San Onofre) Pressurizer nozzle (ANO1) Pressurizer nozzle (Palo Verde 1) Nozzle (Palo Verde 2) Explosive expansion Fessenheim 1 Hard rolling on SG hot leg (Gravelines 6) Hydraulic expansion (Doel 2) Small U-bends Vallourec Small U-bends Westinghouse Sensitive hard rolling on SG hot leg Very sensitive hard rolling on SG hot leg 1300 MW Pressurizer Nozzle Mechanical pluggs French CRDM Nozzles

0.2 0.5 2 0.5 0.5 0.5 0.5 0.5 1 0.5 2 2 2 1 2 0.5 0.5 0.5 0.5 1.1 1.1 1.1 1.1

0.4 0.3 2.2 0.1 0.9 0.3 0.4 1.5 0.4 2.2 0.4 2.2 10 2.2 2.2 3.2 1 1.5 1.5 2.8 2.5 2.5 2.5

1 0.9 0.1 3.3 3.3 3.3 3.3 1.1 1 1 1 0.3 0.3 1 1 3.3 1 0.5 0.5 0.08 0.08 0.08 0.08

0.08 0.14 0.44 0.17 1.49 0.5 0.66 0.83 0.4 1.1 0.8 1.32 6 2.2 4.4 5.28 0.5 0.4 0.4 0.24 0.22 0.22 0.2

80000 80000 48000 56000 8000 84336 33320 25000 75000 30000 30000 30000 6000 20000 8000 8000 40000 80000 26800 72909 48427 58868 90777

NC NC C C C C C C C C C C C C C C C C C C C C C


Ref.

*: NC: non cracked; C: cracked.

9

10

Corrosion issues in light water reactors Influence of time and susceptibility

Cumulated percentage of cracked penetrations

99.99

1300 MW Units: Circles 12 to 14 99 Simulation Inspection results

90 70

Susceptibility C

50 30

Susceptibility B

10

Susceptibility A

1

0.01 20000

50000 80000 Time (hours)

200000

1.2 Results of Monte Carlo simulations of IGSCC in upper head CRDM penetrations of 1300 MWe French PWRs and comparison with inspection results for each class of alloy 600 [8].

significant scatter in the susceptibility indices for different heats about the mean associated with each class. This dispersion in material properties combined with the dispersion of stress values for any particular component gives rise to a distribution of failures with time that can be fitted to an appropriate function such as the Weibull distribution. The main advantage of the Weibull distribution is that it has a linear transform that can be fitted to the early failures in order to give a reliable prediction of the increase in stress corrosion failures with time [7, 8]. Further improvements in estimating the progression of IGSCC failures in Alloy 600 with time as well as the uncertainty in those predictions have come about by applying the Monte Carlo simulation technique of randomly sampling distributions of the input parameters in equation (1.1) [8]. An example of the results using the Monte Carlo approach is shown in Fig. 1.2 in the form of a Weibull distribution comparing the results of these simulations with the inspection results for upper head penetrations in each susceptibility class of Alloy 600. When the Monte Carlo simulations are repeated many times, the dispersion in the resulting Weibull distribution of failure times is relatively small because the number of penetrations considered for each PWR plant series is quite large (over 1000). It can be shown that the progression of the problem for each design series of PWRs has relatively little inherent uncertainty. On the other hand, if the problem is considered on a plant by


11

plant basis, then the statistical uncertainty in predictions of the proportion that will crack in a given operating time is much greater because there are less than a hundred CRDM penetrations per upper head. For a given upper head, this statistical uncertainty can be of the order of ±1 to ±5 on the mean prediction, which is easily demonstrated and quantified in a probabilistic sense using the Monte Carlo simulation technique. Once a stress corrosion crack has been detected by non-destructive examination in a PWR primary circuit component, an essential step in the justification of structural integrity and further operation without repair or replacement of the affected component is an assessment of crack growth during the next few operating cycles. Practical approaches to assessing crack growth by IGSCC in Alloy 600 components have relied on empirical measurements of crack growth rates as a function of crack tip stress intensity, KI, of the form [10, 11]: da = C ◊ ( K – 9) n I dt

( K I in MPa m )

1.4

The values of the coefficients C and n for given practical circumstances vary between different publications but there is a reasonable consensus that the apparent or effective activation energy to be used for adjusting the coefficient C for temperature is ~31 kcal/mole.

1.3

Nickel base alloys on the secondary side of PWR steam generators

The main type of PWR steam generator in general use is the vertical Recirculating Steam Generator (RSG) with tube bundles, depending on age, made from either mill annealed or thermally treated Alloy 600, thermally treated Alloy 690, or Alloy 800. Thermal treatment is carried out at ~700 ∞C with the objective of precipitating dissolved carbon as chromium carbides on the grain boundaries. Sub-cooled primary water flows through inside of the tubes and boils secondary water on the shell side of the tubes. The steam quality of the water-steam mixture entering the steam driers of RSGs is typically 10% and the superheat across the tubes may vary from 10 to 40 ∞C. Vertical PWR steam generators have experienced a variety of corrosioninduced problems and many have been replaced, usually because of corrosion induced cracking of mill annealed Alloy 600 steam generator tubes. Only a very few thermally treated tubes have experienced such problems and they appear to be due to isolated failures of the thermal treatment to ensure an adequate grain boundary carbide microstructure. Some steam generators with mill annealed Alloy 600 tubes have been replaced after only 8 to 12 years of operation, which is well short of the usual initially licensed plant operating period of 40 years. New or replacement RSGs are supplied with thermally

12


treated Alloy 690 or Alloy 800 tubing, which to date have resisted both primary and secondary corrosion problems. Secondary side steam generator tube corrosion problems involving mill annealed Alloy 600 include denting, wastage, intergranular attack, IGSCC, and pitting on the outside surfaces of the steam generator tubes [3, 12]. The evolution of steam generator tube corrosion with time in terms of relative importance of each damage mechanism is shown in Fig. 1.3. Many secondary side corrosion problems with mill annealed Alloy 600 tubes have been associated with the interstices between the tubes and the tube supports. The tube support structures for most of the early units were made of carbon steel, while later units switched to Types 405, 409, and 410 ferritic stainless steels for greater corrosion resistance. Tube support structures of early units used plates with drilled holes, then plates with trefoil or quatrefoil broached holes, initially with concave lands and then flat lands, or lattice bars (egg crates). The objective of the more open tube support designs is to reduce the accumulation of impurities in the interstices by the phenomenon of hideout (see later). Another corrosion sensitive zone for steam generator tubes has been in and just above the tubesheet. In some of the very early RSG designs, the tubes were only partly expanded just above the seal weld with the lower tubesheet face, thus leaving a crevice between the outside diameter of the tube and the inside diameter of the hole in the tubesheet. Later, the tubes were expanded into the tubesheet along nearly their full length in order to close all but the last ~4 mm of the tube to tubesheet crevice. Tube expansion has been achieved by various methods, mechanical rolling, hydraulic, and explosive. Each expansion method generated its own characteristic residual stress fields in the tubes that have influenced subsequent stress corrosion behaviour if, or when, impurities concentrate by hideout either in the tube sheet crevice or under sludge that accumulates on the upper face of the tubesheet. The underlying cause of all forms of localized corrosion observed on the secondary side of steam generators is the phenomenon of hideout of low volatility solutes in superheated crevices with restricted water circulation. Most impurities entering recirculating steam generators in the feed water are relatively insoluble in the steam phase and can concentrate by potentially many orders of magnitude in occluded superheated crevices by a wick boiling mechanism. Due to the potential variety of impurities entering the steam generators, many complex mixtures of concentrated chemicals can be envisaged. This severely complicates the task of understanding the mechanisms of tube attack and defining adequate remedies. Tube damage such as wastage, pitting and denting has been attributed to the local formation of strong acids and, evidently, has been largely eliminated by appropriate management of secondary water chemistry (Fig. 1.3). By contrast, the steadily rising trend in


13

100 90

Other

80 Pitting

Percentage

70 60 SCC (ID)

50 40

Denting Wastage

30 20

SCC/IGA (OD)

10 0

Fretting 73 74 75 76 77 78 79 80 81 82 83 84 85 86 87 88 89 90 91 92 93 94 95 96 Year

1.3 Worldwide causes of steam generator tube plugging [12].

IGA/IGSCC (Fig. 1.3) suggests that counter measures have not been completely effective, probably because the mechanism has not always been correctly identified. The morphology of IGSCC in mill annealed Alloy 600 steam generator tubing consists of single or multiple major cracks with minor-to-moderate amounts of branching that are essentially 100% intergranular. Experience suggests that secondary side IGSCC requires stresses greater than 0.5 yield in order to propagate rapidly. At lower levels, propagation rates may approach zero, or the corrosion may take the form of intergranular attack (IGA). IGA is the second generally recognized form of secondary side corrosion attack of mill annealed Alloy 600 where there is substantial volumetric attack of every grain boundary. Stress is not strictly necessary for IGA to occur, which distinguishes it from IGSCC. Nevertheless, the two are clearly closely related. IGA/IGSCC varies greatly with height of the tube support plate in recirculating steam generators, being much more prevalent at the lower levels where the temperature difference between the primary and secondary fluids is greatest. This is clearly strong evidence for the importance of impurity hideout, which increases as a function of the available superheat on the secondary side. Broached tube support plates minimize the extent of the narrow gap between the tube and its support plate and hence substantially reduce the tendency for impurity hideout in such locations. When the fraction of tubes affected by IGA/IGSCC at tube support plate intersections is plotted as a function of time on Weibull distribution coordinates, it is observed that the slopes of the Weibull plots are rather high, typically

14


between 4 and 9 [7, 13]. This indicates that once IGA/IGSCC starts, its progression to other tubes is rather rapid and relatively consistent between different plants. On the other hand, incubation periods before cracking starts vary considerably. In some cases, IGA/IGSCC has not been observed at all, even on mill annealed Alloy 600 tube bundles after very long periods of operation. There is a tendency to attribute this variability between plants mainly to differences in secondary water chemistry and impurities. However, heat to heat variability in sensitivity of mill annealed Alloy 600 to IGA/ IGSCC is very important in this respect and the proportion of very sensitive heats varies markedly between different plants [4, 13]. Following the retirement of some steam generators with degraded tubing, it has been possible to extract and observe metallographically complete tube/ tube support plate intersections [14]. These studies have revealed that the crevice between the tube and tube support plate is typically plugged at its entrance and exit with a very low porosity ( 1 MeV) with yield stresses typically in the range 800 to 1100 MPa. Point defect trapping at grain boundaries leads to changes of local elemental composition in a zone about ±5 nm wide due to atoms of different elements exchanging at different rates with the diffusing point defects. Typically chromium, iron and molybdenum depletion and nickel and silicon enrichment are observed. More generalized changes in elemental composition may also be caused by nuclear transmutation reactions In the case of the oxygenated coolants of BWRs, the modification of grain boundary composition due to neutron irradiation, particularly chromium depletion, has been shown to be an important precursor of IASCC. Neutron doses exceeding 5 ¥ 1020 n/cm2 (E > 1 MeV) are associated with the occurrence of IASCC in BWRs, this being the dose required to develop sufficient irradiation-induced chromium depletion at grain boundaries. (Note that the maximum end-of-life dose to the core internals of BWRs is about 8 ¥ 1021 n/cm2 (E > 1 MeV), which is about an order of magnitude less than that anticipated for PWRs due to the wider water gaps between the fuel and internals in the former case.) In addition, the formation of oxidizing species, oxygen and hydrogen peroxide, by radiolysis plays an important role in this manifestation of IASCC in BWRs, which is absent in PWRs due to the hydrogen added to PWR primary water. Nevertheless, PWR field experience has also shown that intergranular cracking of highly irradiated core components can occur. Type 304 cladding of control rods and cold worked Type 316 core baffle-former bolts of some first generation (CP0 series) 900 MWe French PWRs have cracked intergranularly in service [21]. Fast neutron doses of >2 ¥ 1021 n/cm2 (E > 1 MeV), strains >0.1 %, and absence of water circulation around the bolt shanks in the affected plants have been implicated in the cracking. Clearly, the absence of oxidizing species, oxygen and hydrogen peroxide, is an obvious environmental difference compared to BWRs that renders grain boundary chromium depletion of no particular consequence in PWR primary water. However, the considerable hardening that occurs very probably plays an important role (as indeed it also does in BWRs) [22]. In addition to the phenomena of radiation induced hardening and changes to grain boundary composition, other radiation damage processes could have an important influence on the development of IASCC. Helium bubble

18


formation, particularly if it occurs on grain boundaries, may have an adverse effect on intergranular stress corrosion resistance. Irradiation creep can relax residual and applied stresses and is independent of temperature in the range of interest to light water reactors. Swelling, hitherto only considered of importance to fast reactors, could in principle also appear at the high neutron doses associated with the second half of life of PWRs and affect the loads applied to components such as baffle bolts due to differential swelling rates between Type 304 and 316 stainless steels. Thus, although significant advances have been made in the understanding of IASCC, much remains to be learned, and it is today a very active field of research in the context of both BWR and PWR plant aging.

1.4.3

High strength fasteners

Precipitation hardened high strength nickel base alloys have already been mentioned earlier but high strength stainless steels are also widely used in PWRs for components such as bolts, springs and valve stems. The main ones are A286 precipitation hardened austenitic stainless steel, A410 and similar martensitic stainless steels, and 17-4 PH precipitation hardened martensitic stainless steel. Small numbers of such components have cracked over the years due to stress corrosion or hydrogen embrittlement and on occasions, loose parts have been generated in the primary circuit. A286, an austenitic, precipitation hardened, stainless steel is strengthened by g ¢, Ni3(Ti, Al), formed during aging at 720 ∞C. Its use is favoured where the expansion coefficient relative to other austenitic stainless steels is an important design factor. Unfortunately, it is susceptible to IGSCC in PWR primary water when loaded at or above the room temperature yield stress, typically 700 MPa. [23, 24]. Cold work prior to aging in combination with the lower of two commonly used solution annealing temperatures of 900 and 980 ∞C has a particularly adverse effect on resistance to IGSCC. Hot heading of bolts, which can create a heat-affected zone between the head and shank, is another known adverse factor. Nevertheless, even if these metallurgical factors are optimized, immunity from cracking cannot be assured unless the stresses are maintained below the room temperature yield stress, which necessitates strictly controlled bolt loading procedures. There is also strong circumstantial evidence that superimposed fatigue stresses can lower the mean threshold stress for IGSCC even further. Finally, impurities, including oxygen introduced during plant shut down and possibly consumed only slowly in confined crevices, may help crack initiation. Once initiated, cracks grow relatively easily even in well-controlled PWR primary water. Components such as valve stems, bolts and tie rods requiring rather high strength combined with good corrosion resistance in PWR primary circuit water have been typically fabricated from martensitic stainless steels such as


19

Type 410 and 17-4 PH. Significant numbers of failures of Type 410 and similar martensitic stainless steels have occurred [25]. In most cases, the affected components have usually entered service too hard due to tempering at too low a temperature. No in-service aging seems to have been involved, however, in the case of Type 410 and similar martensitic stainless steels. An additional problem has been caused by galvanic corrosion with graphite containing materials in the packing glands of valves, sometimes leading to valve stem seizure. The preferred replacement material has often been 17-4 PH with its higher chromium and molybdenum content no doubt conferring better resistance to crevice corrosion. Service failures of 17-4 PH precipitation hardening stainless steel have also occurred in PWR primary water [26, 27]. Initially, intergranular cracking by stress corrosion/hydrogen embrittlement was associated with the lowest temperature aging heat treatment at 480 ∞C (900 ∞F) designated H900. This gives a minimum Vickers hardness value of 435HV clearly in excess of the limit of 350HV commonly observed to limit the risk of hydrogen embrittlement. The 593 ∞C (H1100) aging heat treatment was subsequently widely adopted and normally yields a hardness value below 350HV. Nevertheless, a small number of failures have continued to occur. The origin of these failures appears to be thermal aging in service rather than ‘reversible temper embrittlement’ that is related to the diffusion of phosphorus to grain boundaries at aging temperatures generally above 400 ∞C. Thermal aging of precipitation hardened stainless steels such as of 17-4 PH arises from an intra-granular decomposition of the martensitic matrix into two phases, a which is rich in iron, and a ¢ which is chromium rich. A progressive generalized increase in hardness is observed with corresponding increases in strength and ductile/ brittle transition temperature and loss of fracture toughness. The hardening cannot be reversed without re-solution annealing. French studies have shown that this aging mechanism can occur in 17-4 PH steels on time scales relevant to the design lives of PWRs at temperatures exceeding 250 ∞C and quantitative models for component assessment have been developed [26]. Intergranular failures have been associated with hardness values following in-service aging that have significantly exceeded 350HV and have also been apparently aggravated by impurities coming from valve packing gland materials.

1.5

Low alloy steels

1.5.1

Secondary circuit components

A small number of potentially serious failures caused by transgranular stress corrosion/corrosion fatigue have occurred in low alloy steel steam generator shells and carbon steel feedwater piping that are directly exposed to secondary water. The combination of fabrication and operational factors necessary for

20


such cracking to occur in carbon and low alloy steels in steam-raising plant has ensured that it has in reality been highly plant specific. Extensive circumferential cracking of the upper shell to cone girth welds of all the Indian Point 3 steam generators was found in 1982 following a steam leak through one of more than a hundred circumferential cracks [28]. Subsequently, the steam generator shells of six other plants located in the United States and Europe were also observed to be cracked in the same location. In some cases, cracking recurred after local repairs had been made by contour grinding. The steam generator shell cracking was caused by an environmentally assisted cracking mechanism and has been variously called ‘corrosion fatigue’, ‘stress corrosion cracking’, or ‘strain-induced corrosion cracking’. In fact, the last term seems most appropriate since it recognizes that although the cracking is environmentally controlled, a dynamic strain is necessary to maintain crack propagation [29]. Consequently, crack extension tends to occur intermittently alternating with pitting at the crack tip during quiescent periods. This environmentally assisted cracking (EAC) mechanism observed for steam generator shell materials is well known and characterized both for bainitic low alloy steels as well as for ferritic-pearlitic carbon manganese steels used extensively in both conventional and nuclear steam-raising plant [29–31]. In addition to the dynamic loading requirement usually caused by large thermal transients, cracking has been associated in practice with high residual welding stresses due to poor or non-existent stress relief. The worst affected plants had been weld repaired during fabrication of the final closure weld. In one case, the girth weld had to be completely remade and stress relieved at a higher temperature of 607 ∞C compared to 538 ∞C originally. Water chemistry transients, particularly oxygen ingress, occurring at the same time as dynamic loading have also been strongly linked to the observed cracking. The effect of oxygen was observed to be greatly exaggerated if copper corrosion products (e.g. from brass condenser tubes) were also present [28]. The only metallurgical factors that appeared to play a role were the sulphur impurity content of the steel in the form of manganese sulphide, where the risk of cracking was greater the higher the sulphur content, and possibly also the free nitrogen content via the phenomenon of strain aging [31]. Practical resolution of steam generator shell cracking has been mainly achieved by contour grinding of existing cracks and by ensuring that auxiliary feed water is properly de-oxygenated prior to use, particularly during plant start-up. In addition to these reported incidents of steam generator shell cracking, a very large technical literature exists concerning EAC of carbon and low alloy steels in both nuclear and conventional steam-raising plant [29–31]. The observed cracking is usually transgranular cleavage-like in appearance although can occasionally be intergranular without any obvious involvement of other chemical pollutants.


1.5.2

21

High strength fasteners

High strength martensitic and maraging steels are used in many external fastener applications in nuclear power plants as well as for some internal fasteners in PWR secondary circuits. A significant number of corrosion related failures of external fasteners used for support bolting and pressure boundary flanges have occurred [32]. Failures of low alloy (AISI 4340 and 4140) and maraging steel support bolting have been attributed mainly to hydrogen embrittlement. Steels with ultra high yield strengths greater than 1000 MPa have failed due to a combination of too high applied stresses and humid or wet environments collecting around the bases of components. Pitting often precedes cracking in such cases. Steels with lower yield strengths have also failed due to poor heat treatment or material variability. Hydrogen cracking is usually avoided by specifying an upper bound strength limit (normally defined by a hardness level acceptance criterion of L) of crack traces of depth greater than L versus the length L. In the case of 600 MA alloy, it was not possible to discriminate between a slow and a rapid propagation step, as shown in Fig. 9.4: In addition, the same result was obtained in the case of 600 TT alloy. In conclusion, a rapid propagation step is not determined in the case of alloy 600 in the reference ‘complex’ environment. This result is further confirmed in the reference ‘complex’ environment by an investigation of propagation on WOL specimens. Indeed, no crack deeper than 40 mm was locally obtained, leading to a local crack growth rate of ca 0.02 mm/h, fully consistent with a propagation in the slow regime.

9.3.5

Slow propagation

The slow propagation step was studied in a ‘complex’ environment by using C-ring specimens in 600 MA and TT alloy at 320 ∞C. The effects of several parameters on slow propagation, including chemical conditions, potential and temperature have been studied in detail. 1000

N (DL)

100

10

1 0

50

100 L (mm)

150

200

9.4 Number N (N > l) of cracks traces of depth greater than L versus the length L on CERT specimen in 600 MA alloy in the reference ‘complex’ environment at 320 ∞C (strain rate of 5 ¥ 10–8 s–1).

Stress corrosion cracking of steam generator tubes

125

Organic compounds effect Firstly, the results obtained in a ‘complex’ environment at 320 ∞C (with or without acetate) suggest that the choice of the amine (ammonia or morpholine) is not decisive on IGSCC. Moreover, cracking of alloy 600 is reduced if a large concentration of ammonia is used; acetate alone does not play a significant role in the cracking. Secondly, the results obtained in an AVT environment with or without phosphate confirm their inhibitor effect on cracking at a high concentration (7.8 g/L), whereas they are harmful at an intermediate concentration (2.6 g/ L) [7]. Moreover, phosphates are not necessary in the ‘complex’ environment for the cracking to occur. Thirdly, the effect of carbonate has also been studied. Indeed, carbonate increases IGSCC at high pH (9.2). However, no effect is observed at neutral pH. It has been shown that no cracking is observed when carbonate is replaced by sodium hydroxide (pH = 9.2). Therefore, this effect is likely to be due to carbonate and not to the pH value. Finally, several tests have also been carried in the ‘complex’ environment at 320 ∞C by varying the Al/Si ratio (see Section 9.4), as shown in Fig. 9.5: These tests, conducted in EDF laboratories, show that the Al/Si ratio has little effect on cracking in a ‘complex’ environment. However, some significant detrimental effect was observed in CEA laboratories [6]. This different behavior could be due to different hydrogen concentrations (and potential) used during these tests.

Crack length (mm/h)

150 U581 MA U573 MA U581 MA 100

50

0 0

0.1

0.2

0.3

3191 h test, no break

0.4

0.5 0.6 Al/Si ratio

0.7

0.8

0.9

9.5 Al/Si ratio effect on the crack length on C-ring specimens (600 MA alloy) in a ‘complex’ environment (T = 320 ∞C).

1

126


Redox potential effect Four tests carried out for C-ring specimens (600 MA and TT alloy) at different potential values between 0 mV and 360 mV/Ecorr. These tests strongly suggest that the crack velocity decreases rapidly as the potential value increases from 0 to 400 mV/Ecorr, as shown in Fig. 9.6: Temperature effect Four tests carried out for C-ring specimens (600 MA and TT alloy) at different temperature values between 305 ∞C and 335 ∞C show that the crack velocity varies with temperature, as shown in Fig. 9.7: In particular, the maximum of IGSCC for alloy 600 in the ‘complex’ environment occurs between 312.5 ∞C and 320 ∞C. This effect, which has already been observed in the sulfate environment [9], could be due to a change in the corrosion mechanism (i.e., from IGSCC to general corrosion). In most tests, 600 TT alloy is a bit less sensitive than 600 MA alloy towards IGSCC, but in some conditions (e.g. at 312.5 ∞C, see Fig. 9.7), 600 TT alloy can be more sensitive that 600 MA alloy towards IGSCC. Material effect The IGSCC results obtained for 600 MA alloy do not depend on the heat considered during the tests. Moreover, 600 TT alloy was revealed to be generally a bit more resistant towards IGSCC than 600 MA alloy. In particular,

0.04 600 MA 600 TT

Vmax (mm/h)

0.03

0.02

0.01

0 0

100

200 E (mV/Ecorr)

300

400

9.6 Potential effect on the crack velocity of 600 alloy in the reference ‘complex’ environment at 320 ∞C.


127

0.04 600 MA 600 TT

vmax (mm/h)

0.03

0.02

0.01

0 300

305

310

315

320 T (∞C)

325

330

335

340

9.7 Temperature effect on the crack velocity of 600 alloy in the ‘complex’ environment at 320 ∞C.

the ratio in the same experimental conditions between the crack velocities for 600 MA and TT C-ring specimens in a ‘complex’ environment is 0.45 ± 0.35.

9.4

Discussion

IGA and IGSCC of mill annealed alloy 600 occur under alumino-silicate deposits. Such alumino-silicate deposits are expected to have similar structures to zeolites which are widely used as catalysts or ion exchange products since they correspond to potential Lewis acids (compounds with an electron doublet vacancy). On a structural point of view, zeolites correspond to crystalline inorganic polymers based on a repeating framework of AlO 4– and SiO 4– tetrahedra linked by some common oxygen atoms. These compounds can be chemically described by the formula M 1/n+n (AlO 2 , ySiO 2 ), z H 2 O , where Mn+ corresponds to an alkali, an ammonium or a transition metal cation. As a consequence, these compounds may accept cations from the oxidizing surface of alloy 600, according to the following mechanism [6]: ∑

electron transfer reaction: M 1/n+n (AlO 2 , ySiO 2 ), z H 2 O Æ zeolite-Al[]Lewis acid center + H2O zeolite-Al[]Lewis acid center + Ni Æ Ni2+ + zeolite-Al[:]

∑

ion exchange reaction: Ni 2+ + M 1/n+n (AlO 2 , ySiO 2 ), z H 2 O Æ 2+ Ni 1/2 (AlO 2 , ySiO 2 ), z H 2 O + 1/ n M n+

128


According to this mechanism, the Al/Si ratio should have a strong influence on IGSCC. An increase in the Al/Si ratio in the zeolite framework would occur with an increase in the Al/Si ratio in the alumina and silica sources. However, some of the previous tests (see Section 9.3) show that the Al/Si ratio has little effect on cracking in a ‘complex’ environment. One has to consider that the experimental conditions used here are not optimal for the zeolite compound synthesis. IGA and IGSCC also involve an underlying hydroxide gel rich in chromium, which is brittle and non-protective. Another mechanistic hypothesis considers that the dissolution of Ni could be increased by the cyano coordination of Ni2+, according to the following mechanism: ∑

amid formation: CH3COOH + NH3 Æ CH3COO– NH +4

∑

amid deshydratation (probably via the formation of P2O5 from phosphates): CH3COO– NH +4 Æ CH3CN + 2H2O

∑

Ni dissolution by formation of the complex: [NiII(CN)4]2–

Indeed, IGSCC was detected in an AVT, phosphate and acetic acid environment, i.e. without any alumino-silicate deposits. However, some of the previous tests (see Section 9.3) show that phosphate is not necessary to obtain cracks.

9.5

Conclusions

The effects of several parameters such as chemical conditions, potential value and temperature, on the slow propagation step for alloy 600 in a ‘complex’ environment (T = 320 ∞C) have been investigated in detail. The ‘complex’ environment succeeds in obtaining a cracking at the approximate rate of 0.02 mm/h, similar to field experience. Based on the mechanistic hypotheses detailed in Section 9.3, further investigations are necessary to verify the detrimental effect of aluminosilicate zeolite compounds and cyano ligands on IGSCC of alloy 600. Indeed, the results obtained in this study do not allow discrimination with confidence between the two hypothesis Finally, it seems to be impossible, with the results obtained in this study, to model the slow propagation step of alloy 600 in the ‘complex’ environment. Indeed, the effects of the previous parameters (in particular the chemical conditions) are not definite enough to allow the modeling of IGSCC velocity for alloy 600 in the reference ‘complex’ environment.


9.6

129

References

1. Vaillant F., Prieux B., Cattant F., Stutzmann A., Lemaire P., ‘Saint-Laurent B1: expertises, chimie en fonctionnement et essais de corrosion’, Contribution of Materials Investigation to the Resolution of Problems Encountered in Pressurized Water Reactors, International Symposium, Fontevraud III, 12–16 September 1994, SFEN, pp. 383– 393. 2. Ollar P., Viricel-Honorez L., ‘Better understanding flow-restricted environments from hideout return analyses’, Contribution of Materials Investigation to the Resolution of Problems Encountered in Pressurized Water Reactors, International Symposium, Fontevraud IV, 14–18 September 1998, SFEN, pp. 465–476. 3. Cattant F., Dupin M., Sala B., Gelpi A., ‘Analysis of deposits and underlying surfaces on the secondary side of pulled tubes from a French plant’, Contribution of Materials Investigation to the Resolution of Problems Encountered in Pressurized Water Reactors, International Symposium, Fontevraud III, 12–16 September 1994, SFEN, pp. 469– 480. 4. Sala B. Henry K., Lancha A.M., Dupin M., Combrade P., Erre R., Gelpi A., ‘Analysis of the deposits and surface layers on tubes pulled from PWR French steam generators’, Proc. Eurocorr’96, Nice (France), September 1996, paper IX OR 14. 5. Sala B., Gelpi A., Chevallier S., Dupin M., ‘Complementary investigations concerning the analysis of the deposits and underlying surfaces observed on French PWR steam generator pulled tubes’, Proc. Intern. Symp. Fontevraud IV, September 1998, SFEN, p. 553. 6. de Bouvier O., Vaillant F., Millet L., Scott P. M., Tran Q. T. ‘Duplication in laboratory of deposits, films and IGA/SCC damage observed on pulled steam generator tubes’, Contribution of Materials Investigation to the Resolution of Problems Encountered in Pressurized Water Reactors, International Symposium Fontevraud V, 23–27 September 2002, SFEN, p. 1049. 7. Tran T., Scott P., Vaillant F., ‘IGA/SCC of Alloy 600 in complex mixtures of impurities’, Proceedings of the tenth International conference on environmental degradation of materials in nuclear power systems – Water reactors, South Lake Tahoe (NV), USA, August 5–9, 2001. 8. Vaillant F., Pavageau E. M., Bouchacourt M., Boursier J M., Lemaire P, ‘Modeling the secondary side corrosion of tubings: a help to the maintenance policy of PWR steam generators’, Proceedings of the ninth International conference on environmental degradation of materials in nuclear power systems – Water reactors, Newport Beach (CA), USA, August 1–5, 1999, p. 673. 9. Pavageau E. M., Vaillant F., de Bouvier O., Bouchacourt M., Caire J. P., Dalard F., ‘Secondary side corrosion modeling of alloy 600 for steam generator tubes based on laboratory tests in sulfate environments‘, Proceedings of the tenth International conference on environmental degradation of materials in nuclear power systems – Water reactors, South Lake Tahoe (NV), USA, August 5–9, 2001. 10. Santarini G., ‘Comprehensive interpretation of CERTs: a method for the characterization and the prediction of IGSCC’, Corrosion, vol. 45, no. 5, p. 369.

10 The effect of sulphate and chloride transients on the environmentally-assisted cracking behaviour of low-alloy RPV steels under simulated BWR conditions S. R I T T E R and H. P. S E I F E R T, Paul Scherrer Institute (PSI), Switzerland

10.1

Introduction

The ageing of light water reactor structural materials has been one of the major factors affecting the availability, cost and safety of nuclear power. Along with fatigue and irradiation embrittlement, environmentally-assisted cracking (EAC) is a potential RPV ageing mechanism. The EAC behaviour of low-alloy steel (LAS) pressure boundary components under transientfree, steady-state boiling water reactor (BWR) power operation conditions, which covers the largest part of the lifetime, is well established [1–5]. These investigations have shown very low susceptibility to SCC crack growth under static loading conditions in oxygenated, high-purity, high-temperature water/BWR environments at temperatures around 288 ∞C and also formed the basis for the definition of the BWRVIP-60 SCC disposition lines for SCC crack growth in LAS during BWR power operation (Fig. 10.1, [3]). On the other hand, the EAC behaviour during and after water chemistry transients has hardly been investigated so far, but is also of great practical relevance because BWR operation inevitably involves periodic short-term variations in water chemistry and oxygen/corrosion potential (ECP). Conductivity and oxygen/ECP transients occur during start-up/shut-down and occasionally during steady-state power operation (ion exchanger resin intrusions, condenser leakages, etc.). Nowadays, the extent (magnitude, period) and frequency of such transients is strongly minimised by following the current EPRI BWR water chemistry guidelines (Table 10.1, [6]). The possible effect of relatively short-term water chemistry transients on the transition EAC crack growth behaviour during and, in particular, after a transient under subsequent steady-state power operation is of great interest for safety assessments. Effects of oxygen/ECP transients on EAC crack growth in LAS have already been widely investigated in the context of hydrogen water chemistry/noble metal chemical addition (NMCA) by Andresen et al. in tests with periodical partial unloading (PPU) [7]. There is an obvious lack of qualified and well documented testing with sulphate and chloride 130

The effect of sulphate and chloride transients

131

10–8

da /dtscc (m/s)

BWRVIP-60 SCC DL 1: Stationary power operation BWRVIP-60 SCC DL 2: During and 100 h after transients of water chemistry and load

10–9

10–10

10–11

0

20

40 60 80 Stress intensity factor Kl (MPa · m1/2)

100

10.1 BWRVIP-60 SCC disposition lines [3] for SCC crack growth in LAS. Line 1: Stationary, transient-free BWR power operation (< EPRI action level 1), line 2: During and 100 h after water chemistry transients (> EPRI action level 1). Table 10.1 EPRI water chemistry guidelines for reactor water during BWR/NWC power operation [6] Control Parameter

Action level 1

Action level 2

Action level 3

Conductivity [mS/cm] Sulphate [ppb] Chloride [ppb]

> 0.3 >5 >5

> 1.0 > 20 > 20

> 5.0 > 100 > 100

transients under highly oxidising BWR/normal water chemistry (NWC) conditions at high ECP, where the most distinct and severe long-term/hysteresis effects might be expected. Therefore, the EAC crack growth behaviour of three different low-alloy RPV steels during and after sulphate and chloride transients was investigated under simulated BWR/NWC power operation conditions by tests with PPU and experiments under constant load. These tests should indicate, if hysteresis or long-term effects might occur under these highly oxidising conditions and reveal information on the transition behaviour during and after such a transient (response times, incubation periods, delay times, acceleration of crack growth, etc.). Furthermore, the adequacy and conservative character of the BWRVIP60 SCC disposition lines (Fig. 10.1, [3]) for SCC crack growth in LAS during and after water chemistry transients was evaluated and assessed in the context of the current EPRI BWR water chemistry guidelines (Table 10.1).

132

10.2


Experimental procedure

10.2.1 Materials Three different types of low-alloy, nuclear grade RPV steels with either a low or high sulphur content were investigated (Tables 10.2 and 10.3). The RPV steels had a granular, bainitic microstructure with an average former austenitic grain size of 10 to 20 mm. The spatial distribution and morphology of the MnS-inclusions was fairly homogenous and similar in alloys B and C covering the range from small, spherical to large (up to a few 100 mm), elongated inclusions. In alloy A distinct sulphur segregation zones with large clusters of MnS inclusions were observed. The local sulphur content therefore significantly deviated from the average bulk sulphur content of 0.015 wt.% and varied between 0.003 and 0.053 wt.%.

10.2.2 Specimens 25 mm thick compact tension specimens (1T-C(T)) were used for all experiments. They were manufactured in the T-L or L-T orientation. The specimens were pre-cracked by fatigue in air at room temperature, using a load ratio R of 0.1. The maximal KI at the final load step was £ 15 MPa · m1/2.

10.2.3 Environmental parameters The tests were conducted in modern high-temperature water loops [8]. Water chemistry (oxygen content and conductivity) and flow rate were measured at the autoclave inlet and outlet. Inside the autoclave pressure and temperature were measured. The stainless steel autoclave volume of 10 litres was exchanged three to four times per hour. In the vicinity of the specimens a flow velocity in the range of mm/s was obtained. The concentration of dissolved oxygen (DO) was adjusted by adding an argon-oxygen mixture to the storage tank. After the demineralised water in the storage tank was purified by ion exchangers, active coal and microfilters, the conductivity was controlled by dosing 0.02 M Na2SO4 or NaCl to the high-purity (£ 0.06 mS/cm) water. Conductivity and concentration of DO were controlled at the inlet water. Ionic impurities of the water (grab samples at inlet and outlet) were analysed by Inductive Coupled Plasma – Atomic Emission Spectroscopy (ICP – AES) and Ion Chromatography (IC) several times during each test. The ECP of the specimens and the redox potential (platinum probe) were continuously monitored by use of an external Ag/AgCl/0.01 M KCl reference electrode. The specimens were electrically insulated from the autoclave and from each other by ZrO2 spacers.

Table 10.2 Chemical composition (in wt.%) and heat treatment of the investigated low-alloy RPV steels (WQ = water quenched, FC = furnace cooled, AC = air cooled, PWHT = post weld heat treatment Alloy

C

Si

Mn

P

S

Ni

Cr

V

Mo

Al

N

Heat treatment

A

0.25

0.33

1.54

0.014

0.015*

0.62

0.18

0.024

0.68

0.021

0.004

900 ∞C/9h/WQ, 650 ∞C/34h/AC, 660 ∞C/14h/AC + PWHT

22 NiMoCr 3 7 (∫ SA 508 Cl.2)

B

0.22

0.20

0.91

0.008

0.007

0.88

0.42

0.010

0.53

0.018

0.008

895 ∞C/7h/WC, 645 ∞C/17h/AC + PWHT

SA 533 B Cl.1 (∫ 20 MnMoNi 5 5)

C

0.25

0.24

1.42

0.006

0.018

0.62

0.12

0.007

0.54

0.030

0.006

915 ∞C/12h/AC, 860 ∞C/12 h/WQ, 660 ∞C/12 h/FC + PWHT

*local sulphur content: 0.003–0.053 wt.%)

Table 10.3 Mechanical properties (tensile tests in air, DIN 50125, B5 ¥ 50-specimens, Rp0.2 = yield stress, Rm = tensile strength, A5 = elongation at fracture, Z = reduction of area) 288 ∞C

Room temperature Alloy

Rp0.2 [MPa]

Rm [MPa]

A5 [%]

Z [%]

Rp0.2 [MPa]

Rm [MPa]

A5 [%]

Z [%]

A

512

663

19

56

462

618

17

51

22 NiMoCr 3 7

B

467

605

17

72

400

578

16

70

SA 533 B Cl.1

C

456

618

23

60

412

588

21

55

133

20 MnMoNi 5 5


20 MnMoNi 5 5 (∫ SA 508 Cl.3)

134


10.2.4 Crack growth monitoring and fractographical posttest evaluation Crack advance was continuously monitored using the reversed direct current potential drop (DCPD) method with a resolution limit corresponding to roughly 5 mm [8]. The crack growth increment was calculated by the Johnson formula [9]. The calculated crack length at the end of the experiment was then verified and, if necessary, corrected with regard to the mean final crack length as revealed by post-test fractography [8]. The crack growth rates (CGR) were determined by linear fit of the crack increment versus time curve. After the test, the specimens were broken apart at liquid nitrogen temperature for post-test evaluation. For fractographical analysis in the scanning electron microscope, the oxide film on the fracture surface of one specimen half was removed by galvanostatic reduction in an ENDOX-bath [10].

10.2.5 Mechanical loading Two pre-cracked specimens were investigated simultaneously under the test conditions in oxygenated high-temperature water in a daisy chain. The load was actuated with a screw-driven, electro-mechanical tensile machine with computer control. The KI values were calculated according to ASTM E 399 by the measured load and by the actual mean crack length , derived by the DCPD method and by post-test fractographical evaluation [8].

10.2.6 Test procedure of the sulphate transient experiment (Test 1) In the first test, the effect of a sulphate transient on an actively growing EAC crack in two low-alloy RPV steels with different sulphur contents (alloy A and B) under low-flow and highly oxidising BWR conditions was investigated under PPU (constant load amplitude loading with trapezoid waveform) conditions. The four major experimental phases of this transient test are shown in Fig. 10.2. After achieving the desired environmental conditions, the specimens were pre-oxidised in the test environment (8 ppm DO, kinlet = 0.06 mS/cm) under a small mechanical pre-load. Before applying the sulphate transient, an EAC CGR in the range of the ‘low-sulphur SCC line’ of the GE model [11] was generated in the initially high-purity water by PPU (asymmetrical trapezoid loading at high load ratio R of 0.8 with a rise time DtR of 1000 s and long hold time at maximum load DtH of 5 h). The testing phase in oxygenated, high-temperature water at 288 ∞C consisted on a sequence of three different water chemistry conditions WC 1 to 3 with different sulphate concentrations (Fig. 10.2 and Table 10.4). After 170 h in high-purity water,

The effect of sulphate and chloride transients DtR

Load

2–

Test 1 (368 ppb SO4 ) Test 2 (50 ppb Cl–) Load

135

DtR

Test 3 (20 ppb Cl–) DtR

k (Na2SO4 or NaCl)

(CO2)

O2 (ECP) T Time Pre- Heating Conditioning test phase phase phase

Testing phase

Cooling phase

10.2 Simplified schematic of the test procedures of tests 1 to 3 with the major experimental phases.

the sulphate level was increased within 1 h from < 0.6 to 368 ppb by dosing 0.02 M Na2SO4. The sulphate level was then kept constant for further 310 h (WC 2). Afterwards, the sulphate level was decreased within 2.6 h to a value of < 0.6 ppb and the high-purity water chemistry conditions (WC 3) were maintained for further 240 h before the specimens were unloaded.

10.2.7 Test procedure of the chloride transient experiments (Tests 2 and 3) The effect of chloride transients on the EAC behaviour of three low-alloy RPV steels with different sulphur contents under low-flow and highly oxidising BWR conditions was investigated by two different tests under PPU and constant load conditions. Test 2 (50 ppb chloride, PPU) In this test the effect of a chloride transient on an actively growing EAC crack in two low-alloy RPV steels with different sulphur contents (alloy A and B) under low-flow and highly oxidising BWR conditions was investigated under PPU conditions. The four major experimental phases of this transient test were similar to the sulphate transient test (see above) and are shown in Fig. 10.2. Instead of sulphate, chloride was added to the high-purity water for 40 h (WC 2). The chloride level was increased within 1 h from < 0.4 to

136


50 ppb by dosing 0.02 M NaCl. Afterwards, the chloride level was decreased within 1.6 h to a value of < 0.4 ppb and the high-purity water chemistry conditions (WC 3) were maintained for further 170 h before the specimens were unloaded. Test 3 (20 ppb chloride, constant load) The EAC behaviour of alloy C and B during and after a chloride transient was investigated in a third test under constant load. In this test 20 ppb chloride was added for 290 h under pure constant load under otherwise similar environmental conditions as in test 1 and 2. After 240 h of NaCl dosing a single partial unloading with a load ratio R of 0.7 was performed (DtFall = 200 s, DtRise = 1000 s). Then the load was kept constant for further 1130 h.

10.3

Results and discussion

10.3.1 Effects of a sulphate transient on the EAC behaviour The average EAC CGR of the two specimens (alloy A and B) during the water chemistry phases WC 1 to 3 are summarised in Table 10.4. In material A no effect of the sulphate transient on the EAC CGR was observed. The EAC CGR before, during and after the sulphate transient were almost identical (Table 10.4). The very small increase of the CGR from WC 1 to 3 was primarily the result of the increasing KI,max value. The PPU (R = 0.8, KI,max = 69–85 MPa·m1/2, DtR = 1000 s, DtH = 5 h) under these highly oxidising conditions resulted in stable EAC crack growth in the range of the BWRVIP60 SCC disposition line 2. During the PPU at high KI,max values it was possible to resolve the crack growth in each trapezoid cycle and to qualitatively differentiate between the EAC crack growth during the slow rising load and the constant load part of the cycle (Fig. 10.3). The crack mainly grew during the rising load phase with some very minor crack advance (and very low CGR) during the subsequent constant load phase. The crack arrested either during the constant load or during the unloading part of the cycle and reinitiated again during the rising load part of the next cycle. In accordance to the GE model [11], the absence of an acceleration of EAC crack growth may be attributed to ‘high-sulphur’ crack chemistry conditions, which already existed during the high-purity water chemistry phase WC 1 because of the high ECP and the dissolution of MnS-inclusions in the enclave of the incipient crack. Therefore, an increased bulk sulphate concentration did not result in an acceleration of the crack growth. In material B continuous cessation of EAC crack growth was observed in the high-purity water chemistry phase WC 1. The crack was growing with a

Alloy A

Alloy B

Test phase

WC 1

WC 2

WC 3

WC 1

WC 2

WC 3

<ECP> [mVSHE]

+110

+70

+115

+130

+100

+120

0.07

1.00

0.06

0.07

1.00

0.06

< 0.6

368

< 0.6

< 0.6

368

< 0.6

KI,max [MPa·m ]

68.8–70.9

70.9–78.0

78.0–84.7

61.9

62.0

62.0–65.9

[m/s]

8.1 ¥ 10–10

1.4 ¥ 10–9

1.6 ¥ 10–9

1.4 ¥ 10–11

1.4 ¥ 10–11

1.6 ¥ 10–9

inlet

[mS/cm] [ppb] 1/2


Table 10.4 Summary of the most important environmental parameters and the mean CGR during the three water chemistry phases of test 1 (WC 1–3). WC 2 = sulphate transient

137

138

Corrosion issues in light water reactors 25.06 Load 42

40

25.02

38

Load (kN)

Crack length (mm)

25.04

36 25.00 A: 20 MnMoNi 5 5, 0.015 wt.% S T = 288 ∞C, DO = 8 ppm, ECP = 71 mVSHE WC 2: k = 1.00 mS/cm, 368 ppb SO42–

24.98 710

715

720 Time (h)

725

34

32

10.3 Crack length of alloy A specimen during the sulphate transient. Test 1.

CGR just slightly above the detection limit of the DCPD ( ª 1.4 ¥ 10–11 m/s) during the water chemistry phases WC 1 and 2. The sulphate transient did not result in an acceleration of the very slow EAC crack growth. 65 h after returning to high-purity water, EAC re-initiated. After re-initiation followed by a short transition period, stable and stationary EAC crack growth with a CGR of 1.6 ¥ 10–9 m/s (51 mm/year) in the range of the ‘low-sulphur SCC crack growth curve’ of the GE model [11] was observed in alloy B. The CGR was almost identical to that in alloy A. It is believed, that the reinitiation of fast EAC in alloy B was not the direct result of a delayed effect of the sulphate transient, but rather a result of the probabilistic nature of the EAC cessation/pinning/arrest phenomena and of the re-initiation process [12]. The same crack growth behaviour has also been observed by MPA Stuttgart in the framework of the CASTOC programme [13] and in many other investigations of PSI [1, 4]. In these experiments under comparable environmental conditions (T = 288 ∞C, DO = 0.4 or 8 ppm, ECP = 0 to +150 mVSHE), all investigated RPV steels with different sulphur contents revealed the same crack growth behaviour with comparable CGR under identical loading conditions in high-purity water and in water with very high levels of sulphate (added as H2SO4 or Na2SO4). Even very high sulphate contents of up to 1400 ppb (ca. 10 mS/cm) did not result in an acceleration of EAC crack growth. Fast growing cracks, triggered by cyclic or slow rising loading arrested immediately after switching to constant load at stress intensity factors of up to 53 MPa·m1/2. In spite of the absence of any accelerating effect under highly oxidising BWR conditions, sulphate still remains a harmful species


139

for EAC in LAS, since it has been observed to affect EAC initiation from smooth surfaces in LCF and SSRT tests and to accelerate EAC crack growth under reducing PWR conditions [5]. Furthermore it may help to overcome crack pinning/arrest phenomena and re-initiation problems.

10.3.2 Effect of chloride on the EAC behaviour Test 2 (50 ppb chloride, PPU, [14]) The mean EAC CGR of the two alloys during the individual water chemistry phases WC 1 to 3 are summarised in Table 10.5. In Fig. 10.4 the crack growth behaviour of material A during the chloride transient is shown. An accelerating effect of the chloride transient on the EAC CGR was observed. The PPU (R = 0.73, KI,max = 51–67 MPa·m1/2, DtR = 1000 s, DtH = 12 h) during the high-purity water chemistry phase WC 1 resulted in stable EAC crack growth slightly above the BWRVIP-60 disposition line 2. In the high-purity water chemistry phases WC 1 and 3 the crack growth behaviour was similar to that during the sulphate transient experiment, as described before (Fig. 10.3). 3 h after adding NaCl to the high-purity water, onset of fast EAC occurred during the constant load phase of a PPU cycle (Fig. 10.5). 17 h later, the EAC crack growth reached a stationary state during the rising (da/dtRL = 1.4 ¥ 10–7 m/s) and constant load (da/dtCL = 1.8 ¥ 10–8 m/s) part of the next PPU cycle with a mean CGR of 2.2 ¥ 10–8 m/s (694 mm/year) at stress intensity factor values KI of 55 to 62 MPa·m1/2. After stopping the NaCl dosage, the crack was further growing with the same high CGR under constant load until the next partial unloading (Fig. 10.4). Then the crack growth started to slow down to a mean CGR in the same range (and with the same behaviour: see Fig. 10.3) as observed before the chloride transient. The slightly higher CGR was assigned to the higher KI,max values. The decay of the crack growth after the chloride transient might have been caused by crack closure effects during the partial unloading. Therefore, a possible long-term effect of a chloride transient on the EAC crack growth behaviour under constant load could not be fully excluded based on this experiment. According to the EPRI BWR/NWC water chemistry guidelines (Table 1), the chloride concentration has to be reduced below the action level 2 within 24 h from the time of occurrence by adequate correction actions. Otherwise an orderly shut-down shall be initiated. If it is foreseeable that the parameter will be below the action level 2 value within the time period required to achieve an orderly shut-down, power operation can be maintained. The experimentally observed short incubation period of 3 h for acceleration of EAC is therefore significantly shorter than the maximum allowable time interval for returning to normal operating conditions. The short incubation period for acceleration of EAC in combination with the very high SCC CGR

140

Alloy A

Alloy B

Test phase

WC 1

WC 2

WC 3

WC 1

WC 2

WC 3

<ECP> [mVSHE]

+105

+110

+125

+130

+125

+150

0.06

0.23

0.06

0.06

0.23

0.06

< 0.4

50

< 0.4

< 0.4

50

< 0.4

KI, max [MPa·m ]

51.5–52.9

54.2–62.1

64.8–67.2

49.5–49.6

49.6–49.8

49.8–50.4

[m/s]

1.2 ¥ 10–9

2.2 ¥ 10–8

1.8 ¥ 10–9

2.7 ¥ 10–11

2.5 ¥ 10–9

4.6 ¥ 10–10*

inlet

[mS/cm]

Cl– [ppb] 1/2


Table 10.5 Summary of the most important environmental parameters and the mean CGR during the three water chemistry phases of test 2 (WC 1–3). WC 2 = chloride transient

The effect of sulphate and chloride transients WC 1 (< 0.4 ppb Cl–)

28.0

WC 2 (50 ppb Cl–)

WC 3 (< 0.4 ppb Cl–)

141

0.25

Inlet

m/s 0.15

26.0 0.10 25.0

24.0

Kl = 52 – 67 MPa.m 1.2 ¥ 10–9 m/s

400

1/2

A: 20 MnMoNi 5.5, 0.015 wt.% S T = 288 ∞C, DO = 8 ppm, ECP = 110 mVSHE

450

500 Time (h)

550

Conductivity (mS/cm)

Crack length (mm)

2.2 ¥

10–8

0.20

1.8 ¥ 10–9 m/s

27.0

0.05

0.00

600

10.4 Crack length of alloy A specimen with mean EAC CGR (linear fit) before during and after the chloride transient. Test 2.

24.4

WC 1 (< 0.4 ppb Cl–)

WC 2 (50 ppb Cl–) Load

32

30 24.2

Onset of fast EAC under constant load

28

24.1 26

Load (kN)

Crack length (mm)

24.3

24.0 A: 20 MnMoNi 5.5, 0.015 wt.% S T = 288 ∞C, DO = 8 ppm, ECP = 110 mVSHE k = 0.06/0.23 mS/cm, < 0.4/50 ppb Cl–

23.9 470

475

480 485 Time (h)

490

24 22 495

10.5 Crack length of alloy A specimen at the beginning of the chloride transient. Test 2.

under constant load of almost 1 mm/day arise some safety concern for severe chloride transients under BWR/NWC conditions. Test 3 (20 ppb chloride, constant load) The mean EAC CGR of the investigated alloys C and B during the individual water chemistry phases WC 1 to 3 are summarised in Table 10.6. In both

142

Alloy C Test phase

WC 1

Alloy B WC 2

WC 3

WC 1

WC 2

WC 3

<ECP> [mVSHE]

+140

+140

+140

+145

+145

+145

[mS/cm]

0.06

0.13

0.06

0.06

0.13

0.06

Cl– [ppb]

0.2 mm Da > 0.2 mm 100

12.3 Range of crack growth rates of material A at a frequency of 8.3 ¥ 10–4 s–1 in oxygenated high purity BWR water at 288 ∞C.

one order of magnitude is typically observed in this loading frequency range for RPV steels based on well-behaved crack growth of homogeneous material in a similar environment [9, 12] and all data obtained in the present study for larger crack extensions (Da ≥ 0.2 mm) also falls within this with the exception of one single data point. This ‘outlier’ was related to a highly uneven crack front and a repetition of the test in the corresponding lab revealed a cyclebased crack growth rate da/dN in the expected range. Most of the data obtained in the CASTOC project with low crack growth rates were related either to load drops because of electric power interruptions, highly uneven crack front because of crack pinning, or short crack advances (Da < 0.2 mm) and were believed to be non-representative for this material. EAC growth in simulated BWR environment is affected by the stress intensity, type of loading, the ECP related to the dissolved oxygen concentration, – 2– – the concentration of specific anionic impurities ( SO 2– 4 , Cl , H2S, S , HS , etc.) reflected by the conductivity of the water, the flow rate passing the crack mouth, the steel sulphur content, and the morphology/spatial distribution of the MnS-inclusions [12]. Variations of these parameters can result indifferent crack growth rates. Apart from material properties, all of these factors were fairly similar in all labs. A relevant part of the large data range may therefore be related to the inhomogeneous microstructure in this material: e.g., the

172


investigated specimens showed a strong variation in the amount of MnS inclusions and local islands of intergranular fracture, indicating local segregation. Additionally the nature of the EAC cracking process itself can lead to local crack pinning, crack cessation/arrest and re-nucleation problems, and thus to lower crack growth rates than expected. EAC is a deterministic process, but singular results are influenced by the previously mentioned parameters, thus adding a probabilistic factor. Furthermore, crack growth data, which were derived from test periods with not long enough crack advance (Da < 0.2 mm) with regard to the microstrucutre (e.g., grain size) and to the resolution of fractographic crack length determination and DCPD technique for crack length measurement, in particular in the lower range of da/dN, are associated with a higher uncertainty and thus are less reliable. This has to be considered with regard to practical applications. Several data points, from several specimens representing several locations in a LAS material are needed before representative and conservative crack growth rate values can be derived with satisfactory confidence. From the results of the present work it was concluded, that a large enough (Da ≥ 0.2 mm) crack advance is needed following any type of anomalous testing event before appropriate cyclic crack growth rate data are measured. If not further specified, the following cyclic crack growth rate diagrams therefore only contain data points with crack advances ≥ 0.2 mm. The data were compared with the prediction line as presented in the ASME Boiler and Pressure Vessel Code, Section XI, Appendix A [10]. This comparison clearly reveals that a significant number of data, in particular those which were derived from test periods with sufficient crack advance, exceed the ASME prediction line. It can further be concluded that the lower cycling frequency (8.3 ¥ 10–4 s–1, Fig. 12.3) tends to cause higher crack growth rates, than the higher frequency (8.3 ¥ 10–3 s–1, Fig. 12.2). This is consistent with the general experience that EAC is a strong time dependent process.

12.3.2 Results from cyclic loading and load transients Effect of material Material A was investigated extensively and showed consistently, that most of the da/dN data exceed the ASME XI prediction curve in high purity water (Figs 12.2 and 12.3). Material B, representing optimised nuclear grade, showed significantly lower susceptibility to EAC and in most cases it was difficult to initiate fast EAC crack growth under cyclic load. Therefore, the range of da/ dN data were generally almost one order of magnitude lower than that of material A (Fig. 12.4), but if once fast crack growth was initiated, the crack growth rates of material B, e.g. during slow load transients, were in the same ranges as those of material A [6]. The lower susceptibility to EAC appeared

Crack growth behaviour of low-alloy steels

173

102

Crack growth rate da /dN (mm/cycle)

ASME XI ‘Wet’, R = 0.2 ASME XI ‘Wet’, R = 0.8 BWR water, 240 ∞C 400 mg/kg O2

101

f = 2.2¥10–4 s–1 R = 0.2

100

f = 2.5¥10–3 s–1 R = 0.8

10–1

10–2

10–3

Material A Material B Material B with Da < 0.2 mm 1

10 DKl (MPa÷m)

100

12.4 Comparison of crack growth rates of material A and B under low frequency fatigue in oxygenated high purity BWR water at 240 ∞C.

also in tests with PPU and during the applied chloride transient (compare pages 178–9) and was attributed to the more homogeneous microstructure, lower sulphur content and lower susceptibility to dynamic strain ageing (DSA) in this material. The weld metal and the HAZ of the joint weld of material A showed similar behaviour as the base material A (Fig. 12.5), although due to the limited number of tests and too short crack advances (Da < 0.2 mm) in some cases no clear conclusion for the cyclic crack growth behaviour could be established in the framework of the CASTOC project. Effect of oxygen content The effect of dissolved oxygen content was investigated on material A in high purity BWR water. No significant effect of oxygen on EAC was observed at concentrations between 400 and 8000 mg/kg (Fig. 12.6), which is consistent with the only small differences in the measured ECP values. Effect of temperature In BWR water environment experiments were carried out with material A and B at temperatures of 180 and 240 ∞C in addition to those at 288 ∞C. An

Corrosion issues in light water reactors 102


ASME XI ‘Air’ ASME XI ‘Wet’ 101

BWR water, 288 ∞C 400 mg/kg O2

f = 8.3 ¥ 10–3 s–1 R = 0.7 100

10–1

10–2

0.65 £ R < 1

R = 0.1 10–3 1

Base material A HAZ Weld metal Weld metal with Da < 0.2 mm

10 DKl (MPa÷m)

100

12.5 Crack growth behaviour of weld metal and HAZ compared to data for base material A under low frequency fatigue in oxygenated high purity BWR water at 288 ∞C. 102 ASME XI ‘Air’ ASME XI ‘Wet’


174

101

BWR water, 288 ∞C Material A f = 8.3 ¥ 10–4 s–1 R = 0.7 (0.8)

100

10–1

10–2 0.65 £ R < 1

R = 0.1 10–3 1

10 DKl (MPa÷m)

400 mg/kg O2 8000 mg/kg O2 100

12.6 Effect of oxygen on crack growth rate under low frequency fatigue in BWR water at 288 ∞C.


175

example is presented in Fig. 12.7 for material A, whereas material B showed a similar behaviour. In this temperature range, no clear temperature trend on the crack growth behaviour was observed due to the wide range of data, the limited number of tests, and too short crack advances (Da < 0.2 mm) in some cases. However, literature data in BWR environment reveal an increase in crack growth rates with increasing temperature, sometimes with a maximum at intermediate temperatures (200 to 250 ∞C) in case of materials which show distinct DSA effects [11, 12]. Effect of load transients and frequency By means of the DCPD measurement technique detailed insight into the crack growth behaviour as a function of time was obtained. Tests performed with PPU revealed very clearly that crack advance in high purity BWR water occurs only in the phases of rising load (e.g., Fig. 12.8). Figure 12.9 shows an example of PPU tests in which the hold time was varied from 0 to 105 s (0 to 28 h). As long as EAC under static load during the hold time period at maximum load does not occur, the low frequency corrosion fatigue (LFCF) crack growth rate da/dN depends only on the number of loading events and thus the curve in Fig. 12.9 must result in a horizontal line. Differences appear in the crack growth rate due to different applied load ratios (R = 0.2 and 0.8) and different rise times to maximum load (100, 1000, and 4000 s). With the longer rise times (1000/4000 s), the load increase is slower, causing a lower strain rate in the crack tip region and hence the contribution of EAC is more pronounced and results in a higher crack growth rate da/dN. This behaviour is consistent with the basic mechanism of straininduced corrosion cracking (SICC). In Fig. 12.10 cycle-based crack growth rates of material A are plotted versus the loading frequency and compared to the corresponding ASME XI prediction curves. The crack advance per fatigue cycle da/dN increases with decreasing frequency, whereas sustained, stationary LFCF crack growth was observed down to very low frequencies of 10–5 s–1. The ASME XI prediction curves are significantly exceeded for all investigated loading frequencies under low-flow and oxidising conditions at 240 ∞C. Effect of sulphate The effect of sulphate in BWR water, added as Na2SO4 and H2SO4 respectively, was predominantly investigated on material A. Although the sulphate content was realised up to values far beyond the onset of Action Level 3 of the EPRI Water Chemistry Guidelines according to the sulphate content of up to 1400 mg/kg, no enhanced crack growth rates were observed (Fig. 12.11). This is in good agreement with model predictions and most investigations under BWR/

Corrosion issues in light water reactors 102


ASME XI ‘Air’ ASME XI ‘Wet’ BWR water 400 mg/kg O2

101

Material A f = 8.3 ¥ 10–4 s–1 R = 0.7

100

10–1

10–2 0.65 £ R < 1

R = 0.1 10–3 1

288 ∞C 240 ∞C 180 ∞C 180 ∞C with Da < 0.2 mm 100

10 DKl (MPa÷m)

12.7 Effect of temperature on crack growth rate under low frequency fatigue in oxygenated high-purity BWR water.

23.9

Load

32

da /dt = 1.22¥10–9 m/s

30

23.8 28 DCPD 23.7

26

23.6

24 Material A BWR water, 288 ∞C, 8000 mg/kg O2

23.5

Load (kN)

24.0

Crack length (mm)

176

390

400

410 420 Time (h)

430

22 440

12.8 Effect of load transients (PPU) on the crack advance during the load transient and during the hold time period; material A, high purity BWR water at 288 ∞C.


177

103


Material A, BWR water, 240 ∞C, 400 mg/kg O2 102

58.0

56.0

60.1

DtH = 0 h

101

61.5

58.8

57.5

60.4

57.0 55.0

100

DtH

P 10–1

59.5

58.2

59.8

DtR

R = 0.8, DtR = 1000 s R = 0.2, DtR = 4000 s R = 0.7, DtR = 100 s

PPU DtR = 100/1000/4000s, variation of DtR

10–2

0.01

t

0.1 1 10 Hold time at maximum constant load DtH (h)

12.9 Effect of hold time at maximum load during PPU on crack advance per reloading event; numbers in the diagram indicate the maximum KI during the hold time.

Material A, BWR water, 240 ∞C, 400 mg/kg O2


100

ASME XI ‘Wet’ R = 0.2, DK = 43.4 MPa÷m

10

1

ASME XI ‘Wet’ R = 0.8, DK = 12.5 MPa÷m

R = 0.2 R = 0.8 0.1 10–6

10–5

10–4

10–3 10–2 10–1 Frequency (s–1)

100

101

12.10 Effect of loading frequency on crack growth rate in oxygenated high purity BWR water at 240 ∞C.

178

Corrosion issues in light water reactors 102 ASME XI ‘Air’ ASME XI ‘Wet’


101

BWR water, 288 ∞C 400 mg/kg O2 Material A f = 8.3 ¥ 10–4 s–1 R = 0.7

100

R = 0.1

10–1

10–2

High purity water Sulphate Sulphate with Da < 0.2 mm

0.65 £ R £ 1

10–3 1

10 DKl (MPa÷m)

100

12.11 Effect of sulphate content in BWR water at 288 ∞C on the crack growth rate of material A under low-frequency fatigue, sulphate content up to 150 mg/kg.

NWC conditions known from literature [12]. In some few other experiments under BWR/NWC conditions, an accelerating effect of sulphate was observed, in particular in LFCF tests with low-sulphur steels at lower corrosion potentials, where crack growth rates in high purity water were close to those measured in air [12]. In the CASTOC experiments, however, the effect of sulphate addition was shaded because the investigated material has relatively high sulphur content, so that the dissolved manganese sulphides intersected by the plane of the growing crack strongly control the local water chemistry. The ranges of crack growth rate data without and with additional sulphate overlap completely and there is no tendency that crack growth data corresponding to EPRI Action Level 3 are higher than those corresponding to EPRI Action Levels 1 or 2. Effect of chloride To screen the effect of chloride on the crack growth behaviour of LAS a test with PPU with a hold time of 12 h at maximum load was performed introducing


179

a chloride transient of about 50 mg/kg for a time period of 40 h. According to the chloride content, this water condition corresponds to Action Level 2 of the EPRI Water Chemistry Guidelines, whereas the electrical conductivity which is monitored continuously in the plant, resulted in a value of about 0.27 mS/cm, which is still below the EPRI Action Level 1 limit. Starting in BWR water of high purity but with enhanced oxygen content of about 8000 mg/kg, crack advance was observed typical for PPU tests, i.e., crack growth occurred only in phases of rising load (compare page 175 and Fig. 12.8). After an incubation time of about 3 h, onset of fast crack growth under static load (SCC) occurred in the specimen of material A (Fig. 12.12). The specimen of material B, which was in daisy chain with the specimen of material A in the same experiment, exhibited retarded crack initiation only after 32 h, i.e. after two reloading events in chloride containing water. The detailed analysis of the DCPD signal showed that the chloride transient affected both, the crack growth behaviour under static load da/dt and the crack growth rate under cyclic load da/dN based on the number of reloading events (Fig. 12.13). The crack advance Da during reloading, however, was only slightly increased. This indicates that the mechanism of SICC which is acting during the reloading phase is not much affected by the presence of chlorides. The high value for the cycle-based crack growth rate da/dN has mainly to be attributed to the mechanism of SCC which occurs under static load at the level of the maximum stress intensity factor KImax. After returning to high purity water, crack growth in the phase of constant static load was still observed, however, with strongly decreasing tendency over two more load cycles. This observation indicates that a long-term effect (‘memory effect’) of a chloride transient cannot be fully excluded. Similar behaviour was observed for material B but as already mentioned with retarded crack initiation, lower crack growth rate under static load and less crack advances during reloading.

12.3.3 Results from static loading Based on well-qualified data for simulated BWR environment [13, 14], disposition lines for SCC crack growth in LAS during BWR power operation were proposed by an international group of experts, working within the framework of the EPRI BWRVIP Project, and accepted by the US Nuclear Regulatory Commission as an interim position [13]. The BWRVIP-60 SCC Disposition Line (DL) 1 applies to crack growth in LAS under static loading and transient-free, stationary BWR/NWC or hydrogen water chemistry power operation conditions, whereas the BWRVIP-60 SCC DL 2 may be used for estimating SCC crack growth during and 100 h after transients in water chemistry (> EPRI Action Level 1 limit) or load transients not covered by

180

Corrosion issues in light water reactors < 0.4 mg/kg Cl–

50 mg/kg Cl–

24.4 Load

32

30 24.2

Onset of fast EAC under static load

28

24.1

Load (kN)

Crack length (mm)

24.3

26 24.0

DCPD

24

Material A 23.9 BWR water, 288 ∞C, 8000 mg/kg O2 470

475

480 485 Time (h)

490

22 495

12.12 Crack initiation in a C(T)25 specimen of material A during the phase of static load due to a chloride transient (50 mg/kg) in oxygenated (8000 mg/kg O2) BWR water at 288 ∞C.

da /dt da /dN

Contribution of SCC

1000

2¥10–8

da/dt Decreasing tendency with time 1¥10–8

0

High purity phase 1

da /dN

da /dN

500

da /dt = 0


1500

Material A BWR water, 288 ∞C 8000 mg/kg O2 Kl = 55 to 62 MPa÷m

50 mg/kg chloride phase 2

Crack growth rate da /dt (m/s)

3¥10–8

2000

0 High purity phase 3

12.13 Crack growth behaviour (da/dN and da/dt) of material A before, during and after a chloride transient in oxygenated BWR water at 288 ∞C.

fatigue evaluation procedures. For comparison purposes crack growth rate data of material A obtained from tests in high purity BWR water under constant load are displayed together with these DLs in Fig. 12.14. For stress intensity factors in the range of up to about 60 MPa÷m crack growth was not


181

10–6


BWRMP-60 SCC DL 1 (stationary power operation) BWRMP-60 SCC DL 2 (during and 100 h after transients)

10–7

BWR water, 288 ∞C Material A

‘High-Sulphur line’

10–8 LEFM range 10–9

‘Low-Sulphur line’

all other data under constant load below detection limit < 10–12 to 10–11 m/s

10–10

10–11 20

30

40 50 60 70 80 Stress intensity factor Kl (MPa÷m)

90

100

12.14 Crack growth behaviour of material A under constant static load in high purity BWR water at 288 ∞C, range of plane-strain LEFM validity indicated.

observed in any of the C(T)25 specimens, which were tested in high purity BWR water. These results in which the crack growth rate was below the detection limit of the DCPD method of 10–12 to 10–11 m/s (60 to 300 mm/a) are displayed schematically in Fig. 12.14. When the stress intensity factor was increased far beyond the validity limits for linear elastic fracture mode (LEFM) continuous crack growth was observed even in high purity BWR water obviously due to yielding processes in the crack tip region. Although this plastic deformation occurs under constant static load, the EAC process may mechanistically be attributed to SICC. However, most of the data fall below the DLs. These data in general are not relevant for application to thick-walled pressure boundary components, which are loaded in the LEFM regime during operation, as e.g. the RPV, because the LEFM validity criteria are violated. For thinwalled components, e.g. pipes, a careful assessment of LEFM applicability has to be performed depending on the load and flaw size. The most important result from the investigations under static load was the reproducible observation of immediate crack cessation of growing cracks after changing from cyclic loading to static load for all investigated base materials in oxidising high purity BWR water at stress intensity factors below 60 MPa÷m (e.g., Fig. 12.1). There were only few exceptions in which continuous crack advance was detected under static load. These are described in the following sections.

182


Effect of material As stated above, none of the investigated base materials showed any susceptibility to EAC under constant static load in the relevant stress intensity range and normal BWR water chemistry. This behaviour was also confirmed for the weld metal of the joint weld in material A. A different behaviour was observed in a test with a specimen from the HAZ of the joint weld in material A. Even at a stress intensity factor of KI = 47 MPa÷m, using a C(T)25 specimen, sustained crack growth occurred with a crack growth rate beyond the DL 2 (Fig. 12.15). Based on hardness measurements at room temperature in the base material (235 HV) and the HAZ (340 HV) and also, based on general experience, it may be concluded that the limit for plane-strain LEFM even at elevated temperature is significantly higher for the HAZ as compared to the base material. Therefore, the applied stress intensity factor in this experiment is still supposed to fulfil the LEFM criteria. Effect of environment As already stated (page 175), additional high sulphate content in BWR water did not enhance the crack growth rate under cyclic load. The same is true for static load. In the relevant range of stress intensity, crack cessation occurred in all cases even at a sulphate concentration corresponding to Action Level 3 of the EPRI Water Chemistry Guidelines. 10–6


BWRVIP-60 SCC DL 1 (stationary power operation) BWRVIP-60 SCC DL 2 (during and 100 h after transients)

10–7 BWR water, 288 ∞C

‘High-Sulphur line’ S

10–8

S

E

E ‘Low-Sulphur line’

10–9 S = start of test E = end of test 10–10

10–11 20

HAZ, high purity water Material A, chloride transient Material B, chloride transient 30

40 50 60 70 80 Stress intensity factor Kl (MPa÷m)

90

100

12.15 Summary of conditions under which continuous crack growth was observed under static load in oxygenated high purity BWR and BWR water containing 50 mg/kg chloride.


183

Clear evidence of crack growth under static load was obtained from tests with materials A and B when introducing a chloride transient. According to the chloride content (50 mg/kg), the conditions corresponded to Action Level 2 of the EPRI Water Chemistry Guidelines, whereas the electrical conductivity was still in the range below Action Level 1. Crack initiation occurred during the phase of static load after a relatively short incubation time of 3 h and 32 h respectively (compare pages 178–9). The crack growth rate da/dt significantly exceeds the DL 2 (Fig. 12.15). For material A a crack growth rate of 1.8 ¥ 10–8 m/s (570 mm/a) was determined at a stress intensity factor of about 55 MPa÷m. Material B showed a longer incubation time and lower crack growth rate under static load of 2.5 ¥ 10–9 m/s (80 mm/a). Details are already shown in pages 178–9 and Figs 12.12 and 12.13. After returning to high purity water, a decreasing tendency of the crack growth rate with some delay was observed indicating a kind of ‘memory effect’ due to the residence time of the chloride ions in the crevice.

12.4

Summary and conclusions

The CASTOC project has given more insight into both the phenomenology and the acting mechanisms on corrosion cracking in BWR water environment and the effect of transients. With regard to the application of the results from the CASTOC project for the assessment of components in LWRs, the following aspects should be considered: ∑ ∑

∑

∑

Low-alloy steel base materials for RPV application revealed resistance to SCC crack growth under constant static load up to stress intensity factors of about 60 MPa÷m in BWR/NWC environment. Under certain environmental and material conditions, however, experimental results from tests performed under static load give reason for a more careful consideration of the assessment of components. The screening experiments of this project revealed crack growth under constant load for the following conditions: HAZ of the joint weld of material A, and materials A and B during a water chemistry transient with 50 mg/kg chloride. The proposed BWRVIP-60 SCC Disposition Line 1 [13] for crack growth under steady-state conditions was essentially confirmed for base materials. With respect to the Disposition Line 2 for transients in load and water chemistry, however, further consideration is recommended based on the results of this project. The project has revealed the general trend that the existing prediction curve presented in the ASME Boiler and Pressure Vessel Code, Section XI, Appendix A [10] for da/dN assessment of existing flaws in low-alloy steels may not be conservative under some specific conditions, e.g. low

184


loading frequencies, since the observed crack growth rates at low frequency cyclic loading significantly exceeded the maximum rates of the Code.

12.5

Acknowledgements

The financial support by the 5th Framework Programme of the European Commission under Contract No. FIKS-CT-2000-00048 and the Swiss Federal Office for Education and Science (BBW) is gratefully acknowledged. The valuable contributions of all partners to data generation and discussion of the results were the basis for a successful performance of the project.

12.6

References

1. J. Föhl, U. Ehrnstén, M. Ernestová, D. Gómez-Briceño, J. Lapeña, S. Ritter, A. Roth, B. Devrient, H.P. Seifert, T. Weissenberg, M. Žamboch, ‘Crack Growth Behaviour of Low Alloy Steel for Pressure Boundary Components under Transient Light Water Reactor Operating Conditions – CASTOC’, FISA Conf. on EU Research in Reactor Safety, Luxembourg, Nov. 12–14, 2001. 2. J. Föhl, U. Ehrnstén, M. Ernestová, D. Gómez-Briceño, J. Lapeña, S. Ritter, A. Roth, B. Devrient, H.P. Seifert, T. Weissenberg, M. Žamboch, ‘Crack Growth Behaviour of Low-Alloy Steels for Pressure Boundary Components under Transient Light Water Reactor Operating Conditions – CASTOC’, FISA Conf. on EU Research in Reactor Safety, Luxembourg, Nov. 10–12, 2003. 3. M. Ernestová, M. Žamboch, J. Föhl, U. Ehrnstén, D. Gómez-Briceño, J. Lapeña, S. Ritter, A. Roth, B. Devrient, H.P. Seifert, T. Weissenberg, ‘Crack Growth Behaviour of Low-Alloy Steels for Pressure Boundary Components under Transient Light Water Reactor Operating Conditions – CASTOC, Part II: VVER Conditions’, EUROCORR 2004, Paper No. 241, Nice, France, Sep. 12–16, 2004. 4. U. Ehrnstén, ‘Crack Growth Behaviour of Low Alloy Steels for Pressure Boundary Components Under Transient Light Water Reactor Operating Conditions’, EU Project CASTOC, Technical Report WP1, 2002. 5. J. Lapeña, D. Gómez-Briceño, ‘Crack Growth Behaviour of Low Alloy Steels for Pressure Boundary Components Under Transient Light Water Reactor Operating Conditions’, EU Project CASTOC, Technical Report WP2, 2003. 6. M. Žamboch, J. Föhl, ‘Crack Growth Behaviour of Low Alloy Steels for Pressure Boundary Components Under Transient Light Water Reactor Operating Conditions’, EU Project CASTOC, Technical Report WP3, 2003. 7. A. Roth, B. Devrient, J. Föhl, ‘Crack Growth Behaviour of Low Alloy Steels for Pressure Boundary Components Under Transient Light Water Reactor Operating Conditions’, EU Project CASTOC, Technical Report WP4, 2003. 8. U. Staud, M. Lasch, ‘Die VGB-Richtlinie für Chemie in SWR-Anlagen – aktueller Stand’; VGB-Konferenz ‘Chemie im Kraftwerk 1995’, VGB-Speisewassertagung 1995, Vortrag KKW3. 9. U. Ehrnstén, D. Gómez-Briceño, J. Lapeña, M. Ernestová, M. Žamboch, S. Ritter, H.P. Seifert, A. Roth, J. Föhl, F. Hüttner, T. Weissenberg, ‘Inter-Laboratory Crack Growth Test on Pressure Vessel Steel 20MnMoNi5-5 in Simulated BWR Environment’,


10.

11.

12. 13.

14.

185

11th Int. Conf. on Env. Deg. of Mat. in Nucl. Power Systems – Water Reactors, Stevenson, WA, USA, Aug. 10–14, 2003. ASME Boiler and Pressure Vessel Code, Section XI, ‘Rules for In-service Inspection of Nuclear Power Plant Components’, Appendix A, Article A-4000, ‘Material Properties’. A. Roth, et al., ‘Experimental Investigations Concerning the Possible Effect of Dynamic Strain Ageing in the Environmentally Assisted Cracking of Low Alloy Steels in Oxygenated High Temperature Water’, 29th MPA Seminar, Stuttgart, Germany, Oct. 09–10, 2003. H.P. Seifert, S. Ritter, J. Hickling, Power Plant Chemistry, 6, pp. 111–123, 2004. F.P. Ford, R.M. Horn, J. Hickling, R. Pathania, G. Brümmer, ‘Stress Corrosion Cracking of Low Alloy Steels under BWR Conditions; Assessments of Crack Growth Rate Algorithms’, 9th Int. Conf. on Env. Deg. of Mat. in Nucl. Power Systems – Water Reactors, pp. 855–863, Newport Beach, CA, USA, Aug. 1–5, 1999. J. Heldt, H.P. Seifert, Nuclear Engineering and Design, 206, pp. 57–89, 2001.

13 Crack growth behaviour of low-alloy steels for pressure boundary components under transient light water reactor operating conditions – CASTOC, Part 2: VVER conditions M. E R N E S T O V Á and M. Ž A M B O C H, Nuclear Research Institute (NRI), Czech Republic, B. D E V R I E N T and A. R O T H, Framatome ANP GmbH, Germany, U. E H R N S T É N, VTT Industrial Systems, Finland, J. F Ö H L and T. W E I S S E N B E R G, Staatliche Materialprüfungsanstalt (MPA), Germany, D. G O M É Z B R I C E Ñ O and J. L A P E Ñ A, Centro de Investigaciones Energéticas Medioambientales y Tecnológicas (CIEMAT), Spain and S. R I T T E R and H. P. S E I F E R T, Paul Scherrer Institute (PSI), Switzerland

13.1

Introduction

The ageing of pressure boundary components is one of the main factors controlling the lifetime of nuclear power plants. Under certain circumstances EAC can be one of the major ageing mechanisms of LAS in high-temperature water. The project ‘Crack Growth Behaviour of Low-Alloy Steels for Pressure Boundary Components under Transient Light Water Reactor Operating Conditions’ (CASTOC), was performed within the 5th EC framework programme and addressed the problem of EAC of Western and Eastern type steels used for pressure boundary components [1, 2]. The objective of the CASTOC project was to screen the EAC behaviour of low-alloy reactor pressure vessel (RPV) steels in high-temperature water during load transients and water chemistry transients such as may occur during start-up and shut-down, steady-state operation and the following mode of commercially operating LWRs. This is in contrast to the worldwide activities in the past, which focused mainly on either cyclic loading or static loading and steady-state operating conditions. The main focus of the project was directed at the interactions between static and cyclic loading, which was realised, for example, by low frequency corrosion fatigue (LFCF) tests followed by static load or by periodical partial unloading (PPU) with different rise and hold times. In conjunction with the different load spectra, the effect of water 186


187

chemistry transients was investigated. A more detailed description of the project is available in references [1] and [2]. The results of the project should in particular be assessed with respect to quality, reliability and their application to plant concerns and possible code implementation. This comprises a comparison of the data from the CASTOC project with data from literature and codes and to give indications where the results may be considered in plant life management strategies. The tested materials and water environment conditions were chosen to address the concerns of both BWRs and VVERs. In this chapter the crack growth results obtained in a simulated VVER environment on materials representing VVER RPVs are presented. The main focus of the tests performed in a simulated VVER environment was directed to the crack growth rate (CGR) data during cyclic and static loadings and to investigate the effect of materials, oxygen content and the effect of different constraint situations realised by different specimen size. The tests were performed at NRI within the CASTOC project subsequently after the interlaboratory comparison test. The results of the tests performed under simulated BWR/NWC conditions are summarised in [3]. For a detailed description of all results see [4–8].

13.2


The tests performed within the CASTOC project comply with the current state-of-the-art knowledge of science and technology in laboratory testing of EAC processes. The investigated materials represent nuclear grade materials. With regard to the selected environmental conditions, enveloping parameters were applied.

13.2.1 Materials The ferritic reactors LAS 15Ch2MFA (material C) and 15Ch2NMFA (material D) as RPV steels used at VVER 440 and VVER 1000, respectively, were investigated under simulated VVER conditions. The chemical composition of the materials are summarised in Table 13.1. The steels were quenched and tempered followed by air cooling [6]. Both base materials had a fine-grained bainitic microstructure. Mechanical properties are given in Table 13.2. ∑ ∑

Material C, base material: forged plate 15Ch2MFA (thickness of the plate 140 mm), fabricated according to nuclear grade quality with enhanced sulphur content. Material D, base material: forged plate 15Ch2NMFA (thickness of the plate 320 mm), fabricated according to nuclear grade quality.

188

Base material

Chemical composition [wt.%] C

Mn

Si

P

S

Ni

Cr

Mo

Cu

V

Co

As

Material C

0.15

0.40

0.24

0.013

0.015

0.30

2.78

0.64

0.08

0.29

0.009

0.011

Material D

0.14

0.45

0.25

0.009

0.007

1.23

2.15

0.57

0.05

0.08

*

*

* not determined


Table 13.1 Chemical composition (in wt.%) of the investigated materials


189

Table 13.2 Mechanical properties of the investigated materials Temperature [∞C]

Rp0.2 [MPa] Yield strength

Rm [MPa] Tensile strength

A5 [%] Elongation

RA [%] Reduction of area

Material C

20 350

545 469

651 544

22.6 17.8

75.0 75.0

Material D

22

570

674

19.7

75.3

13.2.2 Specimen preparation The C(T)25 and C(T)50 samples were manufactured from forged plates, material C specimens in L-S direction and material D specimens in S-L direction. The specimen were pre-cracked in air using parameters which fulfill the demand of ASTM E399, 1990 to the final a0 /W ~ 0.5.

13.2.3 Experimental facilities The test equipment at NRI consisted of a heated autoclave vessel with an integrated bellows system to apply the mechanical load, a water refreshing system to adjust desired water conditions, a high pressure pump, and measuring equipment to control the water chemistry [6]. The autoclave had the capacity to install either two C(T)25 or C(T)50 specimens in a daisy chain, where each specimen was electrically isolated using ceramic and mica spacers. The laboratory test unit enables control of the water and loading conditions. The reversed direct current potential drop (DCPD) system for on-line crack length monitoring was used. The detection limit of the DCPD technique is of the order of 10 mm. On-line monitoring was performed on load, pressure and temperature. The outlet conductivity and defining the outlet oxygen concentration as the target value of oxygen were measured continuously. The external Ag/AgCl/deionate reference electrode and a platinum probe were used for continuous measurement of the corrosion and redox potentials.

13.2.4 Testing procedure and environment Each autoclave test was divided into a stabilisation phase, covering the time needed for heating and pressurising the autoclave (about 100 h) and a conditioning phase taking at least 100 h at stable conditions before the test phase was started by fatigue loading followed by constant loading. The main objective of the fatigue loading part of the test was to create an actively growing crack before switching to constant load, and to investigate whether sustained crack growth occurs at desired K value in simulated VVER water. Normal operation of VVER reactors is characterised by very low oxygen

190


Table 13.3 Water chemistry parameters in simulated VVER water Boric acid

Potassium hydroxide

Ammonia

Conductivity (in the outlet)

O2 concentration (in the outlet)

6.8 g/kg

23.5 mg/kg

20 mg/kg

~130 mS/cm

< 20 mg/kg or ~ 200 mg/kg

concentration (< 10 mg/kg). In order to investigate the effect of higher oxygen levels – e.g., as residual oxygen after reactor start-up or as a result of oxygen ingress during power operation – the tests were performed in oxygenated (~ 200 mg/kg) or oxygen-free (< 20 mg/kg) water with boric acid at 288 ∞C to generate conservative data with respect to plant conditions. The VVER water chemistry parameters are listed in Table 13.3. The aim of the project was to obtain the data at stress intensity factors outside the range of linear elastic fracture mode (LEFM). The testing procedure included loading the specimens to a stress intensity ranging from 56 to 88 MPa÷m starting of fatigue loading using a positive saw tooth waveform with a rise time of 1000 s and a decline time of 200 s (f = 8.3·10–4 s–1) and a load ratio of R = 0.1, 0.2 or 0.8. After the crack activation, constant load was applied for at least 300 h. After termination of the autoclave testing the cracks were opened at liquid nitrogen temperature. The pre-crack length a0 and EAC advance were measured at 25/50 (C(T)25/C(T)50) equidistant locations on the fracture surface along the notch. Fracture surface investigations using a scanning electron microscope were performed before and after electrochemical cleaning.

13.3


13.3.1 Results from cyclic loading The CGRs obtained from the experiments were compared with the prediction line of the ASME Boiler and Pressure Vessel Code Case N 643, Section XI, Div. 1 [9]. In this Code Case the prediction line is determined by a threshold value in DKI, the load ratio R and the rise time Dtr of the cycle. Despite the fact that the Code Case was established from Western type PWR water environments, the data obtained from tests in simulated VVER environments within this project are compared with the Code Case prediction line. The crack growth rates in mm/cycle have been calculated using the average crack increment for the cyclic test period, the amount of cycles and the rise time. Effect of material The tested materials differed in sulphur, chromium, phosphorus and vanadium content. Although material C had a higher sulphur content (0.015 wt.%)


191

compared to material D (0.007 wt.%) the CGRs for material D tended to be higher than those of material C under comparable conditions (Figs 13.1 and 13.2). This observation indicates that the sulphur content of the steel was not the sole material parameter controlling EAC growth rates. Effect of oxygen content The effect of dissolved oxygen in VVER water environment on the crack growth behaviour can be established only from the tests performed with C(T)50 specimens. At low loading ratio (R = 0.2 and R = 0.1) the CGRs tended to be slightly higher in oxygenated water than in oxygen-free water. At the higher loading ratio (R = 0.8) the CGRs in water with enhanced dissolved oxygen content were more than one order of magnitude higher than in an oxygen-free environment (Figs 13.3 and 13.4).

1.E+03

VVER water, 288 ∞C < 20 mg/kg or 200 mg/kg O2, f = 8.3 ¥ 10–4 s–1 CT25

1.E+02

CT50

Material C

da /dN (mm/cycle)

1.E+01

1.E+00

R = 0.8 ASME XI Code case N 643 PWR water line R = 0.2 Dtrise = 1000 s

1.E+01

R = 0.2/01

1.E–02

ASME XI air line R = 0.2

1.E–03

1

10 DKl,max (MPa.m1/2)

100

13.1 CGRs of material C at a frequency of 8.3 ¥ 10–4 s–1 in oxygenated and oxygen-free VVER water at 288 ∞C.

192

Corrosion issues in light water reactors 1.E+03

VVER water, 288 ∞C < 20 mg/kg or 200 mg/kg O2, f = 8.3 ¥ 10–4 s–1 CT25

1.E+02

CT50

Material D

da /dN (mm/cycle)

1.E+01

R = 0.8

1.E+00

1.E–01

ASME XI Code case N 643 PWR water line R = 0.2 Dtrise = 1000 s

R = 0.2/0.1

1.E–02

ASME XI air line R = 0.2

1.E–03 1


100

13.2 CGRs of material D at a frequency of 8.3 ¥ 10–4 s–1 in oxygenated and oxygen-free VVER water at 288 ∞C.

Effect of specimen size The effect of specimen size can be evaluated from the tests performed in oxygenated water with a loading ratio of 0.2 and 0.1 (Fig. 13.5). It becomes obvious that the crack growth per cycle in small C(T)25 specimens is about one order of magnitude higher than that of the large C(T)50 specimens for both materials. This might be caused by higher plastic deformation in the crack tip area of the smaller specimens, in particular at the low load ratio R since the prevailing mechanism is strain induced corrosion cracking. With regard to the transferability of laboratory results to large components the smaller specimens obviously provide more conservative data.

13.3.2 Results from constant loading Although the BWRVIP-60 SCC Disposition Lines (DL) [10] apply to BWR/ NWC conditions the da/dt data obtained from the experiments were compared to these DLs. In most of the tests the crack activated during cyclic loading

Crack growth behaviour of low-alloy steels 1.E+03

193

VVER water, 288∞C, f = 8.3 ¥ 10–4 s–1 oxygen-free (< 20 mg/kg O2) CT25

1.E+02

CT50

Closed symbols: material C open symbols: material D

da /dN (mm/cycle)

1.E+01

1.E+00 ASME XI Code case N 643 PWR water line R = 0.2 Dtrise = 1000 s

1.E–01

R = 0.2/0.1

R = 0.8

1.E–02 ASME XI air line R = 0.2

1.E–03

1


100

13.3 CGRs of material C and D at a frequency of 8.3 ¥ 10–4 s–1 in oxygen-free VVER water at 288 ∞C.

arrested after switching to static load. There were only a few exceptions in which continuous crack advance was detected under static load. The crack growth rates, da/dt, were calculated using the average crack increment for the static test period and the whole static period test time. Effect of material The corrosion fatigue crack activated by cyclic loading arrested after the change to constant static load during all tests in oxygenated (~ 200 mg/kg) and oxygen-free (< 20 mg/kg) simulated VVER water performed with C(T)25 and C(T)50 specimens from material C. During the test phases at static load of at least 300 h duration no crack growth could be detected at applied stress intensity factors ranging from 56 to 88 MPa÷m. Concerning the material D, the corrosion fatigue crack activated by cyclic loading, arrested after switching to constant static load during all tests performed on C(T)50 specimens in oxygenated and oxygen-free simulated VVER water. Similar to the behaviour of material C no crack growth was observed during

194

Corrosion issues in light water reactors 1.E+03

VVER water, 288 ∞C, f = 8.3 ¥ 10–4 s–1 oxygenated (200 mg/kg O2)

1.E+02

CT50 Closed symbols: material C open symbols: material D

da /dN (mm/cycle)

1.E+01

R = 0.8

1.E+00 R = 0.2

ASME XI Code case N 643 PWR water line R = 0.2 Dtrise = 1000 s

1.E–01


1.E–03

1


100

13.4 CGRs of material C and D at a frequency of 8.3 ¥ 10–4 s–1 in oxygenated VVER water at 288 ∞C.

the phases of static load. However, crack growth under static load was observed in C(T)25 specimens tested in oxygenated VVER water at stress intensity factors in the range from 58 to 70 MPa÷m. For comparison the results are displayed together with the BWRVIP-60 SCC DLs in Fig. 13.6. For stress intensity factors in the range of up to about 56 MPa÷m no crack growth was observed in any of the specimens of material C or in any of the C(T)50 specimens of material D, which were tested in oxygenated and oxygen-free VVER water. When during the constant load period the stress intensity factor was increased beyond the validity limits for linear elastic fracture mode (LEFM) continuous crack growth was observed on C(T)25 specimens of material D in oxygenated VVER water and the da/ dt data (~10–8 m/s) are higher than the ones expected from DL 2. The investigation of the fracture surfaces on C(T)25 specimens of material D which showed crack growth under static load revealed a small portion of intergranular (IG) cracking during cyclic phase and high portion of IG cracking during the constant load phase (Fig. 13.7).

Crack growth behaviour of low-alloy steels 1.E+03

195

VVER water, 288 ∞C, 200 mg/kg O2, f = 8.3 ¥ 10–4 s–1, R = 0.2/0.1

CT 25

1.E+02

CT 50

Closed symbols: material C open symbols: material D

C(T) 25

da/dN (mm/cycle)

1.E+01

1.E+00 ASME XI Code case N 643 PWR water line R = 0.2 Dtrise = 1000 s

1.E–01

C(T) 50


1.E–03

1


100

13.5 CGRs of material C and D at a frequency of 8.3 ¥ 10–4 s–1 in oxygenated VVER water at 288 ∞C; C(T)25 and C(T)50 specimens.

Effect of oxygen The effect of oxygen could not be demonstrated on materials C and D in VVER water because of crack arrest on most of the specimens. The only possibility of examining the effect of oxygen is to use the da/dt data obtained on C(T)25 specimens of material D in oxygenated (~ 200 mg/kg) VVER water which could be compared with data in oxygen-free VVER water. The additional test with C(T)25 specimens of C and D materials in oxygen-free VVER water is being carried out separately to the CASTOC project and will document the possible effect of oxygen. Effect of specimen size The effect of specimen size on the crack growth rate under static load could not be clearly demonstrated. In tests with material C no crack growth was observed for both specimen sizes (i.e. 25 mm or 50 mm thick) and in material D crack growth was only detected in C(T)25 specimens, this was associated

196


1.E–07

1.E–08

CT25 CT50 Closed symbols: material C Open symbols: material D

da /dt (m/s)

1.E–09

1.E–10 VIP DL 1

1.E–11

VIP DL 2

1.E–12

1.E–13 20

30

40

50 60 Kl (MPa.m1/2)

70

80

90

13.6 Crack growth behaviour of material C and D under constant static load in oxygenated and oxygen-free VVER water at 288 ∞C; the data below detection limit ~ 10–13 m/s.

with IG cracking. It can be concluded that for the smaller specimen size C(T)25 equal nominal stress intensity factors provided more conservative data than the larger one, which suggests that the mechanism was strain induced corrosion cracking because more extended yielding is anticipated in the smaller specimen at high KI values.

13.4

Summary and conclusions

The PWR part of the CASTOC project addresses environmentally assisted crack growth phenomena in the low-alloy steels used for pressure boundary components in Russian-type pressurised water reactors (VVER). The number of tests using sophisticated test facility and measurement technique for the on-line detection of crack advance have provided a more detailed understanding of the mechanism of environmentally assisted cracking and provided quantitative data for CGRs as a function of loading events and time. The work was focussed on the evaluation of crack growth under cyclic load, crack growth and crack cessation under static load, and on determining the effect of oxygen content, constraint and stress state outside the range of LEFM. Nevertheless, regarding the application of the results from the CASTOC project for the assessment of components in LWRs, the following aspects should be considered: ∑

Low-alloy steel base materials for RPV application revealed resistance to crack growth under constant static load up to stress intensity factors


50 mm (a)

50 mm (b)

50 mm (c)

13.7 Typical intergranular attack in C(T)25 specimen of material D tested in oxygenated VVER water environment at 288 ∞C (static loading phase).

197

198

∑

∑

∑


of about 60 MPa÷m in VVER normal water chemistry. The observed crack growth behaviour is consistent with plant experience in general, and especially in those cases, where austenitic stainless steel cladding is not applied or was removed deliberately or by chance. Under certain environmental and material conditions, however, experimental results from tests performed under static load suggest that a more careful assessment of components should be made. The screening experiments carried out in this project revealed crack growth under constant load for the following condition: material D in oxygenated (~ 200 mg/kg) VVER water and stress intensity factor beyond the validity limit for linear elastic fracture mode. The proposed BWRVIP-60 SCC Disposition Line 1 [10] for crack growth under steady-state BWR/NWC conditions was essentially confirmed as well as for steady-state VVER conditions. However, with respect to the Disposition Line 2 for transients in load and water chemistry further consideration is recommended based on the results of this project. The curves provided in the ASME Code Case N 643 for a PWR water environment cover fairly well the data obtained in a VVER environment for Russian type RPV steels, even at increased oxygen content.

The CASTOC results provide an important contribution to the understanding of crack growth behaviour as a function of time and on as a consequence of the number and height of loading events. This is important in evaluating transient events, that may occur in a power plant.

13.5

Acknowledgements

The financial support by the 5th Framework Programme of the European Commision under contract No. FIKS-CT-2000-00048 is gratefully acknowledged. The valuable contributions of all partners to data generation and discussion of the results formed the basis for successful performance of the project.

13.6

References

1. J. Föhl, U. Ehrnstén, M. Ernestová, D. Gómez-Briceño, J. Lapeña, S. Ritter, A. Roth, B. Devrient, H.P. Seifert, T. Weissenberg, M. Žamboch, ‘Crack Growth Behaviour of Low Alloy Steel for Pressure Boundary Components under Transient Light Water Reactor Operating Conditions – CASTOC’, FISA Conference on EU Research in Reactor Safety, Luxembourg, November 12–14, 2001. 2. J. Föhl, U. Ehrnstén, M. Ernestová, D. Gómez-Briceño, J. Lapeña, S. Ritter, A. Roth, B. Devrient, H.P. Seifert, T. Weissenberg, M. Žamboch, ‘Crack Growth Behaviour of Low-Alloy Steels for Pressure Boundary Components under Transient Light Water Reactor Operating Conditions – CASTOC’, FISA Conference on EU Research in Reactor Safety, Luxembourg, November 10–12, 2003.


199

3. S. Ritter, H.P. Seifert, B. Devrient, U. Ehrnstén, M. Ernestová, J. Föhl, D. GómezBriceño, J. Lapeña, A. Roth, T. Weissenberg, M. Žamboch, ‘Crack Growth Behaviour of Low-Alloy Steels for Pressure Boundary Components under Transient Light Water Reactor Operating Conditions – CASTOC, Part I: BWR/NWC Conditions’, Paper No. 281, EUROCORR 2004, Nice, France, September 12–16, 2004. 4. U. Ehrnstén, ‘Crack Growth Behaviour of Low Alloy Steels for Pressure Boundary Components Under Transient Light Water Reactor Operating Conditions’, EU project CASTOC, Technical Report WP1, 2002. 5. U. Ehrnstén, D. Gómez-Briceño, J. Lapeña, M. Ernestová, M. Žamboch, S. Ritter, H.P. Seifert, A. Roth, J. Föhl, F. Hüttner, T. Weissenberg, ‘Inter-Laboratory Crack Growth Test on Pressure Vessel Steel 20MnMoNi5-5 in Simulated BWR Environment’, 11th Int. Conf. on Env. Deg. of Materials in Nuclear Power Systems – Water Reactors, Stevenson, WA, USA, August 10–14, 2003. 6. M. Žamboch, J. Föhl, ‘Crack Growth Behaviour of Low Alloy Steels for Pressure Boundary Components Under Transient Light Water Reactor Operating Conditions’, EU project CASTOC, Technical Report WP3, 2003. 7. J. Lapeña, D. Gómez-Briceño, ‘Crack Growth Behaviour of Low Alloy Steels for Pressure Boundary Components Under Transient Light Water Reactor Operating Conditions’, EU project CASTOC, Technical Report WP2, 2003. 8. A. Roth, B. Devrient, J. Föhl, ‘Crack Growth Behaviour of Low Alloy Steels for Pressure Boundary Components Under Transient Light Water Reactor Operating Conditions’, EU project CASTOC, Technical Report WP4, 2003. 9. Cases of ASME Boiler & Pressure Vessel Code, Case N-643; ‘Fatigue Crack Growth Rate Curves for Ferritic Steels in PWR Water Environment’, Section XI, Division 1, 2000. 10. F.P. Ford, R.M. Horn, J. Hickling, R. Pathania, G. Brümmer, ‘Stress Corrosion Cracking of Low Alloy Steels under BWR Conditions; Assessments of Crack Growth Rate Algorithms’, 9th Int. Conf. on Env. Deg. of Materials in Nuclear Power Systems – Water Reactors, Newport Beach, CA, USA, pp. 855–863, August 1–5, 1999.

14 Effect of yield strength on stress corrosion crack propagation under PWR and BWR environments of hardened stainless steels M. L. C A S T A Ñ O, M. S. G A R C Í A, G. D E D I E G O and D. G O M É Z - B R I C E Ñ O, CIEMAT, Spain

14.1

Introduction

Core components of light water reactors (LWR), made of austenitic stainless steels (SS) and nickel alloys, subjected to stress and exposed to relatively high fast neutron flux may suffer a cracking process termed Irradiation Assisted Stress Corrosion Cracking (IASCC). This degradation phenomenon is a timedependent process in which neutron and gamma radiation are directly implicated in the initiation and propagation of cracking [1]. Although this type of cracking was first recognized in Boiling Water Reactor (BWR), later service failures attributed to IASCC were observed in Pressurized Water Reactor (PWR) components [2]. Among the material modifications induced by neutron irradiation, Radiation Induced Segregation (RIS) and Radiation Hardening have been identified as the main contributors to the susceptibility to IASCC of irradiated stainless steels. Chromium depletion at grain boundaries produce by Radiation Induced Segregation can justify the IASCC response in oxidizing environments, such as BWR normal water chemistry [3]. However, in non-oxidizing environments, such as PWR primary water or BWR hydrogen water chemistry, the role played by chromium depletion at grain boundaries on IASCC behaviour of highly irradiated material seems to be irrelevant [4], and the influence of material hardening is becoming more strongly considered. Radiation hardening can be simulated by mechanical deformation, in spite of the significant difference observed in the microstructure of both types of materials. Furthermore, it is accepted that the study of the SCC behaviour of unirradiated austenitic steels with different hardening levels could contribute to the understanding of IASCC mechanism. In this chapter, crack growth rate data of sensitized 304, 316L and 347 with different hardening levels obtained by cold work will be presented and discussed. Sensitized 304 SS has been tested in BWR conditions to gain some insight into the cracking behaviour of core shroud, in which the cracks appear associated to HAZ. 316L SS will be tested in typical PWR primary 200

Effect of yield strength on stress corrosion crack propagation

201

water and in high lithium primary water to study the behaviour of this material in its use as baffle former bolts (BFB). Tests have been carried out in high lithium primary water to simulate the postulated environment around the BFB due to gamma heating. Niobium-stabilized type 347 SS has been tested in BWR and PWR to study the behaviour of this material in its use as core shroud in the Siemens-KWU plants and in its use as BFB in Westinghouse plants.

14.2


14.2.1 Materials Materials tested were three plates of commercial austenitic stainless steels, type 316L, 347 and 304. The chemical compositions and the mechanical properties, at room temperature, in ‘as-received’ condition, are shown in Tables 14.1 and 14.2, respectively. Stainless steel type 304 SS was subjected to a sensitization treatment of 650 ∞C, 1 hour and air cooling, while no additional heat treatment was applied to 316L and 347 SS.

14.2.2 Hardening process In order to produce materials with different yield strength, several degrees of cold or warm work have been applied to the material by tensile deformation. Large tensile samples were strained in a 100 Tm tensile machine to several Table 14.1 Chemical composition (%wt) of tested materials Material

C

Cr

Mn

Mo

N

316L SS

0.020 17.39 1.28

2.20 0.020 11.49 0.032

347 SS

0.028 18.10 1.79

–

304 SS

0.059 18.02 1.63

0.20 0.084

–

Ni

P

S

Si

Co

Nb

0.001 0.45

0.14

–

10.35 0.022

0.005 0.17

–

0.46

8.10 0.022

0.022 0.44

0.15

–

Table 14.2 Mechanical properties (yield strength, ultimate tensile strength, percentage of elongation and Vickers hardness), at room temperature, in asreceived condition Material

YS MPa (Ksi)

UTS MPa (Ksi)

Elongation (%)

Hv (30 Kg)

316L SS 347 SS Sen. 304 SS*

232 (33.9) 238 (34.8) 240 (35.1)

566 (82.6) 568 (82.9) 707 (103.2)

84 70 86

172 171 194

*Mechanical properties after the sensitization treatment: 650 ∞C, 1 hour, air cooling

202


Table 14.3 Test matrix PWR 340 ∞C

Material 316L CW 347 CW Sen. 304 WW

BWR, 290 ∞C

1200 ppm B, 2 ppm Li 3.2 ppm H2

108 ppm B, 7 ppm Li 3.2 ppm H2

NWC 200 ppb O2

YS (MPa) 542 772 518 639 – –

YS (MPa) 542 819 – – – –

YS (MPa) – – 238 518 240 503

819 705 –

– 639 643

– 705 680

CW: Cold worked, WW: Warm worked,

% of deformation, before necking. Finally the gauge length of the strained samples was used for Compact-Tension (C-T), tensile and metallographic samples fabrication. Stainless steels 316L and 347 were hardened by cold work at room temperature whereas 304 was hardened by warm work at a temperature higher than 200 ∞C.

14.2.3 Test procedure Crack growth test of hardened stainless steels, 316L, 347 and sensitized 304, with different levels of yield strength have been performed in BWR and PWR conditions. Table 14.3 shows the test matrix carried out. Two samples per conditions were tested simultaneously. Crack growth rate tests were performed using 12 mm CT specimens, fabricated according to the ASTM E-399. CT specimens were pre-cracked in air under a triangular wave (22 Hz and R = 0.1). Then, the samples were precracked in high temperature water under a triangular wave (2 ¥ 10–2 Hz and R = 0.6) and, then, constant load with periodic unloading under trapezoidal wave was applied with a holding time of 9000 seconds and 5 and 45 second for the unloading and reloading process. In some of the tests in BWR conditions (347 and sensitized 304 SS with highest YS) the time loading and reloading was 50 and 450 seconds. Nominal stress intensity factor, Ki, was in the range of 25–35 MPam1/2. The apparent crack advance was on-line monitored by Direct Current Potential Drop (DCPD) technique. After each crack propagation test in high temperature water, the CT specimens were opened by fatigue in air, at room temperature, and the fracture surface examined by Scanning Electron Microscopy (SEM). This observation was used to identify the areas of crack propagation, to determine the morphology of the cracking and to verify the accuracy of the DCPD monitoring technique.


14.3

203


14.3.1 Effect of yield strength in PWR conditions Sixteen CT samples were tested at 340 ∞C in primary PWR. After testing, of 230 to 580 hours duration, some samples exhibit a band of intergranular cracking with a regular initiation from the pre-cracking end and almost even crack propagation. However, other specimens show localized initiations along the pre-cracking end, rapid propagation along elongated grains with the presence of ligaments and growing finger-like shapes. Some samples only show localized crack initiation along the pre-cracking crack front, identified as intergranular morphology. In all the cases the fracture morphology was always intergranular. In addition to the crack propagation obtained by fractographic measurements, apparent crack growth rate was available using the DCPD. Agreement between crack growth rates obtained by both techniques has been discussed previously [5]. To establish comparisons among different conditions, maximum crack growth rates (CGR) were obtained by dividing the deepest fractographic intergranular propagation by total testing time. Figure 14.1 plots the crack growth rate of 316L and 347 SS at 340 ∞C obtained under standard and high lithium primary water conditions. All crack growth rate data are normalized to a stress intensity of 30 MPam1/2 [6]. In both materials, as yield strength increases the crack growth rate increases. In the case of 347 SS the crack growth rates obtained are lower than in 316L SS. However, the effect of yield strength seems to be more evident in 347 SS than in 316L SS. Results obtained are consistent with the published data for cold worked 316L SS obtained by Andresen [7], although in this reference tests were performed in hydrogenated pure water instead of primary PWR.

316L Primary water 316L High Li 347 Primary water

CGR* (mm/s)

1E-6

1E-7

1E-8

* Corrected CGR to K30 450

500

550

600

650 700 YS (MPa)

750

800

850

14.1 CGR of CW stainless steels in PWR at 340 ∞C.

900

204


Crack growth rate was ª 3 times higher in high Li primary water than in PWR standard primary water. The SCC susceptibility of cold worked austenitic stainless steels in lithiated solutions at high temperature has been shown previously by Smialowska and co-workers [8]. The possibility of highly alkaline environments by concentration of LiOH in the liquid phase has been considered for some closed crevices, in particular for PWR core baffle bolts [9].

14.3.2 Effect of yield strength in BWR conditions Fourteen CT samples were tested at 290 ∞C in BWR conditions Test temperature, pressure, conductivity, dissolved gases and corrosion potential, using a reference electrode of Cu/Cu2O, were continuously monitored and recorded. After the pre-cracking at high temperature, samples were exposed for around 500 hours to oxidizing environment (200 ppb oxygen). Specimens of sensitized 304 with 240 MPa of YS, used as reference, presented a band of intergranular cracking with regular initiation from the pre-cracking end and almost even crack propagation. However, sensitized and warm-worked 304 SS presents partial crack initiation from the pre-cracking end, rapid propagation along elongated grains, presence of ligaments between intergranular cracking and growing finger-like shapes. In the case of 347 SS, no intergranular cracking was observed in material with yield strength of 238 MPa, but a clear intergranular propagation was observed for higher yield strength. The crack propagation in 347 SS shows similar characteristics to those observed in 304 SS. Following the same criteria as in PWR, maximum crack growth rate has been used for comparisons. Figure 14.2 plots the crack growth rate of sensitized

CGR* (mm/s)

1E-6

Sen. AISI-304 AISI-347

1E-7

1E-8 * Corrected CGR to K30 200

300

400 500 YS (MPa)

600

700

800

14.2 Crack growth rates of sensitized 304 SS and 347 SS tested at 290 ∞C in BWR conditions.


205

304 and 347 SS as a function of yield strength in BWR conditions, at 290 ∞C. According to the fractographic values, the crack growth rate of sensitized 304 SS and 347 SS increases as yield strength increases. However, the capability of increasing the crack propagation is less pronounced in sensitized 304 SS. This material shows significant crack growth rate in the non-warmwork conditions, due to its sensitization treatment. Apparently, the detrimental effect of yield strength could be overwhelmed by the marked and wellknown high crack growth rate of sensitized material in oxidizing environments. A more significant effect of sensitization is reported by M. Spiedel [10], who indicates that crack growth rate of heavily sensitized 304 SS is independent of the stress intensity and yield strength. In fact, in the present study, very similar CGR have been obtained for specimens with yield strength of 503 and 680 MPa. However, it is important to indicate that yield strength of 680 MPa in sensitized 304 was obtained by a two-step deformation and that the resultant increase of hardness was lower than expected. As a consequence, the crack growth rate could be affected by the straining procedure followed to produce the target yield strength. Crack growth rate of 347 SS shows a significant dependence on the yield strength, and crack growth rate as high as 3.1 ¥ 10–7 mm/s was measured for 347 SS with a YS of 705 MPa. These values are consistent with the published data [11]. This material, 347 SS has been tested in PWR and BWR conditions. In spite of the higher test temperature in PWR water, crack propagation was higher in BWR (NWC) water. The influence of yield strength in both environments seems to be quite similar, Fig. 14.3. This similar dependence with the yield strength at both low and high corrosion potential has also been reported by Andresen [11].

CGR* (mm/s)

1E-6

347 BWR 290 ∞C 347 PWR 340 ∞C

1E-7

1E-8 * Corrected CGR to K30 500

600

700 YS (MPa)

800

900

14.3 Crack growth rate of cold worked 347 SS, tested in PWR and BWR conditions.

206


14.3.3 Implications for the IASCC process Hardening induced by neutron radiation is one of the effects observed on structural materials exposed to radiation fields. The potential contribution of radiation induced hardening to the initiation and propagation of cracking process is becoming more strongly considered, especially for environments where other factors, such as microchemistry, have no significant influence. In order to compare the crack growth rate data obtained with hardened and irradiated materials, the correlation between yield strength and dpa shown in Fig. 14.4 has been used [12]. The range of yield strength tested, from 500 to 800 Mpa, could correspond to a radiation damage from 1 to 10 dpa. Figure 14.5 shows crack growth rate data for hardened materials obtained in this work and in other labs and available data for irradiated materials, in PWR conditions. Crack growth rate for irradiated 304 up to 12 and 35 dpa were obtained under constant load with one per day unloading (holding time >80.000 s.) in some steps of the test, at 335 ∞C [13]. Only data from valid K values have been included. Crack growth rate data for irradiated 304 up to 6 dpa under several program conditions were obtained at 340 ∞C [14]. All data were corrected to K30. Irradiated stainless steels at higher yield strength presented a significant

Measured tensile yield strength (MPa)

1000

800

600 304 SS Jenssen 304 SS Jenssen 304 SS Jenssen 304 SS Jenssen 316 SS Jenssen 316 SS Jenssen 316 SS Jenssen 316 SS Jenssen 316 SS Bergenlid 310 SS Kodama 316 SS Kodama 347 SS Kodama 304/316 SS Odette/Lucas Neutron-irradiated 300-Series Stainless Steel

400

200

0

0

1

2

3 4 Square root (dpa)

14.4 Increase of yield strength with dpa [12].

5

6

7


CGR (mm/s)

1E-6

CW 316 CW 347 WW 316 L. Shoji [13] WW 304 L. Shoji [13] WW 304 L. Andresen [14, 15] WW 316 L. Andresen [14, 15] Irrad-304 Halden [16] Irrad-304 CIR-Content Load [17] Irrad-304 CIR-trapez-1000s [17] Irrad-304 CIR-trapez-10000s [17]

207

PWR

1E-7

1E-8 500

600

700

800 YS (MPa)

900

1000

1100

14.5 Comparison of CGR of irradiated material and hardened material at 340–335 ∞C in PWR conditions.

dispersion, SCC behaviour is probably a multi-parameter phenomenon and the influence of testing conditions could modify final results. However, the crack growth rate for irradiated stainless steels and for hardened material show a similar trend when the yield strength increases. Therefore, hardened material seems to reproduce reasonably well the behaviour of irradiated material under stress corrosion cracking conditions, in PWR conditions. Similar comparison has been performed for BWR conditions, Fig. 14.6. Crack growth rate data for irradiated 304 SS (13 dpa), 316NG (1.4 dpa) and 347 SS were obtained in pure water with 6–7 ppm O2 at 288 ∞C [13, 15] Crack growth rate for irradiated 304 SS (1.4 and 3 dpa) and irradiated 316 SS (2 dpa) were obtained in pure water with 300 ppb O2 at 289 ∞C, under several load programs [16]. Only data obtained under constant load or under cycling load with holding time of 7200 were included in the Fig. 14.6. Finally, irradiated 304 SS (12 dpa) was tested in pure water with 900 ppb O2. at 288 ∞C, under constant load [17]. Comparison of crack growth rates of irradiated material and hardened material shows that, in general, crack propagation rates of irradiated material are above crack growth rates of hardened material. In addition, crack growth rates of irradiated material are also above crack growth rates of hardened and sensitized material. These results seem to suggest that hardened materials and even hardened and sensitized materials are not appropriated to evaluate the behaviour of irradiated materials in oxidizing conditions, representative of BWR environments.

208

CGR (mm/s)

1E-5

1E-6

Corrosion issues in light water reactors Se. 304 347 316L Shoji [13] 304 Shoji [13] 347 Shoji [13] 304L Andresen [14, 15] 316L Andresen [14, 15] Irrad. 304, 316, 1.4 and 3dpa NRC [19] Irrad. 304, 12 dpa, Studsvik [20] Irrad. 304, 13 dpa Halden [16, 18] Irrad. 347, 2 dpa Halden [16, 18] Irrad. 316NG, 1.4 dpa Halden [16, 18]

BWR

1E-7

200

300

400

500 600 YS (MPa)

700

800

900

1000

14.6 Comparison of CGR of irradiated material and hardened material at 288 ∞C in BWR conditions.

14.4

Conclusions

In PWR at 340 ∞C, crack growth rates increase as yield strength increases, both in 316L SS and 347 SS. The influence of yield strength is more pronounced in 347 SS that in 316L SS. However the crack propagation obtained is lower in 347 SS than in 316L SS. Crack growth rates are almost ª 3 times higher in high Li and high pH primary water than in PWR standard primary water. The detrimental effect of yield strength on crack propagation, in BWR conditions, was less pronounced in sensitized 304 than in 347 SS. For this material, crack propagation rates are higher in BWR (NWC) than in primary PWR, in spite of the higher temperature of the latter conditions. Radiation hardening and cold work produces similar effects on crack growth in PWR conditions. Hardened material seems to be appropriate to assess the stress corrosion cracking behaviour of irradiated material under this condition. However, crack growth rate data obtained with hardened and sensitized material poorly reproduce the data of irradiated material and therefore seems not to be appropriated to evaluate the behaviour of irradiated materials in oxidizing conditions representative of BWR environments.


14.5

209

References

1. P. L. Andresen. ‘Irradiation Assisted Stress Corrosion Cracking’. Stress Corrosion Cracking Material Performance and Evaluation. R.H. Jones. ASM-210, p 182. (1992). 2. P. Scott, M. Meurier, D. Deydier, S. Silvestre, A. Trency. ‘An analysis of Baffle Former Bolt cracking in French PWRs’. Environmental Assisted Cracking: Predictive Methods for Risk Assessment and Evaluation of Materials Equipment and Structures.ASTM 1410.West Conshohocken, (2000). 3. M. Kodama, R. Katsura, J. Morisawa, S. Nishima, S. Suzuki, K. Takamori. ‘IASCC susceptibility of Irradiated Austenitic Steels under very Low Dissolved Oxygen’. 7th Int. Conf. on Environ. Degradation of Materials in NPS-Water Reactor, NACE, p 1121 (1995). 4. K. Fukuya et al. ‘Stress Corrosion Cracking on cold worked 316 stainless steel irradiated to high fluence’. 10th International Conference on Environmental Degradation of Materials in NPS – Water Reactors. Lake Tahoe (2001). 5. M. L. Castaño, M. S: Garcia Redondo, G. De Diego, D. Gómez-Briceño, ‘Crack Growth Rate in BWR and PWR of Hardened Austenitic Stainless Steels’, 11th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors August 10–14, 2003. 6. P. Andresen, K. Gott, L. Nelson. ‘Stress Corrosion Cracking of Sensitized type 304 Stainless Steels in 288 ∞C Water: A five Laboratory Round Robin’. 9th International Symposium on Environmental Degradation of Material in NPS – Water Reactors. TSM, p 423 (1999). 7. P. Andresen, T. Angeliu, R. Catlin, L. Young, R. Horn. ‘Effect of Deformation on SCC of Unsensitized Stainless Steels’. NACE Corrosion 2000. Paper 203. 8. S. W. Sharkawy, Z. Xia, Z. Szklarska-Smialowska. ‘Stress Corrosion Cracking of AISI304 and 316 Stainless Steels in Lithiated Water at 350 ∞C’. JNM 195, p 184 (1992). 9. P. M. Scott. 200 F. N. Speller Award Lecture: ‘Stress Corrosion Cracking in Pressurized Water Reactors – Interpretation, Modeling and Remedies’. Corrosion, Vol. 56, No. 8, 771–782 (2000). 10. M. O. Spiedel, R. Magdowski. ‘Environmental Degradation Assessment and Life Prediction of Nuclear Piping Made of Stabilized austenitic stainless steels’. Proceedings of the Inter. Symp. Plant Aging and Life Prediction of Corrodible Structures. Sapporo, Japan, p 951 (1995). 11. P. Andresen, T. Angeliu, L. Young, W. Catlin, R. Horn. ‘Mechanisms and Kinetics of SCC in Stainless Steels’. 10th International Conference on Environmental Degradation of Materials in NPS – Water Reactors, Lake Tahoe (2001). 12. S. M. Brummer, E. P. Simonen, P. M. Scott, P. L. Andresen, G. S. Was, J. L. Nelson. ‘Radiation Induced Material Changes and Susceptibility to Intergranular Failure of Light Water Reactor Core Internals’. JNM 274, p 299 (1999). 13. T. Shoji, et al. ‘Quantification of Yield Strength Effects on IGSCC of Austenitic Stainless Steels in High Temperature Waters’. 11th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors August 10–14, 2003. 14. P. L. Andresen. ‘Similarity of Cold Work and Radiation Hardening in Enhancing Yield Strength and SCC Growth of Stainless Steel in Hot Water’. NACE 2002 Paper No. 0509. 15. P. L. Andersen et al. ‘Stress Corrosion Crack Growth Rate Behavior of Various Grades of Cold Worked Stainless Steel in High Temperature Water’. NACE 2002 Paper No. 2511.

210 16. 17. 18. 19.


OECD Halden Reactor Project. Proprietary Information. CIR Proprietary Information. OECD Halden Reactor Project. Proprietary Information. O. K. Chopra, E. E. Gruber and W. J. Shack. ‘Crack Growth behavior of irradiated austenitic stainless steels in high purity water at 289 ∞C’. 11th Int. Con. on Environmental Degradation on Material in NPS – Water Reactors. Skamania Lodge, August 2003. 20. A. Jenssen, P. Efsing, K. Gott, P. O. Anderson. ‘Crack growth behavior of irradiated 304L stainless steel in simulated BWR environment’. 11th Int. Con. on Environmental Degradation on Material in NPS – Water Reactors. Skamania Lodge, August 2003.

15 Corrosion fatigue crack growth behaviour of low-alloy RPV steels at different temperatures and loading frequencies under BWR/NWC environment S. R I T T E R and H. P. S E I F E R T, Paul Scherrer Institute (PSI), Switzerland

15.1

Introduction

Low-alloy steels (LAS) are widely used for the reactor pressure vessel (RPV) of light water reactors (LWR), which is the most critical pressure-boundary component as far as safety and plant life are concerned. The possible effect of environmentally-assisted cracking (EAC) on RPV structural integrity therefore continues to be a key concern within the context of both reactor safety and evaluation/extension of plant service life. The accumulated operating experience and performance of low-alloy primary pressure-boundary components is very good world-wide [1–5]. The current fatigue design and evaluation codes (ASME III and XI) have been quite successful in preventing fatigue cracks and failures in LAS components and would therefore seem to be adequate or conservative under most operating circumstances. Instances of EAC have occurred particularly in boiling water reactor (BWR) service, most often in LAS piping and, very rarely, in the RPV itself [1–5]. Oxidising agents, usually dissolved oxygen (DO), and relevant dynamic straining (e.g., arising from thermal stratification, thermal and pressurisation cycles during start-up/shut-down, etc.) were always involved [1–6]. These cases were attributed either to strain-induced corrosion cracking (SICC) or low-frequency corrosion fatigue (LFCF) (Table 15.1) [2, 4]. Operational experience [1–5] and laboratory background knowledge [7] both indicate, that the current fatigue design curves in Appendix A of the ASME Boiler and Pressure Vessel Code, Section XI, might be non-conservative for certain critical, short-lived, BWR plant transients (start-up/shut-down, hot stand-by, thermal stratification, etc.) and that SICC or very low-cycle corrosion fatigue covers the most important gap in the field of EAC of LAS. There is a relevant lack of quantitative SICC/LFCF crack growth data under these critical conditions (i.e., at slow strain rates or very low cyclic frequencies (< 10–2 Hz), intermediate temperatures (150 to 270 ∞C) and high corrosion potentials (ECP)) [1–9]. Under these combinations of temperature and strain rate, weld and weld HAZ materials could eventually reveal higher CF CGR 211

212


Table 15.1 Basic types of environmentally-assisted cracking Environmentally-assisted cracking (EAC) Mechanism

SCC Stress corrosion cracking

SICC Strain-induced corrosion cracking

CF Corrosion fatigue

Type of loading

Static

Slow monotonically rising or very lowcycle

Cyclic: low-cycle, high-cycle

LWR operation condition

Transient-free, steady-state power Start-up/shut-down, Thermal fatigue, operation thermal stratification thermal stratification, …

Characterisation of crack growth

BWRVIP-60 Disposition Lines

?

ASME XI, Code Case N-643 (PWR)

Characterisation of crack initiation

?

Susceptibility conditions: ECPcrit, de/dtcrit, ecrit

ASME III, Fenv-approach

than the base metal, because of dynamic strain ageing (DSA) [9, 10] or hydrogen-induced EAC (if hardness of HAZ > 350 HV) [11, 12]. An EAC project [9] was therefore started at Paul Scherrer Institute (PSI) to evaluate and assess the adequacy and conservatism of the current reference fatigue crack growth curves in LWR coolant environment in Appendix A of the ASME Boiler and Pressure Vessel Code, Section XI (‘ASME XI wet fatigue CGR curves’) [13] and of existing crack growth models (GE model [3]) under these critical conditions. In this project different RPV steels and weld filler/weld HAZ materials were investigated. This chapter presents some important results and conclusions of this experimental parameter study. Special emphasis is placed on loading frequency, temperature and material effects.

15.2


15.2.1 Materials Five different types of low-alloy, nuclear grade RPV steels (base metal and HAZ) with either a low, medium, or high sulphur (and aluminium) content and a RPV weld filler material were investigated (Tables 15.2 and 15.3) [8, 9, 14]. The weld filler and weld HAZ materials were taken from the circumferential core girth weld of a German pressurised water RPV (Biblis C, 1976), which has not been commissioned. All base materials were quenched and tempered. The weld filler, weld

Material

C

Si

Mn

P

S

Cr

Mo

Ni

V

Al

Cu

20 MnMoNi 5 5

A

0.210

0.25

1.26

0.004

0.004

0.15

0.50

0.77

0.008

0.0130

0.06

SA 508 Cl. 2

B

0.210

0.27

0.69

0.005

0.004

0.38

0.63

0.78

0.006

0.0150

0.16

SA 533 B Cl. 1

C

0.250

0.24

1.42

0.006

0.018

0.12

0.54

0.62

0.007

0.0300

0.15

22 NiMoCr 3 7

D

0.215

0.20

0.91

0.008

0.007

0.42

0.53

0.88

0.007

0.0180

0.04

Weld Filler

E

0.054

0.17

1.19

0.013

0.007

0.04

0.55

0.94

0.006

0.0053

0.06

20 MnMoNi 5 5

F

0.260

0.32

1.44

0.016

0.015

0.15

0.61

0.63

0.020

0.0290

0.17

HAZ of D

G

0.215

0.20

0.91

0.008

0.007

0.42

0.53

0.88

0.007

0.0180

0.04

Corrosion fatigue crack growth behaviour of low-alloy RPV steels

Table 15.2 Chemical compositions of investigated steels in wt.%

213

214

Material

288 ∞ C

S [wt.%]

Al [wt.%]

Nfree [ppm]

Heat treatment

Microstructure

RP [MPa]

DSAindex

20 MnMoNi 5 5 (∫ SA 508 Cl.3)

A

0.004

0.013

30

910–920∞C/6h/WQ, 640–650∞C/9.5h/FC

Bainitic

418

–12.3% +++

SA 508 Cl.2 (∫ 22 NiMoCr 3 7)

B

0.004

0.015

2

900∞C/ 8h/WQ 600∞C/9h/AC

Bainitic/ ferriticpearlitic

396

–16.4% +++

SA 533 B Cl.1 (∫ 20 MnMoNi 5 5)

C

0.018

0.030

10 to 100 Hz), where significant environmental effects on fatigue crack growth can be neglected or excluded, were not defined in ASME XI. For these reasons, a modification of the ‘ASME XI fatigue CGR curves’ or the development of a new code case for BWR/NWC should be taken into consideration. Based on the presented


227

results, the development of more realistic reference curves should consider both the strong effect of loading frequency/strain rate and ECP. A differentiation by material parameters (e.g., steel sulphur content) does not seem to be necessary from an engineering point of view. Different curve sets could be developed for several ECP regimes (e.g., BWR/NWC and BWR/HWC). Any of such procedures would result in more complicated flaw tolerance evaluations than so far, since the loading frequency/strain rate of different transients had to be considered in an adequate way, but it would have the potential to reduce both uncertainty and undue conservatism.

15.3.6 Superposition model and time-domain evaluation of LFCF CGR results A simple linear superposition model [9], which considers both frequency and ECP effects, is briefly outlined in the following paragraphs as one possible way for the development of new reference curves. In this model, the cyclebased CGR in high-temperature water Da/DNEAC is just a simple linear superposition of the cycle-based CGR in air Da/DNAir by pure mechanical fatigue and of the corrosion-assisted CGR Da/DNENV. The first contribution is a purely cyclic-controlled process and independent of loading frequency. The second contribution only occurs during the rising load part of the fatigue cycle and is strongly dependent on crack-tip strain rate de/dtCT (da/dtEAC = A · (de/dtCT)n) and loading frequency. Under cyclic loading conditions it is assumed, that the crack-tip strain rate is proportional to the experimentallyderived and known fatigue CGR in air (de/dtCT a da/dtAir) under otherwise identical loading conditions. For the onset of fast EAC, the fatigue CGR in crit , which is dependent on the air da/dtAir has to exceed a critical CGR da / dt Air ECP. Based on these assumptions, the CF crack growth in high-temperature water can be described by the following equations, which are the basis of the so-called ‘time-domain analysis method’: crit 1. da / dt Air < da / dt Air = f (ECP) :

da/dtEAC = da/dtAir = f (DtR, DK, R)

crit 2. da / dt Air ≥ da / dt Air = f (ECP) :

da / dt EAC C ◊ ( da / dt Air ) m + da / dt Air

C, m = f (ECP) These equations can be easily transformed in a cycle-based form by dividing crit have to be them by the loading frequency. The parameters C, m and da / dt Air conservatively determined by experiments for different ECP regimes, (e.g., crit ,C BWR/NWC-, BWR/HWC- and PWR-conditions). In contrast to da / dt Air crit , and m are expected to be only slightly dependent on ECP. Based on da / dt Air thresholds DKEAC = f (n, R) and critical frequencies ncrit = f (DK, R) for the onset of EAC can be derived. Such a time-domain analysis for a large data base of LFCF tests [9] in

228


oxygenated high-temperature water (simulated BWR/NWC operating conditions) is shown in Fig. 15.12. The air fatigue CGR da/dtAir have been calculated according to Eason [15]. The test conditions covered a wide range of environmental (240 to 288 ∞C, 0.4 to 8 ppm DO, k = 0.06 to 0.25 mS/cm, 0 mVSHE, ≥ 0.4 ppm DO):

15.4.1 Temperature/loading frequency For all frequencies and materials, both cycle-based CGR Da/DNEAC and time-based CGR da/dtEAC increased with increasing temperature from 150 to 250 ∞C, where a maximum/plateau could be observed with further increasing the temperature to 288 ∞C. With decreasing frequency the CF crack advance per cycle Da/DNEAC generally increased for all temperatures and materials. Sustained EAC growth could be maintained down to low frequencies of 10–5 Hz. The time-based LFCF CGR da/dtEAC were in the range of 5 ¥ 10–11 m/s (1.6 mm/a) to 5 ¥ 10–8 m/s (160 mm/a) and decreased with decreasing loading frequency.


229

15.4.2 Material aspects The RPV steels with low and high sulphur content and the weld filler/weld HAZ material showed very similar LFCF CGR over a wide range of loading conditions. A possible effect of steel sulphur content was only observed at 288 ∞C at very low loading frequencies n £ 3 ¥ 10–5 Hz, where Da/DNEAC seemed to increase with increasing sulphur content. The maximum/plateau of CF CGR at/above 250 ∞C and the higher cycle-based CGR Da/DNEAC for the low-sulphur steels with a high DSA susceptibility at intermediate temperatures and very low loading frequencies n £ 3 ¥ 10–5 Hz clearly indicated that DSA might affect the LFCF behaviour. The concentration of ‘free’, interstitial nitrogen and carbon might therefore be just as relevant for EAC susceptibility as the steel sulphur content, at least under conditions where DSA is observed.

15.4.3 Comparison to the GE model and to ASME XI The LFCF CGR of all materials were conservatively covered by the ‘highsulphur line’ of the GE model for all temperatures and frequencies. The transition curves between the ‘low- and high-sulphur line’ seem to be nonconservative under these highly oxidising conditions and the model therefore crit . The current predicts too high critical frequencies ncrit and CGR da / dt Air ‘ASME XI wet fatigue CGR curves’ could be significantly exceeded by cyclic fatigue loading at low frequencies (< 10–2 Hz) for low- and highsulphur steels as well as the RPV weld filler/HAZ materials and low and high load ratios in the temperature range between 150 and 288 ∞C. They do not adequately describe and conservatively cover the experimentally observed CF crack growth behaviour of LAS under BWR/NWC conditions. The development of more realistic reference fatigue crack growth curves for a new BWR/NWC code case should therefore be taken into consideration. A simple superposition model/time-domain evaluation method, which includes both frequency and ECP effects, could be used for that purpose.

15.5

Acknowledgements

The financial support for this work by the Swiss Federal Nuclear Safety Inspectorate (HSK), the Swiss Federal Office of Energy (BFE) and the Swiss Federal Office for Education and Science (BBW) is gratefully acknowledged. Thanks are also expressed to U. Ineichen, U. Tschanz, B. Gerodetti, and E. Groth (all PSI) for their experimental contribution to this work.

15.6

References

1. P. Scott, D. Tice, Nuclear Engineering and Design, 119, pp. 399–413, 1990.

230


2. J. Hickling, D. Blind, Nuclear Engineering and Design, 91, pp. 305–330, 1986. 3. F.P. Ford, ‘Environmentally-Assisted Cracking of Low-Alloy Steels’, EPRI NP7473-L, Electric Power Research Institute, Jan. 1992. 4. Y.S. Garud, S.R. Paterson, R.B. Dooley, R.S. Pathania, J. Hickling, A. Bursik, ‘Corrosion Fatigue of Water-Touched Pressure Retaining Components in Power Plants’, EPRI TR-106696, Final Report, Nov. 1997. 5. H.P. Seifert, S. Ritter, J. Hickling, Power Plant Chemistry, 6, pp. 111–123, 2004. 6. E. Lenz, N. Wieling, Nuclear Engineering and Design, 9, pp. 331–344, 1986. 7. O.K. Chopra, W.J. Shack, Nuclear Engineering and Design, 184, pp. 49–76, 1998. 8. S. Ritter, H.P. Seifert, Power Plant Chemistry, 5, pp. 17–29, 2003. 9. H.P. Seifert, S. Ritter, U. Ineichen, U. Tschanz, B. Gerodetti, ‘Risskorrosion in druckführenden Komponenten des Primärkreislaufes von SWR’, BFE-Final Report, PSI, Switzerland, Feb. 2003. 10. H. Hänninen et al., ‘Effects of Dynamic Strain Aging on Environment-Assisted Cracking of Low-Alloy Pressure Vessel and Piping Steels’, 10th Int. Conf. on Env. Degr. of Mat. in Nucl. Power Systems – Water Reactors, CD-ROM, Paper No. 47, Lake Tahoe, NV, USA, Aug. 6–10, 2001. 11. M. Tsubota et al., ‘Intergranular Stress Corrosion Cracking of Low-Alloy and Carbon Steels in High-Temperature Water’, Proc. 6th Int. Conf. on Env. Degr. of Mat. in Nucl. Power Systems – Water Reactors, pp. 53–58, San Diego, CA, USA, Aug. 1– 5, 1993. 12. J. Hickling, ‘Wasserstoffinduzierte Spannungsrisskorrosion in niedriglegierten Stählen’, 4th MPA-Seminar, Paper No. 7, Stuttgart, Germany, Oct. 4–5, 1978. 13. ASME XI Appendix A-4300, ASME Boiler & Pressure Vessel Code, Section XI, Rules for In-Service Inspection of Nuclear Power Plant Components, Appendix A, Article A-4000, Subsection A-4300: Fatigue Crack Growth Rate, ASME, New York, 1998. 14. S. Ritter, H.P. Seifert, ‘Characterisation of the Lower Shell and Weld Material of the Biblis C Reactor Pressure Vessel’, PSI-Report 02-01, ISSN 1019-0643, Jan. 2002. 15. E.D. Eason et al., Nuclear Engineering and Design, 184, pp. 89–111, 1998. 16. F.P. Ford, P.L. Andresen, ‘Corrosion Fatigue of A533B/A508 Pressure Vessel Steels in 288 ∞C Water’, 3rd Int. IAEA Specialist’s Meeting on Subcritical Crack Growth, 1, pp. 105–124, W. Cullen, Ed., NUREG/CP-0112, Moscow, USSR, May 14–17, 1990.

16 Effect of cyclic loadings on the stress corrosion crack growth rate in alloy 600 in PWR primary water C. G U E R R E, O. R A Q U E T and L. D U I S A B E A U, CEA, France and G. T U R L U E R, IRSN, France

16.1

Introduction

Alloy 600 is widespread in Pressurized Water Reactors (PWR). Results of tests, generally performed with monotonic loadings (constant load, constant deformation or constant extension rate), show that this alloy is sensitive to Stress Corrosion Cracking (SCC) in primary water of PWR. For a few years, there has been a growing interest in assessing the role and the effects of cyclic loadings or stress transients on the Primary Water Stress Corrosion Cracking (PWSCC) behavior of metallic components like Alloy 600. For instance, as French nuclear power plants operate at variable power that may generate cyclic loadings on the components, the determination of the effects of cyclic loadings is of great importance. Congleton et al. [1] concluded there was an accelerating effect of cycling loading on Crack Growth Rate (CGR). Bosch and Vaillant [2–4] agreed with an environmentally assisted cracking for frequencies below 0.01 Hz. Concerning the influence of the wave form, Lidar [5] concluded that the saw-tooth form was the most damaging one. This study takes place in a program with the support of the French Institute for Radiological Protection and Nuclear Safety (IRSN) that focuses on the effect of low frequency and high R ratio cyclic loading conditions on Alloy 600 in order to improve the knowledge of the conditions leading to enhanced CGRs and to assess some testing procedures allowing ‘gentle cycling’ or periodic load discharge to generate steady CGRs. Two heats are tested, the first one (heat 3110439) was air melted by Allegheny Ludlum and the specimens were machined in a hot rolled plate. The second one (heat WL344) was produced by Techphy.

16.2

Materials and specimen

The composition of the heat 3110439 is given in Table 16.1. The plate was cut in a hot rolled larger plate (30% cold-work). The microstructure (Fig. 231

232

RCC-M4101

C

Mn

Si

S

Ni

Cr

Cu

Co

Fe

P

Ti

Mo

N

Al

CB

0.5 MeV was 2.2 ¥ 109 cm–2 and at E > 0.1 MeV – 2.9 ¥ 109 cm–2. While considering and analysing the investigation results, the previous published data [5], obtained from testing the 08Cr-18Ni-10Ti steel wrapper of the emergency assembly removed after 30 years of operation in the VK50 reactor, have been used. The indicated element had failures of corrosion origin in the weld region.

19.3

Experimental results

19.3.1 Metallography Cracks are observed on ring 1 in the weld and heat affected zone (Fig. 19.3a). The outer surface breaking cracks are not revealed. On the outer surface of rings 1 and 2 only the intergranular corrosion is observed. The intergranular corrosion, corrosion cracking (Fig. 19.4a, b) and undersurface corrosion exfoliation are detected on rings 3–5 (Fig. 19.5a, b). The maximum corrosion depth is at the core top level and achieves 3–3.5 mm.

19.3.2 Electron microscopy Electron microscopy of samples cut out from ring 5, having a minimum neutron fluence, showed that the steel structure is characterised by the presence of austenite equiaxial grains where one may find primary precipitates of

276

Corrosion issues in light water reactors (a)

(b)

Magnification ¥ 200

19.3 Cracking and intergranular corrosion at the lower part of the measurement channel (¥ 200): weld and heat affected zone of ring No. 1(a); base metal of ring No. 2(b).

(a)

(b)

19.4 Corrosion cracking (¥ 200): ring No. 3(a), ring No. 4(b).

(a)

(b)

19.5 Under-surface corrosion (exfoliation) of the base metal (ring 5) ¥ 200 (a); ¥ 5 (b).

Corrosion damage to 18Cr-9Ni-Ti steel

277

excess phases up to 1 mm and clusters of linear dislocations. The radiation defects are not revealed. As shown in Fig. 19.6, the precipitation of the secondary phases from few tens of micrometers to some hundreds of nanometers occurs in the 08Cr18Ni-10Ti steel under irradiation. They are mainly located along the grain boundaries and in the adjacent areas (Fig. 9.6 a, b). Their chemical composition is as follows: 37–46 Fe, 11–15 Cr, 32–45 Ti, 5–6 Ni at %. These precipitates in terms of their structure are evidently Me23C6 or Me6C type carbides. The large precipitate, shown in Fig. 19.6a, represents TiC. The fine-dispersed precipitates of the rounded shape, presumably the G-phase, were discovered in the grain body. Their average size is in the order of 8 nm, the density 5 ¥ 1015 cm–3. The average diameter of dislocation loops was 9 nm and their density 1.4 ¥ 1016 cm–3. The clusters of the dislocation loops and linear dislocations are revealed in the samples from rings 2–4. The average size of the dislocation loops is ~10 nm, the maximum one is 20 nm. The loop concentration is estimated as 4 ¥ 1016 cm–3. The crystallographic and elemental analysis of the secondary phase precipitates showed that the precipitates, both large and fine dispersed, are for the most part titanium carbides (TiC) with a FCC lattice.

19.3.3 Fractography Fractographic observations were performed on the samples of 4 ¥ 4 ¥ 1 mm, cut out from rings 2–4. Before the examinations, the samples were mechanically fractured in order to obtain the fresh fracture surface. A region, remarkable for the fracture nature and having all features of brittle fracture, was revealed on the fracture surface at the outer face of the measurement channel. Its depth is ~3–5 mm for ring 2, 30 mm for ring 3 (Fig. 19.7a) and 100 mm for ring 4 (Fig. 19.7b) that is in a good correlation with the steam content in the coolant (Table 19.1).

500 nm

500 nm

(a)

(b)

200 nm

(c)

19.6 Microstructure of the tube material after irradiation – Transmission electron microscopy.

278


(a)

(b)

19.7 Fracture surface of the samples cut out from the ring 3 (a) and 4 (b). Table 19.1 Depth of the corrosion failure of the measurement channel at different fluences and coolant densities Ring No.

Point of cutting out

Neutron fluence E > 0.1 MeV

Coolant density (g/cm3)

Corrosion failure depth (mm)

1 2 3 4 5

130 cm below the core center 90 cm below the core center 3 cm below the core center 90 cm above the core center 550 cm above the core center

3.0¢1010 2.7¢1021 5.1¢1021 1.7¢1021 2.1¢109

1.000 0.665 0.375 0.305 £ 0.305

0.1 0.1 1.0 3.5 1.8

19.3.4 Auger spectroscopy Auger spectroscopy was carried out on the samples cut out from ring 4 in order to evaluate possible modification of the chemical composition of the superficial brittle layer. The choice was governed by the fact that the upper part of the measurement channel is mostly affected by the corrosion attack and has the maximum thickness of the brittle superficial layer. The examinations were carried out on the fresh fragment of the fracture surface of 2 ¥ 2 ¥ 0.3 mm. To obtain the averaged concentration values of various elements on the surface, the primary electron beam of about 3 mm diameter was scanned in a raster to light up a surface point of 200 ¥ 200 mm on the sample. After registering Auger spectra (10–12 atomic layers) from the fracture surface, multiple ion etching of the examinated surface was performed for analysing the elemental composition in depth. The thickness of the remote layer was calculated from the etching rate, equal approximately to 20 Angstroms per minute. Before measuring Auger peak intensities of the alloying elements


279

(Fe, Cr, Ni and Ti) and carbon, ion purification of the surface was carried out. During the operation the concentration of the main alloying elements in the material matrix was not modified, but very high carbon value was observed at the outer face washed by coolant (Fig. 19.8). As shown in Fig. 19.9a, no chrome and nickel and a very high content of titanium is observed on the fracture surface. Titanium value is about 60 times as high as that one in the matrix. The depth of the modified layer varies from 0.2 to 1.2 mm. The carbon content decreases up to the initial state level at the distance of ~1.2 mm, whereas the high level of oxygen is not modified and even at the depth of 1 mm is in the order of 40 at% (Fig. 19.9b). The depth of the sulfur and chlorine layer is approximately 0.2 and 0.4 mm respectively and the copper layer is in the order of 1 mm. The calcium content is strongly decreases in the layer of 0.2 mm thickness, but it is entirely disappeared at a depth of 1 mm (Fig. 19.9c). To provide information on the uniformity or irregularity of various element distribution on the fracture surface, the registration of elemental charts series was performed at the accelerating voltages of 3 and 9 kV. Figure 19.10 presents the elemental charts of the fracture surface of one of the grains located in the center of each photo. As seen in Fig. 19.10a, chlorine is distributed uniformly. In contrast sulfur is concentrated mainly in the clusters of 5–6 mm (Fig. 19.10b). It is especially good seen with lower magnifications at the accelerating voltage of 3 kV (Fig. 19.10c). As a rule, an elevated concentration of iron (Fig. 19.10e), chrome (Fig. 19.10h), as well as nickel is noted in the sulfur clusters. These sulfur clusters appear to represent complex sulfide compounds containing iron, chrome and nickel. Copper is distributed in the form of separate clusters along the edges of the grain selected for examination (Fig. 19.10d). An elevated copper concentration is observed in several points, located along the fracture pattern perimeter, and it is absent at the surface of the fracture itself. The carbon Atomic part of carbon (%)

25 20 15 10 5 0 0

100

200 300 400 500 600 Distance from the outer edge of the tube (mm)

700

19.8 Carbon distribution along the thickness of the measurement channel wall (ring No. 4) in its upper part.

800

Corrosion issues in light water reactors 80 Concentration (%)

Fe 60

40 Cr

20

Ni Ti

0 0

0.2 0.4 0.6 0.8 1 Distance from the failure surface (mm)

1.2

(a)

Concentration (%)

60 O 40

20

C 0 0

0.2 0.4 0.6 0.8 1 Distance from the failure surface (mm) (b)

1.2

4

Concentration (%)

280

3

Cu

2

1 Cl

Ca

S 0 0

0.2 0.4 0.6 0.8 1 Distance from the failure surface (mm)

1.2

(c)

19.9 Concentration of the main added elements (a), minus the light ones and impurities), oxygen and carbon (b), as well as sulfur, chlorine, copper and calcium (c) in the irradiated steel as a function of the distance from the fracture surface.


281

A B

(a) Chlorine

(b) Sulfur

(c) Sulfur

(e) Iron

(f) Oxygen

C (d) Copper

D

(g) Carbon

(h) Chromium

(i) Absorbed electrons

19.10 Charts of the elemental distribution on the fracture surface of the irradiated steel (a, b, d-i ¥ 1400; c ¥ 400; A, B, C, D, – points where the local elemental composition is determined).

distribution chart (Fig. 19.10g), the Auger peak of which has the maximum intensity, in terms of topographic nuances, bears a strong resemblance to a fracture surface image in absorbed or secondary electrons (Fig. 19.10i). There are no particular carbon clusters, unless the areas are available with some lower carbon concentration in the points where an elevated oxygen content is observed (Figs 19.10f, g). The elevated oxygen content in the right lower part of the fracture patter (Fig. 19.10f) is accompanied by an elevated nickel content and a slightly heightened content of chrome (Fig. 19.10h).

19.4

Discussion

During the entire operation period the measurement channel has been subjected to small longitudinal tensile stresses in the range of 0.5–0.7 Mpa, not able to provoke stress corrosion cracking, even at the critical concentration of oxygen and chlorine in the coolant. Accumulation of chlorine and carbon at the metal/coolant interface and their diffusion in the metal thickness is a more serious factor [2]. So, when examinating the emergency assembly wrapper operated during 30 years in the same reactor, a higher chlorine content was detected at a depth up to 0.4 mm from the fracture surface (Fig. 19.9c). The

282


results of this work show that the saturation of the measurement channel surface with carbon up to 20 at.% is observed in the layer of 30–100 mm deep, resulting in the brittle fracture. The carbonisation effect leds to generation of the internal stresses which may be sufficient to initiate and propagate stress corrosion cracking. The properties of one or other section of the measurement channel depend on its location: the farther off the core bottom, the greater is the carbon saturation and the thickness of the brittle layer (p. 2.3–2.4). It evidently relates to the considerably variable coolant density along the axis. The lower is the density (at a high fluence), the greater is the depth of the corrosion failure (Table 19.1). The coolant density in the lower core part of the reactor under operation is 0.735 g/cm3. In this region, only intergranular corrosion is observed on the measurement channel. The steam content increases approaching the core center, the coolant density decreases up to 0.375 g/cm3, and intergranular corrosion and corrosion cracking achieves 1 mm depth. Futher decreasing of the coolant density to 0.304 g/cm3 at the level of the core top is followed by a sharp increase in the corrosion depth. It appears to be related to the elevated oxygen concentration in steam, carbonisation of the tube surface, possibility to concentrate the chloride impurities, when evaporating multiple water portions, and to facilitate oxygen access. The neutron fluence varies along the axis of the measurement channel from 2.1 ¥ 109 cm–2 to 5.1 ¥ 1021 cm–2, however all these areas are revealed to have corrosion failures. As seen in Table 19.1, the corrosion depth of the measurement channel material is not a strong function of the fluence value. The results lead to the following conclusions: ∑ ∑ ∑ ∑

Failure occurs exceptionally along the grain boundaries. Ti segregates and the Cr and Ni content decreases on the grain boundaries. There are harmful corrosive elements such as chlorine, sulfur and copper on the failure surface. Segregation thickness on the failure surface is about 1–1.5 mm; in any case the segregation has practically disappeared after ion etching at a depth ~3 mm (Fig. 19.11).

The elemental charts give the qualitative representation of the elemental distribution. The quantitative information can be provided from analysing the local elemental composition in the separate points with an electron probe diameter of 0.1–0.3 mm (in this study in the points ‘ACE’ in Fig. 19.10, Table 19.2). The same table presents for comparison the integral measurements of the elemental composition in the raster 200 ¥ 200 mm, the surface of which comprises the examined fracture pattern. The results of the local measurements are in a good agreement with the qualitative information provided by the


(a) Sulfur

(b) Iron

283

(c) Absorbed electrons

19.11 Charts of sulfur and iron distribution on the failure surface of the irradiated steel after the ion etching at a depth in the order of 3 mcm ¥ 1500. Table 19.2 Local elemental composition of the failure surface (‘A’, ‘B’, ‘C’ and ‘D’ refer to Fig. 19.10) Element

Fe Cr Ni Ni C O Cl S Cu Ca

Content at % ‘A’

‘B’

‘C’

‘D’

Raster

4.0 0.5 3.8 1.0 75.9 13.0 0.5 0.4 0.8 0

12.2 2.8 5.1 0.6 61.7 13.1 0.4 2.7 1.4 0

1.7 0.8 1.4 0.6 82.5 6.3 0.6 0.1 6.2 0

7.4 1.6 6.5 1.2 47.0 30.8 0.3 0 3.0 2.3

2.4 0.6 0.3 0 91.5 3.9 0.3 0.4 0.5 0

elemental charts. As shown in the table, only chlorine is rather uniformly distributed on the failure surface. The concentration of the other elements varies in a wide range. The references [1–7] and the findings permit us to reveal three groups of factors: (1) chief factors, (2) governing factors and (3) related factors.

19.4.1 Chief factors Tensile stresses The cracks have been observed only in the upper part of the tube in the vicinity of the weld (Fig. 19.3a), where the tensile stresses, according to the work results, can reach 300 MPa. As the cracks have not been observed below, in the base metal, where the neutron fluence was increased, one may consider that the tensile stresses are of importance in the radiation-induced intergranular corrosion cracking and prevail over the radiation-induced processes of segregation/depletion.

284


Sensitisation The weld adjacent zone is heated to 550–800 ∞C under welding, which results in the precipitation of the large Me23C and/or Me6C type carbides of some hundreds nanometers size (Fig. 19.6), containing about 13% chromium, 42% iron, 39% titanium and 6% nickel, on the grain boundaries and subboundary areas. As a result, the grain boundaries are depleted in chrome. Secondary source of depletion is the purely thermal low-temperature sensitisation in the temperature range of 250–350 ∞C. As estimated, the time needed for its proceeding at the temperature of 285 ∞C is 12 years [6] that is two and a half times lower than the service life of this tube. And the third, probably the most powerful source of the grain boundary depletion in chromium, is caused by the radiation-induced processes: (a) rising of the available carbides; (b) radiation-induced depletion of the grain boundary in chromium; (c) precipitation of the fine-dispersed phases containing chromium. In the work [7] the segregation processes are analysed, which occur on the grain boundaries of steel 304 in the initial state and after irradiation at the temperature of 300 ∞C (Table 19.3). As indicated in Table 19.3, the results of the energo-dispersion analysis (EDA) show that at the initial state, a notable segregation of chromium and silicon and a strong segregation of molybdenum and phosphorus are observed on the grain boundaries. According to the results of the Auger electronic Table 19.3 Irradiation effect on the chemical composition modification of the grain body and sub-boundary zones of steel 304 at the temperature of 300 ∞C Element

Fe Cr Ni Mo Mn Si C P S

EDA TEM (mass %) Initial state

Initial state

Grain body

Auger electron analysis (at.%) 5 ¥ 1021 n/cm2, grain boundary Grain boundary

Base 18.54 8.28 0.32 1.52 0.55 0.069 0.023 0.021

62–64 24–25 9 1.7–1.9 n/d 1 n/d 0.8 n/d

61 16 15–16 1 n/d 4.2–4.5 n/d 1.1 n/d

EDA – energo-dispersion analysis n/d – not determined

5 ¥ 1021 n/cm2

Grain body

Grain boundary

Grain body

Grain boundary

69.39 19.58 7.75 0.18 n/d 1.08 n/d 0.04 n/d

74.9 16.4 7.8 n/d n/d n/d n/d 0.7 n/d

68.4 18.7 8.43 1.53 n/d 1.25 n/d 0.06 n/d

60.6 14.2 14.1 0.2 n/d 8.3 n/d 1.9 n/d


285

analysis, the segregation of chrome is not confirmed. However, the data of the energo-dispersion analysis show that after irradiation, the grain boundaries are depleted in chromium, molybdenum and enriched in nickel, phosphorus and especially silicon. The results of the energo-dispersion analysis and Auger spectroscopy are in qualitative agreement. The results from both investigation methods are in a good agreement for chrome. But opposite results are obtained for nickel: in this study the grain boundary is observed to be depleted in nickel up to zero, where as, in work [7] the grain boundary is observed to be enriched in nickel. Thus, one may conclude that the radiation-induced segregation on the grain boundaries is a second significant factor affecting radiation-induced intergranular corrosion cracking.

19.4.2 Governing factors Concentration of oxygen A steam-water mixture as high as 0.2–0.5 mg/l in water and up to 20–50 mg/ l in steam due to radiolysis [1, 2] favors increasing corrosion potential with corrosion rates in intergranular channel peaks. Chlorine precipitation on the boundaries The coolant is probably the principal source of chlorine. In this case the possibility to accumulate chlorine at the steam-water interface favors the initiation of chloride cracking even when the ion-chlorine composition is 1 (Feoxide deposits)

Fe(OH)3 – 42 Fe3O4 – 16 austenite – 42

–(XRD)

–(XRD)

Decontaminated (2001)

2002

7

ª 0.5

1

Fe(OH)3 – 63 Fe3O4 – 24 austenite – 13

–(XRD)

–(XRD)


2002

9

ª 0.6

1

Fe(OH)3 – 46 austenite – 54

NA

NA


2000

11

ª 0.3

Corrosion Issues in Light Water Reactors: Stress Corrosion Cracking (EFC51)

Stress Corrosion Cracking: Theory and Practice

Electrochemistry in Light Water Reactors: Reference Electrodes, Measurement, Corrosion and Tribocorrosion Issues (EFC 49)

Corrosion and Corrosion Control

Corrosion

Corrosion

Corrosion Engineering Handbook (Corrosion Technology)

Corrosion localisée

Developments in Marine Corrosion

Super Light Water Reactors and Super Fast Reactors: Supercritical-Pressure Light Water Cooled Reactors

Corrosion Inhibitors

Corrosion-Resistant Linings and Coatings (Corrosion Technology)

Underground Corrosion

Corrosion Chemistry

Corrosion and Corrosion Control 4th Edition

Principles of Corrosion Engineering and Corrosion Control

Corrosion in Refineries

Principles of Corrosion Engineering and Corrosion Control

Electrochemical Techniques in Corrosion Science and Engineering (Corrosion Technology)

Amine unit corrosion in refineries EFC46 (European Federation of Corrosion)

Oxide Fuels for Light Water Reactors and Fast Neutron Reactors

Radiation Protection at Light Water Reactors

Corrosion of Steel in Concrete

Encyclopedia of Corrosion Technology

Encyclopedia of corrosion technology

Microbiologically Influenced Corrosion

Encyclopedia Of Corrosion Technology

Corrosion-Resistant Piping Systems

Corrosion in reinforced concrete structures

Electrochemistry and corrosion science

Corrosion of Weldments

Corrosion Issues in Light Water Reactors: Stress Corrosion Cracking (EFC51)

Stress Corrosion Cracking: Theory and Practice

Electrochemistry in Light Water Reactors: Reference Electrodes, Measurement, Corrosion and Tribocorrosion Issues (EFC 49)

Corrosion and Corrosion Control

Corrosion

Corrosion

Corrosion Engineering Handbook (Corrosion Technology)

Corrosion localisée

Developments in Marine Corrosion

Super Light Water Reactors and Super Fast Reactors: Supercritical-Pressure Light Water Cooled Reactors

Corrosion Inhibitors

Corrosion-Resistant Linings and Coatings (Corrosion Technology)

Underground Corrosion

Corrosion Chemistry

Corrosion and Corrosion Control 4th Edition

Principles of Corrosion Engineering and Corrosion Control

Corrosion in Refineries

Principles of Corrosion Engineering and Corrosion Control

Electrochemical Techniques in Corrosion Science and Engineering (Corrosion Technology)

Amine unit corrosion in refineries EFC46 (European Federation of Corrosion)

Oxide Fuels for Light Water Reactors and Fast Neutron Reactors

Radiation Protection at Light Water Reactors

Corrosion of Steel in Concrete

Encyclopedia of Corrosion Technology

Encyclopedia of corrosion technology

Microbiologically Influenced Corrosion

Encyclopedia Of Corrosion Technology

Corrosion-Resistant Piping Systems

Corrosion in reinforced concrete structures

Electrochemistry and corrosion science

Corrosion of Weldments

Recommend Documents