Environmentally-Assisted Cracking of Low-Alloy Reactor ...

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P A U L S C H E R R E R I N S T I T U T

CH0200020

PSI Bericht Nr. 02-05February 2002

ISSN 1019-0643

Nuclear Energy and Safety DivisionLaboratory for Materials Behaviour

Environmentally-Assisted Cracking of Low-AlloyReactor Pressure Vessel Steels under BoilingWater Reactor Conditions

H.P. Seifert, S. Ritter

3 3 / 1 3Paul Scherrer InstitutCH - 5232 Villigen PSITelefon 056 310 21 11Telefax 056 310 21 99

Environmentally-Assisted Crackingof Low-Alloy Reactor Pressure Vessel Steels

under Boiling Water Reactor Conditions

Progress Report of the RIKORR Project

H.P. Seifert, S. Ritter

Nuclear Energy and Safety DivisionLaboratory for Materials Behaviour

Structural Integrity Group

Abstract

The present report summarizes the experimental work performed by PSI on theenvironmentally-assisted cracking (EAC) of low-alloy steels (LAS) in the frame of theRJKORR-project during the period from January 2000 to August 2001. Within this project,the EAC crack growth behaviour of different low-alloy reactor pressure vessel (RPV) steels,weld filler and weld heat-affected zone materials is investigated under simulated transient andsteady-state BWR/NWC power operation conditions.

The EAC crack growth behaviour of different low-alloy RPV steels was characterized byslow rising load (SRL) / low-frequency corrosion fatigue (LFCF) and constant load tests withpre-cracked fracture mechanics specimens in oxygenated high-temperature water attemperatures of either 288, 250, 200 or 150 °C. These tests revealed the following importantinterim results:

Under low-flow and highly oxidizing (ECP > 100 mVSHE) conditions, the ASME XI "wet"reference fatigue crack growth curve could be significantly exceeded by cyclic fatigueloading at low frequencies (< 10"3 Hz), at high and low load-ratios R, and by ripple loadingnear to AKth fatigue thresholds.

The BWR VIP 60 SCC disposition lines may be significantly or slightly exceeded (even insteels with a low sulphur content) in the case of small load fluctuations at high load ratios(ripple loading) or at intermediate temperatures (200 - 250 °C) in RPV materials, whichshow a distinct susceptibility to dynamic strain ageing (DSA).

Summary

Within the RIKORR-project, the EAC crack growth behaviour of different low-alloy RPVsteels, weld filler and weld heat-affected zone materials is investigated under simulatedtransient and steady-state BWR/NWC power operation conditions. System conditions, whichchallenge the conservatism of the current ASME XI reference crack growth curves and of theBWR VIP 60 SCC disposition lines, shall be identified by means of an experimentalparameter study.

The strain-induced corrosion cracking (SICC) / low-frequency corrosion fatigue (LFCF) andstress corrosion cracking (SCC) crack growth behaviour of different low-alloy RPV steelsunder simulated BWR/NWC conditions was characterized by slow rising load (SRL) / low-frequency corrosion fatigue (LFCF) and constant load / periodical partial unloading / rippleload tests with pre-cracked fracture mechanics specimens in oxygenated high-temperaturewater at temperatures of either 288, 250, 200 or 150 °C. Modern high-temperature waterloops, on-line crack growth monitoring and fractographic analysis by scanning electronmicroscopy (SEM) were used to quantify the cracking response. These tests revealed thefollowing important interim results:

SICC/LFCF:

Maximum SICC susceptibility (initiation of crack growth from incipient cracks) wasobserved at intermediate temperatures (= 250 °C) and slow strain-rates («5-10~5 s"1). TheSICC growth rates increased with increasing strain-rate and with increasing temperature witha maximum/plateau at/above 250 °C. SICC initiation and the extent of subsequent growthwere crucially dependent upon simultaneously maintaining a positive tensile crack-tip strain-rate and a high sulphur-anion activity in the crack-tip environment. Under low-flow andhighly oxidizing (ECP > 100 HIVSHE) conditions, the ASME XI "wet" reference fatigue crackgrowth curve could be significantly exceeded by cyclic fatigue loading at low frequencies(< 10"3 Hz), at high and low load-ratios R, and by ripple loading near to AKth fatiguethresholds. Stationary LFCF crack growth could be sustained down to low frequencies of10"5Hz. The LFCF growth behaviour of low- and high-sulphur steels was comparable over awide range of loading conditions.

SCC:It is concluded that there is no susceptibility to sustained SCC crack growth at temperaturesaround 288 °C under purely static loading, as long as small-scale-yielding conditions prevailat the crack-tip and the water-chemistry is maintained within current BWR/NWC operationalpractice (EPRI water-chemistry guidelines). However, sustained, fast SCC (with respect tooperational time scales) cannot be excluded for faulted water-chemistry conditions (> EPRIAction Level 3) and/or for highly stressed specimens, either loaded near to Ku or with a highdegree of plasticity in the remaining ligament. The conservative character of the BWR VIP60 disposition lines 1 and 2 for SCC crack growth in low-alloy steels (LAS) has beenconfirmed by this study for 288 °C and RPV base material (< 0.018 wt.% S).

Preliminary results indicate, that these disposition lines may be significantly or slightlyexceeded (even in steels with a low sulphur content) in the case of small load fluctuations athigh load ratios (ripple loading) or at intermediate temperatures (200 - 250 °C) in RPVmaterials, which show a distinct susceptibility to dynamic strain ageing (DSA). In the case ofperiodical partial unloading (PPU), the SCC crack growth rates fall below the BWR VIPdisposition line 2 for long hold periods at maximum (constant) load > 10 h. Furthermore,cessation of crack growth was often observed for long hold periods.

Table of contents

ABSTRACT 2

SUMMARY 3

0. ABBREVIATIONS AND SYMBOLS 6

0.1 ABBREVIATIONS 6

0.2 SYMBOLS AND UNITS 8

1. INTRODUCTION: EAC IN LOW-ALLOY RPV AND PIPING STEELS 10

1.1 TYPES OF EAC 10

1.2 OPERATING EXPERIENCE 11

1.3 EXPERIMENTAL BACKGROUND KNOWLEDGE 11

1.3.1 Corrosion fatigue / strain-induced corrosion cracking 111.3.2 Stress corrosion cracking 12

2. RIKORR RESEARCH PROGRAMME 13

2.1 OBJECTIVES OF THE RIKORR RESEARCH PROGRAMME 13

2.2 WORKING PROGRAMME 14

3. MATERIALS AND EXPERIMENTAL PROCEDURE 15

3.1 MATERIALS 15

3.2 SPECIMEN FORM / PRE-CRACKING 16

3.3 MECHANICAL LOADING 16

3.4 WATER-CHEMISTRY 19

3.5 ELECTROCHEMICAL POTENTIAL 19

3.6 ON-LINE CRACK GROWTH MONITORING 21

3.7 POST-TEST EVALUATION 22

4. RESULTS AND DISCUSSION 23

4.1 STRAIN-INDUCED CORROSION CRACKING / SLOW RISING LOAD TESTS 23

4.1.1 Effect of loading rate 234.1.2 Effect of water-chemistry 254.1.3 Effect of temperature 294.1.4 Effect of steel sulphur and aluminium content 324.1.5 Effect of microstructure 334.1.6 Effect of specimen orientation and location 36

4.2 STRAIN-INDUCED CORROSION CRACKING / VERY LOW FREQUENCY CF TESTS 38

4.3 CORROSION FATIGUE / LOW-FREQUENCY CORROSION FATIGUE TESTS 43

4.3.1 Effect of temperature and loading frequency 434.3.2 Effect of loading conditions 484.3.3 Comparison to literature data 524.3.4 Comparison to ASME XI and to a proposal for an upgrading 534.3.5 Comparison to the GE-model 554.3.6 Comparison to SRL- and vLFCF-tests 58

4.4 STRESS CORROSION CRACKING 60

4.4.1 Summary of the main results ofSpRKII 604.4.2 Effect of temperature 714.4.3 Ripple loading 74

4.4.4 Periodical partial unloading 774.4.5 Effect of microstructure 784.4.6 Comparison with BWR VIP 60 SCC disposition lines and literature data 794.4.7Summary of SCC 81

4.5 DYNAMIC STRAIN AGEING 83

4.5.1 Background 834.5.2 Phenomenological features and metallurgical mechanism ofDSA 844.5.3 Characterization ofDSA response by tensile tests 924.5.4 Characterization ofDSA response by internal friction measurements 984.5.5 Correlation between the observed DSA- and EAC-response 104

5. EXTENDED SUMMARY AND CONCLUSIONS 106

6. FUTURE WORK 110

7. ACKNOWLEDGEMENT I l l

8. REFERENCES 112

0. Abbreviations and symbols

0.1 Abbreviations

AC

ASME BPV

ASTM

ASTM E 399

BNCT

BWR

BWR VIP

C(T)

CERT

CF

CGR

CODLL

DCPD

DO

DSA

EAC

ECP

EPFM

EPRI

FC

FRAD

FW

GE

HAEAC

HAZ

HSK

HWC

ICG-EAC

IF

IG

Air-Cooled

ASME Boiler and Pressure Vessel Code

American Society of Testing and Materials

Test Method for Plane-Strain Fracture Toughness of Metallic Materials

Blunt-Notched C(T) Specimen

Boiling Water Reactor

Boiling Water Reactor Vessel and Internals Project

Compact Tension Specimen

Constant Extension Rate Test

Corrosion Fatigue

Crack Growth Rate

Crack Opening Displacement at Load Line

Direct Current Potential Drop Method

Dissolved Oxygen

Dynamic Strain Ageing

Environmentally-Assisted Cracking

Electrochemical Corrosion Potential

Elastic-Plastic Fracture Mechanics

Electric Power Research Institute

Furnace-Cooled

Film Rupture/Anodic Dissolution Mechanism

Feedwater

General Electric

Hydrogen-Assisted EAC Mechanism

Heat-Affected Zone

Hauptabteilung für die Sicherheit der Kernanlagen

Hydrogen Water-Chemistry

International Co-operative Group of Environmentally-Assisted Crackingof LWR Materials

Internal Friction

Intergranular

Ki-Threshold for SCC

KTA Kerntechnischer Ausschuss, Germany

LAS Low-Alloy Steel

LCF Low-Cycle Fatigue

LEFM Linear-Elastic Fracture Mechanics

LFCF Low-Frequency Corrosion Fatigue (Test)

LWR Light Water Reactor

LWV Labor fur Werkstoffverhalten

MPA Staatliche Materialpriifungsanstalt, University of Stuttgart, Germany

N Normalised

NDT Non-Destructive Testing

NRC National Regulatory Commission, USA

NWC Normal Water-Chemistry

PEER Panel of Experts for Evaluation and Review

PPU Periodical Partial Unloading Tests

PWHT Post-Weld Heat Treatment

PWR Pressurized Water Reactor

RPV Reactor Pressure Vessel

Q Quenched

Q + T Quenched and tempered

S Snoek- Peak in IF Spectra

SCC Stress Corrosion Cracking

SEM Scanning Electron Microscopy

SHE Standard-Hydrogen-Electrode

SICC Strain-Induced Corrosion Cracking

S-K Snoek-Koester-Peak in IF Spectra

SRL Slow Rising Load Test

SS Stainless Steel

SSRT Slow Strain-Rate Test

SSY Small Scale Yielding

TG Transgranular

UTS Ultimate Tensile Strength

VGB Technische Vereinigung der Grosskraftwerksbetreiber, Deutschland

vLFCF Very-Low-Frequency Corrosion Fatigue (Test)

WQ Water-Quenched

YS Yield Strength

0.2 Symbols and units

Symbol Unit Designations

Elongation at fracture

Crack advance

Crack length

Crack advance per fatigue cycle

Time-based crack growth rate: time-derivate of a(t)

Time-based fatigue crack growth rate in inert environment

Time-based CF crack growth rate in high-temperature water

Concentration of dissolved oxygen

Corrosion-assisted crack advance by SCC

Corrosion-assisted crack advance by SICC

Corrosion-assisted crack advance by EAC

Corrosion-assisted crack advance by LFCF

Total crack advance during a test / test period under constant load

SICC crack growth rate in SRL-tests

Crack opening displacement rate at load line

Crack opening displacement

Strain

Crack-tip strain-rate

Stress intensity factor rate

AK = Ki"1^ - Ki™": Total stress intensity factor range

AK-threshold for fatigue

Test period under constant load

Rise time

Potential drop (DCPD)

Resistance change (DCPD)

A5

Aa

a

Aa/AN

da/dt

da/dw

da/dtEAC

DO

AaScc

AaSicc

AaEAc

AaLFCF

AaEAC/AtcL

AaDRK/AtiE

dCODL i7dt

5

e

dEcx/dt

dKi/dt

AK

A K t h

AtcL

AtR

AU

AR

%

|im or mm

mm

|im/cycle

m/s

m/s

m/s

ppb or ppm

(xmormm

|im or mm

(Xm or mm

|im or mm

m/s

m/s

mm/s

[mm]

-or%

s"1

MPa-m1/2/h

MPam1/2

MPa-m1/2

h

h

(iVorV

Symbol Unit Designations

E

ECP

K

K,

K l A S ™

KU

KD

KISCc

V

P

Q"1

R

RP

Rm

a

T

Z

GPa

mVswE

|xS/cm

MPa-m1/2

MPam1/2

MPa-m1/2

MPam1/2

MPa m1/2

Hz

kN

[-]

-

MPa

MPa

MPa

°C

%

E-Modul

Electrochemical corrosion potential

Specific electric conductivity

Stress intensity factor

ASTM E 399 limit for Kj

Ki-value at crack initiation by SICC in SRL-tests

Ki-value at the onset of ductile crack growth in inert environment

Ki-threshold for SCC

Frequency

Load

Internal friction

Load-ratio: R = Kmjn / K ^

Yield stress

Ultimate tensile strength

Stress

Temperature

Reduction of area

1. Introduction: EAC in low-alloy RPV and piping steels

The reactor pressure vessel (RPV) of both boiling water (BWR) and pressurized waterreactors (PWR) is the most critical pressure-boundary component as far as safety and plantlife are concerned [1, 2]. The possible effect of environmentally-assisted cracking (EAC) onRPV structural integrity therefore continues to be a key concern within the context of bothreactor safety and evaluation/extension of plant life. During service, cracks might initiate andgrow in low-alloy primary pressure-boundary components by EAC under the synergisticeffect of the reactor coolant and thermo-mechanical operational loads. To ensure the safeoperation of these structural components, it is essential to identify the conditions, which maylead to EAC. Therefore, reliable quantitative experimental data on EAC initiation and growthunder different LWR operation conditions and a basic knowledge on the underlyingmechanisms are essential to evaluate the possible effects of EAC on the RPV structuralintegrity and to define possible mitigation actions.

Because of the limited resolution of the in-service inspection methods, reactor safety analysisand structural integrity assessments are primarily focused to the growth of postulatedincipient cracks. Reliable quantitative crack propagation data are needed for the assessmentof the flaw-tolerance/safety margins and for the verification/adaptation of the inspectionintervals of the in-service inspection.

1.1 Types of EAC

Currently, no internationally accepted consensus definition for the different basic types ofEAC exists. Stress corrosion cracking (SCC) is therefore sometimes understood either in anarrow sense (EAC under purely, static mechanical loading) or as a part of the broaderspectrum of EAC (including the transition to strain-induced corrosion cracking (SICC) andlow-cycle/low-frequency corrosion fatigue (LFCF)). On the basis of the applied, external,mechanical load, the different types of EAC can be assigned approximately to different LWRoperation conditions (Table 1): SICC and low-frequency corrosion fatigue are characteristicfor transient operating conditions as plant start-up/shut-down. SCC is characteristic fortransient-free, steady-state, power operation, when static loading of the RPV prevails.

Mechanism

Type of loading

LWR operationcondition

Quantitativecharacterisation

SCCStress Corrosion

Cracking

static

steady-statepower operation

BWR VIP 60disposition lines

EAC

SICCStrain-Induced

Corrosion Cracking

slow monotonicallyrising or

very low-cycle

start-up/shut-downthermal stratification

-

CF

Corrosion Fatigue

cycliclow-cycle high-cycle

thermal fatiguethermal stratification

ASME m and XI

Table 1: Basic types of EAC.

10

1.2 Operating experience

The worldwide accumulated operating experience and performance of low-alloy primarypressure-boundary components is very good [3]. Instances of EAC have particularly occurredin BWR service, most often in LAS piping, and, very rarely in the RPV itself. In many cases,EAC has occurred as a result of substantial departures from either design intentions or fromnormal operation practice [3], which are identifiable and can therefore be avoided. Oxidisingagents, usually dissolved oxygen, and relevant dynamic straining (e.g. arising form thermalstratification, thermal and pressurisation cycles during start-up/shut-down, etc.) were alwaysinvolved [3-8]. These cases could be attributed either to SICC or low-frequency CF. Severalincidents were related to unanticipated sources, frequency, and/or severity of thermal stresscycles in critical locations, thus indicating design inadequacies [7]. Cases with major orrelevant contribution of SCC to the total crack advance in properly manufactured and heat-treated LAS primary pressure-boundary components are not known to the authors.

1.3 Experimental background knowledge

EAC initiation and growth are governed by a complex interaction of environmental, loadingand material parameters. It is known from both experimental and field experience that EACmay occur in low-alloy pressure-boundary component steels in oxygenated, high-temperaturewater if the following conditions are simultaneously attained [8]:

- Corrosion potential: ECP > ECPcrit = -200 mVSHE (if K < 0.3 |iS/cm)

- Strain-rate: 0 <ecrit-min < e < eait>mK= 101 %/s

- Strain: e > e ^ = 0.1 - 0.3 % > elastic limit (ocrit > RP)

In particular, slow, positive, dynamic tensile straining with associated plastic yielding appearsto be essential for EAC initiation and growth. The extent of EAC cracking is cruciallydependent both on maintaining a positive crack-tip strain-rate and a high sulphur-anionactivity in the crack-tip environment [9]. Hence, high EAC crack growth rates (CGR) areobserved in steels with a high MnS-inclusion content in highly oxidizing environments withhigh impurity level in the water and, in particular, under dynamic loading conditions. Underhigh-flow conditions, which are characteristic for most RPV locations, crack initiation maybe retarded, or completely suppressed, e.g. in corrosion pits [10]. So far, no majorfundamental discrepancy was observed between the operating experience and the laboratorybackground knowledge [8].

1.3.1 Corrosion fatigue I strain-induced corrosion cracking

The corrosion fatigue crack growth behaviour of LAS in high-temperature water is wellestablished for temperatures in the range of 270 - 320 °C and loading frequencies > 10"3 Hzfor PWR and to a lesser extent for BWR environments. The intensive experimentalinvestigations on EAC have resulted in a revision of the ASME XI curve in 1980. In themeantime, new proposals how environmental effects could be better implemented in theASME m [11-13] and in ASME XI [14, 15] code have been established. Implementation ofthese new proposals to the codes and code upgrading is still under discussion. The proposedprocedures seem to be too complex for practical applications.

11

Under certain critical combinations of environmental, loading and material parameters,experimental EAC data may significantly exceed/fall below the ASME XI wet/ASME Eldesign curve. Nevertheless, under most circumstances the two codes seem to be adequate andconservative, even if environmental effects are considered. Despite the good record of currentfatigue design rules in preventing cracks in carbon and LAS components, a very fewincidents do still occur. Thus identifying specific circumstances that challenge theconservatism in present rules is necessary and very desirable. On the other hand, there is alsoan obvious desire to eliminate overconservatism, especially in the context of plant-lifeextension assessments.

Both, operation experience and laboratory background knowledge indicate that the fatiguedesign codes/curves and usage methodologies might be non-conservative for certain criticalshort-lived BWR plant transients (start-up/shut-down, hot stand-by, thermal stratification,...)and that SICC, or very low-frequency CF behaviour, covers the most important gap in thefield of EAC of LAS [8]. These plant transients have been identified as periods when SICC ismost likely to occur, so intermediate temperatures and slow strain-rates in connection withrelatively high strains are of greater interest. The temperature/strain-rate combinationsrelevant for SICC also draw the attention to possible effects of dynamic strain ageing (DSA)[8 ,16-19] .

The SICC susceptibility of LAS in high-temperature water has been investigated over a widerange of loading (strain, strain-rate), material (sulphur content) and environmental parameters(DO, conductivity), mainly by SSRT and strain controlled low-cycle fatigue (LCF) tests withsmooth specimens [20-25]. Apart from flow rate effects, most influencing factors andconditions, which could lead to crack initiation are now well established. From a safetyperspective, there is still a relevant lack of quantitative SICC crack growth data at slowstrain-rates/very low cyclic frequencies (< 10~3 Hz), intermediate temperatures (100 - 290 °C)and high ECP (> + 100 mVSHE), where susceptibility to SICC is very high and the ASME XIwet curve may be significantly exceeded [8]. Because of possible DSA effects, both, weldfiller and weld heat-affected zone (HAZ) materials have to be included in such investigations[8, 16-19].

1.3.2 Stress corrosion cracking

Notwithstanding the absence of SCC in the field, SCC has been observed in laboratory testsunder simulated BWR conditions with crack growth rates ranging from 30 |im/year to3 m/year even under nominally similar testing conditions [26- 35]. The SCC behaviour oflow-alloy RPV steels in oxygenated high-temperature water and its possible relevance toBWR power operation has therefore been a subject of controversial discussions. Thisproblem has now been largely resolved thanks to some very careful laboratory work carriedout at MPA [32, 33] and PSI [28] using fracture-mechanics specimens.

Recently performed, well-qualified experiments at MPA Stuttgart [32, 33], at PSI [28], in theframe of a European Round Robin test [34] and unpublished testing by ABB and GE inautoclaves attached to the primary coolant systems of BWR revealed a very lowsusceptibility to SCC crack growth under static loading conditions in oxygenated high-temperature water/BWR environments at temperatures around 288°C. Fast and sustainedSCC growth was only observed at high anionic impurity levels (Cl\ SO4

2") in the bulk waterexceeding the EPRI action level 3 or in specimens loaded near to Ku or with gross ligamentyielding [28]. Based on the results of these qualified tests in simulated BWR environment[28, 32, 34 - 36] interim disposition lines for SCC crack growth in low-alloy pressure-

12

boundary component steels for BWR power operation have been proposed by an internationalgroup of experts within the BWR VIP 60 project [35, 36]. The BWR VIP 60 SCC dispositionlines have been recently accepted by the NRC as an interim position [35].

The overwhelming part of SCC tests has been performed with RPV base metal under purestatic loading in the temperature range of 270 - 300 °C [8]. Furthermore, the ECP in most ofthese tests were significantly or slightly lower than the ECP of +150 to +200 mVsHE,estimated to be present on the stainless steel cladding of a BWR RPV [37]. In the feedwatersystem (T = 200 - 220 °C) and in the region of the RPV feedwater nozzle lower temperaturescan be present and small load fluctuations can be observed at least temporary. Theconservative character and adequacy of the BWR VIP 60 SCC disposition lines shouldtherefore be further verified at lower, intermediate temperatures, for small load fluctuations athigh load ratios (ripple loading) and for weld filler and weld HAZ materials. Furthermore, theextent to which these disposition lines also cover water-chemistry or mechanical loadingtransients (which are not covered by the current fatigue evaluation procedures) should beanalysed thoroughly.

2. RIKORR research programme

2.1 Objectives of the RIKORR research programme

In contrast to the low SCC susceptibility under transient-free stationary power operationconditions, preliminary PSI tests with slow rising or low-frequency cyclic loading conditionsrevealed a distinct SICC and LFCF susceptibility of LAS in high-purity, oxygenated high-temperature water under simulated plant operational transients [26, 28].

A new EAC project (RIKORR, 2000 - 2002), co-sponsored by the HSK/BFE, was thereforestarted at PSI to evaluate the most important open questions concerning safety/structuralintegrity in the field of EAC of LAS. The main objective is to quantitatively characterise theSICC and LFCF crack growth behaviour of different low-alloy RPV steels, RPV welds andweld HAZ under transient BWR operating conditions. The special emphasis is placed totemperature/strain-rate effects on crack growth. The second objective is to verify the BWRVIP 60 SCC disposition lines for RPV welds and HAZ, at lower, intermediate temperatures(120 - 290 °C) and under small load fluctuations at high load-ratios (ripple loading).

System conditions, which challenge the conservatism of the current ASME XI referencefatigue crack growth curves and of the BWR VIP 60 SCC disposition lines shall be identifiedby means of an experimental parameter study. The generated data will provide the basis toassess the flaw tolerance/remaining safety margins of the RPV for the case of these criticalsystem conditions/transients. The programme may help to identify critical componentlocations and operation conditions and therefore to improve and optimise in-serviceinspection programmes and operating procedures with respect to EAC.

13

2.2 Working programme

The project includes the following three major experimental tasks [38]:

1. Verification of the current ASME XI reference fatigue crack growth curves for transientBWR/NWC conditions (Section 4.1 - 4.3):

- Effect of loading rate/strain-rate/frequency and temperature on SICC/LFCF crackgrowth.

- Effect of loading parameters (rise time, hold time, R-value, AK, ...).

- SICC/LFCF crack growth susceptibility of RPV weld and weld HAZ. A real RPV girthweld (Biblis C) will be included in these investigations.

The special emphasis is placed to BWR start-up/shut-down conditions and other operationconditions, where thermal stratification typically occurs in horizontal piping and the adjacentRPV nozzles. The testing conditions shall include high ECP, increased conductivity andimpurity levels, intermediate temperatures (150 - 288 °C) and slow strain-rates/lowfrequencies.

2.Verification of the BWR VIP 60 SCC disposition lines for LAS (Section 4. 4):

- SCC crack growth susceptibility of RPV weld filler and, in particular, of weld HAZmaterials.

- Effect of temperature on SCC crack growth.

- Ripple loading (R > 0.95) and periodical partial unloading: Effect of small pressure andtemperature fluctuations on SCC crack growth under steady-state operation conditions.

3. Characterization of the PSA response of the different RPV materials (Section 4.5):

- Tensile tests at different temperatures/strain-rates.

- Internal friction measurements with selected materials.

- Measurement of total and free Nitrogen content.

- Characterization of the lower shell RPV base material and girth weld material betweenthe upper and lower RPV shell of Biblis C (see [39]).

14

3. Materials and experimental procedure

The SICC/LFCF and SCC crack growth behaviour of different RPV materials wascharacterized by slow rising load (SRL) / low-frequency corrosion fatigue (LFCF) andconstant load (CL) /periodical partial unloading (PPU) / ripple loading (RL) tests with pre-cracked fracture mechanics specimens. The tests were performed in modern high-temperaturewater loops (Figure 1) under simulated BWR/NWC conditions, i.e. in oxygenated high-temperature water at temperatures of either 288, 250, 200 or 150 °C. The testing facilities andprocedures are described in detail in [26, 28, 40]. In the following, only the most importantaspects are summarized.

Figure 1: Schematic drawing of the hot-water loop.

3.1 Materials

Four different types of low-alloy RPV steels with either a low, medium or high sulphurcontent between 0.004 and 0.018 wt.% and a RPV weld material from a real RPV (Bibiis C,PWR, 1200 MWe, 1976) were investigated. The heat treatments, chemical composition andmechanical tensile properties are summarized in Table 2 and 3. The RPV base materials had agranular, bainitic (alloy A, B, D) or a mixed bainitic/ferritic-pearlitic structure (alloy C) withan average former austenitic grain size of 10 to 20 Jim. The spatial distribution of MnSinclusions was fairly homogeneous in all of the steels and no distinct sulphur segregationzones were observed by sulphur prints. The weld material had a very fine-grained ferriticmicrostructure with a grain size < 6 fxm. The morphology of the MnS-inclusions in the weld

15

material E was completely different from that in the RPV base material: The weld materialonly contained fine-dispersed, very small, spheric MnS-inclusions. The RPV base materialsrevealed a wide range of morphologies and size distribution of MnS-inclusions, includingsome large, elongated ones (up to some few hundreds of ^m), but no regions with distinctclusters were observed in all materials.

Steel20 MnMoNi 5 5

SA 533 B Cl. 1SA 508 Cl. 2

22NiMoCr3 7S3 NiMo 1

ABCDE

C

0.21

0.25

0.21

0.22

0.05

Si

0.25

0.24

0.27

0.20

0.17

Mn

1.26

1.42

0.69

0.91

1.19

P

0.004

0.006

0.005

0.008

0.013

s0.004

0.018

0.004

0.007

0.007

Cr

0.15

0.12

0.38

0.42

0.04

Mo

0.50.54

0.63

0.53

0.55

Ni

0.77

0.62

0.78

0.88

0.94

V

0.008

0.007

0.006

0.007

0.006

Al0.013

0.03

0.015

0.018

0.006

Cu

0.06

0.15

0.16

0.04

0.06

N,o,[ppm]

706011080110

Nfree[ppm]

30<12316

Table 2: Chemical composition (in wt.%) of low alloy RPV steels and weld materialsinvestigated.

Steel

20 MnMoNi 5 5

SA533BCI. 1

SA508CI. 2

22NiMoCr3 7

S3 NiMo 1

ABC

D

E

Rex

[MPa]

485468

448

467

492

jm temp

Rm

[MPa]648616

611

605

592

eratun

As[%]19.3

21.0

17.9

17.3

17.4

Z

[%]

72.1

43.0

71.0

71.9

73.3

288 °C

RPr0.2

[MPa]418411

396

400

430

Heat treatment

910 oC-920°C/6h/WQ/640 oC-650 oC/9.5h/FC

915 o C/12h/860°C/12h/WQ/635°C/12h/FC

as received

870-905°C/7h/WQ/635-655°C/11.3h/ACPWHT: 540-555 °C / 59 h / 590 - 610 °C/ 21 h / 590 -605 °C /11 h

PWHT: 540-555 °C / 59 h / 590 - 610 °C/ 21 h / 590 -605 °C /11 h

Table 3: Mechanical tensile properties (DIN 50145, B5X50 specimens, mean value of T andL gauge length) and heat treatments (WQ: Water-quenched FC: Furnace-cooled,AC: Air-cooled, PWHT: Post-weld heat treatment) of low alloy RPV steels andweld materials investigated.

3.2 Specimen form / pre-cracking25 mm thick compact tension specimens (1T-C(T)) were manufactured from forged ingots orhot-rolled, nuclear grade steel plates, mainly in the T-L or L-T orientation. All specimenswere pre-cracked by fatigue in air at room temperature using a load ratio R of 0.1 and a final

3.3 Mechanical loadingA screw-driven, electromechanical tensile machine with computer control actuated the loadduring the experiment. With this system, the stroke of the pull rod could be controlled veryaccurately by an electric step motor down to very low displacement rates of 10"9 m/s. Theload was measured by an external load cell and corrected by the actual measured value ofautoclave pressure and an experimental calibration. The friction between the pull rod and theautoclave sealing was less than 100 N. The pressure fluctuations from the high-pressurepump were minimized by an accumulator and pulsation dampener to < 0.1 MPacorresponding to load fluctuations of < 50 N. Besides measuring the pull rod stroke by anexternal LVDT, the crack opening displacement (COD) was measured by internal clip gaugesmounted on the specimens (electrically insulated from specimens by ceramic spacers). Ki-values were calculated according to ASTM E 399 [41] by the measured load and the actual

16

mean crack length derived by post-test fractography and by the reversed, direct-current,potential-drop (DCPD) method. [26, 40]

The different phases of the experiments are shown in Figure 2 and 3. After heating indeoxygenated high-purity water (1), the specimens were pre-oxidised in the test environmentfor > 168 h under a small mechanical pre-load of 9 kN, corresponding to a Ki between 12 and18 MPam1/2 (2). The loading procedure following the pre-conditioning period varied for thedifferent test types as follows:

Slow rising load / low-frequency corrosion fatigue tests (SICC/LFCF):

After the pre-conditioning period, the loading of the specimens in SRL tests and the initialloading in LFCF tests were performed by using a constant rate of the pull-rod stroke. Undermost test conditions, a constant pull-rod stroke rate resulted in constant dP/dt, dCOD/dt anddKi/dt values. The subsequent cyclic loading in LFCF tests was performed under loadcontrol. Constant load amplitude loading with a positive saw tooth waveform (slow loading,fast unloading) was applied.

Load, K,

Pre-load

Temperature

' EnvironmentalO2, K. ...

parameters

! AtRj

Asymmetrical sawtooth loading

Timet

*Timel

t = 0 At Timet

Figure 2: Simplified schematic of test procedure for low-frequency CF tests.

In SRL-tests, the rise time AtR from the initial (12-18 MPa-m1/2) to the final Ki-values (60 -76 MPa-m1/2) was varied between 292 h and 0.2 h corresponding to dCODuydt values of210"7 to 6-10"4 mm/s and to dK/dt values ranging from 0.2 to 320 MPa-m1/2/h.

In LFCF-tests, initial loading was generally performed in the range of maximum SICCsusceptibility (dCODu/dt values of 1-10"6 to 3-10"6 mm/s). The loading parameters werevaried in the following range (the values in parenthesis give the typical values): R-value: 0.2- 0.98 (0.8), v = 2.8-10"6 - 1.410'2 Hz, AK: 2.3 - 64.4 MPa-m1/2 (12 - 13 MPa-m1/2), Ki"^ :50 - 80 MPa-m1/2 (60 - 70 MPa-m1/2). In most cases, the K^-values were below the ASTME 647 [51] limit. In the overwhelming part of the different test phases, stable, stationary EACgrowth and crack growth increment AaEAC ^ 100 - 200 \xm (> 10 X DCPD resolution limitor 10 X mean grain size) were observed.

17

Constant load, periodical partial unloading and ripple loading tests (SCO:

For our standard constant load tests (Figure 3), the increase in loading (3) prior to theconstant load testing phase (4) was carried out rapidly within 0.3 to 2 h at a constant loadingrate (constant rate of the pull-rod stroke). For different tests, the loading rate ranged from 10to 100 kN/h, corresponding to a load-line crack-opening-displacement rate dCODLi7dt of3-10"5 to 3-10"4 mm/s and a dKj/dt value of 10 MPa-m1/2/h to 200 MPam1/2/h (most often «30 MPa-m1/2/h). After reaching the intended Kj value, the load was held constant for ca.1000 h.

Environ mental parameters

Time t

Figure 3: Schematic of test procedure of constant load tests.

The modified constant load tests started either with a slow rising load or cyclic loading phasebefore switching to constant load. The constant load test phase therefore always started withan actively growing EAC crack. Both EAC CGR and crack advance were generally above thecritical values (da/dtEAC > 10"10 m/s, AaEAC > 0.2 - 0.3 mm). The rate of loading during SRL-phase (3) was systematically varied between 0.12 and 161 kN/h, corresponding to dCODuydtvalues of 2-10'7 to 6-10"4 mm/s and to dKr/dt values ranging from 0.2 to 330 MPa-m1/2/hresulting in EAC CGR daEAc/dt between 10"10 to 8 -10"7 m/s. The cyclic loading prior to theconstant load period was performed under load control. Constant load amplitude loading witha positive saw tooth waveform (slow loading, fast unloading) with a rise time AtR of 1000 sand a load ratio of 0.8 was applied. After reaching the intended maximum load (Ki), theduration of the subsequent constant load test phase (4) was held constant for ca. 1000 h. Inthe case of stationary or fast SCC growth the constant load period was reduced to 300 -400 h.

Preliminary tests with periodical partial unloading were performed under load control withconstant load amplitude loading. An asymmetrical trapezoid waveform with a rise time AtR of

18

1000 s and a load ratio of 0.8 was applied. The hold time at maximum constant load wasvaried between 0.1 and 10 h.

Preliminary ripple loading tests at very high load ratios R of 0.96 to 0.98 with a AK-amplitudes between 2 and 3 MPa-m1/2 near to the fatigue-threshold AKth were performedunder load control with asymmetrical saw tooth loading. The rise time AtR (loadingfrequency) was systematically varied between 0.01 and 10 h (2 10~2 and 2-10"5 Hz).

3.4 Water-chemistry

The experiments were performed in refreshed loops with 10 1 stainless steel autoclavefacilities (Figure 1) under "low-flow" conditions (ca. 4 autoclave exchanges per hour). Aflow rate in the range of mm/s is to be expected in the vicinity of the specimens. In the low-temperature/low-pressure part of the facility, the main water-chemistry parameters(conductivity and concentration of dissolved oxygen) were measured in both the inlet and theoutlet water. The values measured at the inlet were used to control these parameters. The inletconductivity was controlled by dosing 0.02 m Na2SC>4 solution into the high-purity(< 0.06 |iS/cm) water. The outlet water is completely cleaned by ion exchangers. The level ofadded SO42" was 1 ppb (0.06 |0,S/cm) to 360 ppb (1.0 |i.S/cm) (Figure 4). Ionic impurities wereanalysed by Inductive Coupled Plasma - Atomic Emission Spectroscopy (ICP-AES) and IonChromatography (IC) of periodically taken grab samples (inlet and outlet water) (Table 4).

InletOutlet

SO42"

<1-2<1-2

lonchci-

<1-2<1-2

romatocN03-<1-2<1-2

iraphP<1<1

y [ppbPO43-<1<1

1SiOz*-<2-10<2-10

Shot

2-33-6

Kto,<1<1

Natot

<1<1

ICPCUtot

<0.5<0.5

-AES [|Few

<0.5-1<0.5-1

>pb]Cfa<11-4

Nhot<0.10.5

COtot

<0.8<0.8

Zntot

0.30.5

Table 4: Typical concentration range of impurities in tests withdifferent DO.

< 0.06 |iS/cm and

3.5 Electrochemical potential

The ECP of the specimens and the redox potential (platinum probe) were continuouslymonitored by use of an external Ag/AgCl/O.Olm KCl-reference electrode connected to anappropriate electronic device with a high input impedance of > 1014 Q [28, 40, 42].Calculation of the ECP versus the Standard Hydrogen Electrode (SHE) and correction of thethermal-liquid-junction potential was performed according to Macdonald [43]. Thespecimens were electrically insulated from the autoclave, from each other and from the clipgauges by Z1O2 spacers. The ECP was set by the concentration of dissolved oxygen (DO) inthe hot water. Depending on DO (< 5 to 8000 ppb) and temperature T (150 - 288 °C) an ECPof -400 up to ca. +250 mVsHE was measured (Figure 5 and 6).

19

400

0.2 0.6 0.8 1.0

Inlet conductivity K [uS/cm]

Figure 4: Relationship between conductivity and sulphate content at autoclave inlet.

200

•"1

100 1000

DO at inlet [ppb]10000

Figure 5: Effect of DO on ECP at 288 °C.

LLJ

w 500

— 400

Xo•o0

DC•D

CCQ .OLU

200

100 DO1*' = 8 ppmK = 0,25 nS/cm, 65 ppb SO4

125 150 175 200 225 250 275 300 325

Temperature [°C]

Figure 6: Effect of temperature on ECP at 8 ppm DO.

20

3.6 On-line crack growth monitoringCrack advance was continuously monitored using the DCPD method. The crack growthincrement was calculated by the Johnson formula [44]. The mean, pre-fatigue crack length<ao> was assigned to the potential drop at the point of crack-growth initiation during initialloading in the test, as determined according to ASTM E 1737 [45] (Figure 7 and 8). Thecalculated crack length at the end of the experiment was then verified and, if necessary,corrected with regard to the mean final crack length <ao + AaEAc> as revealed by post-testfractography [40]. In the case of fairly even SICC crack advance, the difference betweencalculated and fractographically determined increments of crack advance was < 5%. Theevaluated DCPD resolution limit corresponded to ca. 5 - 2 0 |im (depending on testconditions) [26, 40].

separation of 2 * * ^ c r a c k

crack fianks b l u n t i n9 initiation irowtb

AUAR

CMOD P

Figure 7: Schematic of DCPD signal during SRL-tests. Transition from region II to HI =point of crack initiation.

In the case of cyclic loading conditions, cycle-dependent CGR Aa/AN and time-dependent"true" CGR da/dt (time-derivate of a) and "apparent" CGR (Aa/AN) / AtR were calculated(AtR: rise time, see Figure 2). Under slow rising load conditions "true" SICC CGR dasicc/dt(or Aasicc/AtiE) were given. Under static loading conditions "true" CGR dascc/dt (time-derivate of a) and "apparent" CGR (AaEActot/AtcL, Aascc/AtcL) w e r e determined.

21

ECP = + 160 mVouc> 65 ppb SOSICC: <Aaoirr> = 0.73 mm

Crack initiation

ECP = -380 rrA/ , 65 ppb SOno SICC: <Aaoir,^> < 20 urn

SICC

10 20 30 40 50 60 70 80K, [MParn172]

Figure 8: DCPD signal during a SRL-test. a: SICC initiation and growth (8 ppm DO), b: noSICC (< 5 ppb DO).

3.7 Post-test evaluationThe fracture surfaces of specimens broken apart at liquid nitrogen temperature wereinvestigated with a scanning electron microscope (SEM). For fractographic analysis in theSEM, the oxide film on the fracture surface of one specimen half was first removed bygalvanostatic reduction in an ENDOX-bath [46, 47].

22

4. Results and discussion

4.1 Strain-induced corrosion cracking / slow rising load tests

4.1.1 Effect of loading rate

Figure 9 and 10 show the effect of loading rate on SICC initiation and subsequent growth inmaterials A, B and C at a temperature of 288 °C. The tests were conducted with a DO of8 ppm (ECP = + 150 mVsHE) and an inlet conductivity of 0.25 |xS/cm (65 ppb SO4

2"). Thecrack growth initiation stress intensity K^ and SICC CGR <Aasicc/At> are plotted versus thedisplacement rate at the load line dCODLi/dt. Additionally, the corresponding increase in thestress intensity factor dKi/dt and the estimated crack-tip strain-rate d£cr/dt [70] are shown inFigure 9 and 10. Alloy B (0.018 wt.% S) revealed SICC over the whole range of testedloading rates. In this material, Ky decreased with decreasing loading rate, with a possible

1/2 ,-6.minimum of 34 MPa-m at a relatively low displacement rate of = 10" mm/s, correspondingto a crack-tip strain-rate of 5-10"5 to 10"4 s"1. At lower loading rates, Ky seemed to increaseagain.

1/2,

10,-1

dK/dt [MPanV'Th]

10° 101 102 103

60

Estimated decl/dt [s'1]

10'5 10"4 10"3 10'2 10"1

30

55 -

50 -

45 -

40

35

-

-

O SA 533 B0.018 wt.

• SA 508 C

i

Cl. 1%s.2

(0.004 wt.% S)A 20 MnMoNi 5 5

(0.004 wt.

. •

: A

x O O

\ o

%S)

o

o/

AK = 0

so;

•••1 i

i

o // •

o /

/ :

/ ;

= 8 ppm [

+ 150mVSHE:

.25 nS/cm' = 65 ppb !

\-710" 10 10 10 10 10

dCODM/dt [mm/s]

'2

LL/

Figure 9: Effect of loading rate on SICC initiation.

23

The subsequent SICC CGR was constant during the whole rising load phase after a shorttransition period following SICC initiation (Figure 8a). The SICC CGR depended strongly onthe applied loading rate, but not on the actual value of Kj, at least up to a Ki of 60 to70MPam1/2. Here, a simple power law relationship with an exponent of 0.8 was foundbetween SICC CGR and dCODLL/dt and dK/dt (Figure 10). Thus, both SICC initiation andgrowth were strongly dependent on applied loading rate, but with opposite tendency. Theobserved high SICC CGR of 10"9 m/s to 810"7 m/s were thus only observed if initiationoccured and this requires high stress intensities, which would correspond to relatively deepcracks in the RPV.

dK/dt [MPa-m1/2/h]

10"1 10° 101 102

f ,«•LU

ooCO

10-9

10-10

SA 533 B Cl. 1 (0.018 wt.% S)SA 508 Cl. 2 (0.004 wt. % S)20 MnMoNi 5 5 (0.004 wt.% S)

• Linear Fit of O• 95 % Prediction Interval

<Aas|CC>/AtE a (dK/dt)0.8

1 0 ' 7 10"6 10"5 10"4

dCOD/dt [mm/s]

Figure 10: Effect of loading rate on SICC growth.

In alloys A and C with a low sulphur content of 0.004 wt.%, SICC could only be detected byDCPD at the lowest loading rate applied (2 to 310"7 mm/s). The measured SICC CGR werewithin the scatter range of the results of alloy B (Figure 10), whereas the Ky values wereslightly higher (Figure 9). At higher loading rates, the extent of SICC revealed byfractography was restricted to some few, very localized thumbnail-shaped areas along theformer pre-crack front. The mean crack increment was below the DCPD resolution limit ofca. 20 fim. The SICC growth at very slow loading rates in alloys A and C might be associatedwith DSA [16, 17, 48], since both alloys have a very low aluminium content. Preliminarytensile tests at different temperatures/strain-rates revealed DSA in alloy A and C [48 - 50](see also Section 4.5).

24

4.1.2 Effect of water-chemistry

The effects of ECP and bulk SO42" concentration were investigated at two different loadingrates (2 to 3-10"7 and 2 to 4-10"6 mm/s) at 288 °C. Tests at different oxygen levels of < 5, 200and 8000 ppb, corresponding to three different regimes of ECP (- 380, -60 and +150 mVSHE),were performed at a conductivity of 0.25 (xS/cm (65 ppb SO42"). SICC could be detected byDCPD only in tests at the highest DO of 8000 ppb in alloy B with the highest sulphur contentof 0.018 wt.%. In tests at 200 ppb DO, post-test fractography showed that the overall SICCcrack increment had already decreased significantly, even in alloy B. The crack growth hadbecome highly uneven, revealing only some few very localised SICC areas. This finding inthe high-sulphur alloy B tested at an ECP of -60 mVsHE was equivalent to that for the low-sulphur steels A and C at a high ECP of +150 mVSHE- At the lowest ECP of -380 mVSHE,even in the high-sulphur steel B, no SICC was found by post-test fractography up to a high Kivalue of 76 MPa-m1/2 for all loading rates tested. To highlight the strong effect of ECP, thelocal maximum SICC advance in the high-sulphur alloy B revealed by post-test fractographyfrom tests in the loading range of maximum SICC susceptibility (dCODuydt = 2 to 4- 10"6

mm/s) and stopped at comparable final Ki values is plotted versus the ECP in Figure 11.

10

• 1

E

I< 0.1

0.01 ^

2-AlloyB (0.018 wt.% S)K = 0.25 |iS/cm, 65 ppb SO

dCODLL/dt = 2 - 3 • 10'6 mm/s

-400 -300 -200 -100 0 100 200

ECP [mVCH JSHEJ

Figure 11: Effect of ECP on maximum SICC advance.

25

Tests at conductivities of 0.06, 0.25 and 1.0 (iS/cm (<1, 65, 360 ppb SO4 ) were performedin water with a DO of 8000 ppb (ECP = +150 mVSHE). At this high ECP, SICC crackingsusceptibility in the high-sulphur alloy B was not affected by a change in bulk sulphateconcentration from < 1 ppb to 360 ppb SO4

2' and the Ki,j-values and SICC CGR wereessentially the same at all the three conductivities for identical applied load rates. In the low-sulphur steels A and C, SICC crack growth was restricted to some few localized thumbnail-shaped areas and could not be detected by DCPD at displacement rates dCODLi/dt of 2 to4 • 10"6 mm/s up to 65 ppb SO42". However, a further increase of the bulk- SO42" content from65 to 360 ppb resulted in rather uniform SICC growth along the whole former pre-crack frontin alloys A and C (see Figure 12). The measured K^-values were slightly higher than in thehigh-sulphur alloy B, whereas the SICC CGR were comparable for both the high and low-sulphur steels (Figure 14). For the lowest displacement rate (dCOLi/dt = 2 to 3-10"7 mm/s)tested, SICC growth in the low-sulphur steels was fairly uniform down to < 1 ppb SO42" andcomparable for low and high-sulphur steels at all conductivities (Figure 13 and 14).

10 20 30 40 50 60 70 80

CC

0.00

0.20

0.15

0.10

0.05

nnn

'. ECP =

- < A asicc

-

-

-

I

/

/

+ 160 mVSHE, 65 ppb SO42', 0.25 nS/cm '.

> < 20(xm

II

-

-

-

DC<

10 20 30 40 50 60 70 80

K, [MPam1/2]

Figure 12: Effect of SO42" -addition on SICC in the low-sulphur alloy A. An increase in SO42"from 65 (lower diagram) to 365 ppb (upper diagram) was sufficient to induceSICC in the low-sulphur material A at a DO of 8 ppm.

26

0OcCO

</)

cc

12.712.612.512.412.312.212.112.011.911.811.7

ECP = + 160 mVSHE, 8 ppm O.

K = 0.25 nS/cm, 65 ppb SO42'

T = 288 °C

SA 533 B Cl. 1, 0.018 wt.% S, LTa0 = 23.8 mmdCODLL/dt = 2.5E-7 mm/s<Aa.,._> = 0.73 mm

SICC

K^AOMPam"2

Crack initiation

SA 508 Cl. 2, 0.004 wt.% S, LT-

dCODLL/dt = 2.5E-7 mm/s<Aas.cc> = 0.85 mm

K. = 48MPam1S!

5 10 15 20 25 30 35 40 45 50

Load [kN]

Figure 13: Similar SICC behaviour of the low-sulphur alloy C and high-sulphur alloy B atvery slow loading rates with a rise time of ca. 290 h.

In all cases where (uniform) SICC crack growth could be detected by DCPD (e.g., ECP =+150 mVsHE), SICC crack growth seemed to be independent of both the steel sulphur contentand bulk sulphate concentration, i.e., it was mainly governed by the loading rate itself (Figure14). On the basis of these results, the occurrence of SICC requires critical conditions, i.e. ahigh sulphur-anion activity in the crack-tip environment and a slow positive crack-tip strain-rate, to be achieved simultaneously. If these conjoint conditions are not achieved, no, or onlyminor, local SICC growth is observed. However, if they are achieved or exceeded, the SICCCGR seems to be dependent primarily on the applied strain-rate.

These local crack-tip parameters are governed by a set of interrelated corrosion systemparameters such as the ECP, the bulk sulphur-anion concentration, the steel sulphur contentand the loading rate/level. A high SICC susceptibility is favoured by a high ECP and/or ahigh steel sulphur content/bulk sulphur-anion activity and quasi-stagnant flow conditions,which favour the enrichment of sulphur-anions in the crack-tip environment [8, 9]. Possiblesources for sulphur-anions are (i) dissolved sulphur species in the bulk water outside thecrack enclave and (ii) sulphur-anions from the (probably sluggish) dissolution of MnS-inclusions intersected by the growing crack. The higher the ECP, the higher the tendency forsulphur enrichment within the crack enclave. This is due to migration and retention of anionswithin the crack as a result of the potential gradient arising from a low ECP inside the de-aerated crack enclave and a high ECP under oxidising, bulk-water conditions. The synergisticinteraction of strain-rate and sulphur-anion activity can be rationalised by the filmrupture/anodic dissolution model [9]. In this model, crack growth through anodic dissolutionis primarily governed by mechanical rupture of the protective oxide film via dynamicstraining and through the strong effect of sulphur-anion activity on repassivation kinetics. Ahigh sulphur-anion activity significantly delays the formation of a new, protective oxide layerand thus leads to a larger increment of crack advance by anodic dissolution per oxide-rupture

27

event. The observed SICC cracking behaviour supports the basic ideas of the GE-model [9]for these materials. However, our observation that, at very slow loading rates, more SICCcrack-growth initiation sites seem to be activated along the pre-existing crack front, resultingin more uniform crack growth, cannot be explained by this model.

• •

x(0

10"

10-7

o 10"

10-9

95 % prediction intervalLinear fit

Alloy B (0.018 wt.% S)ECP = + 150ITA/

orlb

0.25 |iS/cm, 65 ppb SO42"

Steel sulphur-cone

10-7

•i i i iii

0.06 iiS/cm, 0.018 wt.% S0.06 ^iS/cm, 0.004 wt. % S0.25 |iS/cm, 0.018 wt.% S0.25 jiS/cm, 0.004 wt.% S1.0 iiS/cm, 0.018 wt.% S1.0 jiS/cm, 0.004 wt.% S

10' 10" 10' 10'

dCODM/dt [mm/s]LL

Figure 14: Effect of SO42"-concentration and steel sulphur content on SICC growth.

28

4.1.3 Effect of temperature

Tests at different temperatures of either 288, 250, 200 or 150 °C with the low- and high-sulphur steels A and B were performed in high-temperature water with a DO of 8000 ppb anda S04

2"-content of 65 ppb. The ECP decreased somewhat from +250 at 150 °C to+130 mVsHE at 288 °C (Figure 6). A displacement rate of 4-10"6 mm/s, i.e. in the range of themaximum SICC susceptibility at 288 °C (Figure 9), was used. The effects of temperature onSICC initiation and growth are shown in Figure 15 and 16. At 288 °C, SICC could bedetected by DCPD only in the high-sulphur alloy B. At 250 °C, SICC was observed in bothalloys, with comparable Ky values and SICC CGR. At 200 and 150 °C, only the low-sulphuralloy B revealed SICC by DCPD. As a first trend, a minimum in Ity values was observed atintermediate temperatures around 250 °C, whereas SICC CGR increased slightly withincreasing temperature for the chosen loading rate. As a first estimate, an Arrheniusactivation energy for SICC growth of 32 ±3 kJ/mol was calculated taking the results for bothalloys together. The higher SICC susceptibility of the low-sulphur alloy A at the lowertemperatures of 150 and 200 °C, which was also further confirmed by SCC tests underconstant load [50, 53], is initially surprising and might be regarded as a further indication fora possible DS A effect in this material. Most of the data trends observed here in SRL-tests onfracture mechanics specimens are in good, qualitative agreement with older slow strain-ratetest (SSRT) results using smooth tensile specimens [22 - 25] (Figure 17 - 19).

80

70

- g 60

£ 50

40

30

20

Alloy A (0.004 wt.% S, 0.013 wl.% Al) ;Alloy B (0.018 M.% S, 0.03 wt.% Al) ;Open symbols: Vio SICC detected by qpPD (K.max)

O

100 150 200 250 300

Temperature [°C]350

Figure 15: Effect of temperature on SICC initiation.

29

I 1

110"

(0

oo55 1 0

9

10"10

Alloy A (0.004 wt.% S, 0.013 wt.% Al)Alloy B (0.018 wt.% S, 0.03 wt.% Al)Open symbols: no SICC detedted by DCPDDLinear Fit of • and

DO = 8 ppm, 65 ppb SO 2-

"6dCODLL/dt = 2-4-10"° mm/s

O . . . . O . . P150 200 250

Temperature [°C]

300

Figure 16: Effect of temperature on SICC growth.

I<d

100-

SO-

0-

450de/dt

*

#

*

•• • ' " t ' • " • '

ppb DO= 1E-6 S"1

™" '"™ j~"m" IJ t

0i r•

ISNiCuHoNbS

15Mo3 // ftm / /

ft

»-.~.».»,m^«,.m,.».I/»«....-....m-»».-..J..f....-»». . ^ . . . .

1

# »

itli*

iso 2S0

Temperature

Figure 17: Effect of temperature on SICC susceptibility in SSRT-tests with smooth tensilespecimens [25] (compare to Figure 15).

30

10

450 ppb DOde/dt = 1E-6 s"

\

ISO 250 300 •£temperature •*>

Figure 18: Effect of temperature on SICC growth in SSRT-tests with smooth tensilespecimens [25] (compare to Figure 16).

3 5 -400 ppb DO0.16|LiS/cm

22 NiMoCr 3 7, 0.016 wt.% S

£ 20os 15 :2 Z

E io ~**""* **» lllllt

Q-t1O*8 10"7 10*6 10*5 10*

Strain rate / Us10"

Figure 19: Effect of strain-rate on SICC susceptibility in SSRT tests with smooth specimens[23] (compare to Figure 9).

31

4.1.4 Effect of steel sulphur and aluminium content

So far it was widely accepted that the steel sulphur content is the sole material parameterstrongly affecting the SICC behaviour of LAS. That it can have a major effect on behaviour(at least in the absence of sulphate impurities in the water) has been shown by numerousinvestigations [8]. Figure 20 shows an example, where a low steel-sulphur content resulted ina significantly lower SICC susceptibility in a SRL-test. As shown before, the beneficial effectof a low steel sulphur content at temperatures > 250 °C could be masked by suitablecombinations of corrosion system parameters (e.g. by suitable strain-rate (Figure 13 and21)/temperature (Figure 15) combinations and/or by a high bulk sulphur-anion content + ahigh ECP (Figure 12 and 14)). The results at very slow strain-rates (Figure 13 and 21) andlower temperatures (Figure 15) indicate that DSA might significantly affect EAC crackingbehaviour in susceptible LAS under certain temperature/strain-rate combinations. Thealuminium content and concentration of free, interstitial nitrogen and carbon might thereforebe just as relevant for EAC susceptibility as the steel sulphur content, at least underconditions where DSA is observed (see also Section 4.4 and 4.5).

Pre-load Experimental load

•\5A

ICDO

:an

COCOCD

cc

15.3

15.2

15.1

15.0

14.9

14.8

14.7

14.6

14.5

-

-

-

• i

I

fti i

• • • • i • • • • i • • • . i •

Initial loading in 22.45 hdCODLL/dt = 3 E-6 mm/s

Crack initiation

1

Blunting

n

Crack growth(SICC) Sm - ^

Blunting11

. . . i . .

SA 533 B Cl.0.018 Gew.%

Aasicc> = 1.09

iiiiiii * ~

SA 508 CI. 2

0.004 Gew.%

<AaS]cr> < 20

/

1Smm

Sfim

-

-

10 15 20 25

Load [kN]

30 35

Figure 20: Beneficial effect of a low steel sulphur content in slow rising load tests in2-oxygenated high-temperature water with 8 ppm DO and 65 ppb SO4 ". In contrast

to the high-sulphur alloy B, the low-sulphur alloy C revealed no SICC underthese test conditions.

32

r—i

CDoCOCOCOCD

*o0

cCO

SIo

u

0

0

0

0

n

.u

.4

.3

CM

.1

n

. " " i 11 • • i . . . • • • i • i i i i _

: ECP = + 160 mVSEW, 8 ppm O2

: K = 0.25 nS/cm, 65 ppb SO42"

: T = 288°C

SA 508 Cl. 2, 0.004 Gew.% S, LTa0 = 23.4 mm j :dCODu/dt = 2.5E-7 nun/s / j<AaORK> = 0.85 mm f "

/ '•« :

l Crack initation / j

K. = 48 MPa m1/2 j f f \

-

/ '•

/ •

jtgHJ^iWW^ SA 508 Cl. 2, 0.004 Gew.% S, LT :j J r ^ ao = 26.9mm -

F dCODLL/dt=2.9E-6mtn/s :

1 <AaDRK> < 2 0 ^m :

i Jl i i i i i i ':

10 20 30 40 50 60 70 80

K [MPa-m1/2]

Figure 21: Effect of very slow strain-rates. The low-sulphur alloy C revealed SICC at veryslow strain-rates (rise time ca. 290 h). At higher loading rates no SICC has beenobserved in this material under otherwise identical conditions.

4.1.5 Effect of microstructure

The effect of microstructure has been investigated with the high-sulphur alloy B in theloading range of maximum SICC susceptibility at 288 °C. The tests were conducted inoxygenated high-temperature water with 8 ppm DO and 65 ppb SO,*2'. Differentmicrostructures were generated by a variation of the thermal heat treatment (see Table 5).Besides the bainitic "standard" microstructure (Q+T), which is characteristic for the RPVbase metal, a martensitic (Q) and a ferritic-pearlitic "equilibrium" microstructure (N) wereproduced by austenizing/water quenching and austenizing/ slow furnace cooling. The MnS-morphology was not affected by the applied heat treatments. The different behaviour cantherefore be related to the different microstructure, yield strength level and plasticdeformation behaviour.

As shown in Figure 22/23 and Table 5, the ferritic-pearlitic and the bainitic microstructurerevealed a very similar SICC behaviour at 288 °C. On the other hand, the martensiticmicrostructure resulted in significantly higher transgranular SICC CGR and in lowerinitiation stress intensity factors Ku in SRL-tests under otherwise identical test conditions. Inall microstructure a quasi-cleavage with feather morphology was observed on the fracturesurface, which is characteristic for EAC in LAS in high-temperature water. Based on the veryhigh hardness and yield strength level (Table 5), hydrogen-assisted transgranular (orintergranular) EAC cannot be excluded for the martensitic microstructure [8, 54, 76]. In fact,the martensitic microstructure also revealed stationary, fast SCC crack growth under constantload at 288 °C (see Section 4.4.5). Such a behaviour has not been observed with the other

33

microstructures in high-purity high-temperature water at 288 °C, eventually indicating adifferent cracking mechanism.

Alloy B, SA 533 B Cl. 1, 0.018 wt.% S

T = 288 °C, ,,low flow": 4 autoclave exchanges / h

ECP = + 150 mVsHE, 8 ppm O2, 65 ppb SO42~, 0.25 ( S/cm

Heat treatment

Heat treatment

Microstructure

Vickers hardness

r, 288 °CKP

Grain size

ao

AtR

dCODu/dt

dK,/dt

<asicc>/AtiE

Ki (constant load)

<ascc>/AtcL

SCC growth

Q + T

915°C/12h/860°C/12h/WQ635°C/12h/FQ

bainitic

197 VH10

411 MPa

16 |im

23.3 mm

25.1 h

2.8 E-6 mm/s

1.9MPa-m1/2/h

40 MPa-m1/2

9.3E-9 m/s

61 MPa-m1/2

2.2 E-11 m/s

decaying

N

900 °C / 30 min / FQ

ferritic-pearlitic

260 VH10

577 MPa

15 |im

21.6 mm

22.3 h

2.4 E-6 mm/s

2 MPa-m1/2/h

32.5 MPa-m1/2

7.5 E-9 m/s

59.4 MPa-m1/2

<lE-llm/s

no

Q

900°C/30min/WQ

martensitic

466 VH10

960 MPa

-

20.3 mm

22.3 h

2.3 E-6 mm/s7.5 E-6 mm/s

1.8MPam1/2/h2.6 MPa-m1/2/h

25 MPa-m1/2

l.lE-8m/s1.9E-7m/s

64.7 - 84 MPa-m1/2

1.9E-7m/s

stationary

Table 5: Summary of test conditions and results of SRL-tests with different microstructures.(Q+T: Quenched + tempered, N: Normalized, Q: Quenched).

Our preliminary results support the basic idea that the high-temperature water EAC behaviouraround 288 °C of different carbon and LAS is almost identical as long as their steel sulphurcontent and morphology are very similar and their yield strength levels are roughlycomparable. At normal yield strength levels, which are representative for low-alloy primarypressure-boundary component steels (300 to 500 MPa), no distinct effect of yield strength on

34

SCC was observed [8, 55]. At very high yield strength (> 800 MPa) or hardness levels (< 350VH), a different behaviour is expected, since hydrogen-assisted EAC cannot be excluded,even at temperatures of 290 °C [54]. Here a distinct effect of yield strength has been observedon both SCC growth rates and Kiscc-thresholds [55]. It is stressed that the ferritic-pearliticand bainitic microstructures might show a significantly different behaviour at othertemperature/strain-rates, in particular at intermediate temperatures, because of their differentDSA response (see Section 4.5). Extrapolation from one material to another one shouldtherefore be treated with caution.

1/2.dK/dt [MPam'Th]

10"1 10° 101 102

10-*

CO

3a>2 -8

oO)•g 10"£

SA 533 B Cl. 1 (0.018 wt. % S)ECP = + 150 mVSHE> O2 = 8 ppm

K = 0.25 uS/cm, SO42 ' = 65 ppb

quasi-stationary CGR

CGR afterinitiation

o bainiticA martensitic• ferritic-pearlitic• bainitic

1 0,-7 10"6 1 0 5 10"4 10"3

dCODM/dt [mm/s]LL

Figure 22: Effect of microstructure on SICC growth in SRL-tests.

35

Q_

55

50

45

40

35

30

25

20

SA 533 B Cl. 1 (0.018 wt.% S)

Heat treatment/microstructure:o Q + T, bainitic• N, ferritic-peariitic

- o

02 = 8 ppmECP = + 150mV

K = 0.25 |aS/cmSO/' = 65 ppb

SHE

10-7 v-5

10° 1(T 10" 10

dCODM/dt [mm/s]

-3

LL

Figure 23: Effect of microstructure on SICC initiation in SRL-tests.

4.1.6 Effect of specimen orientation and location

The local sulphur content and morphology of MnS-inclusions depend on the steel making (—>segregation) and fabrication process of the plate (forming: hot rolling, forging,...) and theycan significantly differ in large plates [99]. Sulphur strongly accumulates in the liquid phaseduring the solidification process. MnS is a relatively low melting compound, and hencemigrates in the molten state toward the centre of a solidifying ingot. The concentration ofsulphides is generally higher in the centre of the plate. In segregation zones local sulphurcontent can be several times larger than the mean bulk sulphur content [100]. Sulphides tendto be rod-like in unidirectionally-rolled plates, and disk-like in cross-rolled plates, while inforgings the sulphides tend to be more spherical. The orientation (T-L, L-T, T-S, ...) andlocation of the specimens may therefore be important factors and a further source of scatter ofexperimental SCC/SICC data.

It has been reported, that the T-L orientation produces higher CF CGR than the L-Torientation which, in turn, produces higher CF CGR than the L-S orientation [101], while inan other reference, the rates for the L-S, L-T and T-L orientation were found to be essentially

36

similar [102]. The size and morphology of the MnS-inclusions seem to be as important as thesulphur content itself [3]. Plate-like or rod-shaped sulphides are often reported to be moredetrimental than spherical inclusions. This all suggests that steel making process and productform (plate vs. forging vs. weld) may play a role in EAC susceptibility. The low apparentEAC susceptibility of welds is attributed to the very small, uniformly-distributed sulphides,as well as to the compositional differences between wrought steels and welds [3, 8].

A distinct effect of specimen orientation (T-L, L-T, T-S, S-T, L-S, ...) or of specimenlocation (0T (top), lA T, Vi T (mid thickness), % T, IT (bottom surface)) could be neitherresolved in SRL- nor in LFCF- and CL-tests with all RPV materials investigated. This mightbe related to the highly oxidizing test conditions (8 ppm DO), which favour the enrichment ofsulphur-anions in the crack-tip environment and, to the fact, that the spatial distribution ofMnS-inclusions was fairly homogeneous in all of the steels and no distinct sulphursegregation zones were observed by sulphur prints or by post-test fractography. In SRL-teststhe SICC CGR were essentially the same (Figure 24). Only the high-sulphur alloy B revealedsome differences in the crack initiation stress intensity factors Ky-values for differentorientations. The Ky-values for the T-L- and L-T-orientation were similar, whereas theyseemed to be slightly higher for the T-S-orientation. Based on the few tests, it was notpossible to evaluate if the difference is significant, since the Ky-values are subjected torelevant scatter.

10 10° 10° 10" 10dCODM/dt [mm/s]

-3

LL

Figure 24: Effect of specimen orientation on SICC growth in SRL-tests.

37

4.2 Strain-induced corrosion cracking / very low frequency CF testsVery low-frequency corrosion fatigue (vLFCF) tests were performed in water at 288 °C witha DO of 8 ppm and 60 ppb SO4

2". In vLFCF-tests with a low load ratio R of 0.2 (AK = 51.1 to64.4 MPam1/2) and an extremely long rise time of 94 h, EAC crack growth could be resolvedby DCPD during each individual fatigue cycle (Figure 25 - 27). The crack growth initiationstress intensity factor Kia in each fatigue cycle is plotted for the high-sulphur alloy B and thelow-sulphur material A in Figure 28 and 29. Figure 30 shows the corresponding true time-based EAC CGR for each fatigue cycle for the alloys A and B tested simultaneously. EACoccurred only during the rising load phase of the fatigue cycles.

60

50

1 40

§ 30

20

10

5 0 j

1 ££,30

TJ 20.. (0o

_J 10-

i

\

\

1 1 1 1 1

/

/

/

\/

1//

-20

/

/

/

0 20

/

/

-40 60 80 100 '_

Time [h]

I

I

/

//

/ //

/

----

0 100 200 300 400 500 600 700 800 900 1000

Time [h]

Figure 25: Very low-frequency corrosion fatigue tests with a rise time of 94 h in high-temperature water with 8 ppm DO and 65 ppb SO42".

In the high-sulphur alloy B, EAC initiation was observed by DCPD already during the firstfatigue cycle (Figure 26 and 28). The observed Ki,;-value and EAC CGR were comparable tothe results of the corresponding SRL-tests with the same loading rate (Figure 31, compare toFigure 9 and 10). In the next two cycles Ky-values increased but stabilized after the thirdcycle, probably because of crack closure effects. The EAC CGR remained approximatelyconstant during all the fatigue cycles (Figure 30).

In the low-sulphur alloy A, EAC initiation was first observed by DCPD during the secondcycle (Figure 27 and 29). In the next cycle, the Ki,; value decreased; the EAC CGR increasedand then stabilized after the third cycle (Figure 30). The somewhat different behaviour hereduring the stabilization phase is probably caused by localized crack growth in alloy A duringthe first fatigue cycles, since less crack initiation sites were activated in the initial rising loadphase in this material. After stabilization, the Ky-values and EAC CGR were identical forboth alloys (Figure 30). This result clearly demonstrates, that monotonic SRL-test results at ahigh ECP with only minor, local SICC should not be interpreted as definite evidence of a lowSICC cracking susceptibility under vLFCF conditions.

38

3.0x10'

2.9x10'

2.8x10

§• 2.7x10"4

•J5 2.6x10"

0 2.5x10

°" 2.4x10"

2.3x10'

: >2.exio" •

2.6x10"-

1- I 2.6x10"

I

Cycle 4

400 420 440 460 480 500

Time [h]

SA 533 B Cl. 1 (0.018 wt.% S)= + 180mVcSHE

60 ppb SO4

t = 0: Start of loading cycle 1

0 100 200 300 400 500 600 700 800 900 1000

Time [h]

Figure 26: DCPD signal during vLFCF-test with high-sulphur alloy B.

2.8x10"4

2.7x10'

§" 2.7x10"4

•S 2.6x10"4

I£ 2.6x10"4

2.5x10'

. "

ential

dro

p

• S.

-

i

2*x10" j

2.6x10"

2.5x10"

2.5x10"

2.5x10"

2.5x10"

2.5x10"

>o

-~r—rj Cycle 1

/ no SICC

0 20 40 60 80 100

Time [h]

y

. . . i . . . . . .

JvX

20 MnMoNi 5ECP = +180

60 ppb SO42'

. . . . . . . . . . . . .

yA/

M -

-

5 (0.004 wt.% S) ;m V SHE J

0 100 200 300 400 500 600 700 800 900 1000

Time [h]

Figure 27: DCPD signal during vLFCF-test with low-sulphur alloy A.

39

COQ .

65

60

55

50

45

40

35

-

-

51.6" 0.20

0.58

57.1Q.190.81

- 4 —

i58.30.190.89

—r~

62.20.19

is T -nil I 072^ 0.75 •

0.19 1 x

0.78

SA533BCI. 1 (0.18

ECP = +180mVQHP, 60AtR = 94 h, dCODa/dt = 6.9

4

e:0-200.80 £ £

0.60

wt.% S)

ppb s o ;

* i6 2 ^ -0.210.79 -».peff

-8.1 E-7 mm/sasymmetrical saw tooth

1 1 1 1—H h-

-

—I :

AKR = K /K

mm max

8 10 11

Cycle number [-]

Figure 28: Crack growth stress intensity factorswith high-sulphur alloy B.

y for each fatigue cycle in vLFCF-tests

70

65

60

<C 55Q.

50

45

40

1 ' 1

" o m a X°-2wnoSKXbyDCPD

! iocat?-

I i; 54.5'. 0.19

0.89

• 1 1

55.80.190.77

i

- h -

1

55.6Iiiii

At R =

H -

1 '

57.957.8 O^9

0^9 0.780.78 T

i •20 MnMoNi

ECP = + 180

1

58.80.190.81

5 5(0mV

: 94 h, dCODLL/dt~="

asymmetrical s

1 \~—h

i

AK/ R = K

' 5 7 . 40,200.77

*

.004 wt.°

, 60 ppb

6.7 - 7.4

saw tooth

1

i

/ K •

mm max _

l,t max, N-1

56.40.200.73

\

/oS)

so;

55.0o.2i ;0.77

\ ]

E-7 mm/s ;

I

- 1 —8

Cycle number [-]10 11

Figure 29: Crack growth stress intensity factors Ki,j for each fatigue cycle in vLFCF-testswith low-sulphur alloy A.

40

10

10'

10

,-7

uo

£ 1010

V

10.-11

0 1

- • - Alloy B (0.018 wt.% S)- • - Alloy A (0.004 wt.% S)

Saw tooth loadingdCODLL/dt = 7E-7 mm/s, AtR = 94 h

R = 0.2, AK = 50 - 64 MPa-m1/2

ECP = + 180 mVSHE, 8 ppm O2

58 ppb SO42\ 0.023 nS/cm

T = 288 °C, "low flow"

H 1 1 h H h3 4 5 6 7 8 9 10 11

Cycle number [-]

Figure 30: True SICC CGR <AaSicc>/At for each fatigue cycle in a vLFCF-test with the low-and high-sulphur alloy A and B tested simultaneously.

10-7

LLJ

10-10

1/2,

10-1dK/dt [MPam"/h]

10"

ECP = + 150mV,

10-7

SHE

• SRL, SA 533 B Cl. 1, 0.018 wt.% S— Linear fit for •

95% Predicition interval for •• vLCCF, SA 533 B Cl. 1 (0.018 wt.% S)O vLCCF, 20 MnMoNi 5 5 (0.004 wt.% S)

10'dCOD/dt [mm/s]

LL

Figure 31: Comparison of CGR from vLFCF test with CGR from SRL-tests.

41

After stabilization, EAC occurred only during the last 20 to 25 % of the rising load phase ofthe fatigue cycles, i.e. at an effective load ratio Reff = Ko / Kiimax of ~ 0.8. Since EAC growthonly occurred during a small part of the whole rising load phase, apparent EAC CGR(Aa/AN)EAc/AtR may be significantly lower than true CGR daEAc/dt during each cycle. InFigure 31, the true EAC CGR from vLFCF-tests are compared to SICC results from SRL-tests under identical environmental conditions. Within the typical scatter range, these trueEAC CGR were identical to the SICC CGR for comparable loading rates. Additionally,results from constant load tests [50, 53] are included to Figure 32. These results furtherconfirm the fundamental relationship between EAC CGR and crack-tip strain-rate overalmost 6 orders of magnitude.

77 ioH

" 10

uCO

CO"D

-7

10"

10

10"

-10

Constant load testsVery low-cycle fatigueSlow raising load testsLinear fit of all data95% prediction interval

10" * 107 "* 105 10"4 103

dCOD/dt [mm/s]LL

Figure 32: Comparison of results from vLFCF-experiments with SRL and CL-tests revealingthe fundamental relationship between crack growth rate and crack-tip strain-rate.

42

4.3 Corrosion fatigue / low-frequency corrosion fatigue tests

4.3.1 Effect of temperature and loading frequencyTemperature/strain-rate effects on CF crack growth were investigated by LFCF tests withalloy A and B at four different temperatures (150, 200, 250, 288 °C) and frequencies (105,10"4, 810"4, 2.5-10~3 Hz) in water with 8000 ppb DO and 65 ppb SO4

2\ Positive saw tooth172172loading with a high R-value of 0.8 and a AK of 12 to 13.7 MParn172 was applied. For both

materials and all temperatures, the CF crack advance per cycle AaEAc/AN increased withdecreasing frequency, whereas the measured crack growth rate daEAc/dt decreased withdecreasing frequency (see Figure 33, which shows an example for T = 250 °C and alloy A).

5"5Stable, stationary CF crack growth was observed down to very low frequencies of 10"5 Hz.

100

2

!

10

1<

20 MnMoNi 5 5 (0.004 wt.% S)

T = 250 °C, DO = 8 ppm, 65 ppb SO4

R = 0.8, AK = 12.6 -13. 7 MPa-m1/2

110"9 ftu

i•a

10,-10

10" 5 10" 4

Frequency v [Hz]

10-2

Figure 33: Effect of loading frequency on LFCF growth.

For all frequencies, both cycle-based CGR AaEAc/AN and time-based CGR daEAc/dt increasedwith increasing temperature from 150 - 250 °C (see Figure 34 - 37). No noticeable change inCF CGR was observed by further increasing the temperature from 250 to 288 °C (Figure 36and 37). In the high-sulphur alloy B, the cycle-dependent CGR at 250 °C and 288 °C seemedto saturate or decrease again at the very low loading frequencies below 10"4Hz, eventuallyindicating the existence of a critical frequency. Based on the results from other temperaturesand other materials, it is concluded that this might be rather an experimental artefact thancritical frequency behaviour, since there are many reasons for crack arrest and local crackpinning, which could feign such a behaviour. This result should therefore verified by furthertests. In the temperature range from 150 - 250 °C an Arrhenius EA activation energy between40 and 50 kJ/mol has been calculated for the different frequencies and materials.

43

100

o

0.1

20 MnMoNi 5 5 (0.004 wt.% S)

DO = 8 ppm, 65 ppb SO42"

R = 0.8, AK = 12.0 -13.7 MPa-m

Temperature

250 °C288 °C

ASME XI ,,Wet"

10-5

150°C. . . . i

10 10° 10

Frequency v [Hz]

-2

Figure 34: Effect of loading frequency and temperature on cycle-based CGR in the low-sulphur alloy A.

100

10

o

0.1

SA 533 B Cl. 1 (0.018 wt.% S)

DO = 8 ppm, 65 ppb SO42"

R = 0.8, AK = 12 -13.7 MPa-m1'2

I ASME XI ,,wet"

-5

Temperature \

288 °C

250 °C

150°C

10"a 10 10° 10

Frequency v [Hz]

-2

Figure 35: Effect of loading frequency and temperature on cycle-based CGR in the high-sulphur alloy B.

44

10"

I 10,-9

io- 1 0

10-11

: 20 MnMoNi 5 5 (0.004 wt.% S); R = 0.8, AK=11.7-13.7 MPa-m. DO = 8 ppm, 65 ppb SO

100

= 1000 s• v = 1E-5Hz, At = 100000 s

H

150 200 250

Temperature [°C]

300

Figure 36: Effect of temperature and frequency on time-based CGR da/dtEAC in the low-sulphur alloy A.

10

10"

-7

10-9

O<LU

10"10

TJ

10.-11

SA 533 B Cl. 1 (0.018 wt.% S)R = 0.8, AK=12-13.7MPamDO = 8 ppm, 65 ppb SO4

1/2

100

v = 2.5 E-3 Hz, AtR = 200 s :v = 1E-4 Hz, At = 10000 s !

R

150 200 250

Temperature [°C]300

Figure 37: Effect of temperature and frequency on the apparent, time-based CGR in thehigh-sulphur alloy B.

45

Figure 38 shows a time-domain plot of these LFCF-tests with the low-sulphur alloy A. Thecorrosion-assisted crack growth rates da/dtEAC in oxygenated high-temperature water areplotted against the corresponding fatigue crack growth rate in air daJdtmen under otherwiseidentical loading conditions. The environmental acceleration of fatigue crack growthincreased both with decreasing loading frequency and increasing temperature. In LFCF testsunder these highly oxidizing conditions, environmental acceleration of fatigue crack growthin the range of 1 to 3 orders of magnitude was observed.

TO10

10"

^ 1 0

" 10"

-9

10

10"1 2

10-13

2.5E-3HZ 250 °C "8.3E-4 Hz j ^ - — " " 288 °C "i-

1E-4Hz

10-13 10- 1 2

20 MnMoNi 5 5 (0.004 wt.% S)

DO = 8 ppm, 65 ppb SO42

R = 0.8,AK = 12-13.7MPa-m1/2

10

Inert

10-10 10-9

da/dt . [m/s]

Figure 38: Time domain plot LFCF-tests with low-sulphur material A.

The low- and high-sulphur steels A and B showed a comparable CF crack growth behaviourunder these low frequency loading/highly oxidizing environment conditions (Figure 39 and40). The identical behaviour of simultaneously tested specimens is exemplary shown inFigure 39 for a temperature of 200 °C and different loading frequencies. At 288 °C, thecomparable behaviour under highly oxidizing conditions has been observed over a very widerange of loading conditions with different load ratios R, stress intensity factor amplitudes AK,and loading frequencies v (Figure 39). All the CGR data are within a small scatter band ofless than one order of magnitude. Preliminary tests in high-purity water without addition ofimpurities at 8 ppm or 400 ppb DO fitted very well to this scatter band. Deviations from thedescribed cracking behaviour were sometimes observed in individual specimens, eitherbecause of crack initiation problems or crack pinning/arrest.

46

Jtt)

I30

10

0.3

o

simultaneously tested ("daisy chain")• 20 MnMoNi 5 5 (0.004 wt.% S)O SA 533 B Cl. 1 (0.018 wt.% S)

Q

T = 200 °C

DO = 8 ppm, 65 ppb SO4

R = 0.8, AK = 12.2 -13.2 MPa-m1/2

10-5 10"4 10-

Frequency [Hz]10,-2

Figure 38: Similar cycle-based CGR Aa/ANEAc for the low- and high-sulphur steel A and Bin LFCF-tests in high-temperature water at 200 °C with 8 ppm DO and 65 ppbSO 2-

(fi

Eo<LLJ

T3

10"

T =288 °C= 8 ppm, 65 ppb SO4

2

v = 1E-5-3E-3HZ1/2R = 0.2 - 0.8, AK = 11 - 62 MPa-m T

simultaneously tested ("daisy chain")O SA 533 B Cl. 1 (0.018 wt.% S)• 20 MnMoNi 5 5 (0.004 wt.% S)i . i i ...i . . , i

10-12 10 •11 10-10 10-9 10' 10-7

da/dt. rt [m/s]inert

Figure 39: Comparison of time-based CGR from tests with the low- and high-sulphurmaterial A and B under different loading conditions in a time-domain plot.

47

4.3.2 Effect of loading conditions-6LFCF-tests at different frequencies ranging from 2.9-10" Hz to 1.4-10" Hz and at three

1/21/2)different R (and AK) levels of 0.2 to 0.34 (22. 9 to 64.4 MPa-m1/2), 0.7 to 0.88 (7.8 to18.4 MPam 1 / 2 ) and of 0.96 to 0.98 (2.3 to 3.2 M P a m 1 / 2 ) were performed at 150 - 288 °C inwater with 8000 ppb DO and 65 or < 1 ppb SO4

2". In Figure 40 the measured CF CGRare plotted versus the corresponding fatigue CGR da/dtinert in air under otherwise

identical loading conditions. Air fatigue CGR have been calculated according to Eason [14,15]. Additionally, the corresponding low- and high-sulphur lines of the GE-model [9] areplotted in Figure 40. CF CGR between the low- and high-sulphur line of the GE-model wereobserved. High CF CGR near to the high-sulphur line were achieved in low- and high-sulphur steels under many different loading conditions down to very low loading frequenciesof 2.9-10"6 Hz. Similar rates were also observed in preliminary LFCF tests in high-puritywater (< lppb SO4

2") at a DO of 8000 or 400 ppb. In LFCF tests under these highly oxidizingconditions, environmental acceleration of fatigue crack growth in the range of 1 to 3 orders ofmagnitude was observed. The environmental acceleration of fatigue CGR generally increasedwith decreasing loading frequency and was at maximum for ripple loading near to the AKthfatigue threshold. A noticeable mechanical fatigue crack growth contribution to the total CFcrack growth, and associated fatigue striations on the fracture surface, were only observed atloading frequencies > 10"3 Hz.

10-7

UJ0

10-10

10-11

• D 0.19<R<0.34• O 0.68 < R < 0.88A A 0.96 < R < 0.98

Solid symbols: 0.004 wt.% SfOpen symbols: 0.018 wt.% S

DO = 8 ppm: < 1 or 65 ppb SO4

2"T=150-288°C

:High-sulphurlineX.

;Low-sulphurline

da/dt = da /d ft inert

10-15 10-13 10-11

da.nert/dt [m/s]

Figure 40: Time domain plot for LFCF-CGR under different loading conditions.

48

In Figure 41, CF crack growth increments per fatigue cycle are plotted versus the appliedtotal stress intensity factor range AK and are compared to the corresponding ASME XIfatigue reference crack growth curves [52]. The CF crack growth rate Aa/ANEAC in LFCFtests significantly exceeded the ASME XI "wet" curve by a factor of 2 - 50 for both low- andhigh-sulphur steels and low and high load ratios. Values below the ASME XI "wet" curvewere only observed at loading frequencies > 10"3 Hz.

103r

^ 102

10"

10-2

• • D 0.19 <- • O 0.68 <: A A 0.96 <• Solid symbols: 0Open symbols: 0

"DO = 8 ppm, < 1: T = 150-288°C

A

: AA AA A

: A: A A

A

I A yA /

/

: R < 0.34R < 0.88 -R < 0.98 n :

004 wt.% S :.018 wt.% Sor 65 ppb SO4

2' "

O

1h*JL A rtn jr~ \ / i *

f MOIVIC /\ l

R > 0.65 !

ASME XI"R <0.25 !

10 100AK [MPam172]

Figure 41: Comparison of LCF CGR with ASME XL

The Aa/ANEAc values increased and the corresponding CGR da/dtEAC or (Aa/AN)EAc/AtRdecreased with decreasing loading frequencies (Figure 34 -38). In LFCF tests, CF CGR of theorder of some few up to several ten |im per fatigue cycle were observed. The maximum cyclebased SICC CGR Aa/ANEAc of up to some few hundred fim/cycle were observed under verylow frequency loading conditions. The environmental acceleration factor of fatigue crackgrowth increased with decreasing frequency and increasing temperature and load ratio(Figure 38 and 40). The maximum environmental acceleration of crack growth per cycle wasfound for ripple loading at very high load ratios of 0.96 to 0.98 near to the AKth fatiguethresholds.

49

The effect of the rise time AtR and hold time at maximum constant load AtcL (trapezoidwaveform) on the cycle-based LFCF CGR is shown in Figure 42 and 43 for a high load ratioR of 0.8. The cycle-based CGR Aa/ANEAc and the deviation from the ASME XI wet CGRincreased with both increasing rise time AtR and hold time at maximum constant load AtcL(Figure 42). The increase in cycle-based CGR with increasing hold time at maximumconstant load AtcL indicates, that corrosion-assisted crack growth also occurs at maximumconstant load and not only during the rising load phase. Based on our constant loadexperiments (see Section 4.4), we expect that the corrosion-assisted growth under constantmaximum load decays with time with cessation of crack growth after sufficiently long holdperiods. Therefore a saturation of cycle-based crack growth rates is expected for long holdperiods. It is stressed, that the corresponding time-based CGR da/dt decreased with increasinghold time at maximum constant load At<x (Figure 43). The results should be verified with alow load ratio and long hold times. Because of possible crack closure effects under theseconditions, reduced EAC CGR or complete cessation of EAC growth might be observed.

10"1

I

•»-2

o

f -

10"*

- * R: A At,

: R =

= 0.8,= 0h

= 0.8,.variation

• • R=

: AtH

0.33,= 0h,

,\K =

A t R

R

-12MPa

19.3

= 2.C

0.21

variation of AtP

of At

AK =

12MPa

= 46 MPa= 8.3h

.21 h

/ /

I

R

h

h

h

//

F

= 0.8

^ AtH = 1 0 h

^ A t H = 1h

[,'11

1 20MnMoNiT

ECP = + 150

t = 0.33

s

•^ ASMEXI

"Wet Curve

5 5 (0.004 wt.= 288 °C

mVSHE, 65 ppb

= 0.8 '•

0.33 "

-1 \

% S)]

so/- ;

10 1001/2-,

AK [MPa-nO

Figure 42: Effect of rise time and hold time at constant maximum load on cycle-based CGR.

50

100

.2

10

1

20 MnMoNi 5 5 (0.004 wt.% S), ECP = + 150 mVSHE, 65 ppb SO42

constant load amplitude, asymmetrical trapezoidal loading1/2

AK= 11.7-13.1 MPam , R = 0.8, AtR = 0.21 h

1

"10-10

10" 10"' 10" 10u 101

Hold time at constant maximum load AtCL [h]

Figure 43: Effect of hold time at constant maximum load on cycle-based and time-basedCGR at a high load ratio of 0.8.

51

4.3.3 Comparison to literature data

Only few data under similar test conditions could be found in the open literature [56 -59].Most of the available data is proprietary information and cannot be referenced. As shown inthe time-domain plot in Figure 44, the literature data at higher loading frequencies > =10~3 Hzunder highly oxidizing conditions fall within the scatter band of PSI results. Van der Sluys etal. [56] observed a decrease of the cycle-based CGR below a critical frequency between ca.

•v210 to 10" Hz (Figure 45). The critical frequency decreased with increasing AK anddecreasing temperature. In contrast to that behaviour, PSI had observed stationary EACgrowth and increasing cycle-based CGR down to very low frequencies of 10*5 Hz under verysimilar loading conditions. The ECP have not been reported in the tests of van der Sluys.Based on the DO of 200 ppb a slightly lower ECP than in the PSI-tests can be expected, butthis aspect is compensated to some extent by the significantly higher sulphur content of thesteel (0.025 instead of 0.004 wt.% S). The same temperature trends on CF CGR have beenobserved in oxygenated high-temperature water by KWU [57]. It is stressed that thetemperature trends might be different at lower ECP and turbulent flow rate conditions withcomplete or partial flushing of the crack-tip electrolyte [57, 60].

vt, AKT

1 0 H

110

10"

-7

3UJ

10-9

1T3

10-10

10-11

Envelope of data of van der Sluys200 ppb DO

T = 288 °Cv = 3E-6Hz-10Hz

R = 0.1 - 0.8, AK = 11 - 62 MPa-m1/2

1 0 " 10"12 1 0 " 10"10

da/dtinert

•icr 10"

[m/s]10'7 10'

• PSI, 8 ppm DO, 0.004-0.018 wt.% S• unpublisehd, 8 ppm DO, 0.013 wt.% S< KWU, 8 ppm DO, 0.011 wt.% S> unpublished, 400 ppb DO.0.012 wt.% S

A GE, 10 ppm DO, 0.0013 wt.% s• unpublished, 7-10 ppm DO, 0.01 - 0.021 wt.% S• van der Sluys, 200 ppb DO, 0.025 wt.% S

— "High-sulphur line" "Low-sulphur line"

Figure 44: Comparison of PSI-results with literature data under similar conditions.

52

102

0

I io°

10-1

10

10"

-2

• ASMEXI,,wet"

• A Van: • PSI,

• PSI,

/der Sluys,8 ppm DO8 ppm DO

R = 0

/ ^

1.200 ppb,288 °C,, 250 °C,

.8,AK = 11

DO, 288 °C0.004 wt.%0.004 wt.%

-13.7

, 0.025SS

MPa-m1/2 •

wt.% S 1

10-5 10"4 10'3 10"2 10"1 10°Frequency [Hz]

Figure 45: Comparison of PSI-results to similar investigations of van der Sluys.

The most important experimental finding of the PSI parameter study is the observation ofsustained, fast EAC growth down to very low loading frequencies under highly oxidizing(150 - 250 mVsHE) and low flow conditions. It is stressed that the ECP/water-chemistryconditions in PSI-tests conservatively cover most BWR plant transients. In most of the olderinvestigations, the measured or estimated ECP were slightly or significantly lower than theECP expected to be present on the stainless steel cladding on the inner RPV wall.

4.3.4 Comparison to ASME XI and to a proposal for an upgrading

As already shown in Figure 34, 35, 38, 41 and 45, the CF CGR Aa/ANEAc in LFCF testsunder highly oxidizing (8 ppm DO, +150 to +250 mVsHE) and low-flow conditionssignificantly exceeded the current ASME XI "wet" curve by a factor of 2 - 50 for both low-and high-sulphur steels and low and high load-ratios in the temperature range between 150and 288 °C. The situation was worst for ripple loading at very high load ratios near to thefatigue thresholds AKth. It is more appropriate to compare the ripple load CGR results to theBWR VIP 60 SCC disposition lines [35, 36] than to the ASME XI code (see Section 4.4).Values below the ASME XI "wet" curve were only observed at loading frequencies> = 8-10"4 Hz. Within the investigated parameter range, the excess difference to the ASME XI"wet" CGR increased with decreasing frequency and increasing load ratio and temperature.

The current reference CGR curves in the ASME BPV Code Section XI Appendix A [52] arebased on data obtained prior to 1980. They depend explicitly on AK and R-ratio, but not onother variables that are known to be important, such as loading frequency. In the meantime,several laboratories have conducted extensive testing programmes, including a wide variation

53

of controlled variables. Based on this extended database, a proposal for new reference crackgrowth curves has been worked-out by Eason et al. [14, 15], taking into account the strongeffect of strain-rate/loading frequency. The BWR database for this proposal was relativelysmall and mainly based on tests with an ECP of < + 50 mVsHE, temperatures around 288 °Cand rise times AtR < 1000 s. There is still a relevant lack of experimental data for high ECP,low loading frequencies < 10"3 Hz and intermediate temperatures, but also for very high R-ratios > 0.95. This lack of data therefore results in a relevant uncertainty concerning theconservatism and adequacy of the proposed reference curve for these parametercombinations. In Figure 46 the LFCF time-based CGR of PSI tests are compared with thisproposal. The proposed reference curves could be exceeded at low loading frequencies< 10~3 Hz and highly oxidizing and low-flow conditions.

The non-conservatism of the current ASME XI crack growth curves and of the proposal foran upgrading has been demonstrated for highly oxidizing and low-flow conditions. Theseresults indicate the need for further work and should be verified at lower Kimax- and AK-levels, corresponding to crack depths slightly above the non-destructive testing (NDT)resolution limit during periodic in-service inspection. The test conditions (ECP, T, dKi/dt, Kj,...) should also be further adjusted to correspond more closely to different BWR transientoperating conditions (start-up/shut-down, hot-stand-by, thermal stratification, ...).

<

10-6

10-7

If) 10

1-8

10-11

10-12

10

r

r

r

"2881200:150r

1 • • •

1E5s

vE5.

Proposal for upgradingof ASME XI (Eason et al

1E4s

11E4

1

s h

r

1E3s

- *JBL

1E3s I1 WF

s Xy

r

r20 MnMoNi 5

DO = 8 ppm,

R = 0.8, AK =

Xj^ a / d t E A C =

5 (0.004

r

da/dt.

Wt.%

65 ppb SO42'

12-13.7 MPa

'Anert

S) !

•m1/2 ]-13

1 0 1 4 10- 1 3 1 0 1 2 1 0 1 1 1 0 1 0 1 0 9 1 0 8 1 0 7 1 0 6

da/dtinert

[m/s]

Figure 46: Comparison of results from LFCF-tests with the proposed reference curves for anupgrading of the current ASME XI code.

54

4.3.5 Comparison to the GE-model

In Figure 48 and 49, the CGR results from the PSI-LFCF-tests in oxygenated high-temperature water are compared to the model developed by GE [9, 75]. The GE-model isbased on the film rupture/anodic dissolution mechanism and is of mixed mechanistic -phenomenological nature. In this model, crack growth through anodic dissolution is primarilygoverned by mechanical rupture of the protective oxide film via dynamic straining andthrough the strong effect of sulphur-anion activity on repassivation kinetics. The crackgrowth is mainly controlled by the crack-tip strain-rate and the sulphur-anion activity in thecrack-tip electrolyte, which govern the oxide film rupture frequency and the repassivationbehaviour after the film rupture event. A high sulphur-anion activity significantly delays theformation of a new, protective oxide layer and thus leads to a larger increment of crackadvance by anodic dissolution per oxide-rupture event. Based on the relationship betweensulphur-anion activity and repassivation derived experimentally under simulated crack-tipconditions, a lower and an upper limiting crack growth-crack-tip strain-rate equation could bedefined. The so-called "low-" and "high-sulphur line" represent an lower and upper boundingline for EAC crack growth in LAS in high-temperature water (Figure 47).

10"5

I

1 10"6

cr

O 10"i—

o

fy I \J

CL

10" 10

.-M i ' 0 ' i«T*» ) i i i i I i i i i I i i

,1CTS

li I i I • I • I i I I I I

1 0 (Kl lVE) *> 30 40 80 80^,100 /

MOOmCAHON OF M I I M L PROPAGATION RATE ALGORITHMS TOACCOUNT FOR POTENTIAL-DRIVEN DIFFUSION

/ • STAGNANT-FLOW• 0.02X SULPHUR

I I 11 III! I I I I I III! I I I | mil i mil"

10, -2

10"3

10"4 g

i d " 5

i d " 6

10"8 10~7 10"6 10-5 10"4 10CRACK TIP STRAIN RATE s - 1

Figure 47: Transition curves between low- and high-sulphur lines of the GE-model inoxygenated high-temperature water under quasi-stagnant conditions for differentECP and a LAS with a sulphur content of 0.02 wt.% S [9].

To sustain high-sulphur CGR, a high-sulphur-anion activity has to be maintained in thecrack-tip electrolyte. A high CF CGR is favoured by a high ECP and/or a high steel sulphurcontent/bulk sulphur-anion activity and quasi-stagnant flow conditions, which favour theenrichment of sulphur-anions in the crack-tip environment [8, 9, 58, 59]. If such a highsulphur-anion activity cannot be sustained, the CGR rapidly drop down to low-sulphur rates.For given environmental (ECP, bulk-sulphur-anion content) and material conditions (steel-sulphur content) "transition curves" between the "low-" and "high-sulphur line" can becalculated numerically [9, 61, 62, 75]. As shown in Figure 47, the transition lines are shiftedto higher crack-tip strain-rates (and loading frequencies) with decreasing ECP or decreasingsteel sulphur content. At high ECP, the transition occurs within a small strain-rate interval, atlow ECP the transition occurs over a wide strain-rate range.

55

In Figure 48, the CGR-results of different LFCF-tests in high-temperature water under low-flow conditions at 288 °C with 8 ppm DO (+150 mVSHE) and 65 ppb SO4

2' and LAS with alow and high sulphur content of 0.004 and 0.018 wt% are compared to the GE-model. Theexact transition lines for the given test conditions were not available. The transition lines foran ECP of +200 and +100 mVsHE are based on a steel-sulphur content of 0.02 wt.%, high-purity water without addition of SO42" and quasi-stagnant flow conditions. All CGR data werebetween the low- and high-sulphur line of the GE-model. At high loading frequencies theCGR were reaching near to the high-sulphur line. At the present stage it is not clear if thesustained, stationary EAC growth at very low crack-tip strain-rate around 10"8 s"1 (even inlow-sulphur steels) is in contradiction to the GE-model, since it is not clear how far thetransition lines are shifted to lower crack-tip strain-rates by the addition of 65 ppb SO4

2" inthe case of 8 ppm DO. LFCF-tests in high-purity water were only available at higher crack-tip strain-rates, where the CGR were comparable for 65 ppb and < 1 ppb SO42" at a DO of8 ppm. The sustained stationary SICC crack growth in SRL-tests in high-purity water(K < 0.06 u,S/cm) with a DO of 8 ppm at extremely slow loading rates (with rise times of290 h) would suggest that CF CGR significantly above the low-sulphur line might besustained under very low-frequency loading conditions. The conservatism of the GE-modelshould therefore be verified by further tests in high-purity water at very low loadingfrequencies and highly oxidizing conditions. It is stressed, that the sustained EAC growth atvery slow crack-tip strain-rates/very low loading frequencies might also be the result of theoccurrence of DSA, which has been observed in all investigated LAS to some extent.Possible DSA effects are not considered in the GE-model, and could result in an increase ofthe local crack-tip strain-rate.

10

10"

-7

10

ju 10

1

-9

10-12

"High-sulphur" iine

Transition line0.02 wt. % S

..quasi-stagnant'

= +100mVSHE

O SA 533 B Cl. 1 (0.018 wt. % S)• 20 MnMoNi 5 5 (0.004 wt.% S)

"Low-sulphur" line

1 0 ' 9 10"8 10"7 10"6 10"5 10"

Crack tip strain rate [s1]

Figure 48: Comparison of PSI-LFCF-test results in oxygenated high-temperature water witha DO of 8 ppm and 65 ppb SO4model.

2- (test conditions, see Figure 39) with the GE-

56

In Figure 49 the cycle-based CGR of a LAS with low sulphur content of 0.004 wt.% inLFCF-tests in high-temperature water with a DO of 8 ppm at different temperatures between150 and 288 °C are compared to the GE-model. Additionally, the corresponding ASME XIreference CGR for these loading conditions in an inert ("inert") and in high-temperaturewater environment ("wet") are plotted. The cycle-based CGR were below the high-sulphurline for all temperatures and loading frequencies. Based on a steel-sulphur content of 0.02wt.% and an ECP of +200 mVsHE, the GE-model would predict a critical frequency between10"4 Hz and 10"3 Hz for high-purity water. Above the critical frequency, the cycle-based CGRincrease with decreasing frequency as it has also been observed in the LFCF-tests. Below thiscritical frequency a decrease of the cycle-based CGR with decreasing loading frequency,which drops down to low-sulphur rates, would be expected.

Frequency [Hz]JO'6 10"5 10"4 10"3 10* 10"1 10° 101 102

I 10"10-2

20 MnMoNi 5 5 (0.004 wt.%S)8 ppm DO, 65 ppb SO4

R = 0.8,AK=12MPa-m1/2

10"9 10"8 10"7 10"6 10 ' 5 10"4 10"3 1 0 * 10"1 10°-1-Crack tip strain rate [s ]

Figure 49: Comparison of the cycle-based CGR from LFCF-tests with a low-sulphur steel(0.004 wt.% S) in oxygenated, high-temperature water with a DO of 8 ppm to theGE-model.

At the present stage it is not clear to which extend the critical loading frequency is shifted tolower frequencies by the addition of 65 ppb SO42" and if the GE-model also conservativelycovers the results at the lowest loading frequencies. Nevertheless, there is a lot ofexperimental evidence that sustained SICC and CF crack growth might be observed at slowcrack-tip strain-rates under highly oxidizing conditions and that the current reference crackgrowth curves in the ASME XI code are not conservative under these loading conditions.These loading/environment (ECP, water-chemistry) conditions are characteristic for BWRoperational transients such as plant start-up/shut-down. Plant start-up is also related toincreased global and local conductivities [63]. It is believed, that the PSI-tests conservativelycover these conditions, in contrast to some other laboratory investigations in high-puritywater (K < 0.07 |iS/cm) with DO of 200 to 400 ppb.

57

4.3.6 Comparison to SRL- and vLFCF-tests

The time-based CGR from SRL-, vLFCF- and LFCF-tests under identical environmentalconditions (8 ppm DO, 65 ppb SO4 ", 288 °C) and similar loading conditions (similar dKi/dtand Ki™ ) are compared to each other in Figure 50. The corrosion-assisted CGR are plottedagainst the measured crack-opening displacement rate at load line, which is a rough measurefor the crack-tip strain-rate. In the case of SRL- and vLFCF-tests the true SICC CGR couldbe measured. In the case of LFCF-tests, only apparent, time-based CGR (Aa/ANEAc)/AtRcould be determined.As already mentioned in Section 4.2, the true time-based CGR of SRL- and vLFCF-testswere identical for similar loading rates. The apparent CGR (Aa/ANEAc)/AtR of vLFCF-testswere smaller than the true CGR da/dt, since the corrosion-assisted crack growth onlyoccurred during the final part (the last 10 - 30 %) of the rising load phase. On the other hand,the apparent CGR of LFCF-tests were significantly smaller than the true CGR of SRL- andvLFCF-tests for identical loading rates and Kimax-values. The difference increased withincreasing loading rate, frequency and load ratio R. The power law-relationship betweenCGR and dCODLi/dt for the LFCF-tests had a smaller slope (m = 0. 5 - 0.7) than thecorresponding relationship for SRL-tests (m = 0.8). The difference seems to be quite large,even if the observations of the vLFCF-tests is considered, that the corrosion-assisted crackgrowth occurred only during the final part of the rising load phase. If the same true CGRwould be assumed as in SRL-tests, than the corrosion-assisted crack growth could have onlyoccurred during less than 1% of the rising load phase!

The reasons for the large difference and different slope under these low-frequency loadingconditions (where the contribution of pure mechanical fatigue crack growth can be neglected)are not yet completely clear. The observation that the corrosion-assisted crack growth doesnot occur during the whole rising load phase in vLFCF-tests may be explained at least in partby crack closure effects. But crack closure effects alone may probably not explain thecomplete difference between the apparent CGR of LFCF-tests and the true CGR of SRL- andvLFCF-tests. A convection-induced change of the crack-tip chemistry induced by the cyclicmovement of the crack flanks during fatigue loading could be another plausible explanation[8]. ,,Fatigue pumping" results from the relative displacements of the crack flanks betweenmaximum and minimum stress portions of a fatigue cycle. A certain volume of water couldtherefore be pumped into, and out of, the crack during each fatigue cycle. This can result in adilution of the crack electrolyte (and therefore in a decrease of CGR). Such a process wouldbecome of greater importance as cyclic frequencies increased and with lower stress ratio andhigher flow rates across the crack mouth. Furthermore, the crack-tip strain-rate might bedifferent for cyclic and single slow rising loading conditions in spite of similar dCODLi7dt-and Ki-values, since the cyclic plastic deformation behaviour might be significantly differentfrom that under monotonic loading (cyclic hardening or softening, DSA, ...).

58

10-5

^ 10" 6

10-7

10"

10"

1/2

LCF, constant load amplitude, saw tooth: Aa/AN / AtR

• R = 0.33, AK = 41 - 53 MPam1/2

SA 533 BCI. 1,0.018 wt.% S

• R = 0.33, AK = 43.7 - 50 MPam1

20 MnMoNi 5 5, 0.004 wt.% S

% R = 0.8, AK = 11.7- 12.3 MPam1

20 MnMoNi 5 5, 0.004 wt.% S

O R = 0.68, AK = 7.8 -11.6 MPam1/2

SA 533 B Cl. 1, 0.018 wt.% S

0 R = 0.79, AK= 10.4-11.4 MPam1

SA 533 BCI. 1,0.018 wt.% S

1/2

SRL Tests

A vLCF, R = 0.2SA533 BCI. 1 -0.018 wt.% s

V vLCF, R =0.220 MnMoNi 5 50.004 wt.% S

10-7

10" 10"

dCODLL/dt [mm/s]

Figure 50: Comparison of CGR from SRL-, vLFCF and LFCF-tests in oxygenated, high-temperature water at 288 °C with a DO of 8 ppm and 65 ppb SO4

2".

59

4.4 Stress corrosion cracking4.4.1 Summary of the main results of SpRK II

The following Section summarizes the most important experimental findings of the SpRK IIprogramme [26, 28, 64, 65]. During transient-free, steady-state BWR power operation, staticloading conditions and temperatures between 270 and 290 °C prevail at the RPV. Within theSpRK II programme, the SCC crack growth behaviour of different low-alloy RPV steels(base metal) was therefore investigated under constant displacement and constant load inoxygenated, high-temperature water at 288 °C. The DO and the conductivity of theenvironment were systematically varied between 0.2 to 8 ppm (-50 mVSHE ^ ECP <+200 2and 0.06 to 1.0 uS/cm ( < 1 ppb < SO4

2" < 360 ppb).

The tests under constant displacement did not reveal any indications for SCC growth, evenunder aggressive environmental conditions with 8 ppm DO and 360 ppb SO42". In contrast toconstant displacement tests, EAC could be initiated by active loading within the testenvironment, if an appropriate combination of mechanical, material and environmentalconditions was simultaneously attained. If EAC was observed by DCPD, crack growth hadalways initiated whilst rising the load to the intended value for the subsequent, constant loadexperiment. Following initiation, the crack grew under rising load conditions, and thecorresponding crack advance Aasicc can be attributed to SICC. The crack advance by SICCAasicc was strongly dependent on the loading rate applied during the rising load phase. Aftertransition from rising to static load, the CGR decayed continuously. This is exemplary shownin Figure 51 for a test with a relatively slow rising load phase with a rise time of ca. 25 h.

Loading phase Experimental phase: constant load

[m/s

];

da/

dt

th r

at<

|2

Cra

ck

10"8

109

10"1 0

10"

/•• • I • • • • I • V

SICC .

[ /

r

/ V

10 20

^ Aa a In t

SCC

SA 533 B Cl. 1, 0.018 Gew.% S, LTK, = 73 MPa m1/2

T = 288 °C, O2 = 8 ppm, ECP = + 170 mVSWE

K = 0.25 nS/cm, 65 ppb SO42"

\ ^

• ^ * S N " ^ Hci/rlt ju 1/t

200 400 600 800

Time t [h]

-

• ;

1700

1600

1500

:1400

1200

1000

800

600

400

200n

1000

(0

0)

ua

•o(0

So

Figure 51: SICC during rising load and SCC during constant load. The SCC growth rate ofca. 5 • 10"8 m/s at the beginning of the constant load phase continuously decaysdown to a value < 2 • 10"11 m/s after 943 h of constant load testing.

60

The corresponding crack advance under static load Aascc is assigned to SCC, in accordancewith the definitions in Table 1. Fractographically, only the total crack advance AaEAC, whichis the sum of Aasicc and Aascc> c a n be determined. Apart from the extent to which thefracture surface is oxidized, fractography does not supply any direct, time-related informationon crack-growth initiation and advance. Based on fractography, therefore, only an apparentCGR AaEAc/AtcL can be derived. Based on the development of crack growth monitored on-line (Figure 51), it is evident that the apparent, fractographic CGR is a conservative estimatefor the SCC CGR under static load. The fractographically derived apparent CGR may differby several orders of magnitude, even for identical material and environmental conditions andcomparable Ki levels, depending on the relevant combination of CGR determination method,constant load test duration, and loading procedure prior to testing under constant load. Thusour observations provide a further explanation for the reported scatter in EAC CGR in LASunder BWR conditions.

4.4.1.1 Effect of stress intensity

In our standard tests with a constant load test period AtcL of 1000 h, the specimens wereloaded up at a relatively high rate dKi/dt of > 10 MPam1/2/h within a rise time of typically0.3 to 2 h. This procedure was intended to restrict the initial crack advance by SICC, but,even under these conditions, the crack increment Aascc under subsequent constant load wasalways < 20 % of the total crack advance AaEAC ; at least for tests with Ki < 80 MPa-ml/2.Therefore, the apparent CGR AaEAc/AtcL were generally more than one order of magnitudehigher than the actual SCC CGR under static load.

In Figure 52 , the apparent, maximum CGR AaEAcmax/AtcL from our standard SCC-tests witha constant load testing period of 1000 h are plotted versus both the applied stress intensityfactor Ki and the dimensionless parameter KI/K^ASTM E399- The testing conditions include testswith low- and high-sulphur steels at conductivities up to 0.5 fxS/cm (< 165 ppb SO42") underlow-flow conditions, both at a high ECP of +150 ITIVSHE (8 ppm O2) and at realisticconcentrations of oxidants (200 - 600 ppb O2). With regard to Ki, ASTM, three regimes can bedifferentiated:

Regime I (Ki <

There was no corrosion-assisted crack advance found within this regime, except in one testwith high-sulphur steel (0.018 wt.% S) at a Revalue of 41 MPam1/2 (very close to theKtASTM-value of 42 MPa-m1/2). Here, a maximum crack increment < 100 Jim and a meancrack increment < 20 |im were observed in the SEM; these were too small to be detected byDCPD. On the basis of the high degree of oxidation of the fracture surface, and of ourseparate results quantifying the influence of loading rate on crack-growth initiation (Section4.1), we conclude that the observed crack advance took place while rising the load to theintended value for the constant load experiment.

61

K. [MPa m12]i

20 30 40 50 60 70 80 90 100

E

- io2 ^

: SEM detection limitno corrosion-assisted crack

OC(0•ats

2o

K./K,, ASTM H

Figure 52: Apparent maximum CGR vs. applied Ki for standard tests under constant, active,external loading with 1T-C(T) specimens. T = 288 °C, K < 0.5 nS/cm, < 165 ppbSO4

2\ < 0.018 wt.% S. Open symbols are for tests with 200 - 600 ppb O2 (-50 to+40 HIVSHE), solid symbols for test with 8 ppm O2 (+150 mVSHE)- Low- and high-sulphur line are taken from Ford [9].

Regime II (Ki, ASTM < KJ < 2 Ki, ASTM):

Thanks to the on-line measurement of crack growth, the following, important observationswere made within this regime:

• Crack growth under constant load only took place if the crack advance had already beeninitiated during the period in which the load was increased to the intended value forsubsequent testing.

• If crack growth under constant load was detected, it was of limited duration and crackarrest generally occurred within a period of 1000 h. This observation is indicated bydownward pointing arrows on the data points in Figure 52. Up to a Ki of 60 MPa-m1/2,crack arrest occurred within a very short period (10 - 100 h) after the transition from risingto static load and the total, mean crack advance under static load <Aascc > was very small(< 60 Jim). The SCC crack growth increment over the remaining, constant load test time of

62

900 to 1000 h was below the DCPD detection limit of ~ 20 ^im, corresponding to a CGRof < 6 • 10"12 m/s (< 200 |j,m /year). Because of the extremely low susceptibility to SCCcrack growth under these conditions, no effect of Ki on SCC CGR could be resolvedbelow 60 MP-m1/2.

• The SCC CGR under constant load decayed in a similar way to the strain-rate from low-temperature creep, i.e. following a reciprocal time law:

dascc/dt« 1/t under constant load (3)

Therefore, the crack advance Aascc followed a logarithmic time law:

Aascc <* ln(t) under constant load (4)

This is illustrated by typical, on-line crack growth measurements for two different levelsof Kiin Figure 53.

In region n, at least up to a Ki level of 80 MPam1/2, the apparent, maximum CGR inexperiments with both high- and low- sulphur steels (Figure 52) were often well below theso-called low-sulphur line [9, 62]. Based on the water-chemistry conditions and on thedevelopment of crack growth during the test, it can be concluded that the low-sulphur line isa very conservative, upper-bound estimate for constant load SCC CGR in low-alloy, RPVsteels under conditions relevant to stationary BWR power operation.

250

200

150

100

50

SA 533 B Cl. 1, 0.018 wt.% S, LTT = 288 °C, O2 = 8 ppm, ECP = + 170 mVSHE

K = 0.25 nS/cm, 65 ppb SO4"Initial loading in 25.1 h, dP/dt = 1.24 kN/h

K, = 61 MPa m

100 200 300 400 500 600 700 800 900 1000

Time from initial loading [h]

Figure 53: Time-dependent crack advance under constant load in regime II (KLASTM ^ KI <2 • KLASTM)- The CGR decays continuously with time. Shortly after the transitionto static load, the rate of the crack advance follows a logarithmic time law.

63

Regime III (Ki > 2

In most cases, the rate of crack advance decayed during a test period of 1000 h at Ki levelsabove = 84 MPam1/2. However, sustained, corrosion-assisted crack advance was alsoobserved in some cases and sometimes amounted to several mm within a test period of1000 h (corresponding to a CGR > 10 mm/year). This observation is thought to be due tosustained, low-temperature creep of the remaining specimen ligament under conditions ofgross-ligament yielding, which allows a positive, crack-tip strain-rate to be maintained. Thisconclusion is clearly supported both by low-temperature-creep experiments with pre-crackedlT-C(T)-specimens at 288°C in air [66], and by FEM modelling of visco-plastic deformationbehaviour under the same conditions [70].

45

44

43

42

41

40

39

38

37

?fi

'-

'-

Aa s c , :

/j CODIL ;

4.0

3.5

3.0

2.5

2.0

1.5

1.0

0.5

0.0

-_

[mrr

o \o .w

^ ;

-_

1.3

1.2

1.1

1.0

0.9

0.8

0.7

0.6

0.5

O

40 41 42 43 44 45 46 47 48 49 50

Time from initial loading [h]

E

•—'A

oo

il 2oca

T3(3

O

6

Linear Regression: ^»#^dascc/dt = 1.4E-7m/s > " ^ ^scc .R = 0.9998 \ ^ j r

• ^^^^^^ ^

• ^ T

~r

125

120

115

110

105 E

100 •§

95 x.

90

85

«n42 43 44 45 46 47

Time from initial loading [h]

48

Figure 54: Stationary, fast SCC crack growth in specimen loaded near to Ku. Alloy B(0.018 wt.% S), 288 °C, 8 ppm O2; +170 mVSHE; 65 ppb SO4

2'.

Steady-state crack growth was observed solely in the high-sulphur alloy D (0.018 wt.% S)and only in those cases where the applied Ki approached the Ku-value of 105 MPa-m1/2. Thissituation resulted in continuous opening-up of the specimen as fast as the loading system

64

would allow. The maximum, pull-rod stroke rate was not sufficient to maintain the load leveland continuous unloading of the specimen therefore resulted (Figure 54). Even in thisextreme situation, the (micro-) fractographic appearance of the fracture surface was found tobe that of EAC in LAS. Interestingly, the effective J- or Ki-value of the rapidly growing crackremained approximately constant during this unloading of the specimens. Under suchconditions, CGR in the range of 1 to 3 • 10~7 m/s were measured, which agree with the SICCCGR found in the rising load phase for the same crack-opening-displacement rate (seeSection 4.1, Figure 10). At similar Ki-levels, both the crack-opening-displacement and crack-growth rates were much lower in the low-sulphur material (with a significantly higher Re-value) and continuously decayed during the test period.

4.4.1.2 Effect of environmental parameters

Within the chosen range of environmental and material parameters, no distinct effect of theseon the SCC crack growth under static loading could be found, simply because there is a verylow susceptibility to SCC growth, at least as long as gross-ligament yielding is avoided.However, the following important aspects concerning environment and material can still bederived from our tests:

At the present time, the following environmental margins can be established for SCC crackgrowth under static load in the investigated low-alloy steels:

• No SCC is observed for ECP <+150 mVSHE and conductivity < 0.25 |J,S/cm(corresponding to < 65 ppb sulphate) as long as small-scale yielding according to ASTM E399 prevailed in the specimen ligament. Under these conditions, fast SICC crack growth(up to 10"7 m/s), as triggered by suitable, slow rising load, could not be maintained aftertransition to static loading and crack arrest occurred within a short period (10 - 100 h).Preliminary results indicate that this statement might even be valid up to a conductivity of1.0 nS/cm (corresponding to 360 ppb SO4

2") at a high ECP of +150 mVSHE-

• Sustained, fast SCC crack growth with respect to the operational time scale of a BWRcannot be excluded for high ECP and/or a conductivity > 1 (xS/cm (> 360 ppb SO42"), i.e.under water-chemistry conditions exceeding EPRI action level HI. Since no SCC-testswere performed under these conditions in the present program, this statement is primarilybased on older tests in static autoclaves [30] and tests in refreshed hot-water loops [67]under quasi-stagnant flow conditions. These results indicate that steady-state, fast SCCmight occur in oxygenated high-temperature water, even under purely static loading andsmall-scale yielding conditions, with high concentrations of specific impurities (e.g. SO42"+ Cl" > 500 ppb - 1 ppm).

It is emphasised that the above SCC margins are valid only for purely static loading. EACcan occur at much lower values of both ECP and anion concentration if there is a positivestrain-rate acting at the crack-tip (see Section 4.1).

4.4.1.3 Effect of steel sulphur contentA difference in behaviour between low- and high-sulphur steels under static load was onlyobserved if small-scale yielding conditions were clearly exceeded and the stress intensityfactor in the high-sulphur material approached the Ko-value (more precisely: J —» Ji). Underthese conditions, steady-state, fast SCC was observed in the high-sulphur material. At thesame Kj level, the CGR in the low-sulphur material were much lower and, furthermore, theydecayed during the test period. This observation relates to a higher Ku-value for the low-sulphur steel. In general, in quenched and tempered LAS, Ku decreases with increasing steel

65

sulphur content, yield strength level and grain size. In segregation zones of very high-sulphursteels, the local Ku-value could be equal to or even smaller than the Ki-limit of ASTM E 399(KLASTM)-

With respect to EAC susceptibility, a low sulphur content in the steel may be beneficial inmany cases. The higher toughness of these steels also results in higher allowable/criticalcrack sizes, and hence in a higher safety margin, but as already discussed (Sections 4.1 - 4.3),EAC cracking susceptibility in low-sulphur steel may be as high as in high-sulphur steels forcertain combinations of corrosion, mechanical and environmental parameters relevant toBWR power operation.

4.4.1.4 Comparison to recent literature data and proposed disposition lines

Table 6 gives a summary of the observed crack growth data from all tests performed underactive, external load control. Both standard tests, with relatively rapid initial loading rates,and tests with very low initial loading rates are considered here.

Corrosion-assisted crack growth

K,[MPa-mI/2]

K|<30

30 < K, < 60

60 < Kt < 80

K,>80K] > 2 Ki ASTM

K, ->KV"

Region

SSY

= SSY

transitionregion

ligamentyielding

Initialloading

no

yes

yes

yes

Subsequentstatic loading

no

strictly limited in time=10 'h -=10 2 h

limited in time=102h-=103h

sustained> 103 h, > test period

steady-state crack growth

AaEAC/AtCL

AtCL = 1000 h[m/s]

< 10"*

< io-10

<10"9

> 3 • 10'°< 3 1 0 7

1 -3-10'7

AaSCC /AtCLAtCL = 1000

h[m/s]

< io-"*

<2- 10""

< io-10

> 3 • 10'°<310"7

1-3-KT7

«k/dtDCPD(t)t=1000h

[m/s]

<10"*

< 10"*

<3 10"

> 3 • 10'°< 3 1 0 7

1-3-10"7

Alloy(Table2 + 3)

AtoF

*

dCODLL/dt[mm/sl

2 • 10 7 to 6 • 10"4

< 10-um/s: < detection limit **Kn = (JiE/d-V))"2

dKi/dt[Mra-m^/h]

0.2 - 330

AtR

[h]

0.3 - 290

T

ra288

K

[US/cm]

0.06 to 0.25

so42-

[ppb]

Ito65

ECP[mVSHE]

-50 to + 200

o2[ppm]

0.2 to 8

Table 12: Compilation of data from active loading tests with 1T-C(T) specimens of differentsteels containing 0.004 to 0.018 wt.% S.

For the investigated steels and the bounding environmental conditions (DO < 8 ppm,equivalent to an ECP of ca. < 150 mVsHE, conductivity < 0.25 |lS/cm, equivalent to 65 ppbSO42"), a conservative, stress-intensity-factor threshold Kiscc of = 30 MPa-m1/2 can beestablished (see Section 4.1). This value, which we derived from SRL-tests, is in accordancewith threshold values of 20 - 30 MPa-m1/2 reported in the literature for tests under constantloading conditions [30, 55].

Our CGR data are substantially consistent with other, recently published data from well-qualified experiments. This is shown by a comparison of the PSI-results with both data ofMPA Stuttgart [27, 32, 33] and the results of a European Round Robin test [34]. Todifferentiate these data according to both the level of dissolved oxygen in the water and thesulphur content in the steel, the same results have been plotted in two different diagrams ofapparent CGR versus applied stress-intensity-factor Ki (Figure 55 and 56). Only constant,active, external-load tests using IT and 2T C(T) specimens (with corresponding KI,ASTM

values of ca. 40 and 60MPam1/2) and with the duration of static loading AtcL ^ 1000 h

66

(except if Ki —> Kn) were considered. The rise time AtR of initial loading was typicallybetween 0.3 h and 2.0 h. In most cases, the fractographically derived, local maximum crackadvance AaEAC™* was used for calculation of the apparent CGR AaEAc/Atci,- In a few cases,only the average crack advance <AaEAC> was available to the present authors. In addition, thelow-sulphur line according to Ford [9] is plotted in Figure 55 and 56. The materials taken intoaccount are representative for the RPV of modern, Western LWR (after 1965) and fulfil therequirements of the relevant nuclear codes (ASME BPV, KTA).

10H

3» 109

rf 10"10

LtJ

A "low sulphur": S < 0.009 wt.%# "medium sulphur":

0.009 wt% <= S <= 0.013 wt.%+ "high sulphur": S > 0.013 wt.%

"low sulphur line" ^

104

103

0

10"1

0 10 20 30 40 50 60 70 80 90 100

K. [MPam1/2]

Figure 55: Comparison of PSI-results with recently published data of MPA Stuttgart [32]and of an European Round Robin test [34] for different sulphur contents in thesteels.

67

10"

10'

iff31

10

10

-7

-8

"9

•10

-11

-12

-13

" OA

: •

r

-

i /

o,

o?o2

= 200= 200-= 8000= 8000

-600600ppb,

PPb,

ppbppb,T =T =

-" low sulphur line"

/

V

/

/

\

) 8

8J

,T =T =288240

«

Slh

•

0

= 288 °C 1240 °C %:°C j ^°C

A :

^ ^ A #O

k :

Ao !

DO°O :

i 10J

102 ,_,CO

-i- 10u

- 1 0 1

- 10'

0 10 20 30 40 50 60 70 80 90 100

K, [MParn172]

Figure 56: Comparison of PSI-results with recently published data of MPA Stuttgart [32]and of an European Round Robin test [34] under various environmentalconditions.

Within a broad range of environmental (Figure 55) and material parameters (Figure 56), nosustained SCC crack growth (i.e. CGR < 10"11 m/s, < 300 |im/a) was observed under staticloading up to a Ki-value of approximately 60 MPam1/2. Furthermore, the apparent, EACCGR (i.e. calculated using the total crack advance) were bounded by the low-sulphur line upto the same Ki-value. Above a stress intensity of 60 MPa-m1/2, most data points were stillbelow the low-sulphur line and sustained cracking was observed only in high-sulphur steels.In such cases, the applied Krvalues approached the Ku-value of the material concerned.

Furthermore, for experiments where Ki was below Ku, the biggest part of the EAC crackadvance plotted in Figure 55 and 56 took place during initial loading or at the very beginningof the constant load phase of the experiment. For Ki < 60 MPam1/2 (and sometimes at highervalues), the rate of crack advance during static loading decayed and crack arrest wasgenerally observed within a period of 1000 h. This was shown directly by on-line crack

68

growth monitoring in the PSI-experiments, while the same behaviour has been indirectlyobserved in the MPA-experiments [32].

Based on the results of qualified tests in simulated BWR environments, interim dispositionlines for assessing possible SCC crack growth in low-alloy, pressure-boundary-componentsteels during BWR power operation have been proposed by an international group of expertsin the context of the BWR VIP project [35, 36]. These bounding curves, which have recentlybeen accepted by the US NRC [35], consider both the disposition line proposed by MPAStuttgart [32], and the low-sulphur line [9]. Our results are consistent with both thesedisposition lines as shown in Figure 57. Moreover, the underlying material behaviour isverified by the PSI-experiments. In particular, considering tests of high-sulphur steel(0.018 wt.% S) at conductivities of up to 0.25 jxS/cm (equivalent to a single-impurity level of65 ppb SO42 and > EPRI action level II) in oxygenated, high-temperature water at a highECP-value of +150 mVsHE* it was shown that an active, fast-growing EAC crack sloweddown rapidly under subsequent constant load conditions and arrested within a testing time of1000 h. It should be mentioned that the high ECP values set in the PSI-experiments arerealistic for the RPV during BWR power operation, whereas most of the MPA tests havebeen performed at significantly lower potentials under quasi-stagnant flow conditions.

K, [MPa-m172]10 20 30 40 50 60 70 80 90 100

101

10,-121

K./K,,ASTM H

No corrosion-assisted £ra*k;igrowth^:iSEM-det< ction limit "

104

103

102

101

10°

10-1

year

s.ac

k ad

var

O

Figure 57: Comparison of PSI-results form constant load tests to BWR VIP 60 SCCdisposition lines [35, 36] and to low- and high-sulphur line [9]. The PSI-testsconfirm the conservative character of the disposition lines for static loadingconditions and 288 °C. (T = 288 °C, K < 0.5 nS/cm, < 165 ppb SO4

2\ <0.018 wt.% S. Open symbols are for tests with 200 - 600 ppb O2 (-50 to+40 mVsHE), solid symbols for test with 8 ppm O2

69

From an engineering point of view, the SCC crack growth behaviour of the investigated low-alloy RPV steels under static loading conditions at 288 °C can thus be conservativelydescribed by the following equations:

• Ki < Kiscc = 30 MPa-m1/2: no SCC (daSCC/dt < 1 - 10"11 m/s)

30 < Ki < 60 MPa-m1/2: daSCC/dt < 2 • 10"11 m/s ± f(Ki)

60 < Ki< 2-KI.ASTM (« 80 MPam1/2) or Ki< 0.8 Kn:

dascc/dt < 3.29 • 10"17-Ki4 ("low-sulphur line")

4.1.4.5 Conclusions form the SpRK II programme

At 288 °C, the investigated, low-alloy, RPV steels revealed no sustained, SCC crack growthunder purely static loading conditions over a wide range of loading and environmentalconditions, as long as small-scale yielding conditions prevail at the crack-tip, and the water-chemistry is maintained within current BWR/NWC operating practice (EPRI water-chemistryguidelines).

However, fast and sustained SCC crack growth cannot be excluded:

• for faulted, water-chemistry conditions (> EPRI action level 3), where the guidelinesrecommend prompt shut-down of the reactor.

and/or

• if a positive strain-rate at the crack-tip can be maintained under external, static loading,either by gross-ligament yielding of material with high net-Section stresses, or byoverloading to stress intensities near to Ku.

High SCC CGR of up 3-10"7 m/s, as sometimes reported in older literature, were not observedwithin this parameter study under simulated BWR/NWC water-chemistry conditions.Otherwise, recently published CGR data of MPA Stuttgart, and from a European RoundRobin test, were confirmed by the present study, even though the tests at PSI were performedunder more severe environmental conditions (in particular at higher ECP-values, approachingthose expected under high-flow and/or radiation flux conditions in the RPV). Moreover, itwas demonstrated that even fast, SICC crack growth cannot be maintained during subsequent,static-load testing up to a high Ki level of 73 MPam1/2. The PSI-data thus support theadequacy and conservative character of the interim disposition lines, recently proposed in thecontext of the BWR VIP project, for assessing possible SCC crack growth in LAS duringBWR power operation.

The disposition lines have just been validated for RPV steels (base metal, < 0.020 wt.% S) inthe quenched and tempered (and stress relieved) state, for temperatures around 288 °C andpure static loading conditions. The conservatism of the disposition lines has therefore to befurther verified for weld and weld HAZ materials, for lower intermediate temperatures (100 -250 °C) and for small load fluctuations.

70

4.4.2 Effect of temperature

In the feedwater piping system and the region of the feedwater RPV nozzle lower operatingtemperatures of 200 - 240 °C prevail and at least temporary small load fluctuations cannot beexcluded here. Therefore, both the effect of temperature and small load fluctuations on theSCC growth in LAS was investigated.

The effect of temperature on the SCC growth was investigated in tests under constant loadwith a steel with a low- (alloy A, 0.004 wt.% S) and a high-sulphur content (alloy B,0.018 wt.% S) at temperatures between 150 and 288 °C. The tests were conducted inoxygenated high-temperature water with 8 ppm DO and either 65 ppb or < 1 ppb SO42". TheECP slightly decreased form +250 mVSHE at 150 °C to +150mVSHE at 288 °C. The constantload period started with an actively growing EAC crack (3 mm/a to 300 mm/a), which wastriggered by periodical partial unloading at a load ratio of 0.8. Constant load amplitudeloading with load control and a positive saw tooth waveform was applied for the periodicalpartial unloading.

The SCC growth behaviour of the RPV steel with the high sulphur and aluminium content(alloy B) revealed the same SCC crack growth behaviour over the whole temperature rangeas at 288 °C and as discussed in Section 4.4.1. The SCC crack growth continuously decayedafter switching form cyclic to constant load. Crack arrest occurred either within 100 h or thecracks were further growing with a very small rate of less than 0.6 mm/a. The observedbehaviour is exemplary shown in Figure 58 for tests with high stress intensity factors of 61 to73 MPa-m1/2.

250

200

150

Aoo

<lV

50

0

SA 533 B Cl. 1 (0.018 wt. % S, 0.03 wt. % Al)

DO = 8 ppm ^ ^ ^ ^ T = 288 °c

T = 288 °CK = 61 MPa-m12

200 400 600 800

Time [h]1000

Figure 58: Effect of temperature on the SCC crack growth in the high-sulphur alloy B(0.018 wt.% S, 0.03 % Al).

71

The RPV steel with the low sulphur and aluminium content (alloy A) showed the same,reproducible SCC crack growth behaviour at 288 °C as the high sulphur alloy B, whereas itsSCC susceptibility was even smaller under many conditions. At intermediate temperatures of200 and 250 °C this material revealed a surprising and totally different behaviour (Figure 59).At these temperatures, the low-sulphur alloy A showed sustained, stationary SCC growthwith rates up to 40 mm/a at relatively high stress intensity factors of 60 to 81 MPam1/2. At288 °C, only very small CGR of less than 0.6 mm/a were observed in this material even atmuch higher stress intensities and more aggressive environmental conditions.

Ao8

600

500

400

300

200

100

0

20 MnMoNi 5 5 (0.004 wt.% S, 0.013 wt/

I - 200 °CKi = 69,8 - 80,6 MPa-m1'

DO = 8 ppm

T = 250 °CK, = 71.3-80.7 MPa-m1/2

T = 288 °C, = 74 MPa-m1/£

0 200 400 600 800

Time [h]1000

Figure 59: Effect of temperature on the SCC crack growth in the low-sulphur alloy A(0.004 wt.% S, 0.013 wt.% Al). Sustained, stationary SCC was observed atintermediate temperatures of 200 and 250 °C.

At 200 and 250 °C, the observed SCC CGR were not dependent on the stress intensity factor,suggesting a plateau-behaviour under these conditions. The same SCC cracking behaviourhas also been observed in a preliminary test in high-purity water (< 1 ppb SO42") at 200 °Cwith a DO of 8 ppm and 400 ppb. The SCC crack growth could be stopped by a furtherreduction of the ECP below the critical ECPcrit. At 150 °C, the low-sulphur alloy revealedagain the same behaviour as at 288 °C.

The different behaviour of the low- and high-sulphur alloy is exemplary shown in Figure 60for two specimens, which have been simultaneously tested at 200 °C. The SCC CGR in thelow-sulphur alloy A at 200 ° and 250 °C was 50 times higher than in the high-sulphur alloyB. This behaviour cannot be rationalized by the current mechanistic understanding of EAC inLAS in high-temperature water. Based on existing models, the SCC crack growthsusceptibility/rates in the low-sulphur steel should be lower or at maximum similar as in thehigh-sulphur alloy. That result clearly demonstrates, that the steel sulphur content or themorphology, size and spatial distribution of the MnS-inclusions are not the sole material

72

parameters strongly affecting the SCC growth behaviour of LAS. The higher SCC crackgrowth susceptibility of the low-sulphur alloy A might be related to the occurrence of DSA inthis temperature range. The low-sulphur material A has a low aluminium content and mighttherefore show a distinct susceptibility to DSA. In fact this material showed the typicalfeatures of DSA in mechanical tensile tests. But all the other materials also revealed DSA,whereas to a smaller extent. The possible effect of DSA is further discussed in Section 4.5.

2.5

2.0 -

EE,

AooCO

CO<V

1.5

1.0

0.5

0.0

'. T = 200 °C, DO =

-

1.2 E-9 m/sj

-• 8 ppm, Constant Load ^j&~ '

j S ^ 4.9E-10 m/s

* < '•

i f 20 MnMoNi 5 5 (0.004 wt.% S) "

/ K, = 69.8 - 80.6 MPa-m1/J

/

SA 533 B Cl. 1 (0.018wt.% S) '.K, = 64.8 MPa-m1'2

2.4E-11 m/s n < „/>.

800 1000 1200 1400

0.5

- 0.4

- 0.3

A0.2 8

C6<

- 0.1

0.01600

Time [h]

Figure 60: Comparison of the SCC crack growth in the low- and high-sulphur alloy A and Bat 200 °C. The SCC CGR in the low-sulphur alloy A was up to a 50 times higherthan in the high-sulphur steel B (note the different scale of the Y-axis).

73

4.4.3 Ripple loading

The effect of small, low-frequency load fluctuations at very high load ratios (R > 0.95) nearto the fatigue threshold AKth ("ripple loading") has been investigated with a steel with a low(alloy C) and a high sulphur content (alloy B). The tests were conducted in oxygenated high-purity (< 1 ppb SO4

2\ K < 0.06 uS/cm), high-temperature water at 288 °C of 8 ppm DO. Theload fluctuations corresponded to a load ratio of 0.957 and 0.97 and a stress intensity factoramplitude AK of 3.3 and 2.0 MPam1/2. Constant load amplitude with load control and anasymmetrical saw tooth waveform was applied. The loading frequency was stepwisedecreased according to Table7 for material B.

Phase

AtR [h]

v[Hz]

R [ - ]

AK [MPam1/2]

1

0.01

1.4E-2

0.957

3.1

2

0.1

2.5E-3

0.957

3.2

3

1

2.7E-4

0.957

3.3

4

0.7

3.8E-4

0.970

2.3

5

7

3.9E-5

0.970

2.3

6

10

2.74E-5

0.957

3.3

Table 7: Loading conditions in the ripple load test with material B.

Figure 61 shows the results of the test with the high-sulphur alloy B. Sustained stationarySCC crack growth was temporary observed in the loading frequency range form 10"2 Hz to10"4 Hz (Figure 62). The SCC CGR thereby reached values up to 146 mm/a at stress intensityfactors between 66.7 and 76.3 MPam"2. In some experimental phases temporary crack arrestwas observed, but the crack growth re-initiated again after several fatigue cycles. Thetemporary crack arrest might be the result of crack closure effects. At lower loadingfrequencies around 10"5 Hz, cessation of the SCC crack growth and crack arrest wasobserved. The ripple loading in the frequency region of 10"2 to 10"4 Hz resulted in thismaterial in an acceleration of the SCC crack growth by a factor of up to 150 compared topure static loading conditions at identical Kj-levels. The fracture surface revealed a veryrough, transgranular quasi-cleavage appearance without any indications of striations, which istypical for SCC/SICC in LAS. In contrast to that, the ripple loading in an inert environmentresults in very smooth hair-line cracks.

The low-sulphur alloy C only revealed SCC in the loading frequency range of 10"3 to 10"4 Hz.The resulting SCC was very localized and the corresponding growth rates were slower than inthe high-sulphur alloy B reaching up to 9 mm/year. Even in the case of the low-sulphurmaterial C, the ripple loading resulted in an acceleration of SCC growth of one order ofmagnitude compared to static loading conditions.

In Figure 63, the SCC CGR from ripple loading tests are compared to time-based CGR fromLFCF-tests at a load ratio of 0.8 under otherwise similar testing conditions. At comparabledCODix/dt-rates, the maximum SCC CGR were slightly higher than the corresponding time-based LFCF CGR.

The strong effect of ripple loading on SCC crack growth has to be further verified at higherloading frequencies and lower Kimax- and AK-values.

74

E£A

ooCO

CO

V

CDOcCO>

T3CO

O

6

1.2

1.0

0.8

0.6

0.4

0.2

0.0

-

1 2 3 4

y

5 6

-

0 100 200 300 400 500 600 700 800

Time [h]

Figure 61: SCC crack growth in alloy B during the different test phases of ripple loading (seeTable 7).

E, 0.90

A

CO

V

CD

o£CO

CO

1O

b 0.85oCO

0.80 -

0.75 -

0.70 -

0.65 -

: 2 3

da/dt = 6.65E-10m/s y ^

4

420 440 460 480 500 520

Time [h]

Figure 62: Stationary SCC crack growth in phase 3 of ripple loading (see Table 7).

75

C/)

LU

oo

<]

10-5

10-6

10-7

10- 8

I 10"9

5 10-10

• Ripple Loading• Raising load /vLCF Tests•i LCF-Tets, R = 0.8, AK = 10.4 -12.3 MPa-m-1^

10-7 .-6 •v-5

10° 10° 10" 10

dCODM/dt [mm/s]

-3

LL

Figure 63: Comparison of ripple loading test results with LFCF- and SRL-test results. Theapparent SCC CGR (Aa/ANScc)/AtR from RL-tests are compared to the apparentCF CGR (Aa/ANLccc)/AtR from LFCF-tests with a high load ratio of 0.8 and withthe true SICC CGR Aasicc/AtiE form SRL-tests. The K^x-values were similar inall tests. The CGR from the ripple load test roughly correspond to the CGR fromLFCF-tests at a high load ratio of 0.8 for comparable crack opening displacementrates dCODLL/dt.

76

4.4.4 Periodical partial unloading

The effect of periodical partial unloading on SCC crack growth was studied with the low-sulphur material A, the medium sulphur material D and the high-sulphur material B. The testswere conducted in oxygenated high-temperature water at 288 °C with a DO of 8 or 0.4 ppmand a SCV-content of 65 ppb or < 1 ppb. Constant load amplitude with load control andasymmetrical trapezoid waveform with a rise time of 1000 s and a load ratio of 0.8 wasapplied. The hold time at maximum constant load was varied between 0 h and 10 h. TheEAC CGR decreased with increasing hold time in all 3 materials. This is exemplary shown inFigure 64 for medium sulphur material. In many cases with long hold periods at maximumconstant load of 10 h, cessation of EAC crack growth and crack arrest was observed.

69.6

70.1

V71.3

Periodical partial unloadingR = 0.8, AtR = 1000 s, variation of AtH

20 MnMoNi 5 5 (0.012 wt.% S)

8 ppm DO, < 1 ppb SO42", 288 °C \ 64 9

Cessation of crak growthand crack arrest —

0.01 0.1 10

Hold time at maximum constant load At. [h]H

Figure 64: Effect of hold time at maximum constant load on SCC crack growth in themedium sulphur steel in tests with periodical partial unloading for two differentK^-levels .

It has been postulated, that the EAC CGR at very long hold periods at constant maximumload should become independent of the hold time and approach the values of pure constantload testing. The EAC CGR at long hold periods might therefore be a good and conservativeestimate for the SCC growth rates under pure static load. Periodical partial unloading (or"gentle cyclic loading") has been proposed for corrosion systems with problems to sustainSCC growth under static loading. The PPU may help to overcome crack pinning and crackarrest problems without any relevant cyclic fatigue effects.

In Figure 65, the results of the tests with periodical partial unloading are compared to theBWR VIP 60 SCC disposition lines. For long hold times > = 5 to 10 h, the SCC CGR weregenerally falling below the BWR VIP 60 SCC disposition line 2 ("low-sulphur line")supporting the basic idea that this reference curve is a conservative upper bound for SCCcrack growth under static loading and BWR conditions.

77

CO

10-8

1 10"9

o10'10

10-11

1 ' • I • • • • I • • • • I • • • • I • ' • • I • • • • I

O 8 ppm DO, 65 ppb SO42', 0.018 wt.% S

• 8 ppm DO, 65 ppb SO42", 0.004 wt.% S

• 8 ppm DO, < 1 ppb SO42\ 0.012 wt.% S

• 400 ppb DO, < 1 ppb SO42", 0.012 wt.% S

^ — B W R VIP60DL2BWRVIP60DL1

1 M

10 20 30 40 50 60 70 80 90 1001/2

Stress intensity factor K [MPa-m ]

Figure 65: Comparison of SCC CGR from tests with periodical partial unloading to theBWR VIP 60 SCC disposition lines.

4.4.5 Effect of microstructure

The effect of microstructure on SCC crack growth has been investigated with the high-sulphur alloy B at 288 °C. The constant load tests were conducted in oxygenated high-temperature water with 8 ppm DO and 65 ppb SO42". An actively growing EAC crack wasgenerated by a slow, monotonically rising load (see Section 4.1.5). The constant load phasestarted with an initial Kj-value of 60 - 64 MPam1/2. Different microstructures were generatedby a variation of the thermal heat treatment (see Table5). Besides the bainitc "standard"microstructure (Q+T), which is characteristic for the RPV base metal, a martensitic (Q) and aferritic-pearlitic "equilibrium" microstructure (N) were produced by austenizing/waterquenching and austenizing/ slow furnace cooling. The MnS-morphology was not affected bythe applied heat treatments. The different behaviour can therefore be related to the differentmicrostructure, yield strength level and plastic deformation behaviour.

The specimen with the ferritic-pearlitic microstructure ("equilibrium") revealed the samebehaviour as the specimen with a bainitic microstructure. Cessation of SCC crack growthafter switching from slow rising to constant load and crack arrest within 10 to 100 h wasobserved. The apparent SCC CGR Aascc/AtcL for a constant load period of 1000 h were verysmall (< 0.6 mm/a). On the other hand, the specimen with the martensitic microstructure(with an excessive hardness of 466 VH10) revealed stationary SCC crack growth with a veryfast rate of 6000 mm/a. The SCC crack growth rate did not depend on the stress intensityfactor in the range from 64 to 84 MPam1/2, indicating a plateau-behaviour. In Figure 66, theSCC CGR for the three different microstructures are compared to each other and to the BWRVIP 60 SCC disposition lines.

78

CO

E10-7

oCO

CO" D" DCCO

d

10.-8

10-9

10-10

o 10,-11

CO<

: — BWR VIP 60 SCC DL 2r BWR VIP 60 SCC DL 1• • Bainitic, Q + T: • Ferritc-pearlitic, N

* Martensitic, Q

: Constant load; 288 ° C, 8 ppm DO, 65 ppb SO/_ SA 533 B Cl. 1 (0.018 wt.% S)

r

10 20 30 40 50 60 70 80 90 1001 /9

Stress Intensity Factor K [MPa-m ]

Figure 66: Comparison of SCC CGR of alloy B with different microstructures to the BWRVIP 60 SCC disposition lines.

The SCC CGR in this stress intensity range in the specimen with martensitic microstructurewas 3 to 4 orders of magnitude higher than the values typically observed in different LASwith bainitic microstructures. In the case of a martensitic microstructure and excessivehardness or high yield strength, the BWR VIP SCC disposition lines can be significantlyexceeded by 2 - 3 orders of magnitude. The hardness of weld HAZ in LAS pressure-boundarycomponents is generally limited to values < 350 VH by suitable post-weld heat treatments(PWHT, stress relieving slightly below the annealing temperature). Nevertheless, thereremains some concern for TG/IG hydrogen-assisted SCC in the peak hardness region of weldHAZ and especially in the case of repair welding without PWHT [8, 54, 76]. Furthermore,SCC crack growth might by also affected by DSA-effects especially at intermediatetemperatures (150 - 290 °C) and in welds and weld HAZ, especially in the as-weldedcondition. The SCC susceptibility of welds and weld HAZ should therefore be furtherinvestigated.

4.4.6 Comparison with BWR VIP 60 SCC disposition lines and literature data

The results of the PSI standard constant load tests at 288 °C, of the tests at intermediatetemperatures of 200 and 250 °C with the low-sulphur material A, and of tests with rippleloading at 288 °C are compared to the BWR VIP 60 SCC disposition lines in Figure 67.Additionally, recent results of MPA Stuttgart [32] and of a European Round Robin test [34]are also plotted to Figure 67. In the case of the constant load tests, the very conservativeapparent SCC CGR AaEAc/AtcL (see Section 4.4.1) were used. The true SCC growth rateswere used for the constant load tests at intermediate temperatures with the low-sulphur alloyand for the ripple loading.

79

w1 10

" 10

-6

CO

$

-7

ooCO

10-8

10-9

10-10

o

s

O O2 = 200 - 600 ppb, T = 288 °CA O2 = 200 - 600 ppb, T = 240 °C• O2 = 8000 ppb, T = 288 °CA O2 = 8000 ppb, T = 240 °C

— BWR VIP DL 2BWRVIPDL1

• Ripple loading, SA 533 BCI.1<i Ripple loading, SA 508 Cl. 2* T = 250 °C, DSAif T = 200 °C, DSA

OCO

22

CO

c

1 0

3

LUO

LCO<10 20 30 40 50 60 70 80 90 100

1/?

Stress Intensity Factor K [MPa-m ]

Figure 67: Comparison of SCC CGR form standard constant load and ripple load tests withthe BWR VIP 60 SCC disposition lines for a wide range of loading,environmental (150 - 288 °C, 0.2 - 8 ppm DO, < lppb - 165 ppb SO4

2") andmaterial (0.004 - 0.018 wt.% S) parameters. In the case of constant load tests, thevery conservative apparent, maximum SCC CGR AaEACtotmax/AtcL are plotted.The true SCC CGR were used for the constant load tests at intermediatetemperatures with the low-sulphur alloy and for the ripple loading.

High SCC CGR in the range of several m/a could not be confirmed within this parameterstudy for properly manufactured and heat-treated low-alloy pressure-boundary componentsteels as long as water-chemistry conditions were maintained within current BWR operatingpractice to some extent. Many of the older tests revealing high CGR significantly above theBWR VIP 60 SCC disposition line have been performed under a combination of both grossyielding of the remaining specimen ligament, which is not transferable to the thick-walledRPV structure, and aggressive water-chemistry conditions, which are not representative ofcurrent BWR operating practice [8, 68, 69].

The adequacy and conservative character of the BWR VIP 60 SCC disposition lines for staticloading and small scale yielding conditions at temperatures around 288 °C has beenconfirmed by many independent tests for different RPV base materials (< 0.018 wt.% S) inthe quenched and tempered (and stress relieved) state as long as water-chemistry ismaintained within current BWR operating practice (EPRI guidelines).

80

At 288 °C, sustained and/or stationary SCC crack growth with rates around or above theBWR VIP 60 SCC disposition lines have been observed for

- low-frequency ripple loading at very high load ratios (R > 0.95) (Figure 67).

- periodical partial unloading with short hold periods ( < ~ 5 h) at maximum constantload (Figure 64 and 65).

- SO42- + Cl- > « 500 ppb (e.g. > EPRI action level 3) [30, 67].

- Ki -» Kn wt.% St , R p t , DSAT(N/AlT) -» K n i (Figure 54).

P —> plastic limit load or onet section —> Rp (Section 4.4.1.1).

Excessive hardness (> = 350 VH), martensitic microstructure (Figure 66).

Otherwise, cessation of SCC crack growth and crack arrest or very slow SCC CGR of somefew hundred |im/a were observed.

At intermediate temperatures of 200 - 250 °C, fast and stationary SCC crack growth wasobserved in a RPV steel with a low sulphur and aluminium content, which revealed distinctDSA effects in tensile tests in this temperature range. The simultaneously tested steel with ahigh sulphur and aluminium showed the same behaviour as at 288 °C with cessation of SCCcrack growth and very low slow CGR of less than 0.6 mm/a. These results further confirm,that the steel sulphur content is not the sole material parameter strongly affecting SCC andindicate that DSA might be as important as the steel sulphur content, at least at intermediatetemperatures and in materials which are susceptible to DSA. The SCC CGR in the low-sulphur material were around and above the BWR VIP 60 SCC disposition lines at 200 and250 °C (Figure 67). These results should be verified by further tests at lower Ki-levels.

4.4.7 Summary of SCC

It is concluded that there is no susceptibility to sustained SCC crack growth at temperaturesaround 288 °C under purely static loading, as long as small-scale-yielding conditions prevailat the crack-tip and the water-chemistry is maintained within current BWR/NWC operationalpractice (EPRI water-chemistry guidelines). However, sustained, fast SCC (with respect tooperational time scales) cannot be excluded for faulted water-chemistry conditions (> EPRIAction Level 3) and/or for highly stressed specimens, either loaded near to Kn or with a highdegree of plasticity in the remaining ligament. The conservative character of the "BWR VIP60 Disposition Lines 1 and 2" for SCC crack growth in LAS has been confirmed by thisstudy for 288 °C and RPV base material (< 0.018 wt.% S).

Preliminary results indicate, that these disposition lines may be significantly or slightlyexceeded (even in steels with a low sulphur content) in the case of small load fluctuations athigh load ratios (ripple loading) or at intermediate temperatures (200 - 250 °C) in RPVmaterials, which show a distinct susceptibility to DSA. In the case of periodical partialunloading, the SCC CGR fall below the BWR VIP 60 SCC disposition line 2 for long holdperiods at maximum (constant) load > 5 to 10 h. Furthermore, crack arrest is often observedfor long hold periods.

Based on these preliminary results, the BWR VIP 60 SCC disposition lines seem to be stillconservative and adequate for the RPV (unaffected base metal) for transient-free steady-stateBWR power operation. There is some concern for

- the RPV feedwater nozzle and feedwater piping system with lower operatingtemperatures and the temporary occurrence of small load fluctuations (—» DSA, rippleloading).

81

- Weld (DSA) and weld HAZ (-» DSA, TG/IG hydrogen-assisted SCC) materials.

Materials in as-welded condition without PWHT (—» DSA).

LAS RPV and piping materials with a low Al-content and a high concentration ofinterstitial N and C (-> DSA).

The BWR VIP SCC disposition lines should therefore be validated by further tests

- at intermediate temperatures.

- with weld filler and weld HAZ materials.

- with ripple loading and periodical partial unloading.

The results should be verified at lower Ki- and AK-levels, corresponding to crack depthsslightly above the resolution limit of the methods for periodic in-service inspection. Thewater-chemistry conditions (ECP, SO4

2\ ...) should also be further adjusted to current BWRoperation conditions.

82

4.5 Dynamic strain ageing

4.5.1 Background

There is now increasing experimental evidence that DSA may affect the EAC crackingbehaviour of LAS in high-temperature water. DSA may be very influential in determiningboth EAC cracking strain-rate sensivity and its temperature dependence. It provides analternate possible explanation for strain-rate thresholds and more significantly, it may offer avery attractive explanation for the EAC cracking susceptibility peak at intermediatetemperatures and rationalize the scattered data in this temperature range. Furthermore it mayprovide a possible explanation for the observed quasi-cleavage fracture surface observed inEAC. [8]

DSA may be observed in susceptible LAS during plastic straining at sufficiently slow strain-rates (< ~ 10~2 s"1) in the temperature range from 100 to 350 °C. These temperature/strain-rateconditions are also characteristic for plant transients such as start-up/shut-down, hot-stand-byor operating conditions, where thermal stratification is typically observed. SICC and LFCFare most likely to occur under these conditions. The operating temperatures of BWR, and inparticular of the feedwater piping system, are close to the peak effects of DSA. In the past,only few EAC tests were performed at intermediate temperatures and lower strain-rates/frequencies, where significant DSA-effects can be expected in susceptible LAS. [8]

DSA in LAS is associated with diffusion of interstitial species such as C and N atoms to thecore region of dislocations and their immobilization. The DSA effects increase withincreasing concentration of free, interstitial N and C and are most pronounced if the diffusionrate of C/N and the dislocation velocity are similar. Parameters, which affect theconcentration of free C and N and their diffusion rate have therefore a strong effect on theDSA-behaviour of LAS. The free C and N content are not specified in the relevant nuclearregulations (KTA, ASME BVP) for low-alloy primary pressure-boundary component steels.They can therefore vary between 0 ppm and their solubility limit (or even higher in thesupersaturated state). The free N and C content may relevantly depend on the steel making(killing) and welding process, on the heat treatments applied and on the exact chemicalcomposition. Differences in free N and C content of otherwise identical or similar LAS maybe one important further reason for the relevant scatter of EAC CGR data, and especially forthe observed different trends in temperature dependence of EAC. Because of possible DSA-effects, a higher EAC susceptibility of weld filler and weld HAZ materials compared to theRPV base material cannot be excluded in the DSA temperature/strain-rate range. [8, 48]

It is stressed that EAC in LAS also occurs under temperature/strain-rate combinations or inmaterials where no or only minor DSA effects can be expected to be present. Therefore DSAis not a pre-requisite for the occurrence of EAC in these steels and it is best regarded as afurther additional contribution to the EAC cracking process together with anodic dissolutionand hydrogen-assisted cracking mechanism.

Because of its relevance for BWR operation, the possible effects of DSA on EAC in LAS inhigh-temperature water will therefore be discussed in more detail in the following Sections.The phenomenological features and the metallurgical mechanism of DSA in LAS and thepossible interaction with EAC are briefly discussed in Section 4.5.2. The DSA behaviour ofselected LAS of the RIKORR programme was characterized by mechanical tensile tests atdifferent temperatures /strain-rates (Section 4.5.3) and by internal friction measurements(Section 4.5.4). In Section 4.5.5, the correlation between the observed DSA and EACresponse of the investigated materials is briefly discussed.

83

4.5.2 Phenomenological features and metallurgical mechanism of DSA

Strain ageing occurs in alloys (typically dilute alloys) containing solutes that segregatestrongly to dislocations resulting in strong elastic interactions between solutes and the stress-strain fields of dislocations and in strong dislocation pinning. Static strain ageing is a processwhere ageing takes place after pre-straining and results in a return of Liiders strain. DSA is aprocess where ageing is sufficiently rapid to occur during straining and it produces a varietyof discontinuous, inhomogeneous deformations which are characterized by terms such asPortevin-le Chatelier effect, serrated yielding, jerky or serrated flow, blue brittleness, etc. Theserrations observed in stress-strain curves are generally classified to A...E types (fromregular to more irregular) depending on the amount of strain and strain-rate. In LAS DSAoccurs at temperatures within 100 - 350 °C, where the stress-strain curves often showserrations, being most marked at 250 °C, depending, however, on strain-rate. Yield drops ofeven 30% due to large amplitude serrations can be obtained in the stress-strain curve (Figure68). The effect of strain-rate on DSA temperature range is related to the diffusing atoms tokeep pace with the moving dislocations during deformation allowing to form atmospheresaround dislocations generated throughout the whole stress-strain curve. Other materials inaddition to LAS relevant to LWR known to cause discontinuities in deformation related toDSA include, e.g., Ni-base alloys including superalloys (due to C and in H-chargedcondition), austenitic stainless steels (due to both interstitial and substitutional alloyingelements) as well as Zr-alloys (due to H, C, N and O). [77 - 81]

600 -

05Q_

=^ 550 -

o

500 -CD

4501.0 1.5 2.0 2.5 3.0

Strain e.3.5 4.0 4.5 5.0

"tech

Figure 68: Detail of stress-strain curves for alloy A at 200 °C in air obtained in tensile testsat different strain-rates. At very slow strain-rates irregular serrations appearedwith yield drops of up to 15 MPa.

84

Phenomenological features of DSA in LAS:

DSA has some detrimental effects on LAS. It results in a peak in ultimate tensile strength(Rm), hardness, and strain hardening rate in the DSA temperature range, a minimum ofductility (elongation to fracture A and reduction of area Z) and results in negative strain-ratesensitivity (Figure 69). Yield strength is affected by static strain ageing rather than DSAresulting in plateau or a small peak in yield strength in the temperature range of DSA (Figure70). The temperature of the peak effect decreases with decreasing strain-rate (Figure 71).

IPa]

reng

t

S8

I

<2

650

640

630

620

610

600

deuo),/dt = 1E-5s"1 *dl/dt = 0.03 mm/min /

J

\ a'

.......50 100 150 200 250 300 350

Temperature [°C]

NCOCD

o

50 100 150 200 250 300 350

Temperature [°C]

CO

cc

I5

Strain rate detech/dt

ENCOCDCO

ion

oflu

ct

Gt

69

68

67

66

65

64

63

62

81

/ T = 200 °C

• /

/

d

••

•

•

10" 6 10 5 103 102

Strain rate detech/dt [s"1]

Figure 69: Ultimate tensile strength Rm (a) and reduction of area Z (b) as a function of tensiletest temperature for alloy A with a strain-rate of 1-10"5 s"1. Rm (c) and Z (d) as afunction of strain-rate at 200 °C for alloy A. A maximum/minimum in tensilestrength/reduction of area and a negative strain-rate sensivity of these propertieswas observed.

The hardening effect of DSA may be better characterized by the tensile properties, rather thanonly based on the observation of serrations in the stress-strain curve alone. Ultimate tensilestrength highlights the hardening effect best and distinguishes the effects of strain-rate mostclearly. Furthermore, DSA results in an increase in the ductile-to-brittle transitiontemperature following plastic deformation in the DSA temperature range as well as alowering of the ductile fracture resistance (decrease of tearing modulus) at temperatureswithin the DSA temperature range [82 - 85]. A reduction of low-cycle fatigue resistance ofLAS in air and water environment is observed as the loading strain-rate is decreased.Microscopically, DSA results in an inhomogeneous localisation of deformation, an increaseof dislocation density and increase in planar deformation. Strain localization and shear bandson the surface of the specimens appear as a result of DSA and intensified acoustic emission isoften related to dislocation multiplication events.

85

550-

COQ _

o500-

CD

CO

450-

400-

25 °C

288 °C

SA 533 B Cl. 1Orientation of gauge length: T

0.0 0.5 1.0 1.5Strain

2.0r o /

2.5 3.0

tech

Figure 70: Detail of stress-strain curve near to the yield point from tensile tests in air withalloy B with a strain-rate of 10~3 s"1 at 25 °C and 288 °C. At 25 °C, the materialrevealed a yield plateau (and Liiders strain), which is characteristic for staticstrain ageing. At 288 °C, no yield plateau, but serrations with yield drops of up to20 MPa were observed, indicating the occurrence of DSA.

CO

c

CO

'encCD

660

640 -

DC 620 -

600 -

580 -

56050 100 150 200 250 300 350 400

Temperature [°C]

Figure 71: Ultimate tensile strength Rm as a function of temperature in alloy B for twodifferent strain-rates. The maximum in Rm is shifted to higher temperatures withincreasing strain-rate.

86

Metallurgical mechanism of PSA:

All effects of static and dynamic strain ageing are explained in terms of either segregation ofsolute atoms to dislocations to form condensed Cottrell atmospheres or precipitates, or Snoek(stress-induced) ordering of solute atoms in the dislocation core structure. DSA in carbon andLAS is an elastic interaction between the strain fields of solute interstitials and movingdislocations with edge character. In ferrite interstitial C and N are producing a nearlyidentical lattice misfit strain, which they can reduce by moving to the region of maximumdilatation in the dislocation core (Cottrell atmosphere), resulting in an overall reduction in thetotal strain energy. Consequently dislocations can be locked in position by strings of C or Natoms along the dislocations, thus substantially raising the stress, which would be necessaryto cause dislocation movement. Because of their segregation to the dislocation core, verysmall bulk C and N contents are sufficient to form condensed atmospheres on dislocations inferrite and cause yield point phenomena. For example, with a typical dislocation density of108/cm2 in Fe a C concentration of 10~6 wt-% is sufficient to provide one interstitial C atomper atomic plane along all the dislocation lines present.

If a steel susceptible to DSA is plastically deformed in the DSA temperature range anunusually high dislocation density is observed coupled with immobilization of many of thedislocations by N and/or C atmospheres. In the DSA temperature range (typically 100 -350 °C), interstitials can diffuse during deformation and can form atmospheres arounddislocations generated during deformation. The dislocations generated are quickly pinned sothat others must be generated to allow deformation to continue. As a result, the dislocationdensity at a given strain is higher than when the straining is done outside the DSA range.Large yield drops ("serrations") in strain-controlled tensile tests are related, thus, togeneration of a large number of new dislocations. The repeated breakaway and repining ofmobile dislocations by the interstitials was the original explanation for the serrations in thestress-strain curve. An alternate explanation is that once C and N atmospheres are formed, thedislocations remain locked, and the serrated yield phenomena arise from the generation andmovement of newly formed dislocations, which are then suddenly pinned.

DSA temperature and strain-rate thresholds can be rationalized as follows: Below 100 °C thediffusivity of C and N is too low to keep pace with moving dislocations. Between 100 and350 °C, the C and N can segregate to the core of moving dislocation resulting in theirimmobilization. At higher temperatures (> 400 °C) thermal energy kT can overcome theelastic interaction or binding energy between the solute and dislocation and dislocations canescape from their atmospheres as a result of thermal activation. The DSA effect is mostpronounced, if the diffusion rate (which is dependent on the temperature) of the freeinterstitials and the dislocation velocity (which is dependent on loading rate/ strain-rate/frequency) are similar. This is schematically shown in Figure 72. At low strain-rate orfrequency, diffusion of the interstitials is much faster than the dislocations. At high strain-rates dislocations are too fast to be pinned by the interstitials. So in both cases there are nosignificant DSA effects. With increasing strain rate, higher temperatures are necessary thatthe interstitials can keep pace with the moving dislocations, but the increasing disordered(random) thermal motion of the interstitials also counteracts the segregation process resultingin a DSA-peak at intermediate temperatures. The maximum is a function of strain-rate (ortemperature). The maximum DSA peak is shifted to higher temperatures with increasingstrain-rate. On the other hand, the maximum DSA peak is shifted to higher strain-rates withincreasing temperatures.

87

o

Q

o

J>to

*

• o DiH

fcj ' j , t i g . -l:£nS

| CompletePassivation >

i ^^^

/Corrosion Fatigue \

/ o r \/ Stress Corrosion \

/ Cracking \/ by Slip/Dissolution ^

fk ' Mechanical

\ t Fracture

c oO *3

uo

SlowLow Frequency

Log Strain RateLog Cyclic Frequency

FastHigh Frequency

Figure 72: Correlation between frequency/strain-rate and DSA or EAC response.

In LAS, activation energy for DSA obtained normally from onset temperature of serratedflow with the change of strain-rate is equal to that of N/C diffusion in ferrite. Based on thetemperature of disappearance of serrations it should correspond to that of a sum of activationenergy for diffusion and binding energy of an interstitial atom to the dislocation core [78].However, the observations for LAS for this explanation to be always true show often too highactivation energy values. The temperature for occurrence of DSA increases with increasingstrain-rate and the peak hardening stress due to DSA decreases linearly with logarithmicincrease of strain-rate. The measured activation energy for the onset of serrations in thestress-strain curve is not sensitive to the microstructure or composition of the steel. Theexemption being the Mn content which seems to influence the diffusion process of N and Cby forming Mn-N and Mn-C pairs increasing thus the activation energy for the diffusionprocess [93]. Large differences in the activation energy for the disappearance of DSAreported suggest that it is also a function of steel composition. The reported values ofactivation energy for the onset of DSA are in the range of 75- 85 kJ/mol being comparable tothe common values for N and C diffusion in oc-iron, 65- 85 kJ/mol (the activation energy fordiffusion is higher for C than for N) [78].

Deformation induced vacancies ("vacancy model"), which are the diffusion vehicles for thesubstitutional solute atoms, can play also an important role in the DSA process especially inFCC-alloys. However, strain-induced vacancies are not expected to accelerate diffusion ofinterstitial atoms. Therefore, the importance of prior strain in inducing serrated flow has beenexplained as follows: vacancies form interstitial-vacancy pairs, which order in the stress-strain fields of dislocations and if substitutional-interstitial complexes (Mn-N, Mn-C) areinvolved, vacancies are necessary to increase the mobility of substitutional atoms. Recently,deformation-induced generation of vacancies and their clustering has been considered to bepromoted by hydrogen and to play a primary role in hydrogen trapping and hydrogenembrittlement susceptibility in many FCC- and BCC-metals and alloys. A substantial densityof vacancies can be expected in plastic strain of steels under presence of hydrogen, which isstabilizing the vacancies. Hydrogen is lowering the formation energy of vacancies by theamount of binding energies of trapped hydrogen atoms. Formation of vacancy-solute

88

complexes (such as H, C and N in steels) elevates the annihilation temperature of vacanciesup to 200 °C and higher [89, 90]. Due to the low migration energy of vacancies in Fe(53 kJ/mol) marked diffusion distances in the crack-tip regions are attainable in short time.Extra vacancies are introduced in the crack-tip plastic zone region in addition due tohydrogen in EAC by oxidation reactions producing vacancies at the oxide metal interface,which are injected into the base metal in the crack-tip region, as well. The extra vacanciesproduced by hydrogen uptake and oxidation due to their agglomeration to clusters and smallvoids may be more important in causing the brittle-like EAC fracture than hydrogen itself aswas already proposed in Reference [91]. It has also been observed that if steels susceptible toHE are tested in conditions, which promote DSA, the ductility decreases remarkably by thesimultaneous effects of DSA and HE [92].

Metallurgical parameters affecting DSA in LAS;

DSA in LAS is associated with diffusion of interstitial species such as C and N atoms to thecore region of dislocations and their immobilization. The DSA effects increase withincreasing concentration of free, interstitial N and C and are most pronounced if the diffusionrate of C/N and the dislocation velocity are similar. Parameters, which affect theconcentration of free C and N and their diffusion rate have therefore a strong effect on theDSA-behaviour of LAS.

In ferrite N and C have similar diffusion coefficients and they are producing a nearlyidentical lattice misfit strain, thus, these two elements are expected to produce nearly similarDSA effects in LAS. Therefore, in general, the effects produced by C and N can beconsidered as additive. At room temperature, the residual solubility of N in ferrite is about100 times greater than that of C. N solubility in steel is higher at all temperatures than that ofC. It is generally assumed that N, rather than C, is mainly responsible for DSA. However, athigher temperatures in the DSA range the increasing solubility of carbon may cause DSAeven in the absence of N.

The concentration of free interstitial C and N strongly depend on the steel making (killing) orwelding process, the thermal history or heat-treatment (annealing-, PWHT- or stress relievingtemperature) and on the chemical composition (Ctot, Ntot, Al, V, Ti, Cr, Mo, O, Mn ...) of thesteel. Alloying elements, which have a strong affinity to N or C, such as Al, Cr and Mo, andform nitrides or carbides may result in a reduction of the residual concentration of free C andN. Parameters, which would affect the interstitial diffusion in dilute alloys might also affectthe DSA-behaviour of LAS. The measured activation energy for the onset of serrations in thestress-strain curve is not sensitive to the microstructure or composition of steel. Theexemption being the Mn-content which seems to influence the diffusion process of N and Cby forming Mn-N and Mn-C pairs increasing thus the activation energy for the diffusionprocess [135]. Large differences in the activation energy for the disappearance of DSAreported suggest that it is also a function of steel composition.

Silicon-killed vs. Aluminium-killed low-alloy steels:

Si-killed, air poured carbon or LAS show a much more pronounced DSA effect than Al-killed vacuum ladle degassed steels, since in Al-killed steels, aluminium nitrides are formedand the remaining free interstitial nitrogen content is very low. The Al/N-ratio is therefore animportant parameter governing the free nitrogen content and the resulting DSA response.

Weld material:

During the welding process uptake of small amounts of O and N can occur. Weld metal oftenhave a very little Al- but high O-contents. Because of its very high affinity to O, Al fist

89

combines to O (aluminium oxide precipitates) in welds, which reduces the beneficial effect ofthe Al. During fast cooling after welding, incomplete precipitation of carbides and nitridesoccurs resulting in an increased or even supersaturated concentration of free C and N in theas-welded condition. The precipitation of carbides during PWHT further reduces theconcentration of free C and N. The PWHT-temperature is therefore detrimental. The free Cand N decrease with increasing annealing temperature because of precipitation of carbidesand nitrides.

Welds might reveal increased concentrations of interstitial N and therefore be moresusceptible to EAC in the DSA-temperature /strain-rate range than generally thought so far.LFCF crack growth tests in high-temperature water with weld materials have been performedunder temperature-strain-rate conditions (> 250 °C, > 10"3 Hz) where DSA is absent or onlymoderately active. The better resistance of the weld material under these conditions has beenmainly attributed to the different morphology and chemical composition of MnS-inclusionsin the welds (very small spherical inclusions) [8].

Weld HAZ material:

The free carbon content in LAS can strongly vary depending on the heat treatment and onchemical composition (Ctot, Cr, Mo, V, ...). Alloying elements, which have a strong tendencyto form carbides as Cr or Mo may also affect the residual interstitial carbon content. In theheat-affected zone near to the fusion line, with high peak temperatures slightly below themelting point (partial) dissolution of carbides and (incomplete) re-precipitation duringsubsequent cooling occurs, which can result in a supersaturated state. During the followingPWHT (stress relieving) precipitation of further carbides occurs. The interstitial carbon isstrongly dependent on the stability of the carbides (chemical composition), on the peaktemperature, heating and cooling rates and on the PWHT-treatments applied. The most severeDSA-effects are expected for materials in the as-welded condition without PWHT (e.g. repairwelding).

Possible interaction between DSA and EAC in LAS:

For the case of LAS in high-temperature water, the following potential cracking mechanismshave been mainly discussed in literature [8, 9, 71 - 74]:

• Film rupture/anodic dissolution (FRAD)

• Hydrogen-assisted EAC (HAEAC)

The controlling factors in EAC of LAS are phenomenologically well known. EAC growthfrom incipient crack is governed by the crack-tip strain-rate and the activity of chloride- andsulphur-anions and the pH in the crack-tip environment and by the steel-sulphurcontent/morphology of MnS-inclusions. The direct experimental evidence for a specificmicroscopic crack extension process is still very weak. The exact crack growth mechanism istherefore still under discussion. None of the proposed mechanism can satisfactorily explainall experimentally observed cracking aspects.

Under crack-tip conditions (ECP, pH, anion content) relevant for LAS under LWRconditions, HAEAC and FRAD mechanism are both thermodynamically and kineticallyviable and may therefore be simultaneously active [8]. Both mechanisms may be dependenton oxide film rupture rates, repassivation rates and liquid diffusion rates, since these factorsaffect the charge transfer per time in the FRAD and the hydrogen ad-atom coverage andhydrogen evolution rate in HAEAC models. Furthermore, since both mechanism may becontrolled by the same rate limiting steps (for example oxide film rupture rate, repassivation

90

kinetics, ...), it is very difficult to experimentally differentiate between these two mechanism.[8].

Both mechanisms are able to explain the experimentally observed dominant effect of (crack-tip) strain-rate and of the MnS-inclusions on the EAC cracking behaviour. The HAEACmodel may better explain some special fractographic features but in contrast to the FRADmodel no quantitative predictive formulation exists. In contrast to intergranular SCC, forexample in sensitized stainless steels, there is no simple reason for the directed, anisotropicdissolution behaviour of LAS assumed in the FRAD models. [8].

To date there is a lack of direct experimental evidence for distinct hydrogen effects in low-alloy primary pressure-boundary component steels under LWR operating conditions. Attypical LWR operating temperatures, the diffusivity of hydrogen is very high and because ofthermal activation trapping (at carbides, grain boundaries, dislocations, ...) is not enoughefficient that decohesion might be active. The most striking argument against relevanthydrogen effects in EAC at temperatures above 150 °C is, that HAEAC should be enhancedby increasing the hydrogen fugacity of the environment, by decreasing the corrosion potentialto lower more cathodic potentials and by catalytic surfaces, which is obviously not the case.[8].

The role of DSA in EAC has not been studied much. However, the EAC data stronglysuggest that the susceptibility of LAS to EAC coincides with DSA behaviour, in terms oftemperature and strain-rate (Figure 72) [86-88]. The EAC behaviour of LAS is controlled bythe crack-tip strain-rate and the activity of chloride- and sulphur-anions and the pH in thecrack-tip environment and by the steel-sulphur content/morphology of MnS-inclusions [8]. InEAC, DSA is especially important in dynamic crack-tip plasticity behaviour such as dynamicloading or development of creep strain. DSA affects the yield strength, strain hardeningexponent and creep rate, which are important factors affecting the crack-tip strain and strain-rate. DSA may result in a higher crack-tip strain and strain-rate than outside the DSA-rangeor than in a material, which is not susceptible to DSA. The inhomogeneous localisation ofdeformation, the increase of dislocation density and increase in planar deformation by DSAcan result in a reduction of the local fracture toughness and favour brittle crack extension, butalso in the mechanical rupture of the protective oxide film and therefore crack advance byanodic dissolution/hydrogen embrittlement mechanism. Therefore DSA may synergisticallyinteract with the FRAD- and HAEAC-mechanism to increase EAC cracking susceptibility.EAC in LAS has been observed under temperature/stain rate conditions or in materials, whereno or only minor DSA effects were present. DSA is therefore not a pre-requisite for EAC andbest regarded as an additional contribution to EAC growth.

From a pragmatic view, the observed EAC behaviour of LAS in high-temperature water canbe best rationalized by a combination of these three fundamental cracking mechanism. Atlower temperatures (< 100 °C) and/or high strength levels (Rp > 800 MPa) hydrogen effectsare more pronounced. At high temperatures (> 150 °C) and/or lower yield strength levels (Rp

< 800 MPa) anodic dissolution seems to dominate. Under certain combinations oftemperature and strain-rate DSA may give an additional contribution to the crack growth insusceptible LAS with a high concentration of interstitial N and C. [8]

91

so-ji

80-joo

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a

ca~u

X!csen

60-t

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30

20 i

10

0

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2S0 300 35C 100

Figure 73: Coincidence of susceptibility of LAS to SICC in SSRT-tests (minimum inreduction of area) in high-temperature water with DSA behaviour (maximum oftensile strength) in air in terms of temperature and strain-rate [86].

4.5.3 Characterization of DSA response by tensile tests

The DSA-behaviour of the low-alloy RPV steels A, B, D and of the weld material E (seeTable 2 and 3) was characterized by mechanical tensile tests in air at different strain-ratesbetween 2-10"6 and 10'2 s"1 in the temperature range from 25 to 350 °C. The tests wereperformed in an electro-mechanical tensile machine with cylindrical tensile specimens with agauge length of 50 mm and a diameter of 5 mm. The displacement was measured on thespecimen by an external optical extensometer via a rod system. The tests were conductedwith a constant rate of pull rod stroke (constant extension rate tests, CERT). This resulted inconstant technical strain-rates during the test. It is stressed, that after the start of yielding, thetrue strain-rates continuously decrease with this procedure. The specimens were as close aspossible from the same locations in the plates/forgings to reduce the amount of scatter. Inmost cases, the orientation of the gauge length was in T-direction. It is stressed that thetensile properties are dependent both on the location in the plates/forging and on theorientation and subjected to relevant scatter, sometimes impeding the resolution of weakDSA effects.

The heat treatment and the chemical composition (selected alloying elements, which areknown to have a strong effect on the DSA and EAC behaviour) of the investigated materialsare summarized in Table 8. The tensile properties of the RPV base materials from standardtensile tests were very similar (see Table 3).

The stress-strain curves revealed a complex behaviour over the tested temperature and strain-rate range. All materials revealed indications for DSA, whereas to a different extent. Clearand expected trends were only observed for the ultimate tensile strength Rm and to a lesserextent for the reduction of area Z. The elongation to fracture A5 and the yield strength Rp

showed no clear trends and were subjected to relevant scatter.

92

Material

20 MnMoNi 5 5a

SA533BCI. 1b

SA 508 Cl. 2c

22NiMoCr3 7Biblis C BM

d

S3 NiMo 1Biblis C Weld

e

S{%]

0.004

0.018

0.004

0.007

0.007

c[%]

0.21

0.25

0.21

0.22

0.05

Ntot[ppm]

70

60

110

80

110

Nfree[ppm]

30

<1

2

3

16

0[ppm]

140

20

260

Al[ppm]

130

300

150

180

53

Mn

1*1

1.26

1.42

0.69

0.91

1.19

Heat treatment

910°C-920°C/6h/WQ/640°C-650°C/9.5h/FC

bainitic915°C/12h/860°C/12h/WQ/

635°C/12h/FCbainitic

As received (Q+T)bainitic-ferritic/pearlitic870-905 °C/7h/WQ/

635-655 °C/11h/AQ + PWHTbainitic

PWHT: 540-555 °C/ 59 h/465 °C/590-610°C/21h/465°C/

590-605 °C/11.5 h/AQferritic

Table 8: Overview on investigated materials: Heat treatments and concentrations of selectedalloying and impurity elements, which affect the DSA and EAC behaviour.

At lower temperatures (< 100 - 150 °C) a yield plateau (sometimes with an upper and loweryield stress) was often observed (probably originating from static strain ageing) with a Ludersstrain of up to 2 %. The level of yield plateau and the Luders strain generally decreased withincreasing temperature from 25 to 150 °C and disappeared above 150 °C. The stress straincurves of all materials only revealed serrations at temperatures between 200 and 250 °C (insome few cases also at 288 °C) in the investigated strain-rate range. The serrations typicallyappeared directly after yielding (—» small critical strains). Above 288 °C, neither a yieldplateau nor serrations were generally observed. This behaviour is exemplary shown for alloyD in Figure 74. Figure 75 - 78 summarize the results concerning the occurrence of serrationsin a temperature/strain-rate diagram for the four materials.

The mechanical property variation with temperature at a strain-rate of 10"5 s"1 is exhibited inFigure 79 for the low- and high-sulphur alloy A and B and for the weld material E. Ultimatetensile strength values show a maximum and reduction of area values show a minimumaround 250 °C indicating DSA behaviour. This is also manifested by the negative strain-ratedependence of these values at 200 °C (Figure 80). Typical stress-strain curves of the low-sulphur material B at 200 °C exhibiting serrations at slow strain-rates are presented in Figure68. The serrations start almost immediately on yielding (small critical strain) and exhibit afterirregular beginning a saw-tooth appearance combined with small serrations (type A + B). Thestress drops associated with the serrations vary up to as high as 20 - 30 MPa in the differentmaterials. In general, the magnitude of the stress drops increases with decreasing the strain-rate. Based on the small number of tests it was not possible to determine the activation energyvalues for the onset and disappearance of the serrations in these steels separately. With theappearance of serrations there is a marked increase in the strain hardening rate, ultimatetensile strength and a loss of ductility. Furthermore, the strain-rate sensitivity of the flowstress becomes negative during DSA. The peak stress regions of the stress-strain curves aregenerally not associated with the serrations and the stress peak is shifted to a highertemperature and the increase of ultimate tensile strength Rm is smaller at higher strain-rates(Figure 71).

93

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T = 25 °CDirection of gauge length: T

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600-

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4 6 8 10 12 14 16 18

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T = 320 °CDirection of gauge length: T

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2 4 6 8 10 12 14 16 18

Strain e, . [%]tech L J

Figure 74: Stress-strain curves of alloy D obtained by tensile tests in air with a strain-rate of10"3 s"1 at different temperatures of 25 °C (yield plateau), 288 °C (serrations) and320 °C (no yield plateau and no serrations).

94

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Temperature [°C]

Figure 75: Map of presence of serrations in stress-strain curves of alloy A.

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Temperature [°C]

Figure 76: Map of presence of serrations in stress-strain curves of alloy B.

95

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Figure 77: Map of presence of serrations in stress-strain curves of alloy D.

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Figure 78: Map of presence of serrations in stress-strain curves of weld material E.

96

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640

620

600

580

560

de^/dt^E-Ss1

• 20 MnMoNi 5 5 (A)• SA 533 B Cl. 1 (B)A Weld(E)

0 50 100 150 200 250 300 350

Temperature [°C]

Figure 79: Ultimate tensile strength Rm as a function of temperature at a strain-rate of 10"5 s"1

for the low- and high-sulphur material A and B and for the weld material E. Allmaterials exhibited a maximum in tensile strength at intermediate temperatures.The low-sulphur material A, which revealed stationary fast SCC at 200 and250 °C, revealed the largest shift in tensile strength at intermediate temperatures.

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Figure 80: Ultimate tensile strength Rm as a function of a strain-rate at 200 °C for the low-and high-sulphur material A and B and for the weld material E. The strengthproperties of all materials showed a negative strain-rate sensivity in thistemperature range.

97

All material revealed typical features of DSA (maximum/minimum in Rm/Z at intermediatetemperatures, negative strain-rate sensivity of strength and ductility properties, serrations,...), but to a different extent. Based on changes in Rm and Z (Figure 81 - 83), thesusceptibility to DSA of the studied steels decreases in the following order:

20 MnMoNi 5 5 (A) > SA533 B Cl. 1 (B) > weld (E) > 22 NiMoCr 3 7 (D)

Apart form the high-sulphur alloy B, the DSA results correlated well with the concentrationof free, interstitial N (Figure 84). The DSA response in alloy B might be caused by interstitialC because of its relatively high Ctot content.

4.5.4 Characterization of DSA response by internal friction measurements

Internal friction experiments were performed at Helsinki University of Technology in thetemperature range of 77 to 1000 K on an inverted torsion pendulum with an applied strainrate of 5 • 10"6 s"1 and a heating rate of 1.5 °C/min. The samples (1.5x0.8x50 mm) weremachined in rolling direction of the original plates. The materials were also studied in cold-worked conditions (cold rolling up to 40 % thickness reduction) to simulate the conditions inthe plastic crack-tip zone and to reveal the Snoek-Koester (S-K) internal friction peaks. Theresults of these investigations are discussed in detail in [48].

In a torsion pendulum, the resonance frequency of the system is exited and then the decay ofthe amplitude of the free oscillations is measured. The energy dissipation per cycle (internalfriction IF or Q"1) can be derived from the decay of the amplitude of these free oscillations.Different resonance frequencies can be produced by different pendulum weights. In a typicalinternal friction spectrum, the IF is plotted against the temperature. Information on thebehaviour of free C and N atoms in the lattice and their interactions with dislocations can bederived from the location and the height of the peaks and by the temperature-shifts of peak-maxima at different frequencies (strain rates).

In a common internal friction spectrum of LAS a Snoek peak associated with redistribution offree C and N between equivalent octahedral sites in the lattice is expected at above roomtemperature. This peak may be asymmetric resulting from overlapping of the C peak (locatedat around 39 °C in ot-Fe) and N peak (located at around 24 - 25°C in a-Fe). The otherimportant thing in LAS compared to pure Fe is the presence of alloying elements (such as Cr,Mo and V) affecting especially the diffusivity (jump process) of N and C and thus reducingmarkedly the Snoek peak height. The Snoek peak height can give a measure of free randomlydistributed N or C content in the bulk lattice. Generally a linear relationship between Snoekpeak height and free interstitial content is assumed [95]. This allows a correlation betweenthe observed ductility loss of LAS in tensile tests to Snoek peak height in internal frictionmeasurements to obtain a measure for DSA sensitivity of steel. The broad Snoek-Koester(cold work) peak observed between 150 - 250 °C is due to the mobility of interstitial atoms inthe dislocation stress-strain fields. The height of the Snoek-Koester peak correlates with thedensity of mobile dislocations and the concentration of interstitial C and N along thesedislocations. The exact mechanism of this relaxation is still to be clarified.

98

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d£tech/dt=1E-5s-

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Figure 81: Relative change of the tensile strength Rm from 25 °C to 250 °C at a strain-rate of10"5 s"1 for the materials A, B, D and E.

c>

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-14

-12

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Figure 82: Relative change of the reduction of area Z from 25 °C to 250 °C at a strain-rate of10"5 s"1 for the materials A, B, D and E.

99

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Figure 83: Relative change of the reduction of area Z at 250 °C by increasing the strain-ratefrom 10"3 s"1 to 10"5 s"1 for the materials A, B, D and E.

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Figure 84: Concentration of free, interstitial nitrogen in the materials A, B, C and E.

100

Selected internal friction results are shown in Figure 85-87. Snoek peaks resulting from freeC and N redistribution between octahedral sites in the ferrite lattice are not visible in thespectra of as-received steels, except in the RPV weld metal, which shows a small Snoek peak.This indicates that in the RPV weld metal the bulk content of free interstitials in solidsolution is the highest. The results indicate also that the C and N contents are not in the solidsolution at the detection limit of the present IF technique (estimated to be 10 ppm).

The weak peak observed in the 500 K temperature region corresponds to the Snoek-Koester(cold work) (S-K) peak and is due to the interaction between interstitial C and N anddislocations. The higher S-K peak of SA 533 B Cl. 1 steel is apparently due to the higher Ccontent of this steel compared to the weld metal. The S-K peak height is related to the densityof mobile dislocations and the interstitial content in the vicinity of dislocations. In order tostudy S-K interaction more closely the density of fresh dislocations was markedly increasedby cold rolling (up to 40 % reduction). It can be seen in Figure 85 that already 5 % tensilestrain in the DSA temperature range (280 °C) results in a well-defined S-K peak as well as aSnoek peak. A correlation with S-K peak height (30 % strain at RT) and interstitial content ispresented in Figure 86. This result suggests that the influence of C masks totally the effects ofN. Thus, it seems that free C content (not known exactly) is of crucial importance incorrelation between IF S-K peak height and interstitial content. The IF- and tensile test resultsshow a good correlation concerning the DSA response of the different materials. The S-Kpeak height correlates well to the relative change of reduction of area in tensile tests (Figure87).

oX

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6

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cold worked5 % tensile strain at 280 °C

de/dt = 5 -10"V

as-recieved

0 50 100 150 200 250 300 350 400 450

Temperature [°C]

Figure 85: Internal friction spectra of alloy A in as-received condition and after 5 % tensilestrain at 280 °C.

101

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to t: 7 0

BC,o,: 2500

Nt0,: 60

20MnMoNi5 5 SA533BCI.1 RPV weld

Figure 86: S-K peak heights for materials A, B and E (30 % strain at RT) related to theirC+N contents (wt. ppm).

-2

r-, -4

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20 MnMoNi 5 5SA 533 B Cl. 1RPV weldLinear fit

2.0 2.5 3.0 3.5 4.0 4.5 5.0

C+N S-K peak height X 104 [-]

Figure 87: Excellent correlation of S-K peak heights for materials A, B and E (30 % strain atroom temperature) with the relative change of the reduction of area in tensiletests, if the strain-rate is reduced from 10"3 to 10"5 s"1 at 250 °C.

102

C and N interaction parameters with dislocations were evaluated with 35 % cold rolled20 MnMoNi 5 5 steel (Figure 88). The activation parameters of the C/N S-K relaxationobtained by frequency shift of IF peak maxima for 20 MnMoNi 5 5 steel are the following:H = 1.67 eV and to = 3.7-10" s. Activation enthalpy of S-K relaxation can be expressed asHS-K= HD + HB, where HD is enthalpy of C/N diffusion in solid solution and HB is enthalpy ofC/N binding to dislocations. Using HD = 0.84 eV [96], HB can be evaluated to be 0.83 eV inthis case.

c

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0.8

0.6

0.4

0.2

0.0

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35 % strain at RT

•

1.90 1.95 2.00 2.05

Reciprocal temperature 1000/T [K" ]

2.10-1-

Figure 88: Arrhenius plots for C/N S-K peaks of 35 % deformed (cold rolling) material A, Band E. Activation parameters are obtained for C/N S-K relaxation from frequencyshift of peak maximum.

The absence of Snoek internal friction peaks at about 40 °C does not allow any conclusionsto be made between free interstitials and mechanical properties of these steels. It has beenobserved in general that when the S-K (cold work) peak is present the Snoek peak decreases.The broad S-K internal friction peak at 500 K increases in height with degree of cold work,i.e., with the dislocation density. This peak has been proposed to be due to the movement ofdislocations either in the presence of an atmosphere of interstitial atoms or small particles ofcarbide (Schoeck's theory). It has also been proposed that the basic mechanism of S-Krelaxation is the reorientation of interstitials in the immediate vicinity of dislocations(Koester's theory) or the formation of kink-pairs in the presence of mobile interstitial atoms(C, N, O) (Seeger's theory). [97, 98]

The mechanical tensile test data of the low- sulphur alloy A are summarized in Figure 89 instrain-rate/temperature coordinates. Generally, four regions are observed: yield plateau, yieldplateau + few serrations, serrations, no plateau and no serrations. Boundaries for onset anddisappearance of serrations (dotted lines) have been drawn rather approximately supposingthat the boundary is described by Arrhenius type of dependence taking the values ofactivation enthalpies of S-K peaks obtained from IF measurements. There is a very good

103

agreement between calculated curve (based on IF data) and boundary of disappearance ofserrations for studied RPV steels. These results clearly confirm the main role of interactionbetween mobile interstitials and dislocations in the development of DSA in LAS.

CO,

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Onset ofserrationsi-L = 0.89 eV j

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0 50 100 150 200 250 300 350 400

Temperature [°C]

Figure 89: Map of presence of serrations in stress-strain curves of alloy A. Dotted lines showthe calculated onset and disappearance of serrated flow using the activationenthalpy for S-K IF peak.

4.5.5 Correlation between the observed DSA- and EAC-response

The following experimental indications for possible DSA-effects on EAC in LAS in high-temperature water were found in PSI-tests:

1. Coincidence of susceptibility of LAS to SICC with DSA behaviour in terms of temperatureand strain-rate:

A good coincidence of temperature and strain-rate between DSA (maximum in tensilestrength and loss of ductility) and susceptibility to SICC of RPV steels in SRL-tests underoxygenated high-temperature water conditions has been observed (Figure 15 and 79 and90).

In SRL-tests, the low-sulphur steel behaved significantly worse than the simultaneouslytested steel with 0.018 wt.% S at temperatures < 250 °C. Note that this situation wasreversed at 288 °C (see Figure 15). This behaviour cannot be explained solely by theFRAD- or the HAEAC-mechanism.

At 288 °C, the high-sulphur alloy B revealed a higher SICC susceptibility in oxygenatedhigh-purity high-temperature water than the low-sulphur alloy A over a wide range ofstrain-rates (Figure 15). At very slow strain-rates (rise times > 100 h, strain-rates < 10"5 s"1)the low-sulphur alloy A showed a similar SICC susceptibility. In LFCF tests the low-frequency CF crack growth rates of the low- and high-sulphur steels were very similarover a broad temperature and frequency range, but at 250 and 288 °C the low-sulphuralloy revealed again higher CF CGR at low frequencies < 10"4 Hz than the high-sulphur

104

alloy (Figure 34 and 35). The increasing susceptibility of the low-sulphur alloy A at veryslow strain-rates might also be an indication for a possible DSA effect, e.g. because of itspronounced negative strain-rate sensivity of strength and ductility properties (Figure 80and 83).

90

8 0

"f 70Q.cE 60

50

40

30

20 MnMoNi 5 5 (0.004 wt.% S, 0.013 wt.% Al)• SRL tests (DO = 8 ppm, 65 ppb SO/")

' A Tensile tests in air (DIN 50125)Open symbols: no SICC detectedby DCPD (K™*)

80

7575

70

65

21N

(0

co'•§

I100 150 200 250 300

Temperature [°C]

60350

Figure 90: Coincidence of susceptibility of LAS to SICC in SRL-tests (minimum in Ky) inhigh-temperature water with DSA behaviour (minimum in reduction of area Z) intensile tests in air in terms of temperature at a strain-rate of ca. 10~5 s"1.

2. High SCC susceptibility in the low-sulphur alloy A at intermediate temperatures of 200 -250 °C:

Both the low- and high-sulphur steel A and B revealed a very low SCC susceptibility inoxygenated high-temperature water at 288 °C with CGR < 2-10"11 m/s up to stressintensity factors of 60 MPa-m1/2. Initially fast EAC crack growth triggered by suitableslow rising or cyclic loading suddenly decayed after switching from rising/cyclic toconstant load and crack arrest or very slow crack growth rates < 210"11 m/s were generallyobserved within a short time interval of 10 - 100 h. The high-sulphur alloy B showed thesame behaviour at all temperatures between 150 - 288 °C, whereas the simultaneouslytested low-sulphur alloy A revealed sustained, stationary and fast SCC crack growth at200 and 250 °C (see Figure 59 and 60). The SCC CGR at 200 and 250 °C was two ordersof magnitude higher in the low-sulphur alloy A compared to the simultaneously testedhigh-sulphur alloy B (Figure 58 - 60). Again, this result cannot be rationalized solely bythe FRAD- or the HAEAC-mechanism. The results correlate well with the maximumSICC susceptibility at intermediate temperatures observed in SRL-tests and the DSAresponse of the alloys in the tensile- and IF-tests. The low-sulphur alloy A showed themost pronounced DSA-response with the largest increase in tensile strength and decreasein reduction of area (Figure 79) and the highest C/N S-K-peaks (Figure 86).

All these results clearly demonstrate that the sulphur content is not the sole materialparameter strongly affecting the EAC behaviour and that DSA might probably also relevantlycontribute to the cracking process under suitable temperature/strain-rate combinations.

105

5. Extended summary and conclusions

The SICC/LFCF and SCC crack growth behaviour of different low-alloy RPV steels undersimulated transient and steady-state BWR/NWC conditions was characterized by slow risingload/low-frequency corrosion fatigue and constant load/periodical partial unloading/rippleloading tests. The tests were performed with pre-cracked fracture mechanics specimens inoxygenated high-temperature water at temperatures of either 288, 250, 200 or 150 °C.Modern high-temperature water loops, on-line crack growth monitoring and fractographicanalysis by SEM were used to quantify the cracking response. Additionally the DSA-response of the different steels was characterized by tensile tests (CERT) in air at differenttemperatures/strain-rates and by internal friction measurements. These tests revealed thefollowing important results:

1. Strain-induced corrosion cracking/ very low frequency corrosion fatigue:

Data trends:Maximum SICC susceptibility (initiation of crack growth from incipient cracks) wasobserved at intermediate temperatures (= 250 °C) and slow strain-rates (~ 51O*5 s"1). Withincertain limits, SICC susceptibility increased with increasing ECP, steel sulphur content andbulk-sulphate concentration. Once initiated, SICC CGR increased with increasing strain-rateand, with increasing temperature with a possible maximum/plateau at 250 °C. High SICCCGR of 10"9 m/s (3 cm /a) to 810"7 m/s (25 m/a) were observed in the loading rate rangedlQ/dt of 0.1 to 500 MPam1/2/h for ECP > 150 mVSHE and Ki> 30 MPa-m1/2 by both SRL-and vLFCF-tests.

Critical conditions for SICC:

The occurrence of SICC requires critical conditions, i.e. a high sulphur-anion activity in thecrack-tip environment and a slow positive tensile crack-tip strain-rate, to be achievedsimultaneously. If these conjoint conditions are not achieved, no, or only minor, local SICCgrowth is observed. However, if they are achieved or exceeded, the SICC CGR seems to bedependent primarily on the applied strain-rate. A high SICC susceptibility is favoured by ahigh ECP and/or a high steel sulphur content/bulk sulphur-anion activity and quasi-stagnantflow conditions, which favour the enrichment of sulphur-anions in the crack-tip environment.

Steel sulphur content and PSA:

The beneficial effect of a low steel sulphur content at temperatures > 250 °C could be maskedby suitable combinations of corrosion system parameters (e.g. by suitable strain-rate/temperature combinations and/or by a high bulk sulphur-anion content + a high ECP).The results at very slow strain-rates (< 10"5 s"1) and intermediate temperatures (150 - 250 °C)indicate that DSA might significantly affect EAC cracking behaviour in susceptible LAS.The Al-content and concentration of free N/C might therefore be just as relevant for EACsusceptibility as the S- content, at least in the DSA temperature/strain-rate range.

2. Low-frequency corrosion fatigue:

Data trends:

Temperature/strain-rate effects on CF crack growth were investigated by LFCF tests withdifferent RPV steels at four different temperatures (150, 200, 250, 288 °C) and frequencies(10'5, 10"4, 810"4, 2.5-10"3 Hz) in water with 8000 ppb DO and 65 ppb or < 1 ppb SO4

2". Forall materials and all temperatures, the CF crack advance per cycle AaEAc/AN increased withdecreasing loading frequency, whereas the measured time-based CGR daEAc/dt decreasedwith decreasing frequency. Stable, stationary CF crack growth was observed down to very

106

low frequencies of 10"5 Hz indicating a critical frequency < 10~5 Hz for highly oxidizingBWR conditions. For all frequencies, both cycle-based CGR AaEAc/AN and time-based CGRdaEAc/dt increased with increasing temperature from 150 - 250 °C. No noticeable change inCF CGR was observed by further increasing the temperature from 250 to 288 °C. Theenvironmental acceleration of fatigue crack growth increased both with decreasing loadingfrequency and increasing temperature. In LFCF tests under these highly oxidizing conditions,environmental acceleration of fatigue crack growth in the range of 1 to 3 orders of magnitudewas observed. The time-based LFCF CGR were in the range of 510"10 m/s (1.6 cm/a) to5-10"8 m/s (1.6 m/a) and decreased with decreasing loading frequency and temperature. TheLFCF crack growth behaviour of low- and high-sulphur steels was comparable over a widerange of loading conditions.

Comparison to ASME XI and to new reference curves proposed by Eason:

Under highly oxidizing (8 ppm DO, ECP > +100 HIVSHE) and low-flow conditions, both theASME XI "wet" reference fatigue crack growth curve and the new reference curves proposedby Eason could be significantly exceeded by cyclic fatigue loading at low frequencies(<10"3Hz) for both low- and high-sulphur steels and low and high load ratios in thetemperature range between 150 and 288 °C. Preliminary test results indicate that this mightbe also true for high-purity water with a DO of 400 ppb. The situation was worst for rippleloading at very high load ratios near to the fatigue thresholds AKth. Under very low frequencyloading conditions, cycle-based CGR of up to several few hundred urn/cycle were observed.Values below the ASME XI "wet" reference crack growth curve were only observed atloading frequencies > = 810"4 Hz. Within the investigated parameter range, the deviationform the ASME XI wet reference CGR increased with decreasing frequency and increasingload ratio and temperature.

These results indicate the need for further work and should be verified at lower Ki"18* and AKlevels, corresponding to crack depths slightly above the resolution limit of the NDT methodsof the periodic in-service inspection. The test conditions (ECP, T, dlQ/dt, Ki, ...) should alsobe further adjusted to correspond more closely to different BWR transient operatingconditions (start-up/shut-down, hot-stand-by, thermal stratification,...).

3. Stress corrosion cracking:

Data trends:

High SCC CGR of up 3-10'7 m/s (9 m/a), as sometimes reported in older literature, were notobserved within this parameter study under simulated BWR/NWC water-chemistryconditions. At 288 °C, no sustained SCC crack growth was observed in different low-alloyRPV steels (base metal, 0.004 - 0.018 wt.% S) under static loading conditions at Ki-values< 60 MPa-m1/2 for ECP < +150 mVSHE and for tests at conductivities < 0.25 ja,S/cm(equivalent to < 65 ppb SO42") as long as small scale yielding in the specimen ligamentprevailed. Under these conditions, even fast SICC crack growth (up to 10" m/s) triggered bya suitable slow rising load could not be maintained under subsequent static loading and crackarrest was observed under the static load. However, sustained, fast SCC (with respect tooperational time scales) cannot be excluded for faulted water-chemistry conditions (> EPRIAction Level 3) and/or for highly stressed specimens, either loaded near to KD or with a highdegree of plasticity in the remaining ligament.

107

Comparison to the BWR VIP 60 SCC disposition lines:

The conservative character of the BWR VIP 60 disposition lines 1 and 2 for SCC crackgrowth in low-alloy steels has been confirmed by this study for 288 °C and RPV basematerial (< 0.018 wt.% S) by constant load and periodical partial unloading tests. Preliminaryresults indicate, that these disposition lines may be significantly or slightly exceeded (even insteels with a low sulphur content) in the case of small load fluctuations at high load ratios(ripple loading) or at intermediate temperatures (200 - 250 °C) in RPV materials, whichshow a distinct susceptibility to DSA.

The BWR VIP 60 SCC disposition lines seem to be conservative and adequate for the RPV(unaffected base metal) for transient-free steady-state BWR power operation, but there issome concern for

the RPV feedwater nozzle and the feedwater piping system with lower operatingtemperatures and the temporary occurrence of small load fluctuations (—> DSA, rippleloading).

- Weld filler (DSA) and weld HAZ (-» DSA, TG/IG HAEAC) materials.

- Materials in as-welded condition without PWHT (—» DSA).

LAS RPV and piping materials with a low Al-content and a high concentration ofinterstitial N and C (-» DSA).

The BWR VIP SCC disposition lines should therefore be validated by further tests

at intermediate temperatures,

with weld filler and weld HAZ materials.

- with ripple loading and periodical partial unloading.

The results should be verified at lower Ki and AK levels, corresponding to crack depthsslightly above the resolution limit of the NDT methods of the periodic in-service inspection.The water-chemistry conditions (ECP, SO42", ...) should also be further adjusted to currentBWR operation conditions.

4. Possible effects of DSA on EAC of LAS:

Physical metallurgy of DSA in LAS:

DSA may be observed in susceptible LAS during plastic straining at sufficiently slow strain-rates (< = 10"2 s"1) in the temperature range from 100 °C to 350 °C. It is associated withdiffusion of interstitial species such as C and N atoms to the core region of dislocations andtheir immobilization. The DSA effects increase with increasing concentration of free,interstitial N and C and are most pronounced if the diffusion rate of C/N and the dislocationvelocity are similar. The concentration of free interstitial C and N are not specified in thenuclear codes (ASME BPV, KTA, ...) and strongly depend on the steel making (killing) orwelding process, on the thermal history or heat-treatment (annealing- and PWHT-temperature) and on the chemical composition (Ctot, Ntot, Al, V, Ti, Cr, Mo, O, Mn ...) of thesteel.

Possible interaction between DSA and EAC in LAS:

DSA may be very influential in determining both EAC cracking strain-rate sensivity and itstemperature dependence. It provides an alternate possible explanation for strain-ratethresholds and more significantly, it may offer a very attractive explanation for the EAC

108

susceptibility peak at intermediate temperatures and rationalize the scattered data in thistemperature range. Differences in the concentration of free interstitials N/C of otherwiseidentical LAS might be one further reason for the scatter of EAC data. A higher EACsusceptibility of weld filler and weld HAZ materials in the DSA temperature/strain-rate rangecompared to the un-affected RPV base material cannot be excluded, because of possibleDSA-effects.

In EAC, DSA is especially important in dynamic crack-tip plasticity behaviour such asdynamic loading or evolution of creep strain. DSA may result in a higher crack-tip strain andstrain-rate than outside the DSA-range or than in a material, which is not susceptible to DSA.The inhomogeneous localisation of deformation, the increase of dislocation density andincrease in planar deformation by DSA can result in a reduction of the local fracturetoughness and favour brittle crack extension, but also in the mechanical rupture of theprotective oxide film and therefore in crack advance by anodic dissolution/hydrogenembrittlement mechanisms. Therefore, DSA may synergistically interact with the FRAD- andHAEAC-mechanism to increase EAC cracking susceptibility. EAC in LAS has beenobserved under temperature/stain rate conditions or in materials, where no or only minorDSA effects were present. DSA is therefore not a pre-requisite for EAC and best regarded asan additional contribution to EAC growth.

Results of tensile tests / internal friction measurements:

Tensile tests at different temperature/strain-rates and internal friction measurements revealedtypical features of DSA in all investigated RPV materials. Apart form one alloy, DSAsusceptibility correlated well with the concentration of free interstitial N. The tests resultsalso indicate that the concentration of interstitial C might be at least as important as the freeN content in LAS.

Experimental indications for DSA effects on EAC:

A good coincidence of DSA behaviour (maximum in tensile strength and loss of ductility) inair and susceptibility to SICC of RPV steels in SRL-tests under low-flow oxygenated high-temperature water conditions has been observed in terms of temperature and strain rate.Furthermore, the low-sulphur RPV steel (0.004 wt.% S) with a high concentration of freenitrogen (30 ppm Nf^e) revealed stationary fast SCC growth (~ 30 mm/a) under static load atintermediate temperatures of 200 and 250 °C in the DSA peak region. The simultaneouslytested high-sulphur alloy (0.018 wt.% S) with a low free nitrogen content (< 1 ppm N&ee) onlyshowed minor SCC crack growth (< 0.6 mm/a). All these results clearly indicate that DSAmight also strongly affect the EAC behaviour and that the concentration of free interstitial Nand C might be as relevant as the S-content, at least in the DSA-temperature/strain-rate range.

Our preliminary results clearly demonstrate the need for further investigations at intermediatetemperatures and slow strain-rates/low frequencies including weld filler and weld HAZmaterials.

109

6. Future work

The following experimental investigations are planned until the end of the project:

l.SICC/LFCF:

• Completion of the parameter study on the effect of temperature and strain-rate/loadingfrequency on the SICC and LFCF crack growth behaviour with weld filler and weld HAZmaterial. The material is from the circumferential weld between the lower and upper shellof a real RPV (Biblis C, PWR, 1200 MWe, 1976). In the case of HAZ material, both pre-cracked IT C(T) specimens and blunt notched C(T) specimens with a large notch radiuswill be used. The geometry of the blunt notch C(T) specimens has been optimised by FE-modelling [18]. In addition to the information on the crack growth behaviour ofmechanically long cracks, the tests with the blunt notch C(T) specimen will also provideinformation on the crack initiation process (number of cycles for crack initiation vs. pseudostress amplitude AK/Vp) and on the short crack growth behaviour.

• The effect of loading conditions (AK, R, v, Atn, •••) will be systematically studied with themost susceptible material in the temperature range of maximum SICC/LFCF-crack growthsusceptibility. The loading and water-chemistry parameters will be adjusted to BWR planttransients (start-up/shut-down, thermal stratification, ...). The AK-values shouldcorrespond to cracks with a length between the resolution limit of the ND in-serviceinspection methods and one quarter of the wall thickness of the RPV or feedwater nozzle.

2. SCC:

• The parameter study on the effect of temperature on the SCC crack growth behaviour willbe completed for weld filler and weld HAZ materials. In the case of HAZ material, bothpre-cracked IT C(T) specimens and blunt notched C(T) specimens with a large notchradius will be used. The high SCC crack growth susceptibility of the low-sulphur materialA at intermediate temperatures will be verified at lower Ki-values.

• The results of the tests with ripple loading and periodical partial unloading will be verifiedat lower K ^ - and AK-values at 288 °C and in the DSA temperature/strain-rate range withselected materials.

110

7. AcknowledgementThe financial support for this work by the Swiss Federal Nuclear Safety Inspectorate (HSK)and the Swiss Federal Office of Energy (BFE) is gratefully acknowledged. The authors aregrateful to U. Ineichen, E. Groth and U. Tschanz (all PSI) for their experimental contributionsto this work and to Prof. Hanninen from Helsinki University of Technology for performingthe IF measurements.

I l l

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