* U.S.NRC

United States Nuclear Regulatory Commission

Protecting People and the Environment

NUREG/CR-7137

ANL-10/36

Stress Corrosion Cracking

in Nickel-Base Alloys 690

and 152 Weld in Simulated

PWR Environment - 2009

Office of Nuclear Regulatory Research
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*U.S.NRC

United States Nuclear Regulatory Commission

Protecting People and the Environment

NUREG/CR-7137

AN L- 10/36

Stress Corrosion Cracking

in Nickel-Base Alloys 690

and 152 Weld in Simulated

PWR Environment - 2009

Manuscript Completed: June 2012 Date Published: June 2012

Prepared by

B. Alexandreanu, Y. Yang, Y. Chen, and W. J. Shack

Argonne National Laboratory

Argonne, IL 60439

Darrell Dunn/Appajosula Rao, NRC Project Managers

Office of Nuclear Regulatory Research

Abstract

Alloys 600 and 182 are used as structural materials in pressurized water reactors (PWRs) and have been

found to undergo stress corrosion cracking (SCC). Alloys 690 and 152 are the replacement materials of

choice for Alloys 600 and 182, respectively. The primary objective of this work is to determine the crack

growth rates (CGRs) in a simulated PWR water environment for the replacement alloys. This is an

ongoing effort, and the report presents the initial investigations on Alloys 690 and 152 weld. A second

objective was to obtain tensile data for these alloys for a wide temperature range, from operating to

extreme conditions. To meet the objectives, testing conducted at ANL included Alloy 690 in the asreceived condition and cold-rolled by 26%, as well as a laboratory-prepared Alloy 152 double-J weld in

the as-welded condition. For the SCC CGR measurements, the specimens were pre-cracked under cyclic

loading in a primary water environment, and the cyclic CGRs were monitored to determine the transition

from the fatigue transgranular fracture mode to the intergranular SCC fracture mode. The cyclic CGRs

of the cold-rolled Alloy 690 showed significant environmental enhancement compared with the asreceived alloy. Compared to Alloy 600, the as-received Alloy 690 proved to be more resistant to SCC in

the simulated PWR environment at 320'C (608 'F), with SCC CGRs of 10-12 m/s. However, the SCC

CGRs for the cold-rolled alloy were an order of magnitude higher, about 10-11, m/s, at comparable stress

intensity factors in the range of 22 to 32 MPa-m12 (20 to 29 ksi-in.l/ 2

). This finding is consistent with

results obtained by industry on similar cold-worked alloys. A possible explanation for the relatively

high CGRs in the cold-worked material is the presence of cracked grain boundary carbides, a feature that

was observed for all cold-worked alloys displaying relatively high CGRs. The environmental

enhancement of cyclic CGRs for Alloy 152 was minimal; nevertheless, the transition from transgranular

to intergranular cracking was successful. Weld samples tested from a single heat of Alloy 152 exhibited

SCC CGR rates of 101" m/s in the simulated PWR environment at 320'C (608 'F), which was only about

an order of magnitude lower than typical for Alloy 182. The report also presents tensile property data

for Alloy 690 and Alloy 152 welds in the temperature range of 25-870'C (81-1598°F). These data meet

the minimum requirements of ASME SB- 167, and are in excellent agreement with data on similar alloys.

iii

Foreword

Primary water stress corrosion cracking (PWSCC),is a significant issue that has affected nickel base

alloy pressurized-water reactor (PWR) components. Observation of PWSCC in the past decade has

challenged the assumptions of leak before break analyses that were based on evaluations that did not

consider PWSCC as a degradation mechanism that may contribute to potential sources of pipe rupture

and subsequently a loss of coolant accident. Operational experience has also shown that PWSCC of

penetrations in the reactor pressure vessel head lead to boric acid corrosion of low-alloy pressure vessel

steels.

The U.S. Nuclear Regulatory Commission (NRC) mandated that inspections are required for reactor

coolant system components constructed from PWSCC susceptible materials, including Alloy 600 and its

weld metals, Alloys 182 and 82. These inspections result in occupational exposures and are also both

difficult to conduct and costly to implement. Consequently, the industry has proposed methods, for

PWSCC mitigation such as the use of replacement components manufactured from Alloy 690 and-its

weld metals, Alloys 152 and 52. Limited operational experience with these higher chromium containing

alloys in thick section pressure boundary components has been favorable. However, the effects of

fabrication processes on alloy microstructure and PWSCC susceptibility are not fully understood.

Ongoing industry research is being conducted to evaluate metallurgical and environmental effects that

influence PWSCC resistance. Confirmatory testing sponsored by the NRC is being conducted to verify'

improvement factors and determine acceptable inspection intervals for components manufactured from

the higher chromium-containing alloys.

This report documents work conducted by Argonne National Laboratory (ANL) during calendar years

2008 and 2009 under contract to the NRC. Testing conducted at ANL during this period consisted of

microstructural analyses and crack growth rate tests to measure the PWSCC resistance of cold-worked

Alloy 690. In addition, microstructural analysis and tests to determine mechanical properties and crack

growth rates for an Alloy 152 weld were conducted. Results of these tests indicate that Alloy 690 in the

as-received and cold-worked conditions was more resistant to PWSCC compared to Alloy 600.

However, results obtained by Bettis Atomic Power Laboratory show that some heats of Alloy 690 in the

cold-worked condition have crack growth rates that are greater than those measured for Alloy 600.

While results with the Alloy 152 weld metal show improved PWSCC resistance compared to Alloy 182,

the measured crack growth rates on limited tests with Alloy 152 were greater than the rates reported by

industry and other testing laboratories.

Additional testing will be conducted to evaluate factors that affect PWSCC resistance that may

contribute to the range of measured crack growth rates. These tests will include multiple heats of Alloy

690 and 152 weld metals. In addition, the microstructure of cold-worked materials may be similar to

weld heat affected zones (HAZ). Subsequent tests will be conducted to determine the microstructural

characteristics and crack growth rates in Alloy 690 weld HAZs.

Michael J. Case, Director

Division of Engineering

Office of Nuclear Regulatory Research

U.S. Nuclear Regulatory Commission

v

Contents

ABSTRACT ............................................................... i

FOREWORD ............................................................. v

C O N T E N T S .............................................................................................................................................. vii

F IG U R E S ................................................................................................................................................... ix

T AB L E S .................................................................................................................................................... xv

EX ECU TIV E SU M M A R Y .................................................................................................................... xvii

A CK N O W LED G M EN T S ....................................................................................................................... xix

ACRONYMS AND ABBREVIATIONS ................................... ...... xxi

1. INTRODUCTION ...................................................... 1

2. EXPERIMENTAL ..................................................... 3

2.1 Materials ........................................................ 3

2.2 Specim en D esign ............................................................................................................. 7

2.3 Test Facilities ..................................................... 7

2.3.1 Facility with 6-Liter Autoclave ............................................................................ 8

2.3.2 Facility with a 2-Liter Autoclave ........................................................................ 9

2.3.3 Prim ary W ater Environm ent ............................................................................. 11

2.4 T est Procedure ........................................................................................................................ 14

2.5 Analysis of Crack Growth Rate Data ............................................................................... 16

3. MICROSTRUCTURAL CHARACTERIZATION ................................................................... 19

3.1 Alloy 690 in Plate Form (A690W C) ................................................................................. 19

3.1.1 H ardness Testing ............................................................................................... 19

3.1.2 General Microstructure - Grain Boundary Carbides ......................................... 20

3.1.3 Grain Boundary Character Distribution ............................................................. 22

3.1.4 Carbide and Grain Size Banding ....................................................................... 24

3.2 Alloy 690 TT from CRDM Tubing (C690) ...................................................................... 28

3.3 A lloy 152 Double-J W eld (A 152) .................................................................................... 30

vii
4. STRESS-STRAIN AND CRACK GROWTH RESULTS .......................................................... 35

4.1 Tensile Data for Alloy 690TT and Alloy 152 Double-J Weld .......................................... 35

4.2 Crack Growth Data for Alloy 690 Cold-Rolled by 26% ................................................... 41

4.2.1 Specimen A690WC-SL-1 ................................................................................. 41

4.2.2 Specimen A690WC-ST-1 ................ ................................................................. 47

4.2.3 Specimen A690WC-ST-2 ................................................................................. 50

4.3 Crack Growth Data for Alloy 690 in As-received Condition ............................................ 54

4.3.1 Specimen C690-CR- I ........................................................................................ 54

4.3.2 Specimen C690-LR-2 ........................................................................................ 60

4.4 Crack Growth Data for Alloy 152 Weld ......................................................................... 66

4.4.1 Specimen A152-TS-2 ....................................................................................... 67

4.4.2 Specimen A152-TS-4 .............................................................. 75

5. D ISC U S SIO N .................................................................................................................................. 85

5 .1 A lloy 690 ................................................................................................................................ 85

5.1.1 Cyclic Crack Growth Rates ............................................................................... 85

5.1.2 SCC Crack Growth Rates .................................................................................. 88

5.1.3 Activation Energy .............................................................................................. 89

5.1.4 SCC of Cold-Worked Alloy 690 from a Mechanistic Perspective .................... 91

5.2 Alloy 152 Weld .................................................................................................... 94

5.2.1 Cyclic Crack Growth Rates ............................................................................... 94

5.2.2 SCC Crack Growth Rates .................................................................................. 96

6. SUMMARY ......................................................................... 99

R E FE RE N C E S ........................................................................................................................................ 10 1

viii
Figures

Figure 1. Orientation of Alloy 690 specimens machined from cold-rolled plate ...................... : ............ 3

Figure 2. Orientation of the Alloy 690 specimens cut from the CRDM nozzle ............................... 4

Figure 3. Schematic of the Alloy 152 weld build-up on the interior of the Alloy 690 CRDM tube ....... 4

Figure 4. Schematic of the weld joint design for the (a) Alloy 152 double-J weld and (b) weld

passes and filler heats used. Dimensions are in inch .......................................................... 6

Figure 5. Configuration of the (a) 1-T and (b) 1⁄2-T CT specimens used for this study.

D im ensions are in m m .......................................................................................................... 7

Figure 6. Schematic diagram of the recirculating 6-liter autoclave system ....................................... 9

Figure 7. Layout of the 2-liter autoclave system .................................... 9

Figure 8. Photograph of the specimen load train for the 2-liter autoclave ....................................... 10

Figure 9. Schematic diagram of the recirculating 2-liter autoclave system ........................................... 11

Figure 10. Hydrogen concentration at Ni/NiO phase transition as a fuiction of temperature (data

taken from R ef. 15) ................................................................................................................ 13

Figure 11. Crack growth rates in air and environment for Alloys 600, 690 and Ni-base welds as a

function of predicted rates in air for Alloy 600 (data taken from Refs. 4,7,9) .................. 17

Figure 12. (a) Schematic of the rolling process of the ANL plate identifying the locations (a, b,

and c) at which hardness was evaluated, and (b) hardness along thickness of the Alloy

690 plate cold-rolled by 26% (ANL). Similar data from a plate cold-rolled by 20% at

GE are included for com parison ...................................................................................... 19

Figure 13. Alloy 690 cold-rolled by 26%: (a) general microstructure, (b) grain boundary carbides,

and (c) TiN precipitates. Images were obtained on the mid-plane of the deformed plate ..... 20

Figure 14. TEM micrographs from mid (rolling plane). Alloy 690 deformed by 26% showing the

following: (a,b) high deformation microstructure and (c,d) grain boundary carbides ........... 21

Figure 15. Alloy 690 cold-rolled by 26%, which was CGR tested at ANL ....................................... 22

Figure 16. Alloy 690 cold-rolled by 20%, which was provided by GE ............................................ 22

Figure 17. (a) OIM map, (b) graph of grain size distribution, and (c) graph of grain boundary

character distribution for Alloy 690 in plate form in the as-received condition ............... 23

Figure 18. (a) OIM map, (b) graph of grain size distribution, and (c) graph of grain boundary

character distribution for Alloy 690 in plate form in the 26% cold-rolled condition ........ 24

ix
Figure 19. Schematic showing the positions of the metallography specimens cut from the plate in

the as-received condition .................................................................................................. 25 . I i

Figure 20. Microstructure of specimen #1 showing carbide streaks ................................................. 25

Figure 21'. Grain size banding in (a) Specimen #5 near edge 1 and (b) Specimen #6 near edge 2

(images rotated by 90'C vs. Fig. 19) ............................................................................... 26

Figure 22. Microstructure observed in Specimen #8: (a) near edge 4 and (b) middle of the

specimen (images rotated by 90'C vs. Fig. 19) ............................................................... 26

Figure 23. Carbide streaks in Alloy 690 plate cold-rolled by 26% ................................................... 27

Figure 24. Grain microstructure in Alloy 690 plate cold-rolled by 26% ............................................ 27

Figure 25. Microstructure of Alloy 690 TT from CRDM tubing ..................................................... 28

Figure 26. Microstructure of Alloy 690 TT from CRDM tubing ..................... L ................................ 28

Figure 27. Microstructure of Alloy 690 TT from CRDM tubing showing grain boundary carbides ...... 29

Figure 28. (a) OIM map, (b) graph of grain size distribution, and (c) graph of grain boundary

character distribution for Alloy 690 TT from CRDM tubing ............................................ 30

Figure 29. Microstructtire of Alloy 152 weld using SEM: (a) secondary electron detector and (b)

back-scattered electron detector ...................................................................................... 31

Figure 30. Optical image of the Alloy 152 weld specimen A152 ...................................................... 31

Figure 31. (a-c) TEM images and (d) EDS spectrum of an inclusion in the double-J Alloy 152

w eld (A 2152) ............................................................................................................................ 32

Figure 32. Alloy 152 weld in double-J geometry (A152): (a) OIM map and (b) grain boundary

character distribution ......................................................................................................... 33

Figure 33. Locations in the double-J Alloy 152 weld where the tensile specimens were cut ............ 35

Figure 34. (a) Stress-strain and (b) true stress-true strain curves for Alloy 690 TT as a function of

tem perature ............................................................................................................................. 37

Figure 35. (a) Yield stress and (b) ultimate tensile stress of Alloy 690TT as a function of

tem perature ............................................................................................................................. 38

Figure 36. (a) Stress-strain and (b) true stress-true strain curves for Alloy 152 weld as a function

of tem perature ......................................................................................................................... 39

Figure 37. (a) Yield stress and (b) ultimate tensile stress of Alloy 152 weld as a function of

tem perature ............................................................................................................................. 40

x
Figure 38. Crack length vs. time for Alloy 690 specimen A690WC-SL-1 in simulated PWR

environment during test periods (a) 1-5 and (b) 6-9 .......................................................... 41

Figure 39. Fracture surface of specimen A690WC-SL-1 by (a) optical and (b) electron

m icroscopy ............................................................................................................................. 42

Figure 40. Fracture modes observed on the fracture surface of A690WC-SL-l: (a) transgranular;

(b,c) faceted with elements of IG; and (d) IG showing carbides on the facets. Crack

advance is from bottom to top ........................................................................................... 43

Figure 41. Fracture modes observed on the fracture surface of A690WC-SL-1: (a) IG and TG, (b)

IG fracture, and (c,d) TG fracture with elements of IG (smooth facets). Crack advance

is from bottom to top ......................................................................................................... 44

Figure 42. Fracture modes observed on the fracture surface of A690WC-SL-1: (a,b) IG fracture and

(c) TG fracture with elements of IG (smooth facets). Crack advance is from bottom to

to p ........................................................................................................................................... 4 5

Figure 43. Cross section of specimen A690WC-SL-1: (a) entire (IG and TG) crack front and (b)

IG region at the end of test ................................................................................................ 46

Figure 44. Crack length vs. time for Alloy 690 Specimen A690WC-ST-1 in simulated PWR

environment during test periods (a) precracking-2, (b) 3-7, (c) 8-11, and (d) 12-13 ...... 47

Figure 45. Fracture surface of specimen A690WC-ST-1. Crack advance is from bottom to top ........... 49

Figure 46. Fracture modes observed on the fracture surface of A690WC-ST-1: (a) TG, (b) IG

fracture, and (c) IG fracture detail. Crack advance is from bottom to top ...................... 49

Figure 47. Crack length vs. time for Alloy 690 Specimen A690WC-ST-2 in simulated PWR

environment during test periods (a) precracking-9, (b) 10-23, and (c) 24-27 ................... 51

Figure 48. Cyclic CGR data for Alloy 690 Specimen A690WC-ST-2 at 350'C (6627F) and for

previous two specimens under similar conditions in PWR environment at 320'C

(6 0 8 FF ) ................................................................................................................................... 52

Figure 49. Fracture surface of specimen A690WC-ST-2. Crack advance is from bottom to top ........... 53

Figure 50. Behavior observed on the fracture surface of specimen A690WC-ST-2: (a) TG and IG

regions, and (b,c) the transition between TG and IG. Crack advance is from bottom to

to p ........................................................................................................................................... 5 3

Figure 51. Crack length vs. time for Alloy 690 specimen C690-CR-I in simulated PWR

environment during test periods (a) precracking-4, (b) 5-9, and (c) 10 ........................... 55

Figure 52. Region of fracture surface of specimen C690-CR-1. Crack advance is from bottom to

to p ........................................................................................................................................... 5 6

Figure 53. Fracture surface of specimen C690-CR-l: (a) first half and (b) second half. ................... 57

xi
Figure 54. Specimen C690-CR-1: (a) section of the fracture surface, (b) TG and faceted fracture

with IG cracks, and (c) TG with IG secondary cracks. Secondary IG cracks are

marked with arrows. Crack advance is from bottom to top ............................................ 58

Figure 55. Fracture surface of specimen C690-CR-1: (a,b) regions near the end of the test, and

(c,d) higher magnification micrographs showing the same areas, with smooth facets

and IG secondary cracks. Secondary IG cracks are identified with arrows. Crack

advance is from bottom to top .......................................................................................... 59

Figure 56. Examples of areas with smooth facets and IG secondary cracks on the fracture surface

of specimen C690-CR-1. Crack advance is from bottom to top .................................... 60

Figure 57. Crack length vs. time for Alloy 690 Specimen C690-LR-2 in simulated PWR

environment during test periods (a) precracking, (b) 1-5, and (c) 6-8 ............................. 61

Figure 58. Oxides on the fracture surface of specimen C690-LR-2 ................................................... 63

Figure 59. Fracture surface of specimen C690-LR-2. Crack advance is from bottom to top ............ 63

Figure 60. Examples IG secondary cracks and occasional smooth facets on the fracture surface of

specimen C690-LR-2. Secondary IG cracks are identified with arrows. Crack advance

is from bottom to top ......................................................................................................... 64

Figure 61. Fracture surface of specimen C690-LR-2, regions near the end of the test. Extensive

secondary IG cracks are identified with arrows. Crack advance is from bottom to top ........ 65

Figure 62. Examples of areas with smooth facets and IG secondary cracks on the fracture surface

of specimen C690-LR-2. Crack advance is from bottom to top ....................................... 66

Figure 63. Crack length vs. time for Alloy 152 weld specimen A152-TS-2 in simulated PWR

environment during test periods (a) precracking, (b) 1-3, and (c) 4 .................................. 67

Figure 64. Crack length vs. time for Alloy 152 weld specimen A152-TS-2 in simulated PWR

environment during constant-load test period 4 ............................................................... 69

Figure 65. Fracture surface of specimen A152-TS-2 by (a) optical and (b) electron microscopy.

Crack advance is from bottom to top .............................................................................. 69

Figure 66. Intergranular cracking on the fracture surface of specimen A152-TS-2 between A and

B in Fig. 65. Crack advance is from bottom to top .......................................................... 70

Figure 67. Specimen A152-TS-2: (a) first third of fracture surface, (b) detail at location 1, and (c)

detail at location 2. Crack advance is from bottom to top .............................................. 72

Figure 68. Specimen A152-TS-2: (a) second third of fracture surface, (b) detail at location 1, and

(c) detail at location 3. Crack advance is from bottom to top ........................................... 73

Figure 69. Specimen A152-TS-2: (a) last third of fracture surface, (b) detail at location 1, and (c)

detail at location 2. Crack advance is from bottom to top ................................... : ............... 74

xii
Figure 70. Region of fracture surface of specimen A152-TS-2. Crack advance is from bottom to

to p ........................................................................................................................................... 7 5

Figure 71. Crack length vs. time for Alloy 152 weld specimen A152-TS-4 in simulated PWR

environment during test periods (a) precracking-2, (b) 3-5, (c) 6-19, (d) 20-25, (e) 26-

37, and (f) 38-39 ............................................................................................................ 77

Figure 72. Fracture surface of specimen Al 52-TS-4. Crack advance is from bottom to top ............. 79

Figure 73. Intergranular region at the end of specimen A152-TS-4. The image was broken into

three regions. Crack advance is from bottom to top ........................................................ 81

Figure 74. Fracture surface of specimen A152-TS-4 at the end of the test (arrows). Crack advance

is from bottom to top ......................................................................................................... 82

Figure 75. Ligaments on the fracture surface of specimen A152-TS-4: (a) out-of-plane crack and

(b) interdendritic (left) and TG (right) fracture modes. Crack advance is from bottom

to to p ....................................................................................................................................... 8 2

Figure 76. Fracture surface of specimen A152-TS-4 prior to period 22. Crack advance is from

b ottom to top .......................................................................................................................... 83

Figure 77. Cyclic CGR data for (a) Alloy 690 in SA and TT conditions from previous ANL and

B&W tests in PWR environment and (b) Alloy 690 in high purity water with low DO

as a function of growth rates for Alloy 690 in air. Test temperature is 320'C (608'F) ......... 85

Figure 78. Cyclic CGR data for 26% cold-rolled Alloy 690 and Alloy 600 from Davis-Besse in

PWR environment at 320'C (608'F) .................................... 86

Figure 79. (a) Transition from TG to IG fracture in Alloy 600 from Davis-Besse Nozzle #3, and

(b) predominant TG fracture during transitioning in Alloy 690 cold-rolled by 26%.

Grain contours, some smooth facets, and IG cracks are evident. In both cases, the

sample notch is at the top of the figure, and crack advance is from top to bottom ........... 87

Figure 80. Cyclic CGR data for 26% cold-rolled and as-received Alloy 690 specimens in PWR

environm ent at 320'C (608'F) ........................................................................................ 87

Figure 81. SCC CGR data for cold-worked Alloy 690 in PWR environment at 320-340'C (608-

644'F). The ANL data are shown with solid green symbols ............................................ 88

Figure 82. SCC CGR data for as-received and cold-worked Alloy 690 in PWR environment at

320-340-C (608-644°F) .................................................................................................... 89

Figure 83. Temperature dependence of the CGR data for Alloy 690: (a) 26% and 24% cold-rolled

and as-received conditions obtained at ANL, GE, and Bettis; (b) 26% and 20% coldrolled and as-received conditions obtained at ANL and GE; and (c) 20% and 50% coldrolled conditions obtained at IN SS .................................................................................... 90

xiii
Figure 84. SCC CGRs in PWR water for Alloy 690 heats deformed by 20-32%. The data were

normalized to 320'C (608'F) in de-aerated water using an activation energy of 130

k J/m o l ..................................................................................................................................... 9 1

Figure 85. (a) Temperature dependence of SCC CGR data for 20% cold-rolled Alloy 690 and

creep data for Alloys 690 and 600 and (b) dependence of steady-state creep at 360'C

and percent cracked boundaries in primary water on the fraction of CSLBs in Ni- I6Cr9 F e alloy s ............................................................................................................................... 92

Figure 86. Crack growth rate in Alloy 600 and 690 samples in the MA and TT conditions vs. the

grain boundary viscosity, 7 .............................................................................................. 92

Figure 87. Smooth IG fracture in Alloy 600 cross-rolled by 30%. Crack advance is from left to

rig h t ........................................................................................................................................ 9 4

Figure 88. Cyclic CGR data for Alloy 182 weld in PWR environment at 320'C (608'F) ................. 95

Figure 89. (a) Transition from TG to IG fracture and (b) higher magnification micrograph of the

boxed area showing the tip of a few secondary cracks. Crack extension from bottom to

to p ............................................................................................................................................ 9 5

Figure 90. Cyclic CGR data for Alloy 152 weld in PWR environment at 320'C (608'F) obtained

at A N L and B & W ............................................................................................................ 96

Figure 91. SCC CGR data for Alloy 152 welds as a function of stress intensity in PWR

environm ent ............................................................................................................................ 97

xiv
Tables

Table 1. Chemical composition (wt.%) of Alloy 690 (Heat NX3297HK12) plate ............................. 3

Table 2. Welding process and conditions used to build up the interior of the Alloy 690 CRDM

tube .......................................................................................................... . . .5

Table 3. Chemical composition (wt.%) of Alloy 690 (Heat WP142) in CRDM tubing form ............ 5

Table 4. Welding process and conditions for various weld passes used for fabricating the

A 152 w eld .................................................................................................................................. 6

Table 5. Chemical composition (wt.%) of Alloy 152 weld wire (vendor) and as-deposited

w eld (A N L ) ................................................................................................................................ 6

Table 6. Test conditions for various temperatures and hydrogen concentrations ............................. 13

Table 7. Maximum allowable stress intensity factors for cyclic and constant load testing of

Alloys 690 and 152 weld in 1⁄22-T CT and 1-T CT specimen configurations ....................... 15

Table 8. Tensile properties of Alloy 690TT in the temperature range 25-870'C (77-1598°F) ...... 36

Table 9. Tensile properties of Alloy 152 weld in the temperature range 25-870'C (77-1598°F) .......... 36

Table 10. Crack growth data for specimen A690WC-SL-1 in PWR water at 320'C ......................... 41

Table 11. Crack growth data for specimen A690WC-ST-1 in PWR water at 320'C ............. 47

Table 12. Crack growth data for specimen A690WC-ST-2 of Alloy 690 in PWR water ................... 50

Table 13. Crack growth data for specimen C690-CR- 1 of Alloy 690 in PWR water...................... 54

Table 14. Crack growth data for specimen C690-LR-2 of Alloy 690 in PWR water ..................... ...... 61 61

Table 15. Crack growth data for specimen A952-TS-2 of Alloy 152 weld in PWR water .................. 67

Table 16. Crack growth data for specimen A152-TS-4 of Alloy 152 weld in PWR water ......... 76

xv

Executive Summary

Nickel-base Alloy 600 used as construction material in light water reactors (LWRs) has experienced

stress corrosion cracking (SCC). Such cracking was first observed in steam generator tubes, but it has

also occurred in components such as instrument nozzles and heater thermal sleeves in the pressurizer and

penetrations for control-rod drive mechanisms (CRDM) in the reactor-vessel closure heads. In operating

plants, weld Alloys 82 and 182 have been used with Alloy 600; however, less cracking has been

observed in the weld materials than in the Alloy 600 wrought material. During the last twenty years,

thin-walled steam generator tubing for PWRs has been fabricated from the higher-chromium material,

Alloy 690. Only in the last decade has this material begun to be used for thick-section components,

particularly nozzle penetrations, during the replacement of heads for reactor pressure vessels. This

increased use of Alloy 690 has necessitated the use of the weld Alloys 52 and 152, which also have high

chromium content.

The main objectives of the current work are to determine the SCC and cyclic crack growth rates (CGRs)

in Alloy 690 and its associated weld materials, and to obtain tensile data for these alloys over a wide

temperature range, from operating to extreme conditions. The CGR work is complemented by

metallographic examinations conducted to understand the observed CGR behavior.

The following observations were made from the metallographic examinations. The microstructure of

Alloy 690 plate cold-rolled by 26% contained TiN inclusions and cracked boundary carbides. The areas

close to the two edges of the plate were decorated with carbide streaks - regions with thick grain

boundary carbides. In addition, grain size banding was observed near both the edges. The mid-plane of

the plate - the region where the material for CGR test specimens was obtained - was free of carbide

streaks and grain size banding. The SCC-resistant coincident site lattice boundary (CSLB) fraction was

approximately 28%, which is much less than is typical for a Ni-base alloy in this condition. Hence, the

proportion of SCC-susceptible random boundaries was approximately 50% higher than in a typical alloy.

The microstructure of the second heat of Alloy 690, actually a piece of CRDM tubing, had equiaxed

grains and dense grain boundary carbide coverage. The grain boundary character distribution was

typical of a conventionally fabricated Ni-base alloy component. The CSLB was approximately 53%, of

which a large proportion was twins. The microstructure of Alloy 152 was also found to be typical of a

Ni-base alloy weld. Transmission electron microscopy revealed Al-oxide inclusions. The grain

boundary character distribution was also typical of a Ni-base alloy weld. The CSLB fraction was

approximately 42%, of which a large proportion was low angle boundaries.

The tensile tests, which were conducted in the temperature range 25-870'C (77-1598°F), showed that the

tensile properties of the alloys meet the minimum requirements of ASME SB-167, and that they are in

good agreement with the properties measured elsewhere on similar alloys. As such, the data set for

Alloy 690 is in good agreement with control data on Alloy 600. Similarly, the data set for Alloy 152 is

in good agreement with the Alloy 82/152 tensile data obtained elsewhere.

The cyclic CGRs for Alloy 690 cold-rolled by 26% showed significant environmental enhancement,

comparable to that observed for a heat of Alloy 600 from the Davis-Besse reactor. In the as-received

condition, the cyclic CGRs for the alloy showed some environmental enhancement, although much less

than that for the cold-rolled alloy. Moreover, the cyclic CGRs for the as-received alloy appeared to be

smaller than the corresponding rates for Alloy 600. Hence, it is expected that Alloy 690 in the asreceived condition is more resistant to SCC than Alloy 600.

xvii
The SCC CGRs in the PWR environment at 320'C (608'F) for Alloy 690 cold-rolled by 26% were

1-3 x 10-11 m/s for moderate stress intensity factors, about an order of magnitude lower than those

typical for as-received Alloy 600. The SCC CGRs for as-received Alloy 690 were about an order of

magnitude lower, 2-4 x 10-12 m/s The data obtained at ANL are in good agreement with data obtained on

similar alloys with cold-working in the 20-32% range due to either cold-rolling or tensile pre-straining at

other laboratories. The cracked grain boundary carbides appear to be a common feature of all coldworked alloys that display relatively high CGRs. These results imply that while Alloy 690 in the asreceived condition is more resistant to SCC compared to Alloy 600, cold work can increase SCC

susceptibility.

The available data on the temperature dependence of SCC CGRs obtained at Argonne National

Laboratory and elsewhere were analyzed to investigate the activation energy for SCC growth rates. The

results indicated that the activation energy for Alloy 690 is strongly affected by the deformation level of

the alloy. Alloys in the as-received condition and at deformation levels of up to 20% cold work have an

activation energy close to that for Alloy 600 (130 kJ/mol). At higher deformation levels (24-26%),

CGRs are not strongly dependent on temperature. This observation is consistent with the results of other

testing laboratories, which show that with 50% cold work, the temperature appears to have little effect

on SCC CGRs.

The cyclic CGR data for the two Alloy 152 specimens tested suggest that specific loading conditions can

lead to environmental enhancement. However, the enhanced rates were always lower than the

corresponding rates for Alloys 82/182. Hence, it is expected that Alloy 152 is more resistant to SCC

than Alloys 82/182. The measured SCC CGRs in PWR environment at 320'C (608'F) for as-welded

Alloy 152 were 1.6-5 x 10-11 m/s for moderate stress intensity factors, about an order of magnitude lower

than those typical for Alloy 182.

xviii
Acknowledgments

The authors thank E. Listwan and W. Soppett and their contributions to the experimental effort. This

work is sponsored by the Office of Nuclear Regulatory Research, U.S. Nuclear Regulatory Commission,

under Job Code N6519; Program Managers: Darrell Dunn and Appajosula Rao.

xix
/
Acronyms and Abbreviations

ANL Argonne National Laboratory

ASTM American Society for Testing and Materials

B&W Babcock and Wilcox

BPR Back Pressure Regulator

BWR Boiling Water Reactor

C Circumferential

CGR Crack Growth Rate

CL Constant Load

CRDM Control Rod Drive Mechanism

CSL Coincident Site Lattice

CSLB Coincident Site Lattice Boundary

CT Compact Tension

DO Dissolved Oxygen

ECP Electrochemical Potential

EDS Energy Dispersive X-ray Spectroscopy

EPRI MRP Electric Power Research Institute Materials Reliability Program

GBCD Grain Boundary Character Distribution

GTA Gas Tungsten Arc

GTAW Gas Tungsten Arc Welding

HAB High-Angle Boundary

HAZ Heat Affected Zone

HX Heat Exchanger

ICP-OES Inductively-Coupled Plasma Optical Emission Spectrometry

IG Intergranular

L Longitudinal

LWR Light Water Reactor

MA Mill Annealed

NRC Nuclear Regulatory Commission

OIM Orientation Imaging Microscopy

PWR Pressurized Water Reactor

PWSCC Primary Water Stress Corrosion Cracking

R Radial

RA Reduction in Area

S Side

SA Solution Annealed

SCC Stress Corrosion Cracking

xxi
SEM Scanning Electron Microscopy

SHE Standard Hydrogen Electrode

SMA Shielded Metal Arc

SMAW Shielded Metal Arc Welding

SS Stainless Steel

T Transverse

TC Thermocouple

TEM Transmission Electron Microscopy

TG Transgranular

TT Thermally Treated

UTS Ultimate Tensile Strength

YS Yield Strength

xxii
1 Introduction

Nickel-base Alloy 600 used as construction material in light water reactors (LWRs) has experienced

stress corrosion cracking (SCC). Such cracking was first observed in steam generator tubes, but it has

also occurred in components such as instrument nozzles and heater thermal sleeves in the pressurizer and

penetrations for control-rod drive mechanisms (CRDM) in the reactor-vessel closure heads. In operating

plants, weld Alloys 82 and 182 have been used with Alloy 600; however, less cracking has been observed

in the weld materials than in the Alloy 600 wrought material. 1 During the last twenty years, thin-walled

steam generator tubing for PWRs has been fabricated from the higher-chromium material, Alloy 690.

This material has only in the last decade begun to be used for thick-section components, particularly

nozzle penetrations, during the replacement of heads in reactor pressure vessels. This increased use of

Alloy 690 has necessitated the use of the weld Alloys 52 and 152, which have equivalent high chromium

content.

A program is under way at Argonne National Laboratory (ANL) to evaluate the resistance of Ni alloys

and their welds to environmentally assisted cracking in simulated coolant environments for light water

reactors. To date, CGR tests have been performed on Alloys 600 and 690,2-7 laboratory-prepared deepgroove and double-J welds of Alloy 182,8 Alloy 600 from Davis-Besse CRDM Nozzle #3, Alloy 182

from a J-groove weld in Davis-Besse Nozzle #11, and Alloys 182 and 82 from a hot-leg nozzle-to-pipe

weld in the V. C. Summer reactor coolant system.9 During the last three years, the research has focused

on understanding the cracking behavior of Alloy 690 and its welds.

The main objectives of the current work are to determine the SCC and cyclic crack growth rates (CGRs)

in Alloy 690 and its associated weld materials, and to obtain tensile data for these alloys over a wide

temperature range, from operating to extreme conditions. The CGR work is complemented by

metallographic examinations conducted to understand the observed CGR behavior. Two heats of Alloy

690 were procured, and one Alloy 152 weld was produced specifically for this program. The CGR testing

approach consists of precracking in the test environment, followed by a loading sequence conducted with

the purpose of transitioning from the transgranular (TG) fatigue fracture mode to an intergranular (IG)

SCC fracture mode. In essence, the cyclic CGRs are monitored for environmental enhancement, that is,

for conditions where the measured CGRs are larger than the rates expected in air under the same loading

conditions. An enhanced CGR signifies that in addition to the mechanical component, the measured CGR

also has an environmental component (corrosion fatigue or SCC). The SCC components are isolated by

introducing increasing hold times and by setting the specimen at constant load. In this approach,

precracking in the environment is an important step as these early data points provide a baseline against

which environmental enhancement is subsequently measured.

The complementary metallographic examinations were focused on features known from previous Alloy

600/182 experience to affect the SCC behavior. These examinations consist of scanning electron

microscopy (SEM) for the analysis of elemental distribution and general microstructure, orientation

imaging microscopy (OIM) for the determination of grain boundary characteristics, and transition electron

microscopy (TEM) with emphasis on the nature of grain boundary carbides.

1

2 Experimental

2.1 Materials

The CGR tests were performed on two heats of Alloy 690 and one of Alloy 152 weld. One heat of Alloy

690 was in plate form, and the other was in CRDM tubing form. The Alloy 152 weld was prepared in a

double-J geometry at ANL Central Shops.

Alloy 690 (Heat NX3297HK12) was received from Nuclear Alloy Corp. in a plate form that was 6.4-cm

(2.25-in.) thick x 7.6-cm (3-in.) wide x 86.4-cm (34-in.) long. The designation for the metallurgical

condition of the as-received plate was MIL-DTL-24802. To reach this condition, the alloy was vacuuminduction-melted, electro-slag-remelted, hot-rolled, de-scaled, and annealed at 1038°C (1900 0F) for 2 h,

then air-cooled. The chemical composition provided by the vendor, as well as that determined at ANL by

inductively-coupled plasma optical emission spectrometry (ICP-OES), is reported in Table 1.

Table 1. Chemical composition (wt.%) of Alloy 690 (Heat NX3297HK12) plate.

Alloy ID (Heat) Analysis C Mn Fe S P Si Cu Ni Cr Ti Nb Co

A 690WC (NX3297HK12) Vendor 0.03 0.20 9.9 <0.001 - 0.07 0.01 59.5 29.5 - - -

ANL 0.04 0.33 8.53 0.001 0.003 0.02 0.04 59.67 30.82 0.47 0.01 <0.01

For the CGR tests, the plate was further cold-rolled at room temperature at Special Metals Corp. to reduce

the thickness by 26%. The desired reduction was achieved in three passes (10, 8, and 8%). The test plane

of interest was the mid-plane of the plate because of its uniform distribution of cold work. The

dimensions of this plate and desired orientations necessitated the use of the small

('/-T) compact tension (CT) specimens. Four such specimens were machined in two orientations, side

longitudinal (SL) and side transverse (ST), as shown in Fig. 1. These specimens are designated

A690WC-orientation-number.

A. Figure 1.

Approx 1.5" Orientation of Alloy 690 specimens

Wa L-I• • machined from cold-rolled plate.

3,5"

1 =A690WC-SL-1

2 =A 69OWC-SL-2

3 =A69OWC-ST-1

4 = A69OWC-ST-2

The second heat of Alloy 690 used for CGR testing was received in CRDM tubing form. The piece of

CRDM tubing was marked "2541 - Heat WP142" and was supplied by Valinox Nucleaire (Fig. 2). As

3
with the plate material, the thickness of the tube necessitated the use of the small 1⁄2/ 2-T CT specimens. In

order to obtain a thickness large enough to fabricate these specimens, a ring was cut from the tubing, and

the interior surface was built up with Alloy 152 by the shielded metal arc (SMA) method, as shown in

Fig. 3 and Table 2. Four 1⁄2/ 2-T CT specimens were machined in two orientations, circumferential radial

(CR) and longitudinal radial (LR), as shown in Fig. 2. These specimens are designated C690-orientationnumber.

Figure 2.

Orientation of the Alloy 690

specimens cut from the CRDM

nozzle.

Figure 3.

Schematic of the Alloy 152 weld

build-up on the interior of the

Alloy 690 CRDM tube.

K .375 w

V

A

1/8

4
Table 2. Welding process and conditions used to build up the interior of the Alloy 690 CRDM

tube.

Weld Filler Heat Code Filler Current, Voltage, Travel Speed,

Pass Process Metal Size, in. A V in./min

1-2 SMA Alloy 152 WC43E9 3/32 55-60 21 -23 6-7

(A5.1 IM-97 ENiCrFe-7)

The chemical composition of Alloy 690, determined by an ICP-OES analysis conducted at ANL, is

shown in Table 3.

Table 3. Chemical composition (wt.%) of Alloy 690 (Heat WP142) in CRDM tubing form.

Alloy ID (Heat) Analysis C Mn Fe S P Si Cu Ni Cr Ti Nb Co

C690 (WP142) Vendor 0.018 0.311 10.02 0.0006 0.006 0.296 <0.05 59.71 29.15 0.240 <0.005 0.007

ANL - 0.28 9.26 - <0.05 0.40 <0.05 54.0 27.3 0.3 <0.07 <0.05

The Alloy 152 was fabricated at ANLCentral Shops by shielded metal arc welding (SMAW) in a doubleJ geometry (Fig. 4) following a procedure developed to be consistent with the American Society of

Mechanical Engineers (ASME) Boiler and Pressure Vessel Code, Section IX; however, this procedure

was not qualified by non-destructive examination, tensile tests, and side bend tests. The weld was

produced by joining two Alloy 690 plates (Heat NX3297HK12, described previously) and was prepared

by 96 weld passes. Root passes 1 to 5 were made by gas tungsten arc welding (GTAW) with Alloy 52

filler, and the other passes were made by SMAW with Alloy 152 filler. The conditions for each weld

pass are listed in Table 4. During welding, the maximum inter-pass temperature was 120°C (250'F).

The weld surfaces were cleaned by wire brushing and grinding, and were rinsed with de-mineralized

water or alcohol. The chemical composition of the Alloy 152 weld wire supplied by the vendor and that

obtained by ICP-OES analysis of the as-deposited weld are given in Table 5.

The size of the double-J weld allowed for larger (1-T) CT specimens. These were cut from the bottom

weld shown schematically in Fig. 4 in the transverse side (TS) orientation, with the notch about 5 mm

(0.2 in.) into the weld. This specimen geometry ensured that the CGR tests were conducted in the region

of the weld made with a single Alloy 152 heat, WC04F6. These specimens are designated A152-

orientation-number.

5
A

'2.25ý SO CK,

75,

(a)

3/4 (PASSES 17-51 HEAT WC96D8)

A

(2.25) STOCK

5/8 (PA'

V

1/4 PASSES 1-16 WITH BACý

S - Y Y

SES 51-63 HEAT WC43pX

-A / 45/8 (PASSE

'GOUGE HEAT NX1681JK)

V

S 64-96 HEAT WCO4F6)

A

(b)

Figure 4. Schematic of the weld joint design for the (a) Alloy 152 double-J weld and (b) weld

passes and filler heats used. Dimensions are in inch.

Table 4. Welding process and conditions for various weld passes used for fabricating the Al 52

weld.

Weld Filler Heat Code Filler Current, Voltage, Travel Speed,

Pass Process Metal Size, in. A V in./min

1-16 GTA Alloy 52 NXI681JK 3/32 185-215 21 -22 6-7

(A5.14M-97 ERNiCrFe-7)

17-51 SMA Alloy 152 WC96D8 1/8 100- 110 23-25 7-9

(A5.1 I M-97 ENiCrFe-7)

51 -63 SMA Alloy 152 WC43E9 3/32 55-60 22-23 6-7

(A5. 11 M-97 ENiCrFe-7)

64-96 SMA Alloy 152 WC04F6 1/8 100-110 22-25 7-9

(A5.1 I M-97 ENiCrFe-7)

Table 5. Chemical composition (wt.%) of Alloy 152 weld wire (vendor) and as-deposited weld

(ANL).

Alloy ID (Heat) Analysis C Mn Fe S P Si Cu Ni Cr Ti Nb Co

A52 (NX168IJK) Vendor 0.02 0.24 10.40 <0.001 0.005 0.12 0.01 59.30 28.68 0.52 <0.01 0.00

A152 (WC96D8) Vendor 0.04 3.88 9.16 0.007 0.008 0.54 0.01 55.88 29.35 0.07 1.81 0.01

A152 (WC43E9) Vendor <0.03 3.95 10.25 <0.003 <0.005 0.51 0.01 54.16 28.60 0.08 1.93 0.01

A152 (WC04F6) Vendor 0.048 3.48 10.39 0.003 0.003 0.41 <0.01 55.20 28.70 0.09 1.54 <0.005

ANL - 3.88 9.56 - <0.08 0.52 <0.04 53.70 28.40 0.10 1.80 <0.04

6
2.2 Specimen Design

All crack growth rate tests on Ni alloys and Ni-base welds were conducted in simulated PWR

environments at temperatures of 290-350'C in accordance with ASTM E-647, "Standard Test Method for

Measurement of Fatigue Crack Growth Rates," 10 and ASTM E-1681, "Standard Test Method for

Determining a Threshold Stress Intensity Factor for Environment-Assisted Cracking of Metallic Materials

under Constant Load."'11 Depending on the dimensions of the available materials, the tests were

performed on either 1-T or 1⁄2-T CT specimens; the configurations of the CT specimens are shown in

Fig. 5.

(a)

DETAIL 'MV

M R

(b)

Figure 5. Configuration of the (a) 1-T and (b) 1

/2-T CT specimens used for this study.

Dimensions are in mm.

2.3 Test Facilities

The CGR tests were conducted in three test facilities - two equipped with 6-liter stainless steel (SS)

autoclaves and one with a 2-liter SS autoclave. Each system has a suite of calibrated instrumentation,

including digitally controlled hydraulic loading and load cells, and an independent water loop with water

chemistry monitoring. Also, each system has a DC potential CGR measuring system.

7
2.3.1 Facility with 6-Liter Autoclave

/

The 6 -liter autoclave facility is typically used for CGR tests at temperatures not exceeding 32 0 'C

(608'F). The facility consists of a closed-loop electro-hydraulic material test system equipped with an

extended column load frame rated at 89 kN (20,000 lb) maximum. The Type 316 SS autoclave has a 175-

mm (6.875-in.) outside diameter and is rated for a working pressure of 5050 psig (35 MPa) at 343°C

(650'F). The load frame is accessorized with an Instron Model 8800 control console and a hydraulic

pump. The autoclave is connected to a recirculating water system. The test facility also consists of a

temperature control unit, a DC potential control console, a DC potential measurement unit, and a data

acquisition and recording system. The autoclave, mounted within the load frame, has been modified to

permit a z19-mm (0.75-in.) dia shaft to load the test specimen through a Bal-Seal® gland in the top of the

autoclave cover.

The autoclave is continuously supplied with simulated primary water solution from a feedwater tank.

Figure 6 shows a schematic diagram of the water system. The water system consists of a feedwater

storage tank, high-pressure pump, regenerative heat exchanger, autoclave preheater, test autoclave,

electrochemical potential (ECP) cell, regenerative heat exchanger, back-pressure regulator, and return line

to the feedwater tank. The storage tank has a hydrogen cover gas to maintain a desired dissolved

hydrogen concentration in the water. In the once-through mode, the return line is connected to the drain.

Water is circulated at relatively low flow rates (15-25 mL/min).

,27 26

5

6

22

124

15

12

32

V13

Figure 6. Schematic diagram of the recirculating 6-liter autoclave system.

8
1. COVER GASS SUPPLY TANK 20. HEAT EXCHANGER (HX)

2. TWO-STAGE HIGH-PRESSURE REGULATOR 21. DRAIN

3. LOW-PRESSURE REGULATOR 22. HX FEEDWATER OUTLET THERMOCOUPLE (TC)

4. FLOW METER 23. AUTOCLAVE & ECP CELL PREHEATER

5. COMPOUND VACUUM & PRESSURE GAUGE 24. HEAT EXCHANGER

6. PRESSURE RELIEF VALVE 25. COMMERCIAL AUTOCLAVE

7. VENT TO AIR & FLASH ARRESTOR 26. THERMOWELL

8. FEEDWATER STORAGE TANK 27. "BAL SEAL" RETAINER

9. SPARGE TUBE 28. ECP CELL

10. FEEDWATER FILL PORT 29. HX EFFLUENT INLET TC

11. WATER SAMPLE PORT 30. ECP CELL INLET TC

2. SYSTEM BLEED PORT 31. AIR-COOLED COIL

.3. SOLENOID VALVE 32. BACK-PRESSURE REGULATOR (BPR)INLET TC

14. HIGH-PRESSURE PUMP 33. BPR

15. CHECK VALVE 34. PRESSURE RELIEF VALVE

16. PRESSURE TRANSDUCER 35. CONDUCTIVITY METER

17. RUPTURE DISC 36. PH METER

r8. ACCUMULATOR 37. ION EXCHANGE BED

19. HIGH-PRESSURE GAUGE 38. ION EXCHANGE BED BYPASS LINE

Figure 6. (cont.)

2.3.2 Facility with a 2-Liter Autoclave

The 2-liter autoclave test facility was built by Cortest Inc. and allows test temperatures of up to 350'C

(662°F). The servohydraulic test frame consists of a load cell, load train, autoclave, and furnace. The

hydraulic actuator is mounted on bottom of the test frame, with the test train components suspended

above it. The load cell is located at the bottom of the pull rod. An Instron Model 8800 control console is

used to load the specimen. A 2-liter autoclave is installed inside the furnace for conducting tests in

simulated LWR environments. The heater bands are mounted on the autoclave vessel. Figure 7 is a

photograph showing the entire test system.

Figure 7. Layout of the 2-liter autoclave system.

9
The load cage that contains the test specimen consists of the cover plate of the autoclave and a thick

bottom plate separated by four compression rods (Fig. 8). The lower two-piece clevis assembly is

fastened to the bottom plate of the cage, and the upper piece is connected to the pull rod. A 1⁄2/ 2-T CT or

dog-bone tensile specimen can be mounted between the clevises. The specimen and clevises are kept

electrically insulated from the load train by using oxidized Zircaloy pins and mica washers to connect the

clevises to the rest of the load train. The crack length is monitored by the same method as for the

previous two facilities. Water is circulated through a port in the autoclave body, which serves both as

inlet and outlet. A schematic diagram of the recirculating water system is shown in Fig. 9.

Figure 8.

Photograph of the specimen load train for the

2-liter autoclave.

28

Figure 9. Schematic diagram of the recirculating 2-liter autoclave system.

10
1. COVER GASS SUPPLY TANK 19. ACCUMULATOR

2. TWO-STAGE HIGH-PRESSURE REGULATOR 20. RUPTURE DISC

3. FLASH ARRESTOR 21. HEAT EXCHANGER (HX)

4. LOW-PRESSURE REGULATOR 22. DRAIN

5. FLOW METER 23. SYSTEM BLEED PORT

6. CHECK VALVE 24. HEAT EXCHANGER OUTLET TC

7. COMPOUND VACUUM & PRESSURE GAUGE 25. AUTOCLAVE PREHEATER

8. PRESSURE RELIEF VALVE 26. PREHEATER OUTLET TC

9. VENT TO AIR & FLASH ARRESTOR 27. COMMERCIAL AUTOCLAVE

10. FEEDWATER STORAGE TANK 28. THERMOWELL

11. SPARGE TUBE 29. "BAL SEAL" RETAINER

12. WATER SAMPLE PORT 30. ECP CELL

13. FEEDWATER FILL PORT 31. AIR-COOLED COIL

14. FEEDWATER TANK RECIRCULATION PUMP 32. WATER COOLED HEAT EXCHANGER

15. SOLENOID VALVE 33. BACK-PRESSURE REGULATOR (BPR) INLET TC

16. HIGH-PRESSURE PUMP 34. BPR

17. PRESSURE TRANSDUCER 35. PH METER

18. HIGH-PRESSURE GAUGE 36. CONDUCTIVITY METER

Figure 9. (cont.)

2.3.3 Primary Water Environment

The simulated PWR feedwater contains 2 ppm Li as LiOH, 1000 ppm B as HBO 3, z2 ppm dissolved

hydrogen (z23 cm3/kg), and less than 10 ppb dissolved oxygen (DO). 12 It is prepared from the laboratory

supply of deionized water by first passing this water through a local filtration system that includes a

carbon filter, an Organex-Q filter, two ion exchangers, and a 0.2-mm capsule filter. The DO in the

deionized water is reduced to <10 ppb by bubbling/sparging pure H2 through the water. To speed

deoxygenation of a fresh tank of water, a vacuum may be applied to the feedwater tank at the vent port.

The PWR water is prepared by dissolving boric acid and lithium hydroxide in 20 L of deionized water

before adding the solution to the supply tank. The dissolved hydrogen in the water is calculated from the

hydrogen pressure and temperature of the supply tank.

Water samples are taken periodically to measure pH, resistivity, and DO concentration both upstream and

downstream from the autoclave. An Orbisphere meter and CHEMetricsTM ampoules are used to measure

the DO concentrations in the supply and effluent water. The redox and open-circuit corrosion potentials

are monitored at the autoclave outlet by measuring the ECPs of platinum and an Alloy 600 electrode,

respectively, against a 0.1 M KC1I/AgCI/Ag external (cold) reference electrode. The measured ECPs,

E(meas) (mV), were converted to the standard hydrogen electrode (SHE) scale,

E(SHE) (mV), by the polynomial expression1 3

E(meas) + 286.637 - 1.0032. (AT) + 1.7447 x 10-4. (AT) 2 - 3.03004 x 10-6 -(AT) 3 , (1)

where AT(0C) is the temperature difference of the salt bridge in a 0.1 M KC1/AgCI/Ag external reference

electrode (i.e., the test temperature minus ambient temperature).

11
The effect of the proximity of the corrosion potential for the SCC test to the potential at the Ni/NiO phase

transition on crack growth rates is well established. 14-

16 While the mechanistic origin is unclear, there

are consistent data showing that CGRs of Ni-base alloys peak near the Ni/NiO phase boundary.14-16

Thus, to isolate the effect of temperature, the tests are conducted so that the difference between the

corrosion potential of the specimen and that of the Ni/NiO phase transition was maintained approximately

constant. The Ni/NiO phase transition is very close to the H2/H20 phase transition, and the latter controls

the corrosion potential of the simulated primary water environment. As such, the corrosion potential

difference reduces to the difference of two hydrogen electrodes: 16

AE = Etest - ENi/NiO = - t, [ H2]tt i (2)

2F ý[H2]Nih

where: Etest = potential at test temperature T,

ENi/NiO = potential of the Ni/NiO transition at test temperature T,

R = universal gas constant

= 8.314 x 10-3 kJ/mol K (1.103 x 10-3 kcal/mol'R),

F = Faraday constant

= 96485 C/mol,

[H2]test = hydrogen concentration at test conditions, and

[H2]Ni/NiO = hydrogen concentration at the Ni/NiO transition.

For two tests, conducted at temperatures T1 and T2 and satisfying the constant corrosion potential

difference to the Ni/NiO line, Eq. 2 transforms into the following relationship between hydrogen

concentrations and test temperatures:

RT2ln [H 2]test 2 l [H2]test]

2F [H2]Ni/NiO test2 ) _, 2F [H2]Ni/NiO-testl (3)

where: [H211.t2 = hydrogen concentration at test temperature T2,

[H2]testl = hydrogen concentration at test temperature T1,

[H2]Ni/NiO-tcst2 = hydrogen concentration at Ni/NiO transition for test 2, and

[H2]NMiO-test1 = hydrogen concentration at Ni/NiO transition for test 1.

The hydrogen concentration at the Ni/NiO transition as a function of temperature was determined by

fitting Attanasio et al. 15 data with a fourth-order polynomial (Fig. 10). The best fit is given by the

following expression:

[H2]Ni/NiO (T)=-8783.1+53.803-T-0.11961.T 2+0.00011238-T 3 -3.628899x10- 8 .T4 (4)

12
30.0 .l . . l . ' . l. . . .1 . . . . .

C

0

C

Cu

0

2

U,

Cu

C

0

Cu

C

0

C.)

C

0

U

I

25.0

20.0

15.0

10.0

5.0

y=- 8783-1 + 53.803x - 011961x2+

+ 0.0001123803 - 3,6288E_8X4

Figure 10.

Hydrogen concentration at NI/NiO

phase transition as a function of

temperature (data taken from Ref. 15).

0.0 r

520

I ý , 1 l I I l I l . . I

0 540 560 580

T (K)

600 620 640

Equation 3 can be solved for the hydrogen concentration at test temperature T2. This concentration

maintains the potential difference to the Ni/NiO line similar to that of a test conducted at temperature Ti.

For the purpose of this calculation, test conditions "1" are considered those typical of tests conducted in a

PWR environment at 320'C (6080F, 593 K), with [H 2 ]tsl = 23 cm 3/kg. 12 The results for several

temperatures of interest are shown in Table 6. For each temperature, the hydrogen concentration at the

Ni!NiO transition was calculated by Eq. 4, then the concentration needed to maintain the corrosion

potential difference constant was calculated by Eq. 3.

Table 6. Test conditions for various temperatures and hydrogen concentrations.

Test Temperature [H2]NiNiO, [H 2

]t~t, AE,

°C (-F) K cm 3

/kg cm 3 /kg mV

290(554) 563 3.0 11.01 31.5

300(572) 573 3.9 13.81 31.3

305(581) 578 4.5 15.60 31.1

310 (590) 583 5.2 17.71 30.8

320 (608) 593 7.1 23.00 30.1

350 (662) 623 17.8 47,59 26.4

13
2.4 Test Procedure

The CGR tests are conducted in the load-control mode using a triangular, sawtooth, or trapezoidal

waveform with load ratio R of 0.3-0.7. The CT specimens are fatigue precracked with a triangular

waveform in the test environment at load ratio R = 0.3, frequency zl Hz, and maximum stress intensity

factor (Kmax) of 20-25 MPa-m112 (18-23 ksi-inl/ 2). Afterz0.5 -mm (20 mils) extension, R is increased

incrementally to 0.5-0.7, and the loading waveform was changed to a slow/fast sawtooth with rise times

of 30-1000 s and unload time of 12 s. This loading sequence is considered to result in reproducible

CGRs. 17 For sawtooth loading, the CGRs are calculated on the basis of the rise time. During individual

test periods, Kmax was maintained approximately constant by periodic load shedding (less than 2%

decrease in load at any given time).

Crack extensions are monitored by the reversing-DC potential difference method. The current leads are

attached to the holes on the top and bottom surfaces of the specimen (Fig. 5), and potential leads are

either welded on the front face of the specimen across the machined notch but on diagonal ends (1-TCT

specimens, Fig. 5a) or attached to pre-machined taps (1⁄22-T CT specimens, Fig. 5b). Also, to compensate

for the effects of changes in resistivity of the material with time, an internal reference bar of the same

material being tested is installed near the test specimen. The CT specimen and reference bar are

connected in series, and the DC potential across the specimen as well as the reference bar is monitored

continuously during the test. The results for the reference bar are used to normalize potential drop

measurements for the CT test specimen.

During crack growth tests in high-temperature water, environmental enhancement of CGRs does not

occur from the start of the test. Under more rapid cyclic loading typically applied during precracking and

initial periods of cyclic loading the crack growth is dominated by mechanical fatigue. In general,

environmental enhancement is typically observed under loading conditions that would lead to CGRs

between 10-11 and 10-9 m/s in air.

The stress intensity factor range (AK) was calculated in accordance with ASTM E 168111 and E 64710 as

follows:

AK = AP W)± (5)a-)

AP = P. - P.mn for R > 0, (6)

f(-ha = 0.886 + 4.64(2) - 13.32(L)2 + 14.72 - 5-6( -)4 (7)

where Pmax and Pmin are maximum and minimum applied load, a is crack length, W is the specimen

width, and the effective thickness Beff = (B BN) 0 .

5 . The applied K and crack size limits for the tests met

the ASTM E 1681 and E 647 criteria. These criteria are intended to ensure applicability' and

transferability of the cracking behavior of a component or specimen of a given thickness under a specific

loading condition to a crack associated with a different geometry, thickness, and loading condition. The

14
K/size criteria require that the plastic zone at the tip of a crack be small relative to the specimen geometry.

For constant load tests, ASTM E 168111 requires that

Beff and (W - a) >2.5 (K!ays) 2 , (8)

and for cyclic loading, ASTM 64710 requires that

(W - a) ?(4/nt) (K/ays) 2, (9)

where K is the applied stress intensity factor, and ays is the yield stress of the material. In

high-temperature water, in circumstances where environmental enhancement occurs, and corrosion

fatigue and SCC overtake mechanical fatigue as the primary mechanism for crack growth, Eq. 8 is

probably the more appropriate criterion. For high-strain hardening materials (i.e., ultimate-to-yield stress

ratio ault/ays >1.3), both criteria allow the use of the flow stress defined as caf = (ault + cys)/2 rather than

the yield stress. Table 7 provides the maximum allowable stress' intensity factors meeting the criteria set

by Eqs. 8 and 9 for cyclic and constant load testing of Alloy 690 and 152 in the as-received and the aswelded condition, respectively. Flow stresses used for the calculations are 402 MPa (58 ksi) for Alloy

690 and 474 MPa (68.7 ksi) for Alloy 152 weld (detailed tensile property data are presented in Section 4.1

of this report). Both '/2-T CT and 1-T CT specimens (Fig. 5) are considered. In addition, typical total

crack extensions of 1.5 mm (0.06 in.) for a 1

/2-T CT and 10 mm (0.4 in.) for a 1-T CT are assumed. As

such, the table shows that in a case of a test involving a 1⁄2-T CT specimen of Alloy 690 in the as-received

condition, a stress intensity factor as high as 39 MPa-m1 /

2 (35.5 ksi-in. 1/

2) can be employed to fatigue

precrack the specimen at the beginning of the test. At the end of the test, after an assumed total crack

advance of 1.5 mm (0.06 in.), constant load with a stress intensity factor not exceeding 26 MPa m1/ 2 (24

ksi.in.1/ 2) may be used. The table also shows that 1-T CT specimens are more versatile in both crack

advance and sampling of microstructure, as well as the range of stress intensity factors that may be used.

However, such specimens are not alway's possible due to the physical size of the material available.

Table 7. Maximum allowable stress intensity factors for cyclic and constant load testing of

Alloys 690 and 152 weld in 1

/2-T CT and 1-T CT specimen configurations.

Alloy 690

Kcyclic, Kscc,

Specimen type a, mm MPa.m12 MPa.m"/2

(in.) (ksi-in. 1

1 2

) (ksi'in. 11

2 )

1

/-T CT 12 (0.47) 39.1 (35.6) 27.9 (25.4)

13.5 (0.53) 36.6 (33.3) 26.1 (23.8)

1-T CT 12.7 (0.50) 69.6 (63.3) 49.6 (45.1)

22.7 (0.89) 59.8 (54.4) 42.6 (38.8)

Alloy 152

K-cyclic, Kscc,

Specimen type a, mm MPa.m"2 MPa-m12

(in.) (ksi-in .1/2) (ksi-in .l/2)

1/2-T CT 12 (0.47) 39.1 (35.6) 27.9 (25.4)

13.5 (0.53) 36.6 (33.3) 26.1 (23.8)

1-T CT 12.7 (0.50) 69.6 (33.3) 49.6 (23.8)

22.7 (0.89) 59.8 (54.4) 42.6 (38.8)

15
After each test is completed, the specimen is fractured in liquid nitrogen, and the fracture surfaces are

examined by optical and electron microscopy to measure the final crack length using the 9/8 averaging

technique; that is, the two near-surface measurements are averaged, and the resultant value is averaged

with the remaining seven measurements.

2.5 Analysis of Crack Growth Rate Data

The experimental approach relies on monitoring the cyclic rates for environmental enhancement. As

such, only those conditions that show environmental enhancement are chosen for transitioning from TG

fatigue to IG SCC. Under cyclic loading, the CGR (m/s) can be expressed as the superposition of the rate

in air (i.e., mechanical fatigue) and the rates due to corrosion fatigue and SCC (aCF and ýScc,

respectively), given as 2,4

aenv =a air +aCF + ýSCC" (10)

The cyclic CGRs for Ni alloys and welds in air were determined from correlations developed earlier at

ANL: 2,4,7,9

Alair = d-----/tr =+.C(1-0.82.R y2.2 . (AK )4"l1]/tr, (1

where da/dN is the growth rate per cycle, tr is the rise time for the loading cycle, R is the load ratio (i.e.,

ratio of the minimum and maximum stress intensity factors Kmin/Kmax), AK is Kmax - Kmin in MPa-m1

/

2,

and the constant C depends on the material and temperature. For Alloy 600 the constant (CA600) is a

third-order polynomial with respect to temperature T (°C), 2,4

CA600 = 4.835x 10-14_ (1.622 x 10-17)T+(1.490x10-18)T2 -(4.355x10-21)T. (12)

In LWR coolant environments, the CGRs of Alloy 600 show frequency-dependent enhancement under

cyclic loading conditions. In high-DO water [i.e., normal water chemistry in boiling water reactor

(BWR)], the environmental enhancement of the growth rates does not appear to depend on the material

composition (e.g., C content) or material heat treatment. In contrast, environmental enhancement of

CGRs of Alloy 600 in low-DO water does seem to be strongly dependent on material conditions. In the

literature,18- 22 such variability has been attributed to thermomechanically controlled parameters, such as

yield strength and grain boundary coverage of carbides, although the evidence for this dependence is

more substantial for steam generator tubing than for structural components.

For Alloy 690, the constant CA690 is given by a third-order polynomial of temperature T (°C), expressed

as:

2

CA690 - 5.423 x 10-'4+ (1.83 x 10-16)T - (1.725 x 10-'8 )T2 + (5.490 x 10-21 )T3 . (13)

For Ni-alloy welds, the constant CNiweld is given by a fourth-order polynomial with respect to temperature

T (°C), expressed as: 9

CNiweld = 8.659 x 10-1- (5.272 x 10- 17 )T + (2.129 x 10-"8 )T2 (14)

- (1.965 x 10-2 )T3+ (6.038 x 10-23)T4

16
In earlier ANL work, correlations were developed to estimate the enhancement of cyclic CGRs in LWR

environments relative to the CGRs in air under the same loading conditions. In the absence of any

significant contribution of SCC to growth rate, the cyclic CGRs for Alloy 600, either in the solution

annealed (SA) condition or the SA plus thermally treated (TT) condition, in z300 ppb DO water at 289'C

are given by the expression4

AenvA600 =a air,A600 +4.4x10- 7 " (airA600 33 (15)

In low-DO environments (e.g., hydrogen water chemistry for the BWR or PWR environment) at 320'C

(608'F), some alloys show little enhancement, while others show enhancement comparable to that

predicted by Eq. 15.

Similarly, correlations describing the cyclic CGRs of Ni-alloy welds (e.g., Alloys 182 and 82) have been

developed. 9 Under similar loading conditions, the CGRs of Ni-alloy welds are a factor of 2-3 higher than

those of Alloy 600. The analysis 9 indicated that the cyclic CGRs of Ni-alloy welds in PWR water are

represented by the expression

aenvNi-weld adairNi-weld + 0.018- (dair,Ni-weld )0.78. (16)

Figure 11 summarizes the expected behaviors for cyclic rates in air and environment for Alloys 600, 690

and Ni-base welds as described by Eqs. 11-16. All data is plotted as a function of the predicted behavior

for Alloy 600 in air.

10-7

,CGR env Ni-weld , .',: , ,

10-8 ........ ................ ........... ,;...... .... . '. ----........

CGR env 600 Figure 11.

Crack growth rates in air and

E ---------- environment for Alloys 600, 690 and Ni-

•> base welds as a function of predicted

0 rates in air for Alloy 600 (data taken 10-1• 6..--•'•;':'" ?from Refs. 4,7,9).

S 'a' eld

..... CGR env 600 (NUREG-4667)

10-11 - .... --- CGR air 690 (NUREG-6383)

CGR air weld (NUREG-6921)

-----. CGR env weld (NUREG-6921)

10-11 10-10 10-9 10-8 10-7

CGRair (m/s)

17
The SCC growth rate data for Alloy 600 and its weld metals have been reviewed by White and Hickling

(MRP-5523 and MRP-1 1524) to determine the effects of critical parameters such as stress intensity factor,

temperature, material heat treatment, cold work, and water chemistry on growth rates. For Alloy 600, the

CGR (m/s) under SCC conditions is represented by the expression, 23

tA600 = exp[--Q(1-- I j(K -Kth)0, (17)

where: Q = activation energy for crack growth

= 130 kJ/mol (31.1 kcal/mol) for Alloy 600,

R = universal gas constant

= 8.314 x 10-3 kJ/mol K (1.103 x 10-3 kcal/mol.°R),

T = absolute operating temperature in K (or 'R),

Tref = absolute reference temperature used to normalize the CGR data

598 K (1076.67 0

R),

(X -- crack growth amplitude (2.67 x 10-12 at 3250C),

K crack tip stress intensity factor (MPa-ml/ 2),

Kth = crack tip stress intensity factor threshold (9 MPa'm1/

2, 8.2 ksi-in.1 /2) and

0 = exponent 1.16.

The effect of K on the SCC growth rate for Ni-alloy welds in PWR environments has been represented by

a modified24 version of the above relationship. Unlike the CGR relationship for Alloy 600, the

relationship for Ni-alloy welds has no threshold value for the stress intensity factor K,

aNiweld = a exp[- f ]KO, (18)

where Q, R, T, and Tref are the same as in Eq. 17; the crack growth amplitude a is 1.5 x 10-12 at 325 0C;

and exponent 0 is 1.6. Also, unlike Alloy 600, for which a reliable value for the activation energy for

crack growth was determined, no such number is available for the Ni-weld alloys. Thus, for Ni-weld

alloys, the activation energy is assumed to be the. same as that for Alloy 600.

18
3 Microstructural Characterization

The alloys used in the present project were subjected to an extensive metallographic analysis. This

section presents the findings for Alloy 690 in plate form and from CRDM nozzle tubing and the

laboratory-prepared Alloy 152 weld. The characterization was focused on features known to affect the IG

SCC behavior. As such, special attention was paid to grain boundary carbides and grain boundary

character distribution. Grain boundary carbides have a profound impact on the SCC behavior of Alloy

600, affecting both crack initiation and propagation. Several early studies reported correlations between

high grain boundary carbide coverage and increased cracking resistance, 25-

26 and subsequent work has

isolated and established the beneficial effect of carbides on SCC behavior. 27-

28 Likewise, grain boundary

character distribution has been shown to affect crack initiation in several austenitic alloys in high

temperature water environments. 29-

36 As such, grain boundaries in special orientation relationships have

been found to be more resistant to SCC initiation than random boundaries. 30-

31 Moreover, SCC was

found to propagate preferentially along random boundaries in Alloys X-750 35 and 182 weld.8

3.1 Alloy 690 in Plate Form (A690WC)

The metallographic analysis of Alloy 690 was focused on the 26% cold-rolled condition - as this was the

form in which this alloy was tested. The analysis included measurements of hardness along the sample

thickness, as well as SEM/OIM and TEM analyses, and focused on the grain boundary carbides and, more

recently, on the existence of banding in this alloy.

3.1.1 Hardness Testing

The cold-work distribution along the thickness in the 26% cold-rolled plate was evaluated by measuring

the microhardness at three positions (a, b, and c, as shown in Fig. 12a).

32.0 . .. VI ,

- ANL

24.0 * UW ANL 26%CR

* GECRD, 20%CR

16.0 -------------

8& 00 ---- -- .. -- . . -. . . . . . . ........ _ 1 .......... . . .

8.0

Rolling direcfion 8 0.00 ...... .................. .

0

a.fl -8.00

-160 ........... . A

. -•. . .. ........... ..... ..... . ..... c.. .....' = - .....

-240. .... .. - - -------- - -- .-------

-32.0 , , , , , , , , i ,

260 280 300 320 340

Hardness HV

(a) (b)

Figure 12. (a) Schematic of the rolling process of the ANL plate identifying the locations (a, b,

and c) at which hardness was evaluated, and (b) hardness along thickness of the

Alloy 690 plate cold-rolled by 26% (ANL). Similar data from a plate cold-rolled by

20% at GE are included for comparison.

19
The hardness measurements were performed by Vickers indentation with a load of 1 kg-f and load time of

15 s at the three positions. The results from two sets of measurements (by ANL and University of

Wisconsin) are shown in Fig. 12b. For comparison, the results obtained on a different plate (heat B25K)

cold-rolled by 20% at GE to a final thickness of 67 mm (2.6 in.) are also included. The data show the

same general trend: the mid-plane has a 10-12% lower hardness compared with that for the positions that

are closer to the rolling surfaces. Also, because the GE plate was initially thicker and cold-rolled by 20%

(vs. 26% for the ANL plate), it had a lower hardness at the mid-plane.

3.1.2 General Microstructure - Grain Boundary Carbides

The microstructure of the Alloy 690 plate that was cold-rolled by 26% was examined by SEM and TEM.

Figure 13 shows that the microstructure of the mid-plane has grain boundary carbides and matrix TiN

precipitates. Surprisingly, the grains do not appear elongated in these pictures. Figure 14 shows TEM

images from specimens prepared from the mid-plane of the plate. Figures 14a and 14b show a highly

deformed microstructure, as expected for cold-worked material. Figures 14c and 14d also show several

grain boundary carbides, as indicated by the white arrows. Other than the absence of elongated grains,

the microstructure of this material did not appear to be unusual.

tap, (b)

Figure 13.

Alloy 690 cold-rolled by 26%: (a) general

microstructure, (b) grain boundary carbides, and

(c) TiN precipitates. Images were obtained on

the mid-plane of the deformed plate.

kc)

20
A_ X

(c) (d)

Figure 14. TEM micrographs from mid (rolling plane) Alloy 690 deformed by 26% showing the

following: (a,b) high deformation microstructure and (c,d) grain boundary carbides.

A second round of TEM examinations was prompted by the finding that the fracture surfaces of alloys

tested at both ANL and Bettis Atomic Power Laboratory (Bettis) 37 appeared to contain cracked grain

boundary carbides. Figure 15 shows images of carbides obtained from the mid-plane of the cold-rolled

plate. Some cracked grain boundary carbides are apparent. Figure 16 shows similar images obtained

from the mid-plane of a 20% cold-rolled Alloy 690 plate provided by GE. The carbides appear somewhat

different in shape (compare Fig. 15a and Fig. 16a); however, cracked grain boundary carbides were also

found in this alloy (Fig. 16b).

21
(a) Mo)

Figure 15. Alloy 690 cold-rolled by 26%, which was CGR tested at ANL.

(a)

Figure 16. Alloy 690 cold-rolled by 20%, which was provided by GE.

3.1.3 Grain Boundary Character Distribution

The grain boundary character distribution (GBCD) is essentially a classification of boundaries according

to the coincident site lattice model as either coincident site lattice boundaries (CSLBs) or high-angle

boundaries (HABs). The CSLBs are formed when the neighboring grains are in specific orientation

relationships, and they are described in terms of the fraction 1/1 of coincident atomic lattice sites. A

22
boundary that contains a high density of coincident lattice points is expected to have low energy because

of a good atomic fit. These boundaries have been shown to possess an increased resistance to IG SCC

over HABs. 29 -

3 6 The GBCD is obtained by OIM.

Figures 17 and 18 summarize the findings of the OIM analysis for the Alloy 690 in plate form in the asreceived and 26% cold-rolled condition. Figure 17a shows a typical OIM map with equiaxed grains. The

resulting average grain size is 36 gtm (Fig. 17b), and the CSLB fraction is 27.5%, of which a large

proportion is twins (1=3). The latter result was expected; however, the former indicates a relatively

small CSLB fraction for this alloy. The typical fraction for a Ni-base alloy in the as-received solution

annealed condition is approximately 50%.31 Similar results were obtained for the alloy in the 26% coldrolled condition (Fig. 18). Again, the relatively small proportion of CSLBs should be noted.

0.2

0

.o

I0.1

0.0

I 10

Grain Size Diameter (w±n)

(b)

100

0

LL

SZ

CM

(.)

Alloy 69OWC 1

0.2

0.1

0.0

N% 1 b 11 N S,% S1% WNS"MIS& bkb0N

CSL E

(a) (c)

Figure 17. (a) OIM map, (b) graph of grain size distribution, and (c) graph of grain boundary

character distribution for Alloy 690 in plate form in the as-received condition.

23
0.20

0.15

0o

0.10 ULL

0.05

0.00

1 10 100

Grain Size Diameter (pam)

(b)

0.3

A690 26%CW

Sum(3_<•<__) = 28.0%

C 1= 0%

=0.2

U

EU,

.k 0.1

Om1

Co

0.c ,-r • T~_ ,• - • . . .

2CSL E

(a) (c)

Figure 18. (a) OIM map, (b) graph of grain size distribution, and (c) graph of grain boundary

character distribution for Alloy 690 in plate form in the 26% cold-rolled condition.

3.1.4 Carbide and Grain Size Banding

Carbide and grain size banding have been advanced by Andresen et al.38 as a possible cause for the

relatively high CGRs of several heats of Alloy 690 measured in the cold-rolled condition. For this

investigation, the Alloy 690 plate at ANL (Heat NX3297HK12) was examined not only in the mid-plane

(where the CGR specimens were cut), but also along the thickness. Both the as-received and the 26%

cold-rolled conditions were examined. For the examination, specimens #1, 5, 6, and 8 were cut from the

as-received Alloy 690 plate (Fig. 19). These specimens were ground to a 1 -Pm surface finish,

electropolished, and etched with Marble's reagent.

24
Figure 19.

Schematic showing the positions of the

metallography specimens cut from the

plate in the as-received condition.

11 12 Edge 4

Ed e3

Edge 2 Center

Figure 20 shows the images obtained from specimen #1. This specimen was decorated with short streaks,

which were further revealed as areas with thick grain boundary carbides.

#1 Edge I

Center

Figure 20. Microstructure of specimen #1 showing carbide streaks.

Grain size banding was found in specimens #5 and #6 close to edges I and 2, respectively (Fig. 21).

Specimen #8, taken from the middle of the plate, showed neither carbide streaks nor grain size banding

(Fig. 22).

25
0.5MM

(a) (D)

Figure 21. Grain size banding in (a) Specimen #5 near edge 1 and (b) Specimen #6 near edge 2

(images rotated by 90°C vs. Fig. 19).

0.5mm 0.5nim

(a) (b)

Figure 22. Microstructure observed in Specimen #8: (a) near edge 4 and (b) middle of the

specimen (images rotated by 900C vs. Fig. 19).

The microstructure of Alloy 690 cold-rolled by 26% is shown in Figs. 23 and 24. The areas close to

edges 1 and 2 (Fig. 19) are decorated with carbide streaks (Fig. 24), while the center is nearly free from

the streaks. In addition, grain size banding was observed near the same edges, as illustrated in Fig. 24.

26
0.m

Figure 23.

Carbide streaks in Alloy 690 plate

cold-rolled by 26%.

/ V

Figure 24. Grain microstructure in Alloy 690 plate cold-rolled by 26%.

In summary, carbide streaks and grain size banding occurred near edges 1 and 2 (Fig. 19) of the Alloy

690 plate. The mid-plane appears to be free of carbide streaks and bands. These features were preserved

by cold-rolling.

27
3.2 Alloy 690 TT from CRDM Tubing (C690)

The Alloy 690TT was examined by SEM, TEM, and OIM. The objectives were to observe the

microstructure, focusing on grain boundary carbides, and to determine the GBCD. As noted in the

previous section, both of these features are known to influence the SCC behavior.

Figure 25 shows an SEM micrograph of the Alloy 690TT microstructure, consisting of equiaxed grains

having an average size of 70 gtm. Figure 26 shows a TEM image of the microstructure. This alloy has a

moderate dislocation density.

Figure 25.

Microstructure of Alloy 690 TT

from CRDM tubing.

I 0Oin EHT =10.00 WV Signal A S E2 Dafte A2 Apr 2008

Mg 10X WD= 12ramn Photo No. 275 Thne:17:0:35

Figure 26.

Microstructure of Alloy 690 TT from

-CRDM tubing.

28
Figure 27 shows additional TEM images, with a focus on the grain boundary carbides. The alloy has a

dense carbide decoration on the grain boundaries. Hence, it is expected to possess superior SCC

resistance.

Figure 27.

Microstructure of Alloy 690 TT from CRDM

tubing showing grain boundary carbides.

(c)

Figure 28 summarizes the findings of the OIM analysis. Figure 28a shows a typical OIM map with

equiaxed grains. The resulting average grain size is 73 grm (Fig. 28b), about double the grain size found

in the plate alloy (Fig. 17). The CSLB fraction in this material, Fig. 28c, is approximately 53%, of which

a large proportion is twins (Y=3), and 6% is low-angle boundaries. While this GBCD is typical of a Nibase alloy, 29-32 it should be noted that, again, this fraction is about double that measured for the plate

29
alloy. The significance of this finding is that the plate alloy has 50% more crack-susceptible boundaries

than the CRDM alloy.

Average Dia (over area) Alloy C690

72.9S.m

0.2

50 .

Grain Size Diameter (p1m)

(b)

0

C

U.

C3

Alloy C690

0.4 Sum(35E•29)-47.0%

0.3

0.2

0.1

0.0

CSL iE

(a) (c)

Figure 28. (a) OIM map, (b) graph of grain size distribution, and (c) graph of grain boundary

character distribution for Alloy 690 TT from CRDM tubing.

3.3 Alloy 152 Double-J Weld (A152)

The Alloy 152 double-J weld was examined by SEM, TEM, and OIM. The objective was to observe the

microstructure, focusing on inclusions, and to determine the GBCD.

Figure 29 shows the microstructure of the Alloy 152 double-J weld, which has elongated, dendritic

grains. The microstructure was also examined by optical microscopy and TEM. Several inclusions are

visible in the images.

30
(a) (D)

Figure 29. Microstructure of Alloy 152 weld using SEM: (a) secondary electron detector and (b)

back-scattered electron detector.

Figure 30 shows an optical micrograph where the micron-sized inclusions mentioned previously are

visible. These appear to be of nearly identical size.

A152

Figure 30.

Optical image of the Alloy 152 weld specimen

A152.

The microstructure of the weld was also examined by TEM with the purpose of determining the nature of

the inclusion noted previously. The TEM images are presented in Fig. 31. The microstructure appears to

show a high density of dislocation loops and dislocation cells. Several inclusions were examined by

energy dispersive x-ray spectroscopy (EDS). Results are given in Fig. 31 d. These inclusions appear to be

oxides, with the major components being Al and 0.

31
0

E

0

Z

Al

Cr

eTi ¶CrcrFe Ni

Nb N i

A.A A. A ~- a

1 2 3 4 5 6 7 8 9 10

KeV

(d)

Figure 31. (a-c) TEM images and (d) EDS spectrum of an inclusion in the double-J Alloy 152

weld (A152).

Figure 32 summarizes the findings of the OIM analysis for the Alloy 152 weld. Figure 32a is a typical

OIM map showing elongated dendritic grains, as noted previously. The resulting GBCD (Fig. 32b) is

typical of Ni-base weld alloys.8 As such, the CSLB fractions are approximately 36%. As expected, this

weld also has a large proportion of low-angle boundaries (1= 1).

32
^ A

Alloy A152

Sum(3U129)=0 11

0.3 5'1=0.31

o0.2

U.

C,)

) 0.1

0.0 . . . . . . . . . . . . . . . . . . 7-1 . . .I ..I

CSL Z

(b)

Figure 32.

Alloy 152 weld in double-J geometry (A152): (a) OIM map and

(b) grain boundary character distribution.

(a)

33
q
4 Stress-Strain and Crack Growth Results

This section presents the tensile property data for Alloy 690TT (thermally treated) and 152 weld in the

temperature range of 25-870'C (77-1598°F), as well as the CGR data for specimens in simulated PWR

water. The CGR tests were conducted on Alloy 690 plate in the 26% cold-rolled condition, Alloy 690

CRDM tubing in the as-received condition, and Alloy 152 in the as-welded condition. Three tests have

been conducted on Alloy 690 in the 26% cold-rolled condition, two tests on Alloy 690 in the as-received

condition, and two tests on Alloy 152 weld.

4.1 Tensile Data for Alloy 690TT and Alloy 152 Double-J Weld

Tensile property data were obtained according to ASTM E8/E2139 on Alloy 690 and Alloy 152 weld

specimens as a function of temperature. The tests were conducted in air in the temperature range of 25-

870°C (77-1598-F). Two specimens of each alloy were tested at each temperature in the test matrix. For

these tests, Alloy 690 Heat NX3297HK12 was given an additional thermal treatment at 720'C (1328 0F)

for 10 h. Companion Alloy 600 specimens were also included in the test matrix. The Alloy 152 round

bars were cut to ensure that the tensile specimens consist of only Alloy 152, not the adjacent Alloy 690

plate or the Alloy 52 weld root (Fig. 33). The available Alloy 152 weld piece allowed for 12 tensile

specimens.

Figure 33.

Locations in the double-J Alloy 152 weld

where the tensile specimens were cut.

The results of the tensile tests are shown in Table 8 for Alloy 690TT and in Table 9 for Alloy 152 weld.

The data for both alloys meet the minimum requirements of ASME SB- 167.

35
Table 8. Tensile properties of Alloy 690TT in the temperature range 25-8700C (77-1598°F).

Specimen Temp., UTS, 0.2% YS, Elong., RA, Modulus, Ult. Load, 0.2% Yield, Orig. Dia., Final Dia.,

ID 0

F 0 C ksi MPa ksi MPa % % msi lb-f lb-f in. in.

A690-1 77 25 96.8 667.4 42.7 294.4 50 65 29.5 9384 4143 0.3514 0.2079

A690-2 77 25 96.6 666.0 42.1 290.3 49 64 23.3 9368 4080 0.3513 0.2105

A690-3 302 150 88.2 608.1 36.1 248.9 46.5 62 31.6 8508 3482 0.3505 0.2156

A690-4 302 150 88.4 609.5 36.0 248.2 46.5 64 29.2 8521 3473 0.3503 0.21

A690-5 572 300 83.4 575.0 33.7 232.4 47 57 24.5 8086 3273 0.3514 0.2303

A690-6 572 300 83.7 577.1 33.5 231.0 48 60 25.8 8063 3231 0.3503 0.2221

A690-7 842 450 80.7 556.4 33.1 228.2 47 62.5 27.1 7780 3193 0.3504 0.2152

A690-8 842 450 80 551.6 32.8 226.1 48 60.5 25.3 7715 3159 0.3504 0.22

A690-9 1112 600 70.4 485.4 31.5 217.2 33.5 34 24.1 6826 3057 0.3513 0.285

A690-10 1112 600 72.6 500.6 29.7 204.8 38.5 41.5 26.2 7045 2880 0.3514 0.2693

A690-11 1382 750 38.4 264.8 28.3 195.1 49.5 49 20.8 3716 2740 0.3512 0.2502

A690-12 1382 750 38.4 264.8 28.4 195.8 48.5 46.5 21.1 3724 2752 0.3512 0.2574

A690-13 1598 870 18.3 126.2 15.8 108.9 43.5 76 16.2 1770 1536 0.3514 0.1728

A690-14 1598 870 17.4 120.0 15.0 103.4 75.5 75 17 1688 1450 0.3514 0.1754

Note: UTS = ultimate tensile strength, YS = yield strength, RA = reduction in area

Table 9. Tensile properties of Alloy 152 weld in the temperature range 25-870"C (77-1598"F).

Specimen Temp., UTS, 0.2% YS, Elong., RA, Modulus, Ult. Load, 0.2% Yield, Orig. Dia., Final Dia.,

ID 'F °C ksi MPa ksi MPa % % msi lb-f lb-f in. in.

A152-11 77 25 93.8 646.7 61.7 425.4 44 54 19.4 4677 3073 0.2519 0.1705

A152-12 302 150 85.9 592.3 53.7 370.2 42 50.5 19.6 4275 2671 0.2517 0.1769

A152-13 572 300 85.6 590.2 59.6 410.9 39 48 23.5 4275 2975 0.2522 0.1815

A152-14 572 300 81.6 562.6 51.7 356.5 41 52 22.5 4061 2575 0.2518 0.1746

A152-21 842 450 76.7 528.8 47.9 330.3 45 51 15.8 3824 2386 0.2519 0.1761

A152-22 842 450 76.7 528.8 50.4 347.5 42 44.5 18.4 3827 2513 0.252 0.1877

A152-23 1112 600 70.8 488.1 50.2 346.1 34 38.5 16 3532 2507 0.2521 0.198

A152-24 1112 600 67.4 464.7 45.3 312.3 41 48.5 18.1 3360 2258 0.2519 0.1807

A152-31 1382 750 38.7 266.8 38.4 264.8 38 43 13.3 1934 1918 0.2522 0.1907

A152-32 1382 750 39.2 270.3 39.1 269.6 37 36.5 13.5 1953 1951 0.2519 0.2009

A152-33 1598 870 22.6 155.8 22.2 153.1 21 25.5 17.6 1125 1105 0.2519 0.2172

A152-34 1598 870 19.7 135.8 19.4 133.8 25 24.5 10.9 984 968 0.252 0.2187

Note: UTS = ultimate tensile strength, YS = yield strength, RA = reduction in area

The resulting stress-strain and true stress-true strain curves as a function of temperature for Alloy 690TT

are shown in Fig. 34. As expected, the yield strength decreases with increasing temperature. Figure 35

shows the 0.2% YS and UTS as a function of temperature. Also included are data obtained on companion

Alloy 600 specimens. Both 0.2% YS and UTS decrease with temperature. The data set for Alloy 690TT

is in good agreement with control data on Alloy 600.

36
lflflI .

80

Co

Co

ci,

in 1 n .g.!. . . i . . .! .! .: .! . .i . . . . . . '-- 

Alloy 690TT

..... ....... i .. t...... : i il.. . ............. :: ......... .. ............ .............. [.....

Tem'Tperature, oF C)

77 (25) .... ... . • l - ............. i...............= .............. ........ " ....

842 (450)

......-..-----1-- 1112 (600)

" .......1382 (750)

1598 (870)

. . . ... . . i . . . . . . . . . . . . . . . . . . . i . .

20

0

)

0 0.1 0.2 0.3 0.4 0.5 0.6 0.7

Strain

(a)

640

560

480

R' 400 a320

240

160

80

0

1000

800

600 •

Co

400 '•

200

a

Cot

160

140

120

100

80

60

40

20

0

Alloy 690T. . .................... .......... ........--- ........ ..........---...... .....------------- ................... .

..................... .................... • •. -- . . ---- . . --- . . . . ...... ............. , .................... .

,• .. --.. • ,v" .. '-- .. ------------- :: .. "il .... .. ... .. .................... . -------- ------- Temperature, ------- ------- 'F (00)---- ---- _--

• ~ ~ ~ ........ ~~------------ ... i................... ............... •1 4 ;0 .. ' 77(25) -

............. - ------- - - - - '- ---....... 842 ý450) --

1112 (600)-

.. -- .. -- .. -- .. -- .. -- .. -- -- - -- -- -- -- -- -- ---- - 1382 (750) -

1598 (870) . . I . . . . . . . . . , . .

0.4

0

0 0.1 0.2 0.3 0.5

Strain

(b)

Figure 34. (a) Stress-strain and (b) true stress-true strain curves for Alloy 690 TT as a function

of temperature.

37
C')

C)J

0-

70

60

50

40

30

20

10

0

I =1 ii I I I I -

AJloy 690TT - - - 0.2% YS A690

0 0.2% YS A600

* ...- ~. --- "4--B ..

..................... ......... ................................. T ................... .....................

~~~~~..................... ,..................... J..................... . ........ . . . . . .- - ............... ......--......------------------------------------------------

......-.- .....------------- £-------------a---- ---------- --- ----------------

.

I I *

480

400

320 a240

160

80

0

0 200 400 600 000

Temperature (°C)

(a)

800 l4

I ~-----------I- I i•EB

100

50

50

-~~~~ --- -IoT UTS A690

0 UTS A600

-~I --- - -

.. -.. --.. -. --.. ----- -- -- --- --- -- ----- --- -- -- --- -- --- --- -- --- -

1000

800

600 cc0-

F400 •

200

U 00-

0 200 400 600 800 1000

Temperature (°C)

(b)

Figure 35. (a) Yield stress and (b) ultimate tensile stress of Alloy 690TT as a function of

temperature.

38
The resulting stress-strain and true stress-true strain curves as a function of temperature for Alloy 152

weld are shown in Fig. 36. As expected, the yield strength decreases with increasing temperature. Figure

37 show the 0.2% YS and UTS as a function of temperature. Both the 0.2% YS and UTS decrease with

temperature. Also included are data on Alloys 52/152 and 82 from other sources.40 The data sets for

Alloy 152 are in good agreement with the Alloy 82/152 data obtained elsewhere.40

100

80

60

(/)

&5 40

20

Alloy 152 Weld

......... ..... ...... - ----- ------ . . . . . .............. ---- -- -- -_ --

Temperature, 'F (°C)

77(25)

* :4;' ~ )(45o

........... ".------_-........ - ----------------------- --------- I 1112 (600)

1382(750)

1598 (870)

, , * I * I ,,, , I . , , I , , . I . .

0

0 0.1 0.2 0.3 0.4 0.5 0.6 0.7

Strain

(a)

640

560

480

400 a.

320 C'

0/)

240

160

80

0

1000

800

CL

600 2

400 ol

200

U)

IOU

140

120

100

80

60

40

20

0

Alloy 152 Weld

..............................- -------------------. ....... . ....... - --------------.-. .

Temperature, 'F (°C) -

77 (25) ..... •••... -........',..................... l ,..................... ;"............... Y ...

.......................................... I ..................... .............. 842 (450 1 ....

1112(600)

.... . . .1382 (750) ....

1598 (870)

0 0.4

_0

0.1 0.2 0.3 0.5

Strain

(b)

Figure 36. (a) Stress-strain and (b) true stress-true strain curves for Alloy 152 weld as a

function of temperature.

39
4 ArI~P Ivuu

Alloy 152 Weld

* , | I , ,

- -i- - 0.2% YS A152

0 0.2 YS MLTS-24Cr

80 A 0.2% YS MLTS-27Cr

60

OR

20

20

E 0.2% YS MLTS-30Cr

A 0.2% YS EN52

......... ----- -- --.........---- -------------. A 0.2% YS EN82H --.

----- -- . .- -- -- -- m ------

- - - - - - I ---- -------- I - - I-- - I -

640

560

480

a400 -

U) 320 >-

240 6

160

80

0

0 200 400 600

Temperature (°C)

(a)

800 1000

150

100

Z'

Alloy 152 Weld -- 0U- -UTSA152

* UTS MLTS-24Cr

A UTS MLTS-27Cr

0 UTS MLTS-30Cr

------------- --------------------- A UTS EN52--------

A A A UTS EN82H • - - - 5l i _ . . -

: "A

.... .. ..... .... .. ......--------- --------1- ---------

1000

800

600 CU)

F400 •

200

50

0 200 400 600 800 1

10

000

Temperature (°C)

(b)

Figure 37. (a) Yield stress and (b) ultimate tensile stress of Alloy 152 weld as a function of

temperature.

40
4.2 Crack Growth Data for Alloy 690 Cold-Rolled by 26%

Three CGR tests in simulated PWR environments were completed on Alloy 690 cold-rolled by 26%.

These tests were conducted on 1/2-T CT specimens designated A690WC-SL-1, A690WC-ST-1, and

A690WC-ST-2.

4.2.1 Specimen A690WC-SL-1

This test was a "screening" test at a time when no other similar data were available. The specimen

orientation was chosen to be SL, as this was expected to be the most susceptible to SCC. Standard

transitioning procedures developed for Alloys 600 and 182 were applied. During the constant load

portion of the test (Period 9), no periodic unloading was used.

The testing conditions for Specimen A690WC-SL-1 are given in Table 10, and the changes in crack

length and Kmax with time are shown in Fig. 38. Surprisingly, transitioning from TG to IG SCC occurred

readily. The measured CGRs at constant load were between 1.5 and 2.8 x 10-11 m/s for Kmax = 26.7-

28 MPa-m1/ 2 (24.3-25.5 ksi-in.1/2).

Table 10. Crack growth data for specimen A690WC-SL-1 in PWR watera at 320°C.

Test Load Rise Down Hold Estimated Crack

Test Time, Ratio Time, Time, Time, Kmax, AK, CGRenv, CGRair, Length,

Period h R s s s MPa-m"/2 MPa-m1/ 2 m/s m/s mm

1 48 0.30 0.50 0.5 0 19.4 13.6 2.44E-08 1.89E-08 11.772

2 62 0.30 50 50 0 19.9 13.95 1.45E-09 2.12E-10 11.973

3 80 0.30 0.5 0.5 0 21.0 14.7 4.96E-08 2.67E-08 12.731

4 105 0.50 50.0 12 0 25.2 12.6 2.57E-09 2.39E-10 12.958

5a 170 0.50 500 12 0 25.8 12.9 5.45E-10 2.62E-11 13.081

5b 216 0.50 500 12 0 25.9 12.9 5.33E-10 2.71E-11 13.173

6 346 0.50 500 12 0 26.2 13.1 3.63E-10 1.41E-1 1 13.318

7 2190 1.00 0 0 0 26.2 0.0 1.46E-11 - 13.447

8 2234 0.50 500 12 0 27.8 13.9 7.15E-10 3.58E- 11 13.561

9 2831 1.00 0 0 0 28.0 0.0 2.83E-11 - 13.627

10 2854 0.50 50 120 0 29.7 20.8 4.31E-09 1.09E-09 13.983

aSimulated PWR water with 2 ppm Li, 1000 ppm B, DO<I10 ppb. Conductivity was 21 + 3 jiS/cm, and pH 6.4.

13.20

13.00

?12.80

E

S12.60

-j 12.40

012.20

12.00

11.80

Alloy 690

Specimen # A69OWC-SL-1 Crack Length

320"C, Simulated PWR Water

10 2 u 3 5a 3.29 x 10-01 Wes

3.36 x 10-t° mis 25.3 MPa min5

25.2 MPa mO 5 0150, Rise time 500

S050, Rise time = 500s

K._

50

45

40

35

30

25

20

15

10 00

E

r

0c

a.

40 60 80 100 120

Time (h)

140 160 180 21

(a)

Figure 38. Crack length vs. time for Alloy 690 specimen A69OWC-SL-1 in simulated PWR

environment during test periods (a) 1-5 and (b) 6-9.

41
80

14.00

13.80

E

-13.60

05

-J

-• 13.40

T

13.20

13.00

Alloy 690

Specimen # A69OWC-SL-1

320°C, Simulated PWR Water

6

3.63x 10-10 rm/s

26.2 MPa m°n5

D.50, Rise time = 1000 s

1.46 x I 0' rn/s

26.7 MPa m°n5

Constant Load

.1 Crack Length

Kmux

'10

2.83 x 10-1' mls

8 28.0 MPa mrnI

Constant Load

70

60

50 E

40

30

20

10

400 800 1200 1600

Time (h)

(b)

2000 2400 2800

Figure 38 (cont.)

The resulting fracture surface was examined by both optical and electron microscopy. Figure 39 shows

the fracture surface of the specimen by both optical and electron microscopy. The images indicate that

the testing procedure resulted in a straight, uniform fracture surface. Two fracture modes can be

distinguished: TG, corresponding to mechanical fatigue, and IG SCC. The images shown in Fig. 39 were

used to determine the correction factors for the DC potential data. No correction was necessary for the

mechanical fatigue periods, and a factor of 2.24 was applied to the SCC periods.

"0

C3

Cu

0)

0

W,

(D)

Figure 39. Fracture surface of specimen A69OWC-SL-1 by (a) optical and (b) electron

microscopy.

42
Figure 40 illustrates the fracture modes that were observed: (a) TG, corresponding to the mechanical

fatigue periods, and (b) faceted IG, almost "granular," corresponding to the SCC periods. Figures 40c

and 40d show additional images. Of interest are the secondary cracks and, especially, the cracked

carbides on the facets.

(c) (d)

Figure 40. Fracture modes observed on the fracture surface of A69OWC-SL-1: (a) transgranular;

(b,c) faceted with elements of IG; and (d) IG showing carbides on the facets. Crack

advance is from bottom to top.

Figure 41 shows an additional set of micrographs from the fracture surface of specimen A690WC-SL-1

with the purpose of comparing the features observed on the TG and IG fracture surfaces. Figure 41 a is a

low magnification micrograph showing both fracture morphologies: TG, corresponding to the mechanical

fatigue test periods, and faceted IG, corresponding to the SCC periods. Figure 41b shows the IG region

where secondary cracks can also be observed. Figures 41 c and 41 d show the TG fracture mode; however,

elements of IG cracking - granular appearance and smooth facets - are also apparent.

43
MILWA

A,,'ý V Magri'% 1,0 J0 kV 46"IX 'InfNOWC 'ýI I

A~xý WMIlý 

(c) (d)

Figure 41. Fracture modes observed on the fracture surface of A690WC-SL-1: (a) IG and TG, (b)

IG fracture, and (c,d) TG fracture with elements of IG (smooth facets). Crack

advance is from bottom to top.

To examine the fracture characteristics in more detail, the surface oxides were removed by standard

metallographic procedures. Figure 42 shows a set of micrographs after the oxides were removed. Figures

42a and 42b show the IG region. The fracture mode appears to be a combination of "granular" and

smooth IG, and a higher number density of narrow secondary cracks is now visible. Figure 42c shows the

TG fracture mode; however, as noted previously, elements of IG cracking - granular appearance, smooth

facets, and narrow secondary cracks - are visible.

44
t ID)

Figure 42.

Fracture modes observed on the fracture surface of

A690WC-SL-1: (a,b) IG fracture and (c) TG fracture

with elements of IG (smooth facets). Crack

advance is from bottom to top.

(c)

Figure 43 shows a cross section of specimen A690WC-SL-1. The side of the specimen was mechanically

polished by finishing with a 1-jim diamond paste, then was electrochemically etched in 10% nital solution

to show grain boundaries. Figure 43a shows the entire crack length, and arrows mark the beginning of the

IG portion of the crack and the end of the test. The overall crack appearance is straight, with no

departures from the initial test plane. Figure 43b shows the IG region at a higher magnification. As noted

previously, because there is no significant deviation from the initial test plane, the crack does not appear

to favor any particular microstructure. Crack branching or deep secondary cracks were not observed on

the cross section.

45
(a)

(b)

Figu Figure 43. Cross section of specimen A690WC-SL-1: (a) entire (IG and TG) crack front and (b) IG region at the end of test.
4.2.2 Specimen A69OWC-ST-1

This test was conducted on a companion specimen to the one tested previously. As with the previous

specimen, the crack front was in the mid-plane of the 26% cold-rolled plate, while the direction was

normal to the rolling direction. As in the previous test, standard transitioning procedures were applied,

and no hold times or periodic unloading was used.

The testing conditions for Specimen A690WC-ST- I are given in Table 11, and the changes in crack

length and Kmax with time are shown in Fig. 44. The measured CGRs at constant load were between

1.31 and 3.3 x 10-11 m/s for K..ax = 31.3 M!Pa'm1 2 (28.5 ksi'in.112

), in good agreement with the previous

specimen.

Table 11. Crack growth data for specimen A690WC-ST-1 in PWR watera at 3200C.

Test Load Rise Down Estimated Crack

Test Time, Temp., Ratio Time, Time, Kmax, AK, CGRenv, CGRair, Length,

Period h °C R s s MPa-mI/ 2 MPa-mI/ 2 m/s m/s mm

Prea 27 319.8 0.28 0.5 0.5 20.4 14.7 4.02E-08 2.51E-08 12.062

Preb 44 319.9 0.28 50 50 20.6 14.8 2.10E-09 2.62E-10 12.129

Prec 49 319.8 0.28 0.5 0.5 22.4 16.1 5.60E-08 3.66E-08 12.618

1 75 319.8 0.47 50 12 23.2 12.3 1.44E-09 1.98E-10 12.746

2 139 319.8 0.46 300 12 24.1 13.0 1.00E-09 4.06E-11 12.977

3 212 319.8 0.45 600 12 24.3 13.3 5.65E-10 2.18E-11 13.111

4 309 319.7 0.45 1000 12 24.9 13.7 3.75E-10 1.48E-11 13.260

5 407 319.7 0.67 1000 12 25.6 8.4 9.81E-11 4.21E-12 13.287

6 476 319.7 0.45 1000 12 25.9 14.2 3.76E-10 1.71E-11 13.366

7 714 319.7 0.45 4350 12 26.1 14.3 1.18E-10 4.06E-12 13.471

8 810 319.7 0.48 1000 12 26.7 13.9 5.59E-10 1.69E-11 13.648

9 1413 319.6 1.00 0 0 27.2 0.0 2.08E-12 13.674

10 1816 319.4 1.00 0 0 31.3 0.0 1.31E-1l1 13.697

11 1816 319.4 1.00 0 0 29.9 0.0 4.76E-12 13.711

12 1816 319.4 0.49 300 12 30.7 15.6 2.07E-09 9.41E-11 13.880

13 4000 319.5 1.00 0 0 31.4 0.0 3.27E-11 14.040

'Simulated PWR water with 2 ppm Li, 1000 ppm B, DO <10 ppb. Conductivity was 21 _ 3 jiS/cm, and pH 6.4.

50

45

40

35 E

30 0A

25 E

20

15

10

20 40 60 80 100 120 140

Time (h)

(a)

Figure4. Crack length vs. time for Alloy 690 Specimen A69OWC-ST-1 in simulated PWR

environment during test periods (a) precracking-2, (b) 3-7, (c) 8-11, and (d) 12-13.

47
14.00

p13.60

E

31 13.20

12.80

12.40

1

14.40

14.20

.14.00 E

E

•13.80

13.60

T 13.40

13.20

13.00

12.80

V

Alloy 690

Specimen # A69OWC-ST-1

Period 6

3.76 . 10-' Wne

5U

45

0.47 Riseflme = l00s

Crack Length 40 ,Period 4 : Perio( 2 :3.75. x1-'°mf. •3

S 24.9 MPa m'°' :. - 0: Ms. ýtime=10ý00. __ •Kn,: ... s ° -

,Period 3 ,Period 7 20

5i5 x 10or-,' rlis 1.18x 10-o nm/s

24.3 MPa mr-5 26.9 MPa mr' 15

0.47: Rise time j 600s 0.47 Rise time = 4350s

10

00 200 300 400 500 600 700 800

Time (h)

(b)

4

Alloy 690 Period 11

Specimen # A69OWC-ST-1 4.76 x 10-12

320"C, Simulated PWR Water 2819 MPa R~o

Period 10 ',Constant Load '

Period 9 1.31 x 10-11 m/s

2.1 x 10-12 Wes 31.3 MPa min5 Crack Length

27.2 MPa rn" Constant Load

Constant Load 7 I

K..

,,. . . ... . . , , , , i

55

50

45

40

35

30

25

20

800 1200 1600

Time (h)

(c)

2000 2400

60

55

50

45

E

40 a.

35 9

30

25

20

2600 2800 3000 3200 3400

Time (h)

(d)

Figure 44. (cont.)

3600 3800 4000

48
The resulting fracture surface was examined by both optical and electron microscopy. Figure 45 shows

the fracture surface of specimen A690WC-ST-1 as determined by optical microscopy. The image

indicates that the testing procedure resulted in a straight, uniform fracture surface. Two fracture modes -

TG, corresponding to mechanical fatigue, and IG SCC - can be distinguished. The image shown in

Fig. 45 was used to determine the correction factors for the DC potential data. No correction was

necessary for the mechanical fatigue periods, and a factor of 2.15 was applied to the SCC periods.

Figure 45. Fracture surface of specimen A690WC-ST-1. Crack advance is from bottom to top.

Figure 46 illustrates the fracture modes observed. Figure 46a shows the TG fracture mode, corresponding

to the mechanical fatigue periods, and Figs. 46b and 46c show the faceted IG - almost "granular" -

fracture mode, corresponding to the SCC periods.

(pA/ Iýnjun

A 0 kV 6 19. M"i AIOOWCý 'A I

(b)

Figure 46.

Fracture modes observed on the fracture surface

of A690WC-ST-1: (a) TG, (b) IG fracture, and (c) IG

fracture detail. Crack advance is from bottom to

top.

(c)

49
4.2.3 Specimen A690WC-ST-2

This test was conducted to determine the activation energy for SCC crack growth in Alloy 690 cold-rolled

by 26%. The approach was to perform a test on a single specimen and observe the effects of changes in

temperature on a propagating crack.

The test conditions for Specimen A690WC-ST-2 are given in Table 12, and the changes in crack length

and Kmax with time are shown in Fig. 47.

Table 12. Crack growth data for specimen A690WC-ST-2 of Alloy 690 in PWR water.a

Test Load Rise Down Hold Estimated Crack

Test Time, Temp., Ratio Time, Time, Time, Kmax, AK, CGRenv, CGRair, Length,

Period h 0C R s s s MPa-mI/ 2 MPa-m1/ 2 m/s m/s mm

Pre a 510 350.4 0.29 1 1 0 18.2 12.9 7.65E-09 9.26E-09 11.755

Pre b 647 350.0 0.29 100 100 0 18.4 13.0 3.23E-10 9.71E-11 11.828

Pre c 654 350.0 0.29 1 1 0 18.8 13.4 1.57E-08 1.07E-08 12.032

Pred 671 349.7 0.29 50 50 0 18.9 13.4 8.40E-10 2.17E-10 12.059

1 743 350.0 0.60 50 12 0 18.9 7.6 1.08E-10 5.07E-11 12.084

2 841 349.9 0.62 50 12 0 19.9 7.6 2.80E-10 5.38E-11 12.131

3 959 349.2 0.62 300 12 0 19.9 7.6 8.48E-10 9.16E-12 12.167

4 1,014 349.6 0.62 600 12 0 20.0 7.6 6.16E-11 4.62E-12 12.178

5 1,183 349.2 0.62 1000 12 0 20.1 7.6 5.56E-11 2.80E-12 12.211

6 1,384 349.5 1.00 0 0 0 20.2 0.0 5.57E-12 - 12.217

7 1,473 349.9 0.49 300 12 0 24.2 12.3 2.90E-10 4.37E-11 12.244

8 1,487 349.8 0.49 600 12 0 24.2 12.4 1.78E-10 2.20E-11 12.266

9 1,511 349.5 0.49 1000 12 0 24.2 12.4 9.42E-11 1.31E-11 12.271

10 1,917 350.0 0.49 300 12 0 24.6 12.5 3.42E-10 4.69E-11 12.370

11 1,990 349.7 0.49 600 12 0 24.7 12.6 1.81E-10 2.39E-11 12.418

12 2,085 349.4 0.49 1000 12 0 24.8 12.7 1.29E-10 1.46E-11 12.462

13 2,189 350.3 0.48 300 12 0 25.0 13.0 2.41E-10 5.28E-11 12.534

14 2,194 352.2 0.48 600 12 0 24.9 13.0 2.33E-10 2.64E-11 12.539

15 2,210 321.5 0.43 300 12 0 23.1 13.2 6.43E-10 3.97E-11 12.570

16 2,234 321.1 0.43 600 12 0 23.2 13.2 2.63E-10 2.02E-11 12.592

17 2,282 321.3 6.43 1000 12 0 23.4 13.3 1.63E-10 1.24E-11 12.621

18 2,404 320.3 1.00 0 0 0 23.4 0.0 4.04E-12 - 12.624

19 2,410 320.4 0.49 300 12 0 24.9 12.7 6.94E-10 4.02E-11 12.634

20 2,428 320.2 0.49 600 12 0 24.9 12.7 3.02E-10 2.04E-11 12.652

21 2,536 319.7 0.49 1000 12 0 25.1 12.8 1.83E-10 1.25E-11 12.727

22 2,698 320.7 0.49 1000 12 3600 25.5 12.9 6.45E-11 1.33E-11 12.769

23 2,938 319.9 0.49 1000 12 7200 25.6 13.0 6.15E-11 1.35E-11 12.820

24 3,799 320.1 1b 0 0 0 25.8 - 2.54E-11 12.890

25 4,154 311.9 1b 0 0 0 26.0 2.63E-11 12.925

26 4,747 319.7 1b 0 0 0 25.9 2.84E-11 12.993

27 5,090 300.0 1 b 0 0 0 26.4 3.77E-11 13.039

aSimulated PWR water with 2 ppm Li, 1000 ppm B, DO <10 ppb. Conductivity was 21 ± 3 [S/cm, and pH 6.4.

bIncludes periodic unloading by 50% (12 s/I 2 s) every 8 h.

The test was initiated at 350'C (662°F) with a dissolved hydrogen concentration of 50 cc/kg. This

concentration maintains approximately the same potential difference between the specimen and the

Ni/NiO line (Table 6) as the standard testing environment used at ANL (320'C or 608'F, 23 cc/kg

dissolved hydrogen). However, the environmental enhancement of cyclic rates was less than in the

previous two tests (Fig. 48). While the reason for the observed behavior is unknown, it was decided to

return to the standard testing environment to transition the crack from TG to IG. Given the relatively low

SCC CGRs initially measured for this specimen, it was also decided to continue load cycling and

introduced hold times to determine the SCC CGR component by superposition. Periods 22 and 23 had

the same loading as period 21, but hold times of 3600 s and 7200 s, respectively. Using the superposition

model (Eq. 10), SCC CGR components of 2.97 x 10-11 m/s and 3.52 x 10-11 m/s were calculated for

periods 22 and 23, respectively. These rates are in good agreement with those measured previously for

50
this alloy condition. Next, an unload/reload cycle was introduced. The sample was initially cycled every

2 h, then the hold period was increased to 4 h and 8 h. At approximately 3 cycles per day (period 24), the

rate appeared to have stabilized to 2.54 x 10-11 m/s again in good agreement with previous determinations

on this alloy condition and temperature. The temperature was decreased to 31 00C (5907F); no significant

effect on the CGR (2.63 x 10-"1 m/s) was observed. The temperature was increased back to 3200C

(608'F), and again the CGRs were essentially unchanged (2.86 x 10-11 m/s). With a more drastic change

in temperature, to 300 0

C (572'F), the rate still stayed unchanged. In summary, the behavior suggests that

for this cold-rolled alloy, once a steady-state SCC CGR is achieved, changes in temperature do not affect

that rate.

, J J• • •60

Alloy 690

12.40 Specimen # A69OWC-ST-2 7 8.9

350'C, Simulated PWR Water

2 C rack Length. . .5

12.20 4: 5 Prec C.] " ; ::4

12.00,

1 10 20

11.60 , " I • • • I • • • I • • • I ' '• ' '• , 10

600 800 1000 1200 1400 1600

Time (h)

(a)

12.90 - - , - 60

"Alloy 690

Specimen # A690WC-ST-2 22 23

12.80 -350'C, Simulated PWR Water 17 18 20 21

14 16- 50

12.70 10 11 12 13

12.60 E.0 E 40

C Crack Length A,

30 E

o

20

123 Km30

12.302"0 . .• .10

2000 2200 2400 2600 2800 3000

Time (h)

(b)

Figure 47. Crack length vs. time for Alloy 690 Specimen A69OWC-ST-2 in simulated PWR

environment during test periods (a) precracking-9, (b) 10-23, and (c) 24-27.

51
13.

13.

E13.

012.

12,

Alloy 690

Specimen # A69OWC-ST-2 27

320°C, Simulled PWR Water 26

10 (, 25 R

4I ), 2.5x 10 - m/s

24 24a 24b 8h t'j' 26MPam°"5

2 h 4 h _! '0:9 CL w/periodic unloading

00: ,

2.8 x to-, /' 4. ~E:

26 MPa, m0 •- 1

90 9 :Periodc Unloading: E OWE SE iCrack Length

60

50

~19

40 0E

a..

30 •

20

10

80

- S

3200C 310 0C

a I

300°C

3000 3500 4000

Time (h)

(c)

4500 5000

Figure 47. (cont.)

10-8

.Da

0>

n0

10-9

26% CR Alloy 690 .........................

Simulated PWR Water

Beet-Fi Curve for Ai 600 iy 

T CGR,' 4.4 x 10-7(CGR 10)

0

,

33 

-- ----- 

. ...

-- -- -.. ............... .............. ..

A A69OWC-ST-1,320C (23 cc/kg)

A690WC-ST-2, 350C (50 cc/kg)

Figure 48.

Cyclic CGR data for Alloy 690

Specimen A690WC-ST-2 at 3500C (6620F)

and for previous two specimens under

similar conditions in PWR environment at

320°C (6080F).

10-11

10-11 10-10 10-9

CGRair (m/s)

108 10-7

The resulting fracture surface was examined by SEM (Fig. 49). The image indicates that the testing

procedure resulted in a straight, uniform fracture surface. Two fracture modes - corresponding to

mechanical fatigue and SCC - can be distinguished. The image shown in Fig. 49 was used to determine

the correction factors for the DC potential data. In this case, the DC potential method underestimated the

crack advance by 14%. The good agreement is not surprising as a cycling component was almost always

present.

52
Figure 49. Fracture surface of specimen A69OWC-ST-2. Crack advance is from bottom to top.

Figure 50 illustrates the fracture modes observed. Figure 50a shows a region of the fracture surface with

both TG and IG fracture modes. Figures 50b and 50c focus on the transition area from TG to IG, which is

marked with arrows in the figures. Unlike previous specimens where elements of IG (smooth facets, IG

cracks, etc.) were found on the fracture surface in the TG region (see, for example, Fig. 41), the current

specimen appears to show a more abrupt transition from a predominantly TG fracture mode. Note that

the fracture mode started to change to IG after approximately 1 mm (0.04 in.), i.e., after transitioning at

320°C (608°F). This result further substantiates the earlier observation regarding the low enhancement of

the cyclic rates at 350'C (6627F), as shown in Fig. 48.

(a) (b)

Figure 50.

Behavior observed on the fracture surface of

specimen A690WC-ST-2: (a) TG and IG regions,

and (b,c) the transition between TG and IG. Crack

advance is from bottom to top.

kc)

53
4.3 Crack Growth Data for Alloy 690 in As-received Condition

Two CGR tests in the simulated PWR environment were completed on as-received Alloy 690. These

tests were conducted on '/2-T CT specimens designated C690-CR-I and C690-LR-2.

4.3.1 Specimen C690-CR-1

The test conditions for Specimen A690WC-SL-1 are given in Table 13, and the changes in crack length

and Kmax with time are shown in Fig. 51. The first measured SCC CGR at constant load (period 5 in

Table 13) was approximately 5 x 10-12 m/s for Kmax = 2.4.9 MPa-m1/

2 (22.7 ksi-in.1/ 2). After introducing a

periodic unloading component, the estimated SCC CGRs were found to be similar to that at pure constant

load, on average ; 5 x 10-12 m/s. Also, in period 9 the CGRair component was about an order of

magnitude smaller than the measured CGR in the test environment; thus, one can safely assume that the

measured CGR is entirely due to SCC. Finally, the specimen was set at constant load one more time to

reconfirm the 10-12 m/s rate.

Table 13. Crack growth data for specimen C690-CR-1 of Alloy 690 in PWR water.a

Test Load Rise Down Hold Estimated Crack

Test Time, Temp., Ratio Time, Time, Time, Kmax, AK, CGRenv, CGRair, Length,

Period h oC R s s s MPa-m1/ 2 MPa-m1/2 m/s m/s mm

Pre a 33 321.2 0.30 1 1 0 23.2 16.2 9.06E-09 2.01E-08 11.716

Pre b 48 321.2 0.30 100 100 0 23.3 16.3 3.14E-10 2.03E-10 11.723

Pre c 50 321.2 0.30 0.5 0.5 0 23.5 16.4 1.60E-08 4.20E-08 11.756

Pre d 53 321.2 0.30 2 '2 0 23.6 16.5 8.38E-09 1.07E-08 11.808

Pre e 57 321.2 0.30 5 5 0 23.6 16.5 5.40E-09 4.26E-09 11.844

1 80 321.2 0.48 50 12 0 23.6 12.3 4.61E-10 2.06E-10 11.855

2 195 321.1 0.48 300 12 0 24.0 12.5 1.17E-10 3.65E-11 11.935

3 313 321.1 0.49 600 12 0 24.1 12.3 7.63E-1I 1.78E-11 11.969

4 480 321.8 0.48 1000 12 0 24.3 12.6 4.25E-11 1.16E-I 1 11.993

5 1,681 320.7 1 - - 0 24.9 0.0 5.40E-12 - 12.027

6 2,568 321.2 0.48 1000 12 7200 21.5 11.2 8.87E-12 8.48E-13 11.054

7 2,882 321.4 lb 0 25.6 13.3 1.72E-1 1 12.088

8 3,340 321.3 Ic 0 25.6 13.3 9.36E-12 12.101

9 3,912 321.5 1d 0 24.8 12.9 5.50E-12 12.111

10 4,678 320.0 1 0 24.5 0.0 1.26E-12 12.118

aSimulated PWR water with 2 ppm Li, 1000 ppm B, DO <10 ppb. Conductivity was 21 ± 3 jiS/cm, and pH 6.4.

blncludes periodic unloading by 50% (12 s/I 2 s) every 2 h.

clncludes periodic unloading by 50% (12 s/12 s) every 4 h.

dlncludes periodic unloading by 50% (12 s/12 s) every 8 h.

54
50

45

4

40

35

igth E

30O(

25 E

20

15

, *10

100 200 300 400 500 600

Time (h)

(a) 4•

1

12.,u ..........

12.20

fE12.10

a) -J

" 12.00

U

11.90

14 an

• ov:,: . . .• •" '. . . . . . . . . . . . . . .' . . . . . . Alloy 6'90*

Specimen # C690-CR-1 8 9 320*C, Simulated PWR Water 7 8

6

Period 5 6.0 X 1()-12 m/s ,•

24.9 MPa ms 1 5,0%,12/12 Constant Load i, • ; 5% 21 50%.12/12 8h

-0.5, 1000/12 ' Constant lofd 4

2h periodic unloading

:0.5,1lOOO/112

Crack Length 2h

Km_(

80

70

60

50 E

40 C

E

30

20

4N

500 1000 1500 2000 2500

Time (h)

(b)

3000 3500 4000

11.

7a

r

0

Alloy 690

Specimen # C690-CR-1

320'C, Simulated PWR Water

.99

98

I ,,, .1i 1 o1il =1 aIl

Period 10 Crack Length

.97 2 x 10-12 m/s

24.4 MPa m".

Constant Load

Km_

80

70

60

50 E

40 x

30

11.

,1t

20

3800 4000 4200 4400 4600

Time (h)

1U

4800 5000 5200

(c) Figure 51. Crack length vs. time for Alloy 690 specimen C690-CR-1 in simulated PWR

environment during test periods (a) precracking-4, (b) 5-9, and (c) 10.

55
The resulting fracture surface was examined by SEM. Figure 52 shows a small region of the fracture

surface, and Fig. 53 shows the full fracture surface of the specimen, broken into two halves. The images

indicate that the testing procedure resulted in a straight, uniform fracture surface. The full fracture

surface, shown in Fig. 53, was used to determine the correction factors for the DC potential data. The

correction was small (4%) as a cycling component was almost always present, and growth under constant

load was limited.

Figure 52 shows that the fracture mode appears to be predominantly TG for the entire test. However, as

the transitioning ends (approximately halfway through the test), the fracture surface appears to become

smooth. This and other features were examined in detail, and will be discussed next.

Figure 52.

Region of fracture surface of specimen

C690-CR-1. Crack advance is from

bottom to top.

56
(a)

(b)

Figure 53. Fracture surface of specimen C690-CR-1: (a) first half and (b) second half.
I K

(c)

Figure 54.

Specimen C690-CR-1: (a) section of the

fracture surface, (b) TG and faceted fracture

with IG cracks, and (c) TG with IG secondary

cracks. Secondary IG cracks are marked

with arrows. Crack advance is from bottom to

top.

(a)

Figure 54 shows an additional set of micrographs from the fracture surface of specimen C690-CR-1 at

higher magnification. Figure 54a shows a section of the fracture surface of specimen C690-CR-1, and as

the transitioning ends (approximately halfway through the test), the fracture surface appears to become

faceted and smooth in appearance. Secondary IG cracks are also present in the second part of the test,

and several are identified by arrows in the figure. Figure 54b shows a region near the end of the test,

again with IG cracks and smooth facets visible. Figures 54c is a micrograph from the TG region where

secondary IG cracks begin to develop.

58
In support of the previous observation, Fig. 55 shows pairs of low and high magnification micrographs

taken near the end of the test. Areas with smooth facets and IG secondary cracks are visible. Figure 56

shows additional examples of areas with smooth facets and IG secondary cracks on the fracture surface of

specimen C690-CR- 1.

I,- Iv

I t) I

(c) (d)

Figure 55. Fracture surface of specimen C690-CR-1: (a,b) regions near the end of the test, and

(c,d) higher magnification micrographs showing the same areas, with smooth facets

and IG secondary cracks. Secondary IG cracks are identified with arrows. Crack

advance is from bottom to top.

In summary, the fracture surface of specimen C690-CR-1 did not show the smooth IG fracture mode

characteristic of fast propagating SCC cracks. Secondary IG cracking on the fracture surface was

extensive. Occasionally, those were accompanied by smooth facets.

59
to~

ira

(e) (t)

Figure 56. Examples of areas with smooth facets and IG secondary cracks on the fracture

surface of specimen C690-CR-1. Crack advance is from bottom to top.

4.3.2 Specimen C690-LR-2

The test on specimen C690-LR-2 complements that on specimen C690-CR-1, but in a different

orientation. The objective was twofold. First, the SCC CGR at 320'C was needed in order to determine

the effect of orientation. Second, once steady-state SCC was established, the temperature was increased

to 350'C (662'F) in an attempt to determine the activation energy for SCC crack growth in Alloy 690.

The premise was that if the SCC CGR for this specimen was comparable to that of C690-CR-1, then an

60
activation energy close to that for Alloy 600 (130 kJ/mol) would cause the SCC CGR at 350'C (662'F) to

be in the low-10-11 m/s range. To further facilitate SCC, the hydrogen concentration was maintained at 23

cc/kg, that is, close to a value that makes the specimen potential close to that for the Ni/NiO transition

(Table 6). The testing conditions are given in Table 14, and the changes in crack length and Kmax with

time are shown in Fig. 57.

Table 14. Crack growth data for specimen C690-LR-2 of Alloy 690 in PWR water.a

Test Load Rise Down Hold Estimated Crack

Test Time, Temp., Ratio Time, Time, Time, Kmax, AK, CGRenv, CGRair, Length,

Period h 0 C R s s s MPa'm 1

/

2 MPa-m"/2 m/s m/s mm

Pre a 79 321.5 0.29 5 5 0 18.3 13.0 3.81E-10 1.59E-09 11.614

Preb 94 321.5 0.29 50 50 0 18.3 13.0 6.14E-12 1.58E-10 11.616

Pre c 102 321.6 0.29 5 5 0 18.3 13.0 1.42E-11 1.58E-09 11.618

Pred 118 321.6 0.29 50 50 0 18.3 13.0 6.04E-11 1.58E-10 11.616

Pree 126 321.8 0.29 1 1 0 18.3 13.0 1.19E-09 7.89E-09 11.631

Pre f 194 321.8 0.29 50 50 0 18.3 13.0 6.74E-11 1.58E-10 11.639

Pre g 199 321.9 0.29 1 1 0 18.4 13.0 2.44E-09 7.98E-09 11.662

Preh 215 321.6 0.29 50 50 0 18.4 13.1 1.32E-10 1.60E-10 11.667

Pre i 222 321.9 0.30 1 1 0 18.6 13.0 4.OOE-09 8.05E-09 11.717

Prej 239 321.5 0.30 50 50 0 18.6 13.0 1.90E-10 1.62E-10 11.728

Pre k 248 321.6 0.30 1 1 0 18.8 13.2 6.76E-09 8.53E-09 11.833

Pre 1 263 321.5 0.30 50 50 0 18.9 13.2 3.35E-10 1.72E-10 11.845

Prem 272 321.6 0.29 1 1 0 19.3 13.7 1.00E-08 9.70E-09 12.003

Pren 287 321.6 0.30 50 50 0 19.3 13.5 5.15E-10 1.88E-10 12.017

Preo 293 321.6 0.30 1 1 0 19.6 13.7 1.36E-08 LO.E-08 12.173

Prep 359 321.6 0.29 50 50 0 19.8 14.1 6.86E-10 2.18E-10 12.258

Pre q 367 321.8 0.29 1 1 0 20.5 14.5 1.71E-08 1.24E-08 12.489

Prer 383 321.8 0.29 50 50 0 20.5 14.6 9.15E-10 2.51E-10 12,519

Pres 386 321.9 0.30 1 1 0 20.9 14.6 2.06E-08 1.30E-08 12.631

1 391 321.9 0.49 50 12 0 20.9 10.7 5.62E-10 1.20E-10 12.641

2 450 321.7 0.49 300 12 0 21.0 10.7 1.14E-10 2.01E-11 12.663

3 507 322.1 0.49 600 12 0 21.0 10.7 5.09E-10 LOIE-I1 12.692

4 559 322.0 0.49 1000 12 0 21.1 10.8 3.45E-11 6.21E-12 12.694

5 1,125 319.7 0.49 1000 12 14,400 21.2 10.8 5.95E-12 4.06E-13 12.724

6 1,411 321.4 1 0 0 0 21.3 0.0 3.78E-12 12.728

7 2,013 350.3 1 0 0 0 21.2 0.0 8.OE-12 12.731

8 3,726 349.8 0.48 1000 12 14,400 21.5 11.2 8.19E-12 5.53E-13 12.775

'Simulated PWR water with 2 ppm Li, 1100 ppm B, DO <10 ppb. Conductivity was 21 _ 3 tS/cm, and pH 6.4.

I 'I A.......................

12.80

Alloy 690

Specimen # C690-LR-2

320°C, Simulated PWR Water

d e f

h ,i, j

k I m n 0 p, q ,rs,

Crack Length

,m

50

45

40

35

E0

30 (L

25 E

20

15 11.80i

L

,b

11 Rfl

80 - - * 10

120 160 200 240 280 320 360 400

Time (h)

Figure 57.

(a)

Crack length vs. time for Alloy 690 Specimen C690-LR-2 in simulated PWR

environment during test periods (a) precracking, (b) 1-5, and (c) 6-8.

61
50

400 500 600 700 800

Time (h)

(b)

21.2 MPa m°'5 45

onstant load with periodic unloading;

40

35~ Crack Length E

30D.

25 E

20

Km.x 15

, I I . . . 10

900 1000 1100 1200

50

1.3 MPa m° I

Ale + hold (8h) 41- " 45

40

Crack Length

E

30 a25 E

20

Km.x 15

J ,10

E12.75 :

12.74

L 12.73

12.72

12,71

12.70L

Figure 57. (cont.)

Time (h)

(c)

The SCC CGRs were determined during both cycle + hold and constant-load conditions. For example,

period 5 (Table 14, Fig. 57b) consists of a cycle + hold, with a hold time of 4 h. The resulting rates (5.9-

6.6 x 10-12 m/s) are in good agreement with those obtained on the previous specimen in the as-received

condition. Since the mechanical contribution due to cycling is approximately an order of magnitude

lower than the measured rate, the measured crack growth rates are due to SCC. Next, the periodic

unloading was removed in period 6, and the SCC CGR measured at constant load was 3.8 x 10-12 m/s,

again, consistent with previous observations. After approximately 200 h, the temperature was increased

to 350°C (662°F). The SCC CGR rate at constant load increased to 8.0 x 10-12 m/s, then apparently

stalled. Cycle + hold (8 h) was introduced in an attempt to reactivate the crack, and a similar CGR of 8.8

x 10-12 m/s was measured. The rate appeared to decrease again, followed by another increase to 7.0 x

10-12 rnMs. While the rates seem consistent with each other, the cause for this occasional stalling is

unknown. Nevertheless, the SCC CGR rate in this as-received alloy appears to respond to temperature

changes, although not to the extent expected based on Alloy 600 experience.

62
The resulting fracture surface was examined by SEM. One observation was that the fracture surface was

covered with oxide filaments (Fig. 58). Such oxides have been observed previously by Mintz and Devine

during corrosion tests of Alloy 600 in the vicinity of the Ni/NiO line.4 1 Figure 59 shows the full fracture

surface of the specimen, imaged such that there is a contrast between the oxides and the metal surface.

This image was used to determine the correction factors for the DC potential data. The correction was

small, 10%, although not as small as that for specimen C690-CR-1, and was probably a reflection of the

fracture surface not being as straight. Nevertheless, in both cases the fact that a cycling component was

almost always present helped with DC potential accuracy.

Figure 59 shows that the fracture mode appears to be predominantly TG for the entire test. However,

features similar to those observed on the fracture surface of specimen C690-CR-1 were observed, and

these will be discussed next.

Figure 58.

Oxides on the fracture surface of

Specimen C690-LR-2.

Figure 59. Fracture surface of specimen C690-LR-2. Crack advance is from bottom to top.

Figure 60 shows micrographs from the fracture surface of specimen C690-LR-2. As with the previous

specimen, the surface has secondary IG cracks, which are identified by arrows in the figure. Occasional

smooth facets are also visible.

63
I I

Figure 60.

Examples IG secondary cracks and occasional

smooth facets on the fracture surface of specimen

C690-LR-2. Secondary IG cracks are identified with

arrows. Crack advance is from bottom to top.

(c)

Figure 61 shows additional examples of secondary cracking 50-70 prm from the crack front at the end of

the test. By comparison with the seemingly isolated cracks observed on the fracture surface of specimen

C690-CR-1 (see, for example, Fig. 55), the cracks here are extensive and may be due to the increased test

temperature during the portion of the test when this cracking occurred.

64
(b)

Figure 61.

Fracture surface of specimen C690-LR-2, regions

near the end of the test. Extensive secondary IG

cracks are identified with arrows. Crack advance

is from bottom to top.

.il

(c)

Finally, Fig. 62 provides additional examples of areas with smooth facets and IG secondary cracks on the

fracture surface of specimen C690-LR-2. At these magnifications, several facets indicate striations

consistent with the cycle + hold loading (Figs. 62b, 62c, and 62d).

In summary, the fracture surface of specimen C690-LR-2 was similar to that of specimen C690-CR-1,

and it did not show the smooth IG fracture mode characteristic of fast propagating SCC cracks.

Secondary IG cracking was present, and during the test period at 350'C (6627F), appeared to be more

extensive than what was observed on the fracture surface of specimen C690-CR-1. Occasionally,

secondary IG cracking was accompanied by smooth facets.

65
I-N

IC P II i A I

ke) kI) Figure 62. Examples of areas with smooth facets and IG secondary cracks on the fracture

surface of specimen C690-LR-2. Crack advance is from bottom to top.

4.4 Crack Growth Data for Alloy 152 Weld

Two CGR tests in simulated PWR environment were completed on Alloy 152 in an as-welded condition.

These tests were conducted on 1-T CT specimens designated A152-TS-2 and A152-TS-4.

66
4.4.1 Specimen A152-TS-2

This test was a "screening" test performed at a time when no other similar data were available, and was

conducted at the same time as the test on Alloy 690 cold-rolled by 26% (specimen A690WC-SL-1). The

weld specimen orientation was chosen to be TS, along the dendritic grains, as this orientation was

expected to be the most susceptible to SCC. Standard transitioning procedures developed for Alloy

182/82 were applied, and no hold times or periodic unloading was used.

The testing conditions for Specimen A152-TS-2 are given in Table 15, and the changes in crack length

and Kmax with time are shown in Fig. 63. Surprisingly, transitioning from TG to IG SCC occurred

readily. The measured CGR at constant load was 5.0 x 10-11 m/s for Kmax = 30.2 MPa-m1 /

2 (27.5

ksi'in. 1

/

2).

Table 15. Crack growth data for specimen A152-TS-2 of Alloy 152 weld in PWR water.a

Test Load Rise Down Hold Estimated Crack

Test Time, Temp., Ratio Time, Time, Time, Kmax, AK, CGRenv, CGRair, Length,

Period h °C R s s s MPa-m" 2 MPa-m"/ 2 m/s m/s mm

Pre a 48.0 322.5 0.3 1 1 0 20.3 14.2 2.26E-08 2.78E-08 12.683

Pre b 115.0 322.2 0.3 100 100 0 20.3 14.2 6.41E-10 2.82E-10 12.751

Pre c 132.0 322.1 0.3 1 1 0 20.7 14.5 2.83E-08 3.07E-08 13.183

Pred 134.0 321.9 0.3 10 10 0 21.0 14.7 5.01E-09 3.21E-09 13.333

Pree 150.0 322.0 0.3 1 1 0 21.9 15.3 4.09E-08 3.81E-08 14.174

Pre f 162.0 321.8 0.3 10 10 0 22.0 15.4 7.07E-09 3.91E-09 14.303

Pre g 172.0 321.9 0.3 1 1 0 22.8 16.0 4.67E-08 4.53E-08 14.999

Pre h 173.0 321.8 0.3 10 10 0 23.0 16.1 7.26E-09 4.69E-09 15.183

Pre i 186.0 321.7 0.3 1 1 0 23.9 16.7 4.92E-08 5.45E-08 15.887

Prej 194.0 321.7 0.3 10 10 0 24.2 16.9 7.86E-09 5.73E-09 16.115

Pre k 210.0 321.8 0.3 1 1 0 24.8 17.4 6.65E-08 6.42E-08 16.701

1 216.0 321.8 0.5 300 12 0 29.5 14.8 2.85E-10 1.88E-10 16.783

2 384.0 321.8 0.5 600 12 0 29.7 14.8 1.95E-10 9.55E-11 16.847

3 721.0 322.0 0.5 1000 12 0 29.8 14.9 1.54E-10 5.84E- 11 17.020

4 1724.0 321.5 1 - - 0 30.3 0.0 5.07E- I I - 17.336

aSimulated PWR water with 2 ppm Li, 1000 ppm B, DO <10 ppb. Conductivity was 21 - 3 pS/cm, and pH 6.4.

E

C

5)

17.6 -" III.I I..11 60

Alloy 152 SMA Double-J Weld Pre

16.8 Specimen # A152-TS-2 p ,

PrePre 

PrPe

320°C, Simulated PWR Water Pre f g h

16.0 • Pre Pre e Crack Length 50 Caciegt

15.2 Pre d

Pre 40

14.4- a b4

13 .6 -' , 3

30

12820

11.2

E

CL

E

ýc

50 100 150

Time (h)

(a)

200 250

Figure 63. Crack length vs. time for Alloy 152 weld specimen A152-TS-2 in simulated PWR

environment during test periods (a) precracking, (b) 1-3, and (c) 4.

67
18.0 Alloy 152 SMA DObIGj Weld

SPe0m0 #0 A I52-TS-2 

60

17.5 320'C, Simulatedj pWR Wate" 

t.53X( 10'onv.

Wa 

5 0

? 1 0 1 ' 2Crack Length 298m am

S16.0 

2 

.0 lein = 0 0 

0 .

40 E

~~~~16.0o 

6oIO

15.5 

30

15.0200 300 400 

400 600 7001

Time (h)

18.00 

(b)

Alloy 152 SMA Doubla-j Weld

32O*C, Simulatedi PWR Water 

5

E 17.50 

s

5 .T x 10-s o ni/ 

45

Crack Length 

30.2 M pa ino 

n

Consta lowd

T 

~ 012 0 

140 

18 460 02

power switch to the pump. The malfunctionig switch caused the pump to turn on and off internhlttently

every few seconds. Figure 64 shows that, as the pump began the on/off cycle at around 1450 h (the stress

intensity factor K reflects that behavior), the crack growth first displayed a jump (consistent with the

breaking of ligaments), then appeared to follow the oscillations in the stress intensity factor until the

switch was replaced. After the switch was replaced, the constant load conditions were restored; however,

the subsequent CGR doubled. The cause for the observed behavior is not understood.

68

8 00 2 0 4 01 6 02
34

32 Km"x

-E17.30,

IS17.20 X

w E 0J 28

17.10 5.0)c 10-)_1 m/s Crack Length 30.2 MPW rnm5

Constant load 26

17.00 .. . .I . . . . I . .

1300 1400 1500 1600 1700

Time (h)

Figure 64. Crack length vs. time for Alloy 152 weld specimen A152-TS-2 in simulated PWR

environment during constant-load test period 4.

The resulting fracture surface was examined by optical microscopy and SEM. Figure 65 shows images of

the two halves of the specimen. These indicate that the testing procedure resulted in a straight, uniform

fracture surface. These images were also used to determine the correction for the DC potential data. The

correction was small (16%) as growth under constant load (where ligaments that confound the DC

potential measurements are expected to form) was small when compared with the total crack advance.

The cracking seems to have propagated on four fronts on slightly different planes, resulting in an arched

appearance (Fig. 65b). Letters A-D indicate the intersections of these planes. The ductile fracture mode

observed at the intersections of the planes (e.g., at locations A-D) most likely occurred when the

specimen was fractured at the end of the test, and this finding substantiates the previous observation that

the cracking occurred on different planes. Nevertheless, the fracture mode was predominantly IG for

each of the four arches. For example, Fig. 66 shows the region between locations A and B with

predominant IG cracking and a band-like appearance. Overall, the engagement of the IG fracture mode

was 82% for this specimen.

*~ ~

~,

(a)

Figure 65. Fracture surface of specimen A152-TS-2 by (a) optical and (b) electron microscopy.

Crack advance is from bottom to top.

69
(b)

Figure 65. (cont.)

Figure 66. Intergranular cracking on the fracture surface of specimen A152-TS-2 between A and

B in Fig. 65. Crack advance is from bottom to top.

The portion of the fracture surface after the transition to a predominantly IG mode is shown in additional

detail in Figs. 67-69. For this purpose, the crack front is divided in three regions. In each region, several

features of interest are shown at higher magnification and discussed. While the fracture mode appears to

be predominantly IG, of particular interest were areas that appeared to be more resistant to IG SCC.

Figure 67a shows the first of the three regions. Figures 67b and 67c are high magnification micrographs

taken at locations 1 and 2 in Fig. 67a. While the fracture mode appears to be predominantly IG in Fig. 67

a, Fig. 67b shows a region of TG fracture. The ductile fracture found in this area (location 1) suggests

that the cracks may be propagating on slightly different planes. Also, there are secondary cracks

perpendicular to the direction of the crack front, but no secondary IG cracks. By contrast, Fig. 67c shows

a region where the fracture mode is mostly IG with an interdendritic appearance. The grain in the center

of the micrograph has interdendritic features, although not as prominent as those of its neighbors.

Moreover, as indicated by arrows, apparently all grain boundaries adjacent to this grain are cracked.

Figure 68a shows the second of the three regions, and Figs. 68b and 68c are high magnification

micrographs taken at locations 1 and 3 in Fig. 68a. As with the previous example, while the fracture mode

appears to be predominantly IG in Fig. 68a, Fig. 68b shows a region of TG fracture. This region appears

similar to that shown in Fig. 67b; however, in addition to the secondary cracks, dendritic elements can

also be observed. A common feature is the absence of secondary IG cracks. Figure 68c was taken at

location 3 in Fig. 68a and corresponds to location C of Fig. 65. The ductile fracture suggests that the

cracks propagated on slightly different planes and became connected only when the specimen was

fractured at the end of the test.

70
Finally, Fig. 69a shows the third of the three regions, and Figs. 69b and 69c are high magnification

micrographs taken at locations 1 and 2 in Fig. 69a. These show areas with an apparent mixed fracture

mode. Nevertheless, secondary IG cracks are observed. Figure 69a shows a case where what appears to

be a smooth faceted fracture is actually preceded by an interdendritic fracture, and illustrates how the IG

crack indicated by the arrow turns into an interdendritic fracture. In a similar fashion, the two IG cracks

shown in Fig. 69c turn into an interdendritic fracture.

Figure 70 further illustrates the role of secondary IG cracks. These types of cracks appear to develop

during transitioning and signal the onset of IG/interdendritic cracking. For example, the IG crack

indicated by an arrow seems to have initiated well ahead of the constant load region; however, it

apparently propagated sideways, toward the side groove of the specimen.

In summary, the engagement of the IG fracture mode was 82% for this specimen. However, areas of TG

and mixed TG/IG were found. A common feature for the TG areas is secondary cracking across the crack

propagation direction. Also, secondary IG cracks were observed in some cases, suggesting that the IG

engagement would have probably been even higher if this (screening) test had lasted longer.

71
(a)

t'J fracture

V Y, ------ d 10011m

(b) (c)

Figure 67. Specimen A152-TS-2: (a) first third of fracture surface, (b) detail at location 1, and (c) detail at location 2.

Crack advance is from bottom to top.
(a)

--4

(b) (c)

Figure 68. Specimen A152-TS-2: (a) second third of fracture surface, (b) detail at location 1, and (c) detail at location 3.

Crack advance is from bottom to top.
(a)

(b) (c)

Figure 69. Specimen A152-TS-2: (a) last third of fracture surface, (b) detail at location 1, and (c) detail at location 2. Crack

advance is from bottom to top.
Region of fracture surface of specimen A152-TS-2. Crack advance is from bottom to

top.

4.4.2 Specimen A152-TS-4

This test was conducted on an Alloy 152 weld specimen similar to the one described previously. The test

conditions for specimen A152-TS-4 are given in Table 16, and the changes in crack length and Kmax with

time are shown in Fig. 71.

75
Table 16. Crack growth data for specimen A152-TS-4 of Alloy 152 weld in PWR water.a

Test Load Rise Down Hold Estimated Crack

Test Time, Temp., Ratio Time, Time, Time, Kmax, AK, CGRenv, CGRair, Length,

Period

Pre a

Pre b

Pre c

Pre d

Pre e

Pre f

Pre g

Pre h

Pre i

1

2

3

4

5

6

7

8

9

10

11

12

13

14

15

16

17

18

19

20

21

22

23

24

25

26

27

28

29

30

31

32

33

34

35

36

37

38

39

40

h °C R s s s

32 319.8 0.3 0.5 0.5 0

96 319.7 0.3 50 50 0

104 319.7 0.3 0.5 0.5 0

119 319.7 0.3 50 50 0

128 319.3 0.3 0.5 0.5 0

144 319.4 0.3 50 50 0

152 319.4 0.3 0.5 0.5 0

168 319.5 0.3 50 50 0

173 319.4 0.3 1 1 0

200 319.4 0.49 50 12 0

264 319.8 0.49 300 12 0

336 319.1 0.49 600 12 0

626 319.4 0.49 1000 12 0

817 319.3 0.69 1000 12 0

960 319.6 0.49 1000 12 0

1293 319.3 1 0 0 0

1,651 319.4 0.49 300 12 0

2,141 320.0 1 0 0 .0

2,170 320.3 0.48 300 12 0

2,375 319.8 1 0 0 0

2,448 319.5 0.48 300 12 0

2,616 319.6 0.48 300 12 3600

2,856 319.4 1 0 0 0

2,888 320.0 0.48 50 12 0

2,952 320.1 0.48 300 12 0

3,023 319.5 0.48 600 12 0

3,120 319.7 0.48 1000 12 0

3,694 319.8 1 0 0 0

3,722 320.3 0.48 50 12 0

3,816 320.4 0.48 300 12 0

3,828 320.9 0.22 1 1 0

3,840 320.6 0.22 50 50 0

3,870 320.3 0.48 50 12 0

3,894 320.1 0.48 300 12 0

4,017 320.5 0.48 600 12 0

4,064 320.5 0.48 1000 12 0

4,399 320.1 1 0 0 0

4,655 319.4 0.48 300 12 7200

4,828 319.6 1 0 0 0

4,969 319.3 0.70 600 12 0

5,136 319.6 0.70 600 12 3600

5,812 319.0 0.70 600 12 7200

5,863 319.1 1 0 0 0

5,966 318.9 0.7 600 12 7200

6,081 319.2 0.7 600 12 0

6,103 319.1 0.5 600 12 0

6,152 319.2 0.5 1000 12 0

7,657 319.1 1b 0 0 0

7,665 319.2 0.5 12 12 0

MPa-mI/2 MPa-m'/2 m/s m/s mm

19.4

19.5

20.5

20.5

21.8

21.8

23.9

24.0

25.0

25.3

25.3

25.3

25.3

25.3

25.4

25.4

25.6

25.6

30.0

30.1

30.2

30.3

30.3

30.7

30.9

31.0

31.1

31.0

31.4

31.7

26.0

26.1

34.4

34.5

34.9

35.0

35.3

35.2

35.2

35.4

35.5

35.5

35.5

35.5

35.4

35.8

36.0

36.2

36.5

13.6 4.61E-08 4.51E-08 13.111

13.7 1.17E-09 4.57E-10 13.235

14.4 5.94E-08 5.62E-08 14.130

14.4 1.52E-09 5.68E-10 14.173

15.3 8.15E-08 7.29E-08 15.327

15.3 1.76E-09 7.36E-10 15.399

16.7 1.18E-07 1.06E-07 17.148

16.8 3.03E-09 1.09E-09 17.267

17.5 7.71E-08 6.42E-08 18.056

12.9 1.36E-09 6.12E-10 18.141

12.9 3.22E-10 1.03E-10 18.225

12.9 8.63E-1 I 5.07E- 11 18.255

12.9 5.85E-11 3.07E-11 18.303

7.8 7.41E-12 8.03E-12 18.306

13.0 7.92E-1I 3.12E-11 18.355

0.0 negligible - 18.355

13.0 3.89E-10 1.06E-10 18.480

0.0 5.07E-12 - 18.490

15.6 5.18E-10 2.18E-10 18.554

0.0 negligible - 18.554

15.7 5.51E-10 2.22E-10 18.682

15.8 6.69E-11 1.73E-1 1 18.725

0.0 negligible - 18.725

16.0 2.17E-09 1.42E-09 18.936

16.1 5.42E-10 2.43E-10 19.067

16.1 2.67E-10 1.23E-10 19.146

16.2 2.92E-10 7.54E-11 19.221

0.0 6.09E-12 - 19.224

16.3 2.34E-09 1.56E-09 19.399

16.5 5.19E-10 2.71E-10 19.569

20.2 6.64E-08 9.85E-08 20.752

20.3 4.49E-09 2.OOE-09 20.832

17.9 3.57E-09 2.29E-09 21.164

17.9 8.07E-10 3.83E-10 21'.233

18.1 5.OOE-10 2.02E-10 21.452

18.2 3:35E-10 1.23E-10 21.505

0.0 8.28E-12 - 21.515

18.3 7.19E-11 1.66E-11 21.561

0.0 negligible - 21.561

10.6 1.12E-10 4.80E-11 21.602

10.6 1.41E-II 6.96E-12 21.607

10.7 1.81E-II 3.74E-12 21.659

0.0 negligible - 21.659

10.7 1.02E-1I 3.74E-12 21.662

10.6 1.24E-10 4.84E-11 21.704

17.9 5.30E-10 1.99E-10 21.753

18.0 3.54E-10 1.22E-10 21.803

18.1 3.08E-11 - 21.957

18.2 7.99E-09 1.1OE-08 22.087

aSimulated PWR water with 2 ppm Li, 1000 ppm B, DO <10 ppb. Conductivity was 21 ± 3 [tS/cm, and pH 6.4.

bAlso includes periodic unloading by 50% every 7200 s. Unload/reload times were 12 s/I 2 s.

76
40

C-)

35

30

25

20

15

40 80 120 160 200 240

Time (h)

(a)

ion . .. .. . .. ................ ,4. n

'E 18.5

E

..J

• 18.0

Alloy 152 SMA Double-J Weld

Specimen # Al152-TS-4

320°C, Simulated PWR Water

5 35

4 Crack Length

32i 30

•. ' 25

Km

E

20

17• 15

240 320 400 480 560

Time (h)

(b)

640 720 800 880

E

E

-,

0i

LO

L)

19.5

19.0

18.5

18.0

17.5

Alloy 152 SMA Double-J Weld 14 1 ,16j1

; 18: 19

Specimen # Al 52-TS-412 13

320=C, Simulated PWR Water 10 11

6'

Crack Length :: : :-

Kmc : , ', .: ... . .

wu

50

40 E

EL

30 Y

20

1000 1500 2000 2500 3000 3500

Time (h)

(c)

Figure 71. Crack length vs. time for Alloy 152 weld specimen A152-TS-4 in simulated PWR

environment during test periods (a) precracking-2, (b) 3-5, (c) 6-19, (d) 20-25, (e) 26-

37, and (f) 38-39.

77
,w . . . . . . . . . . .

Alloy 152 SMA Double-J Weld

Specimen # A152-TS-4

320'C, Simulated PWR Water

21.0-

f 20.0

-j

L 19.0

18.0

18

19

Crack Length

11

22: , 'i24 25

21

20

2O ,- y -

60

50

40 E

E

30 c"

20

K_

1 I -Ji Wel i I I l 37: I

3000 3200 3400 3600

Time (h)

(d)

3800 4000

22.00

21.80

E

21.60

- 21.40

21.20

21.00

2•

Alloy 152 SMA Doublt

Specimen # A152-TS320°C, Simulated PW

27 28

.6:

e-J Weld 37:'

.4

-4R Water 31 3233 34 :-,35:

29 30 313 ,• qe -

• , , Crack Length '

60

50

40 E

0-

30 N

20

Krna

4000 4500 5000 5500 6000 6:500

Time (h)

(e)

LL .OU

22.40

22.20

E

E

22.00 0l

.• 21.80

21.60

21.40

Alloy 152 SMA Double-J Weld

Specimen # A152-TS-4

320°C. Simulated PWR Water

Period 39

3.Ox 10-" mrs

36.2 MPa mr0

Constant Load with PeriodicU

K._

80

70

60

50 E

a.

40 1

30

20

21.20'

6000

-j 10

6500 7000 8000

Time (h)

(f)

7500

Figure 71. (cont.)

78
The resulting fracture surface was examined by optical microscopy and SEM. Figure 72 shows a

photograph of the entire fracture surface. Transitioning to IG was attempted in two microstructural

regions approximately 1 mm (0.04 in.) apart. Between these two regions, the crack was advanced by

fatigue (period 22), and this period appears lighter on the fracture surface. The image indicates that the

testing procedure resulted in a straight, uniform fracture surface. This image was used to determine the

correction factors for the DC potential data. The correction was small (9.5%), as a cyclic component was

almost always present and, thus, ligament formation was kept to a minimum.

Figure 73 shows the fracture surface during the final periods of the test when the SCC CGRs were

determined. To provide additional detail, the fracture surface was divided into three regions along the

width. The fracture mode was almost entirely IG with interdendritic features. Overall, the engagement of

the IG fracture mode was 92%, which is larger than that for the previous specimen; this difference is not

surprising given the much longer time allotted to this test.

Figure 72.

Fracture surface of specimen

A152-TS-4. Crack advance is

from bottom to top.

The approach to transitioning from TG to IG was similar to that used in the previous test. However, the

environmental enhancement was initially lower than that observed previously (compare, for example,

periods 4 and 6 of Table 16 vs. period 3 of Table 15). Hence, the initial attempts to measure a CGR at

constant load were unsuccessful.

A degree of environmental enhancement comparable with specimen A152-TS-2 was obtained only later

in the test. However, the inability to sustain a CGR at constant load consistent with the previous

specimen (;10-11 m/s) suggested the formation of ligaments. To attempt to break the ligaments, in the

second part of the test (after period 22), alternative loading conditions - slow/fast cycling with hold times

or constant load with periodic unloading - were chosen. These loading procedures were successful in

keeping the formation of "fingers" and ligaments at a minimum (hence, the small correction factor

reported for this specimen). The SCC CGRs were calculated during several slow/fast cycling tests with

hold times or periodic unloading (see Table 16). Typically, a well-characterized, environmentally

enhanced slow/fast condition was chosen, and a hold period was introduced. The expectation was that the

SCC CGR component would grow during the hold period. The SCC CGR was then calculated by

79
superposition (Eq. 10). As such, based on the results of periods 21 and 29 (Table 16) the SCC CGR was

calculated to be 2.80 x 10-11 rn/s for period 29 (2-h hold). In a similar fashion, using the crack growth

rate measured in period 31, the SCC CGR was calculated to be 1.60 x 10-11 m/s once period 33 (2-h hold)

was completed. Note that by increasing the hold times from 3600 s in period 32 to 7200 s in period 33,

the measured crack growth rates stayed largely the same (,z 1.5 x 10-11 m/s). This observation further

substantiates the premise that at that stage of the test the SCC CGR component (and IG fracture) was

dominant, although some degree of cycling appears necessary to sustain the crack growth over long

periods of time. Finally, the SCC CGR was measured at constant load with periodic unloading. During

period 39 a steady growth at a rate of 3.08 x 10-11 m/s was measured. The unload/reload cycle was

relatively fast (12-s rise, 12-s down, R= 0.5) compared with the hold time (2 h), and by itself does not

typically result in an IG fracture. As expected, when the hold time was removed (period 40), the

measured CGR was equivalent to that expected in air, and the fracture mode reverted to TG for the final

100 pim of the test (Fig. 74). Therefore, the CGR measured in period 39 is due to IG SCC growth that

occurred primarily during the hold periods at constant load. In summary, the SCC CGRs measured at

constant load with periodic unloading or calculated from cycle + hold periods are consistent with those

measured at constant load.

80
(a)

2/3

(b)

3

(c)

Figure 73. Intergranular region at the end of specimen A152-TS-4. The image was broken into three regions. Crack

advance is from bottom to top.
Figure 74. Fracture surface of specimen A152-TS-4 at the end of the test (arrows). Crack

advance is from bottom to top.

At the end of the test, the fracture surface was also examined for evidence of ligaments to provide

additional validity to the rationale behind the test management described previously. Two examples are

shown in Fig. 75. Figure 75a shows a predominantly IG fracture mode, with a high density of secondary

IG cracks. In the center of the image, the IG crack went out of plane, forming the ligament indicated by

the arrow. Figure 75b shows a high magnification micrograph of the two fracture modes: IG on the left

side and TG on the right side. Interdendritic features are also recognizable on the TG side, as well as

secondary cracks similar to those observed on the previous specimen (Fig. 67b), suggesting a crackresistant grain. Striations are also visible, consistent with a fatigue-driven crack advance. In other words,

the cyclic component broke through such grains (and ligaments) and enabled a CGR to be measured in

real time.

Figure 75. Ligaments on the fracture surface of specimen A152-TS-4: (a) out-of-plane crack

and (b) interdendritic (left) and TG (right) fracture modes. Crack advance is from

bottom to top.

Upon the destructive examination of the specimen, IG fracture features were also observed prior to period

22 (Table 16), as shown in Fig. 76.

82
Figure 76.

Fracture surface of specimen

Al 52-TS-4 prior to period 22. Crack

advance is from bottom to top.

83

5. Discussion

The section summarizes the crack growth rate results for Alloy 690 and 152 weld and compares these

results with similar data obtained earlier at ANL and elsewhere. The significance of the results is

discussed in the context of data obtained previously on Alloy 600 and 82/182 welds. The activation

energy for crack growth in Alloy 690 is also discussed.

5.1 Alloy 690

5.1.1 Cyclic Crack Growth Rates

Cyclic CGRs for Alloy 690 obtained in previous testing programs at ANL have been evaluated to develop

correlations for estimating the cyclic CGRs in Alloy 690 as a function of loading and environmental

conditions. The data were obtained from tests conducted with a triangular or slow/fast sawtooth

waveform. The load ratio R was in the range of 0.2-0.9, Kmax was in the range of

30-50 MPa-m1/

2 (27-46 ksi-in. 11

2), and the rise time was 12 s for most of the tests.2-4,7

Figure 77 shows cyclic CGR data for Alloy 690 in the mill annealed (MA), solution annealed [SA,

1093°C (1999°F) for 1 h], and thermally treated [TT, 715'C (1319'F) for 5 h] conditions from previous

ANL 7 and Babcock & Wilcox42 (B&W) tests. The cyclic rates measured in primary water are plotted vs.

the cyclic rates predicted for Alloy 690 under the same loading conditions in air (see Eqs. 11 and 13).

Figure 77a shows excellent agreement between the measured and predicted rates. It also indicates that

under the testing conditions employed, no environmental enhancement occurred. Previous ANL cyclic

data on Alloy 690 obtained in high purity water at 320'C (608'F) with low DO content are shown in

Fig. 77b. The data indicate that under certain loading conditions, the rates in the environment begin to

drop - an effect consistent with the existence of a threshold, AKth, for cyclic rates.

1 0-5

10-6

Alloy 690 Simulate•

I 111.1n1 M i• 11 1 ii 1 1 1[ 1 I~ 11 l 11 1

0

0

10-8

10-9

d PWR Water at 320"C

... .. . •11 .. , .. ... ,.........o.......... ........

NX8662HG-33 MA+TT

U NX8662HG-33 MA+TT

NX8244HK 1B SA

B&W MA+TT

10-7

10-8

E 1 0 -9

CD

L) 10.10

10-11

HP- Alloy 1""- 1' '' I "" ' ' '.. "'".. 1..! I I I'" . "1 ' ý '"'"1.. y 690 -- - --I ---------.. .. .. .---

Water at 320'C, DO<5 ppb

0 NX8662HG, MA+715'C/5 :

M NX8662HG, SA+715'C/5h . NX8244HK SA 1093.C/lh .

NX8244HK. SA 982 C/lh

* *

.N: .... 44HK SA 982'01 h

A

• A

.. . . .-. ..... - --------------

10-10

10-11

10-1 2 10-12 .v

10-12 10-11 10-10 10-9 10-8 10-7 10-6

CGRair (m/s)

10-5 10-12 10-11 10-10 10-9

CGRair (m/s)

10-8 10-7

(a)

Figure 77. Cyclic CGR data for (a) Alloy 690

(b)

in SA and TT conditions from previous ANL and

B&W tests in PWR environment and (b) Alloy 690 in high purity water with low DO as

a function of growth rates for Alloy 690 in air. Test temperature is 320°C (6080F).

85
Figure 78 shows the cyclic CGR data for Alloy 690 in the 26% cold-rolled condition and data for

Alloy 600 from the Davis-Besse CRDM nozzle #3.9 Also included is a best fit curve for the Davis-Besse

Alloy 600 data, as the rates for this alloy were generally higher than those for other heats of Alloy 600

tested previously. The plot shows that in the mechanical fatigue regime (; 1-3 x 10-8 mI/s), both data sets

are in good agreement. The test protocol employed for the purpose of transitioning from TG to IG led to

environmental enhancement for the two cold-worked specimens, and these rates were typically higher

than predicted by the corrosion fatigue curve for Davis-Besse Alloy 600. The exceptions are two data

points obtained on specimen A690WC-ST- I that fall on the corrosion fatigue curve. These correspond to

periods 5 and 7 of Table 11, which were conducted with the intention of briefly exploring the effect of a

larger load ratio and were not used as final transitioning conditions. The relatively large environmental

enhancement of cold-rolled Alloy 690 suggests that transitioning to IG occurs readily, and relatively high

SCC CGRs are expected.

The high levels of environmental enhancement observed in both Davis-Besse CRDM nozzle Alloy 600

and 26% cold-rolled Alloy 690 are reflected in the respective fracture surfaces (Fig. 79). A unique

feature of the Davis-Besse Alloy 600 was the predominantly IG fracture even during loading, leading to

the high rates of mechanical fatigue crack growth. For example, the fracture surface of the Davis-Besse

specimen N3CL-1 changed to IG near the machine notch, where the first grain boundary is encountered

(Fig. 79a). The IG growth taking place in a regime dominated by mechanical fatigue (which normally

results in TG cracking) suggests a weakness of the grain boundaries. Similarly, grain contours, some

smooth facets, and IG cracks were observed on the fracture surface of the cold-rolled Alloy 690 specimen

(Fig. 79b). Such features observed in an inherently more resistant alloy suggest a similar weakness at the

grain boundaries. While the reason for the weakness of the grain boundaries in the Davis-Besse Alloy

600 is unknown, it may be a result of grain boundary segregation during fabrication. By contrast, the

cracked grain boundary carbides observed in the cold-rolled Alloy 690 indicate, in a straightforward

fashion, a weakness at grain boundaries. This possible effect of cracked grain boundary carbides on SCC

will be addressed.

10-7

Alloy 690 andý Davi's-Besse N#3 Alloy 600 ASimulated PWR Water

10-8 1 08 .CGR, Best-Fit Curve 6.6 x for 10-7{CG-R,,)O33 Davis-Besse N#3 A600 ... .... .... ...-

o 10-9. .. . ........... Figure 78.

_ [Cyclic - CGR data for 26% cold-rolled Alloy

CD '-.. ... .690 10 .. .10...-- and Alloy 600 from Davis-Besse in PWR -- - --.---.-- - -- - -

10-10 .. environment at 320"C (608"F).

112-T T OB N3CC-3

1.4.T CT DB N3CL . .

A A690WC-ST-1

10-11 10-10 10.9 10-8 10-7

CGRair (m/s)

86
I

(a) kt))

Figure 79. (a) Transition from TG to IG fracture in Alloy 600 from Davis-Besse Nozzle #3, and (b)

predominant TG fracture during transitioning in Alloy 690 cold-rolled by 26%. Grain

contours, some smooth facets, and IG cracks are evident. In both cases, the sample

notch is at the top of the figure, and crack advance is from top to bottom.

Figure 80 shows the cyclic CGR data for Alloy 690 CRDM tubing in the as-received condition and, for

comparison, the Alloy 690 plate in the 26% cold-rolled condition. The rates for the alloy in the asreceived condition appear to show some environmental enhancement, although much less than that for the

cold-rolled alloy. Moreover, the cyclic rates for the as-received alloy also appear to be lower than the

corrosion fatigue curve for Alloy 600. Hence, it is expected that as-received Alloy 690 is more resistant

to SCC than Alloy 600.

10-7

10-8

'a- 10-9

L9

Alloy 690

Simulated PWR Water at 320"C

Best-Fit Curve for Alloy 600 .... : ......... . .........

CG + , 44 X 1 7

(CGR)033

.5 .............. • A . . . . . . . -. -......... - - - - -- ......... - - - - -

26% "A CR A690WC-ST-1

.... C.•O-CR, -

* C690-LR-2

Figure 80.

Cyclic CGR data for 26% cold-rolled and asreceived Alloy 690 specimens in PWR

environment at 3200C (608°F).

10.11 10-10 10-9 10-8 10-7

CGRair (m/s)

87
5.1.2 SCC Crack Growth Rates

Figure 81 summarizes the SCC CGR data for Alloy 690 at various levels of cold work. Since the

activation energy for SCC growth is not yet known for Alloy 690, the data presented in the figure were

restricted to the temperature range 320-340'C (608-644'F), close to the operating temperature of PWRs.

Likewise, the data included were obtained at dissolved hydrogen levels of 23-30 cc/kg. The data

presented are from ANL for 26% cold-rolled material, Bettis Atomic Power Laboratory for 24% coldrolled and 32% tensile prestrained alloys,3 7 and Institute for Nuclear Safety System INSS for 20% and

50% cold-rolled alloys. 4 3 The proposed 75th and 5 0th percentile disposition curves for Alloy 600 (EPRI

MRP-55 2 3) are included for comparison.

The SCC CGR data developed at ANL on 26% cold-rolled Alloy 690 are in excellent agreement with

those from 20% cold-rolled material developed at INSS,43 especially for the tests conducted at a similar

temperature (320'C, 608'F). These data are lower than those from 24% cold-rolled developed at Bettis3 7

and from 50% cold-rolled material developed at INSS.43 It is worth noting that the Bettis CGRs on the

24% cold-rolled alloy tested at 338°C (640'F) were the highest in their matrix, which also included tests

at 360'C (680°F). Despite the different methods used to produce the deformation (cold rolling vs. tensile

prestraining) and different deformation levels, the results show, in general, good agreement. The Alloy

600 7 5 percentile tl curve (EPRI MRP-55 2 3 ) seems to bound most of the data.

C,

C-,

C

a)

2

w

Alloy 690 20-50% CW

Simulated PWR environment

io-9 315-340°C, 23-30 cc/kg H2 --------------------------............. .... ..... .. * CGR Curve Alloy 600

: 75th Percentile a

1o-ll ------.---------.---- ----........

/ ::Bettls 24/% CR ST (23 cc/kg) Bttls 24/% CR SL (23 cc/kg)

' 10-1 Beettis 32% prestrain SL Bettis 32% prestrain ST

', INSS 20% CR TL (30 cc/kg) )INSS 50% CR TL (30 B ANL cc/kg) 26% CR A690WC-SL-1 (23 cc/kg) - - - ANL 26% CR A690WC-ST-1 10-13 : (23 cc/kg) ANIL 26% CR A690WC-ST-2 (23 cc/kg) 1 :0 . . . . ;.. ......... ... .... . . . AN 26 . . CR . 69 W -T2(3c/g

Figure 81.

SCC CGR data for cold-worked

Alloy 690 in PWR environment

at 320-340°C (608-644°F). The

ANL data are shown with solid

green symbols.

10 20 30 40 50 60

Stress Intensity K (MPa-m1 /

2

)

Figure 82 shows the data obtained for as-received Alloy 690 (solid black symbols). In order to put this

data set in perspective, the plot includes data on cold-worked alloys as well as the proposed disposition

curves for Alloy 600. The CGR rates measured for the alloys in the as-received condition are about an

order of magnitude less than those obtained on the 26% cold-rolled alloy.

88
10-9

" 10-10

0

0

10-11 E

x

w

10-12

10-13

- r ,,, , , ,

Alloy 690

Simulated PWR environment

CGR Curve Alloy 600

75th Percentile a

-.. ANL 26% CR A69IWC-SL-1 (23 c--kg)

A ANL 26% CR A69OWC-ST-1 (23 cc/kg6

* ANL 26% CR A690WC-ST-2 (23 cc/kg)-

ANL C690-CR-1 5 as-received (23 cc/kg)

... ANL Ce90-LR-2 as-received (23 cc/kg)

- . -- -.-- - . . . . .. ....... i

Figure 82.

SCC CGR data for as-received

and cold-worked Alloy 690 in

PWR environment at 320-3400 C

(608-644°F).

10 20 30 40 50 60

Stress Intensity K (MPa.rn112)

5.1.3 Activation Energy

Researchers at ANL and elsewhere have attempted to determine the activation energy for Alloy 690 in the

cold-worked and as-received conditions. While scarce, the available data are summarized in Fig. 83.

An early attempt to determine the activation energy for 24% cold-rolled Alloy 690 was reported by

Bettis37 (Fig. 83a), and there was a rather weak dependence of SCC CGR on temperature. The ease of

crack initiation and moderately high CGRs measured on the Alloy 690 plate cold-rolled by 26% made it

attractive to attempt to determine the activation energy for SCC crack growth. The approach was to

observe the effect of temperature on a crack that is already propagating. Figure 47c shows that

temperature changes in the range 300-320'C (572-608 0F) have little effect on the CGR (:3 x 10-11 m/s

that was initially established at 320'C (608'F). Similar behavior was observed in tests at GE on this 26%

cold-rolled alloy,38 albeit at a higher CGR [the GE tests were designed to maximize the initial SCC CGRs

by starting the test at 360'C (680'F) at a potential near the Ni/NiO line]. While the SCC CGRs measured

at ANL and GE seem to be in agreement via an activation energy close to that established for Alloy 600

(Fig. 83a), the lack of temperature effect observed in the individual tests, as well as in the Bettis data, is

still not understood. Unlike the 24-26% cold-rolled alloys, the alloys cold-rolled by 20% responded to

temperature changes in both GE38 (Fig. 83b) and INSS43 (Fig. 83c) tests, yielding activation energies

close to that for Alloy 600 (130 kJ/mol). Alloy 690 in the as-received condition tested at ANL (Fig. 83c)

also responded to temperature change in a manner similar to Alloy 600 (Fig. 83b). By contrast, the SCC

rate in Alloy 690 cold-rolled by 50% was not found to be dependent on temperature in INSS tests (Fig.

83c). 43

89
Temperature (°C)

352 333 315 298

Temperature ('C)

283 352 333 315 298 283

1 0-9

10"10

CD

AIloy 690 24-26%-CR ----------- --............... PWR Water 22-26 cc/kg H2

* A -

Q 165 ---------- =/mole kJ10 /mole Kmax = 27.5 MPa m1

/2

--------. .--------- .- .-------- -............ .. .... .... .

GE (C372, C400)-SL

ANL-ST, SL

A - - Bettis-24% CR ST

10-10

10-9

Z

E

XD0)

Alloy 690 PWR Water 22-26 cc/kg H2 ......---........... --

Q ='165 kJ/mole

max =27.5 MPa m 1 12

Q = 148 kJ/mole

-K 27E MPa mV2

26% CR

CR -'2o% (GE) ------------------------

0 100 kJ/mole AR

K max= 21 MPa m 1

1

2

maI xI

10-11

10-12

1.5

10-9

10-11

10-12

5

IIIlIIIllIllltIIlIlIll[I

1.60 1.65 1.70 1.75

1000/T (°K)

1.80 1.55 1.60 1.65 1.70

1000ff (°K)

(b)

1.75 1.80

(a)

Temperature (°C)

352 333 315 298 2. 13

10-10

0D02

Alloy 690 TT

....PW R W ater 30 cc/kg H 2 -------- - ----------

A - . . Q 86 kJ/mole

* * Km 3OM= 30Pam 12

* ~A:

- 20% CW TT690 (TL)

Figure 83.

Temperature dependence of the CGR data for

Alloy 690: (a) 26% and 24% cold-rolled and asreceived conditions obtained at ANL, GE,3 8

and Bettis 37 ; (b) 26% and 20% cold-rolled and

as-received conditions obtained at ANL and

GE 3 8 ; and (c) 20% and 50% cold-rolled

conditions obtained at INSS. 4 3

10-11

-12 10 . . . . . . . . . . . . . . . .

1.55 1.60 1.65 1.70 1.75 1.80

1000/T (-K)

(c)

In summary, despite the large scatter in CGR data, the results to date indicate that the activation energy

for Alloy 690 is strongly affected by the deformation level of the alloy. Although there are substantial

uncertainties in the present activation energies, alloys in the as-received condition and at deformation

levels of up to 20% cold work appear to have an activation energy for SCC growth similar to that for

Alloy 600 (130 kJ/mol). At higher deformation levels (24-26%), the effect of temperature seems to

weaken, and by 50% cold work, the temperature appears to have no effect on SCC CGR.

Figure 84 explores the effect of normalizing the available SCC CGR data on cold-rolled alloys to 320'C

(608'F) when using an activation energy of 130 kJ/mol. The plot contains the same data set shown in

Fig. 81, except for the data obtained on the alloy cold-rolled by 50%. Data obtained at GE on 20-26%

deformed alloy at 360'C (680'F) were normalized and included. The figure shows reasonable agreement

among the data sets obtained at different laboratories. For this data set (Alloy 690 deformed by 20-32%),

the Alloy 600 (MRP-5523) 50th percentile curve seems to bound most of the data.

90
Z,

E

L)

E

k6

6,

U,

Alloy 690 20-32% CW

All data normalized to 320*C using Q = 130 kJ/mole

Simulated PWR environment 1 F 23-30 cc/kg H2 CGR Curv'e-Aloa 6-00 ....... ........

75th Percentilea

10-1° 0hPrcni

--- -- --- -- .... •...........

S Bettis 24% CR ST (23 ccikg)

------ Bettis 24% CR St. (23 cr,;kg)

Bettis 32% prestrain SL

- Betis 32% prestrain ST

INSS 20% CR TL (30 cc/kg)

* GE 26% CR SL (20 cclkg)

GE 20% CR (20 cclkg)

" ANL 26% CR A69OWC-SL-1 (23 cc/kg)

A ANL 26% CR A69OWC-ST-1 (23 ccfkg)

1013 .................... ANL 26% CR A690WC-ST-2 (23 cc/kg)

: .. ... . .

Figure 84.

SCC CGRs in PWR water for

Alloy 690 heats deformed by

20-32%. The data were

normalized to 320=C (608°F) in

de-aerated water using an

activation energy of 130 kJ/mol.

10 20 30 40 50 60

Stress Intensity K (MPa0 1 2 )

5.1.4 SCC of Cold-Worked Alloy 690 from a Mechanistic Perspective

The fundamental processes responsible for primary water SCC in Alloy 600 are not fully understood.

Nevertheless, some key dependencies are known, and the current section attempts to discuss the SCC data

for Alloy 690 based on the Alloy 600 experience. Specifically, the discussion considers a grain boundary

deformation mechanism for SCC and attempts to explain the relatively high SCC CGRs measured on

cold-worked Alloy 690.

This discussion is prompted by recent data by Arioka et al.,4 3 at INSS, who found an excellent correlation

between creep rates and susceptibility to SCC of Alloy 690 (Fig. 85a). The data also suggest that the

activation energy for creep and SCC is similar; therefore, creep is a limiting factor for SCC.

Fundamentally, the finding is very similar to that of Was et al.44 for Alloy 600 (Fig. 85b), who showed

that the steady-state creep rate and the percent cracked boundaries both correlate well with the CSLB

fraction. The manner by which the creep correlates with cracking via the CSLB fraction suggests that

creep and cracking are intimately linked. Since the macroscopic creep rate also reflects the rate of

deformation at grain boundaries, the increased resistance of CSLBs to cracking may stem from their

increased resistance to creep deformation. The similarity of the two data sets implys a fundamentally

identical cracking mechanism for the two alloys and warrants a discussion of SCC behavior of Alloy 690

from the perspective of the proposed mechanisms for SCC of Alloy 600.

91
Temperature (°C)

352 333 315 298 283 10-7

!0-9

10-10

(D0

. . I . . . . I . . . . I . . . . ! . . . . ! . . . . Afloy 690 Tr

PWR Water 30 cclkg H2 .........

Q 86 U/mole

Ka 30 MPa m112

..... L.......... - - -n l .--- ...... --- -------... .. SA

AA

Q 103 kJ/mole

* 20% CW TT 690 (SCC)

20% CW MA 600 oi.ieep A 20% CW TT 690 (creep)

N 20% CW MA 690 (creep)

ZI

S10 -8

GC

7

5i

4

21

I

0

10-11

10-12 1 0 9 .- 0.. 2. .... .. .4 0 . .L

0.10 0.15 0.20 0.25 0.30 0.35 0.40 0.45 0.50 1.30 1.40 150 1.60

1O00/T (°K)

1.70 1.80

Fraction of CSLBs

(a) (b)

Figure 85. (a) Temperature dependence of SCC CGR data for 20% cold-rolled Alloy 690 and

creep data for Alloys 690 and 600 43 and (b) dependence of steady-state creep at

360 0C and percent cracked boundaries in primary water on the fraction of CSLBs in

Ni-16Cr-9Fe alloys."

Some of the most convincing evidence that grain boundary deformation is a critical element in IG SCC

stress comes from a study by Vaillant et al.,4 5 in which they were able to correlate IG SCC in Alloys 600

and 690 in the MA and TT conditions tested in primary water at 360'C (680'F) to grain boundary

deformation data from tests in Ar at the same temperature. Their results are summarized in Fig. 86,

which provides a relationship between the crack growth rate, 6 (as measured in constant extension rate

experiments, not fracture mechanics tests), and the grain boundary viscosity, q. The grain boundary

viscosity is a function of several factors, including stress, grain size, grain boundary carbide coverage,

and most importantly, grain boundary sliding rate. A direct cause-and-effect relationship between grain

boundary deformation and SCC initiation for Alloy 600 was demonstrated by Alexandreanu and Was.46

The mechanism was further supported by an analysis of the SCC crack path in Alloy 182 weld. 8

I1-

0.1-

da/dt

(pzmfh)

0.01-

*600 M"

600 TIT

6.242 9M690 MAA a `242262 a 30L I

6*L 6 1.2

0 61'

9G4

61'

*dweCiat limit

(* (IPm In2 000 h) V9.7,9

Figure 86.

Crack growth rate in Alloy 600 and

690 samples in the MA and TT

conditions vs. the grain boundary

viscosity, q/.45

f i ! . . ... .. ...... ........ . .. ....... .. .. .......

106

TI (MPajs)

lp 1010

92
In the framework of a mechanism for IG SCC induced by grain boundary deformation, one can envision

that those factors that significantly affect this deformation (chromium, grain boundary carbides, etc.) also

affect SCC behavior. While the chromium content of the alloy is known to affect a number of processes

that bear on SCC susceptibility through corrosion mechanisms, chromium is also a solid solution

strengthener. The addition of chromium increases the stacking fault energy and decreases the dislocation

absorption rate in the grain boundary by hindering dislocation motion. As such, in terms of mechanical

behavior, the primary effect of chromium would be to suppress grain boundary cavitation and sliding,

hence, SCC initiation. A high density of grain boundary carbides would reduce SCC susceptibility by

hindering grain boundary sliding directly.

Cold-worked Alloy 690 showed relatively high CGRs, comparable to those for Alloy 600 in some

instances (Fig. 81). In this context, the comparison drawn by Arioka et al. 43 between creep in cold-rolled

Alloy 690 and Alloy 600 and CGRs in Alloy 690 and Alloy 600 is extremely significant. The parallelism

indicates a similar behavior for the two alloys and, further, suggests that the mechanisms developed for

Alloy 600 have a direct application to Alloy 690. It is also important to note that the increased chromium

content was believed to suppress grain boundary cavitation; however, this was observed in cold-worked

Alloy 690. In fact, grain boundary cavitation forms the basis of the mechanism proposed by Arioka et

al.43 to explain the high CGRs in cold-worked Alloy 690.

Based on Alloy 600 experience, the relatively high CGRs observed for cold-worked Alloy 690 can be

understood in the framework of a grain-boundary-deformation-induced mechanism for SCC. The

microstructural features described previously - TiN inclusions, cracked grain boundary carbides, carbide

"streaks," and grain size banding - all can affect the susceptibility to grain boundary deformation. Also,

TiN inclusions can impede intra-granular dislocation motion and, thus, improve creep properties. While

these inclusions would benefit the overall IG SCC resistance, they are not expected to affect the grain

boundaries directly. It is not clear how the "streaks" of thick grain boundary carbides might affect the

SCC behavior of an alloy. In one scenario they could be expected to benefit IG SCC behavior by

hindering grain boundary sliding. However, in a heavily cold-worked microstructure, those carbides

could crack, creating regions with a large amount of damage both at grain boundaries and in the matrix.

This effect coupled with a strain differential due to grain size banding could provide a path -for SCC,

perhaps in a mixed IG/TG fracture mode. Nevertheless, the one feature that would be detrimental to the

IG SCC behavior is the cracked grain boundary carbides in the cold-worked materials. One can readily

envision that a grain boundary with cracked carbides is more susceptible to deformation (sliding) and,

therefore, more susceptible to cracking. All microstructures for which high CGR rates were measured

had cracked grain boundary carbides, including the alloys cold-rolled by ANL and GE (Figs. 15 and 16).

Such cracked grain boundary carbides were also observed by Arioka et al.43 in the INSS 20% cold-rolled

alloy. Finally, carbides on the grain facets on the fracture surfaces of these specimens (similar to those

shown in Fig. 40 for a specimen tested at ANL) were observed by Paraventi et al.37 in several Alloy 690

heats deformed to various levels of cold work by both cold-rolling and tensile pre-straining. It is

important to note that cracked carbides on the fracture surface seem to be a unique feature of cold-worked

Alloy 690 alloys, and were not observed on cold-worked Alloy 600 specimens. For example, Fig. 87

shows the fracture mode of Alloy 600 cross-rolled by 30% and tested in a similar PWR environment. The

fracture mode observed for this specimen was smooth IG, with no evidence of cracked carbides of the

facets.

93
Figure 87.

Smooth IG fracture in Alloy 600

cross-rolled by 30%. Crack

advance is from left to right.

The metallographic examinations reported to date show that a cold-work level as low as 20% (INSS Alloy

690)43 can cause grain boundary carbides to crack. However, the extent of carbide cracking in these

alloys is not clear. Unlike the tensile pre-strained alloy where the deformation is uniformly distributed in

the material, in cold-rolled alloys, depending on the cold-work level and number of passes undertaken to

achieve that level, the deformation is likely to be nonuniform. In fact, hardness measurements performed

on cold-rolled alloys of interest for this program have indicated that the deformation concentrates near the

surfaces in contact with the rolls (Fig. 12). Similar measurements are not available for other alloys that

displayed relatively high SCC CGRs, so the microstructures and extent of grain boundary carbide

cracking cannot be readily compared. Likewise, based on the limited number of metallographic

examinations available, the extent of carbide cracking and SCC CGRs cannot be directly compared at this

time.

5.2. Alloy 152 Weld

5.2.1 Cyclic Crack Growth Rates

A database of cyclic CGRs for Ni-alloy welds in simulated PWR environments is available from120

tests conducted on Alloys 182, 82, 152, and 52 at 243-3450C (469-653°F).42,

47-49 The loading conditions

for these tests include R = 0.1-0.75, Kmax = 20-100 MPa-m112 (18.2-91 ksi-in 1

i

2), and rise time of 0.5-

5000 s. The results indicate little effect of PWR environment on the cyclic CGRs of these welds.

However, only about 10% of the data was obtained under conditions for which significant environmental

enhancement would be expected.

As described previously, the approach taken in ANL tests is to first determine those conditions that

display environmental enhancement, and then use those conditions to transition from a TG to an IG

fracture mode. For example, Fig. 88 shows the cyclic data for an Alloy 182 specimen tested previously in

the program. Environmentally enhanced conditions were identified, and this specimen was transitioned

successfully to an IG fracture mode. Figure 89 focuses on this transition. The images indicate that IG

secondary cracks apparently preceded the interdendritic IG fracture. Similar behavior was observed on

Alloy 152 specimen A152-TS-2 (Fig. 70). The IG crack (indicated by an arrow in Fig. 70) apparently

initiated well ahead of the main crack front; however, it propagated sideways, toward the side groove of

the specimen.

94
10o-7

10-8

0~ 10-10

-Alloy18 2 Weld -- ------

Simulated PWR Water at 320°C

Best Fit Curve for Alloy 182

CGR~r + 0018 (CGR.,)07 U -, . .. . ..

------- - ---- -- - --- - -- -- -- - --

Figure 88.

Cyclic CGR data for Alloy 182 weld in PWR

environment at 320°C (6080F).

10-12

10- 12 10-11 10-1w 10-9 1ao -8 10-7

CGRair Ni-w~eld alloys (m/s)

ka) oD)

Figure 89. (a) Transition from TG to IG fracture and (b) higher magnification micrograph of the

boxed area showing the tip of a few secondary cracks. Crack extension from bottom

to top.

Figure 90 shows the cyclic CGR data for the two Alloy 152 weld specimens tested in a PWR environment

at 3200C (6080F). For comparison, similar cyclic CGR data obtained by Van der Sluys et al.42 of B&W

are included. The cyclic CGR data sets obtained from the two different Alloy 152 welds are in good

agreement. The data resulting from the two specimens tested at ANL suggest that there exist conditions

of R and rise time that lead to environmental enhancement. The conditions that resulted in environmental

enhancement of cyclic CGRs in Alloy 152 are load ratios of R z 0.5 and rise times of 600-1000 s. At R

values beyond 0.5 the environmental enhancement begins to diminish. When R is 0.7 and the rise time is

1000 s, the rate in the environment equals that in air, that is, the environmental enhancement appears to be

completely lost. This behavior is consistent with that observed for Alloy 690 in previous tests (Figs. 78

and 80) and suggests that it may be difficult to maintain a CGR under true constant load.

95
I I I I I ... I I

10-7 -Aloyl52 Weld -.........

Simulated PWR Water at 320*C

1 0 -8 . .. . .. . . .. . .. . .. . .. . .. . .. . . ... . .. ... ..

Best Fit Curve for Alloy 182 A

CGRar + 0.018 (CGRa r)078 Z

E 10_9 - - - ------ ----------- Figure 90.

Cyclic CGR data for Alloy 152 weld in PWR

0 10-10 01-------------environment "- " ....... . ..... . . ..........J ...... ..... ...... ..... at 320°C (608°F) obtained at

ANL and B&W.42

1 0 -11 ......... ....... . A152-TS-2

A152-TS-4

A A152 (Van Der Sluys)

10-12 10-11 10-10 10-9 10-8 10-7

CGRair Ni-weld alloys (m/s)

5.2.2 SCC Crack Growth Rates

The available SCC CGRs obtained on Alloy 152 weld are plotted as a function of stress intensity factor K

in Fig. 91. The proposed 75t and 25" percentile disposition curves for Alloys 182/82 (EPRI MRP-1 1524)

are also included for comparison. The plot shows the data obtained at ANL from the two tests described

previously in this report. The SCC CGR data were measured at constant load (CL), measured at constant

load with periodic unloading, or calculated from cycle + hold conditions (est CL). For the data obtained

elsewhere,4°, 38,

50 the hold periods are provided in the legend. To analyze those data points that may

include mechanical fatigue components, CGRai curves were also included. These were obtained from Eq.

11 by noting that for a constant frequency, hold time, and R, CGR8 ir is a function of only Kmax and

temperature.

The CGRs obtained at KAPL40 are similar to those obtained at ANL, although the nominal K values are

much higher. This similarity is remarkable given the fact that the KAPL tests did not involve any period

of transitioning. The alloys used in KAPL tests contained 27.3% and 29.8% Cr, and the results suggest

that this modest decrease in Cr content resulted in a two-fold increase in CGRs. The reported increase in

IG SCC engagement was ten-fold. The differences in IGSCC engagement may make comparisons of

CGRs tentative, but the difference in engagement may be a better measure of relative SCC susceptibility.

For comparison, the Alloy 152 filler that was used for the ANL weld had a Cr content of 28.7%, and an

analysis found the Cr content of the deposited weld to be 28.4%. The uncertainty in Cr content (measured

by inductively coupled plasma mass spectrometry) is 5%, which is fairly large compared to the measured

differences in Cr concentration. Nevertheless, these observations taken together suggest a strong

dependence of SCC susceptibility on Cr content in the 27-30% range.

Figure 91 includes additional data from the experimental programs at GE38 and Pacific Northwest

National Laboratory (PNL). 50 The testing protocol used by the two laboratories was nearly identical and

always used R = 0.7, f = 0.001 Hz, and 2.5-h hold as a transitioning step. For the purpose of analyzing

these data sets, the air fatigue curves, CGRair, calculated for these conditions at two temperatures, 320°C

(608°F) and 350°C (662°F), are included in the plot. These CGRair curves are shown with dashed lines.

With one exception, the CGR data with 2.5-h hold (purple open symbols) do not seem to exhibit the

dependence on K expected for SCC CGRs (e.g., the disposition curves24), but appear to be parallel to the

96
CGRaff curve. This trend suggests that SCC is not the main driving force for these rates, which appear to

be driven primarily by corrosion or mechanical fatigue. Further increases of hold times to 24 h appear to

decrease the rates accordingly - a fact that substantiates the previous point. Finally, constant-load

conditions confirm that the SCC components are minimal. This observation is consistent with the lack

of,38 or minimal, 50 IG engagement reported for these tests.

The proposed 25h percentile disposition curve for Alloys 182/82 (EPRI MRP-1 1524) appears to bound

most of the data.

A A152-TS-2 (CL) Ni-Alloy Weld Metal CGR Curve 0 A152-TS-4 (est CL)

320-350'C Ni-Welds (

A KAPL MLTS-27Cr (24 h hold)

PNL (2.5 h hold)

GE (2.5 h hold)

PNL (24 h hold)

CD

10"11 .......... .......... "32 °

350'C

a)

a)/s " " 0,' \CGRai (mis) C. CGR Curve CG ai

.- NR = 0.7, f = 0.001, 2.5 h hold

25 Percentile A

A& 10-13

/'p

10 20 30 40 50

Stress Intensity K (MPa.m11 2)

Figure 91. SCC CGR data for Alloy 152 welds as a function of stress intensity in PWR

environment.

97

6. Summary

Tensile and CGR tests were performed on Alloy 690 in the cold-rolled and as-received conditions and on

Alloy 152 weld in the as-welded condition. Samples were also subjected to metallographic examination.

The principal findings are summarized as follows:

1. The microstructure of Alloy 690 plate cold-rolled by 26% contains TiN inclusions and cracked

boundary carbides. The areas close to the two edges of the plate are decorated with carbide

streaks - regions with thick grain boundary carbides. In addition, grain size banding occurs

near both edges. The mid-plane of the plate - the region where the CGR tests were conducted

- is free of carbide streaks and grain size banding. The CSLB fraction is approximately 28%,

which is much less than is typical for a Ni-base alloy in this condition. Hence, the proportion

of SCC-susceptible random boundaries is approximately 50% higher than that in a typical

alloy.

2. The microstructure of CRDM Alloy 690 is typical of that of a Ni-base alloy, with equiaxed

grains and dense grain boundary carbide coverage. The grain boundary character distribution is

also typical of a Ni-base alloy. The CSLB fraction is approximately 53%, a large proportion of

which is twins.

3. The microstructure of Alloy 152 is also typical of that for a Ni-base alloy weld, with Al-oxide

inclusions. The grain boundary character distribution is also typical of a Ni-base alloy weld.

The CSLB fraction is approximately 36%, a large proportion of which is low-angle boundaries.

4. The tensile properties of the alloys tested in the program over a wide temperature range 25-

870°C (81-1598 0F) meet the minimum requirements of ASME SB-167, and agree well with

data obtained elsewhere on similar alloys. As such, the data set for Alloy 690 is in good

agreement with control data on Alloy 600. In addition, the data set for Alloy 152 is in good

agreement with the Alloy 82/182 tensile data obtained elsewhere.

5. The cyclic CGR data for Alloy 690 cold-rolled by 26% indicate significant environmental

enhancement, comparable to that observed for Davis-Besse Alloy 600. The cyclic CGR data

for the as-received alloy also indicate some environmental enhancement, although much less

than that for the cold-rolled alloy. Moreover, the cyclic rates for the as-received alloy are

smaller than the corrosion fatigue curve for Alloy 600. Hence, Alloy 690 in the as-received

condition is expected to be more resistant to SCC than Alloy 600.

6. The SCC CGR in the PWR environment at 320'C (608'F) for Alloy 690 cold-rolled by 26% is

1-3 x 10-11 m/s for moderate stress intensity factors. These rates agree well with data obtained

elsewhere on similar alloys deformed in the 20-32% range by either cold-rolling or tensile prestraining. The cracked grain boundary carbides are a common feature of all cold-worked alloys

that display relatively high CGR. The SCC CGR for as-received Alloy 690 is about an order of

magnitude lower. These results imply that while Alloy 690 in the as-received condition is more

resistant to SCC compared to Alloy 600, cold work can increase SCC susceptibility.

7. The available data on the temperature dependence of' SCC CGRs obtained at ANL and

elsewhere indicate that the activation energy for Alloy 690 is strongly affected by the

deformation level of the alloy. Alloys in the as-received condition and at deformation levels of

99
up to 20% cold work appear to have an activation energy close to that for Alloy 600 (130

kJ/mol). At higher deformation levels (24-26%), the effect of temperature seems to weaken,

and by 50% cold work, the temperature appears to have no effect on SCC CGRs.

8. The cyclic CGR data for the two Alloy 152 specimens tested suggest that loading conditions

(load ratios of R = 0.5 and rise times of 600-1000 s) lead to environmental enhancement. At R

values beyond 0.5, the environmental enhancement begins to diminish. When R is 0.7 and the

rise time is 1000 s, the rate in the environment equals that in air, that is, the environmental

enhancement appears to be completely lost. The enhanced rates are always lower than the

corrosion fatigue curve for Alloys 82/182. Hence, it is expected that Alloy 152 is more

resistant to SCC than Alloys 82/182.

9. The SCC CGRs in PWR environment at 320'C for as-welded Alloy 152 are 1.6-5 x 10-11 m/s

for moderate stress intensity factors. The 25th percentile of the proposed CGR curve for Alloys

82/182 (EPRI MRP-1 15) bounds the available data.

100
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105

NRC FORM 335 U.S. NUCLEAR REGULATORY COMMISSION 1. REPORT NUMBER

(12-2010) (Assigned by NRC, Add Vol., Supp., Rev.,

NRCMD 3.7 and Addendum Numbers, If any.)

BIBLIOGRAPHIC DATA SHEET

(See instructions on the reverse) NUREG/CR-7137

2. TITLE AND SUBTITLE 3. DATE REPORT PUBLISHED

Stress Corrosion Cracking in Nickel-Base Alloys 690 and 152 Weld in Simulated PWR MONTH YEAR

Environment - 2009 June 2012

4. FIN OR GRANT NUMBER

5. AUTHOR(S) 6. TYPE OF REPORT

B. Alexandreanu, Y. Yang, Y. Chen, and W. J. Shack Technical

7. PERIOD COVERED (Inclusive Dates)

5/2006 to 9/2009

8. PERFORMING ORGANIZATION - NAME AND ADDRESS (If NRC, provide Division, Office or Region, U. S. Nuclear Regulatory Commission, and mailing address; if

contractor, provide name and mailing address.)

Argonne National Laboratory

Nuclear Engineering Division

9700 S. Cass Avenue

Argonne, IL 60439

9. SPONSORING ORGANIZATION - NAME AND ADDRESS (If NRC, type "Same as above", if contractor, provide NRC Division, Office or Region, U. S. Nuclear Regulatory

Commission, and mailing address.)

Division of Engineering,

Office of Nuclear Regulatory Research,

U.S. Nuclear Regulatory Commission,

Washington, DC 20555-0001

10. SUPPLEMENTARY NOTES

11, ABSTRACT (200 words or less)

Alloys 690 and 152 are the replacement materials pressurized water reactors (PWRs) for Alloys 600 and 182 that have been found to

undergo stress corrosion cracking (SCC). The objectives of this work were to determine the crack growth rates (CGRs) in a simulated

PWR water environment and obtain tensile data as a function of temperature for the replacement alloys. Compared to Alloy 600, the

as-received Alloy 690 proved to be more resistant to SCC in the simulated PWR environment at 320'C (608 IF). However after cold

rolling, the SCC CGRs increased by an order of magnitude compared to as-received Alloy 690. The environmental enhancement of

cyclic CGRs for Alloy 152 was minimal and SCC CGR rates for the higher chromium weld metal in the simulated PWR environment

at 3200C (608 °F), was about an order of rnagnitude lower than the typical SCC CGR measured for Alloy 182. Tensile test data for

Alloy 690 and Alloy 152 welds in the temperature range of 25-870'C (81-1598°F) meet the minimum requirements of ASME SB-167,

and are in agreement with data on similar alloys.

12. KEY WORDS/DESCRIPTORS (List words or phrases that will assist researchers in locating the report.) 13. AVAILABILITY STATEMENT

alloy 690 unlimited

alloy 152 14. SECURITY CLASSIFICATION

primary water stress corrosion cracking (This Pege)

crack growth rate unclassified

tensile test data (rhis Report)

pressurized water reactor unclassified

15. NUMBER OF PAGES

16. PRICE

NRC FORM 335 (12-2010)
Federa Rocycllng Program

UNITED STATES

NUCLEAR REGULATORY COMMISSION

WASHINGTON, DC 20555-0001

OFFICIAL BUSINESS
NUREG/CR-7137 Stress Corrosion Cracking in Nickel-Base Alloys 690 and 152

Weld in Simulated PWR Environment 

- 2009

June 2012