Degradation of LWR Core 

Internal Materials due to

Neutron Irradiation 

Office of Nuclear Regulatory Research 

NUREG/CR-7027
 

Degradation of LWR Core

Internal Materials due to

Neutron Irradiation

Manuscript Completed: January 2010 

Date Published: December 2010 

Prepared by

O.K. Chopra

Argonne National Laboratory

9700 South Cass Avenue 

Argonne, IL 60439

A.S. Rao, NRC Project Manager 

NRC Job Code N6818

Office of Nuclear Regulatory Research 

NUREG/CR-7027
NUREG/CR-7027

ANL-10/11

Degradation of LWR Core Internal 

Materials due to Neutron Irradiation 

Manuscript Completed: January 2010

Date Published: December 2010

Prepared by

O. K. Chopra 

Environmental Science Division

Argonne National Laboratory

Argonne, IL 60439

Appajosula S. Rao, NRC Project Manager

Prepared for

Division of Engineering 

Office of Nuclear Regulatory Research

U.S. Nuclear Regulatory Commission

Washington, DC 20555

NRC Job Code N6818
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Abstract

Austenitic stainless steels are used extensively as structural alloys in the internal components of 

light water reactor (LWR) pressure vessels because of their relatively high strength, ductility, and fracture 

toughness. However, exposure to neutron irradiation for extended periods changes the microstructure and 

microchemistry of these steels and degrades their fracture properties. This report presents a critical 

assessment of the susceptibility of LWR core internal materials to irradiation effects such as irradiationassisted stress corrosion cracking (IASCC), neutron embrittlement, void swelling, and irradiation-induced 

stress relaxation. The existing data, in the open literature as well as from research carried out by the U.S. 

Nuclear Regulatory Commission (NRC) and industry, have been evaluated to establish the effects of 

material parameters (such as composition, thermo-mechanical treatment, microstructure, microchemistry, 

yield strength, and stacking fault energy) and environmental parameters (such as water chemistry, 

irradiation temperature, dose, and dose rate) on these processes. Differences in radiation-induced 

degradation in material properties between light water reactors and fast-reactor irradiations are also 

discussed. The results are used to (a) define a threshold fluence above which irradiation effects on 

materials properties are significant, (b) develop disposition curves for cyclic and IASCC growth rates for 

reactor core internal materials, and (c) assess the significance of void swelling and irradiation creep 

relaxation on the structural and functional integrity of reactor internal components. The report also 

identifies potential deficiencies or knowledge gaps in the existing experimental data on degradation of 

LWR core internal materials due to neutron irradiation. 

Paperwork Reduction Act Statement

This NUREG does not contain information collection requirements and, therefore, is not subject to 

the requirements of the Paperwork Reduction Act of 1995 (44 U.S.C. 3501 et seq.).

Public Protection Notification

The NRC may not conduct or sponsor, and a person is not required to respond to, a request for 

information or an information collection requirement unless the requesting document displays a current 

valid OMB control number.
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Foreword

The internal components of light water reactor (LWR) pressure vessels are fabricated primarily 

with austenitic stainless steels because of their relatively high strength, ductility, and fracture toughness. 

During normal reactor operational conditions, the internal components are exposed to high-energy 

neutron irradiation and high-temperature reactor coolant. However, prolonged exposure to neutron 

irradiation changes both the microstructure and the microchemistry of these stainless steels such that 

strength typically increases while ductility and fracture toughness decrease. Irradiation and the 

environment also may decrease the resistance of core internal materials to both fatigue crack propagation 

and stress corrosion cracking. These mechanical property changes are caused by a complex interaction 

among irradiation, stress, and corrosion. As nuclear power plants age and neutron irradiation dose 

increases, the degradation of the core internal components becomes more likely. This represents an 

important safety issue associated with long-term operation of nuclear power plants. The U.S. NRC has

been sponsoring research on reactor core internals to (a) define threshold fluence levels above which 

irradiation begins to affect material properties, (b) develop disposition curves for cyclic fatigue and 

irradiation assisted stress corrosion cracking (IASCC) growth rates, and (c) assess the significance of void 

swelling and irradiation stress relaxation/creep on the structural and functional integrity of these 

components. 

The objective of this report is to summarize the current knowledge associated with the effect of the 

reactor environment and neutron irradiation on the mechanical properties of the reactor core internals and 

to identify gaps in this knowledge that could affect the long-term performance of these components. 

Significant gaps should be addressed in future industry- or NRC-sponsored research. This report briefly 

reviews the results of the existing open literature data (on IASCC, neutron embrittlement, void swelling, 

and irradiation-induced stress relaxation/creep) that have been published by the NRC, industry, academia,

and other research agencies. The report also identifies potential deficiencies or knowledge gaps in the 

existing experimental data with respect to defining the threshold fluence levels above which irradiation 

begins to affect material properties; the disposition curves for cyclic fatigue and IASCC growth rates; and 

the significance of void swelling and irradiation stress relaxation/creep on the structural and functional 

integrity of these components.

The results of this final report may be used to (i) support evaluation of the acceptability of reactor 

internal aging management programs in license renewal applications, (ii) support revision of the Generic 

Aging Lessons Learned (GALL) report, (iii) determine appropriate inspection and flaw disposition 

procedures for reactor internals for use in developing ASME code and for determining needed conditions 

within Title 10, Section 50.55a, “Codes and Standards,” of the Code of Federal Regulations (10CFR 

5055a), and (iv) identify issues related to long-term operation (i.e., greater than 60 years) in nuclear 

power plants. This report will also be used by the NRC to evaluate the acceptability of planned industryand NRC-sponsored research to address gaps identified in this report.

Michael Case

Director Division of Engineering

Office of Nuclear Regulatory Research

U.S. Nuclear Regulatory Commission
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Contents 

Abstract ....................................................................................................................................................... iii

Foreword ..................................................................................................................................................... v

Contents....................................................................................................................................................... vii

Figures......................................................................................................................................................... ix

Tables .......................................................................................................................................................... xiii

Executive Summary.................................................................................................................................... xv

Acknowledgments ...................................................................................................................................... xxiii

Acronyms and Abbreviations .................................................................................................................... xxv

Symbols....................................................................................................................................................... xxvii

1 Introduction .......................................................................................................................................... 1

2 Irradiation Assisted Stress Corrosion Cracking ................................................................................. 3

2.1 IASSCC Susceptibility........................................................................................................... 3

2.1.1 Microstructure ........................................................................................................ 6

2.1.2 Microchemistry ...................................................................................................... 9

2.1.3 Radiation Hardening.............................................................................................. 12

2.1.4 Tensile Properties .................................................................................................. 14

2.1.5 Yield Strength ........................................................................................................ 22

2.1.6 Corrosion Potential ................................................................................................ 24

2.1.7 Silicon Segregation................................................................................................ 26

2.1.8 Sulfur Content........................................................................................................ 28

2.1.9 Stacking Fault Energy ........................................................................................... 29

2.1.10 LWR vs. Fast Reactor Irradiations..................................................................... 31

2.2 Crack Growth Rates ............................................................................................................... 31
viii

2.2.1 Quality of Crack Growth Rate Data ..................................................................... 33

2.2.2 Stress Corrosion Cracking Growth Rates............................................................. 40

2.2.3 Fatigue Crack Growth Rates................................................................................. 53

2.3 Initiation of Irradiation-Assisted Stress Corrosion Cracking............................................... 57

3 Neutron Embrittlement ........................................................................................................................ 63

3.1 Fracture Toughness of Irradiated Austenitic Stainless Steels.............................................. 64

3.1.1 Fracture Toughness JIc .......................................................................................... 65

3.1.2 Fracture Toughness J-R Curve.............................................................................. 67

3.2 Fracture Toughness Trend Curve .......................................................................................... 72

3.3 Synergistic Effect of Thermal and Neutron Irradiation ....................................................... 76

4. Void Swelling...................................................................................................................................... 81

4.1 Background Information ........................................................................................................ 81

4.2 Assessment of Void Swelling in PWR Core Internal Components .................................... 84

4.3 Embrittlement due to Void Swelling..................................................................................... 87

5. Stress Relaxation and Creep............................................................................................................... 89

6. Summary.............................................................................................................................................. 93

6.1 Susceptibility to IASCC......................................................................................................... 93

6.2 IASCC Growth Rates............................................................................................................. 96

6.3 Fatigue Crack Growth Rates.................................................................................................. 98

6.4 IASCC Initiation..................................................................................................................... 99

6.5 Neutron Embrittlement........................................................................................................... 100

6.6 Synergistic Effects of Thermal and Neutron Embrittlement ............................................... 102

6.7 Void Swelling ......................................................................................................................... 102

6.8 Irradiation Creep Relaxation.................................................................................................. 103

7. Gaps in Available Data ....................................................................................................................... 105
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References................................................................................................................................................... 109

Figures

1. The effect of fast neutron fluence and average water conductivity on susceptibility of 

austenitic stainless steels to IASCC shown in field correlations of the core component 

cracking............................................................................................................................................. 3

2. Dependence of IASCC susceptibility of austenitic stainless steels on fast neutron fluence and 

corrosion potential as measured from SSRT tests in 288°C water. .............................................. 4

3. Susceptibility of irradiated austenitic SSs to IGSCC as a function of fluence in highdissolved-oxygen water. .................................................................................................................. 4

4. Schematic illustrating mechanistic issues that influence IASCC in austenitic stainless steels... 5

5. The change in density and size of interstitial loops as a function of irradiation dose during 

LWR irradiation at 275-290°C. ....................................................................................................... 7

6. Results from the BOR-60 irradiation experiment with the data for austenitic stainless steels 

irradiated in a BWR. ........................................................................................................................ 8

7. Radiation-induced grain boundary Cr and Si concentration in 300-series stainless steels 

irradiated in LWRs at 275-290°C.................................................................................................... 10

8. Dependence of grain boundary Cr concentration for several 300-series austenitic stainless 

steels irradiated at about 300°C on neutron dose. .......................................................................... 10

9. Comparison of radiation-induced segregation of Cr, Ni, andSi in LWR and BOR-60 fast 

reactor irradiated stainless steels. .................................................................................................... 11

10. Engineering stress-strain curves at 289°C for Type 304 SS and Type 316 SS steel irradiated 

up to 3 dpa in the Halden reactor at 289°C..................................................................................... 12

11. The effects of neutron dose on measured tensile yield strength of 300-series stainless steels, 

irradiated and tested at about 300°C. .............................................................................................. 13

12. Effect of irradiation temperature on the yield strength and uniform elongation of CW 

Type 316 stainless steel irradiated in fast reactors at temperatures between 290-500°C and 

tested at the irradiation temperature................................................................................................ 13

13. Change in yield strength, ultimate tensile strength, uniform elongation, and total elongation 

as a function of neutron dose for solution annealed Type 304, 304L, and 347 stainless steels 

at elevated temperature. ................................................................................................................... 15
x

14. Change in yield strength, ultimate tensile strength, uniform elongation, and total elongation 

as a function of neutron dose for CW Type 316 stainless steel at elevated temperature............. 17

15. Change in yield strength, ultimate tensile strength, uniform elongation, and total elongation 

as a function of neutron dose for Type 308 and 309 weld metal, Type 304 HAZ, and CF-8 

cast stainless steel at elevated temperature..................................................................................... 19

16. Effect of yield strength and martensite content on the SCC growth rate of stainless steels and 

Alloy 600 at 288°C in high-purity water containing 2000 ppb O2 and 95 or 1580 ppb H2. ....... 22

17. Crack growth rates of nonirradiated, cold worked, 300-series austenitic SSs in NWC BWR 

environment as a function of yield strength. .................................................................................. 23

18. Change in SCC growth rates of nonirradiated, sensitized and nonsensitized, stainless steels as 

a function of yield strength in NWC BWR environment at 289°C and PWR environment at 

340°C. ............................................................................................................................................... 23

19. Change in SCC growth rates at 289°C of LWR-irradiated austenitic stainless steels as a 

function of yield strength in NWC and HWC BWR environments.............................................. 23

20. Effect of yield strength on percent IGSCC in SSRT tests on irradiated 300-series stainless 

steels in high-DO water. .................................................................................................................. 24

21. Experimental CGR under constant load in NWC and HWC BWR environments at 288°C for 

austenitic stainless steels irradiated to 1.0-2.5 dpa, 3.0-4.0 dpa, and 13 dpa, plotted as a 

function of the steel ECP versus a standard hydrogen electrode. ................................................. 25

22. Crack growth rates versus Si concentration in custom heats of Type 304L stainless steel 

simulating expected grain boundary microchemistries for high and low corrosion potentials 

and at various stress intensity factors.............................................................................................. 27

23. Effect of bulk S content on IASCC susceptibility of commercial heats of Type 304, 304L, 

316, and 316L tested in high DO water at 289°C. ......................................................................... 29

24. Range of bulk S and C contents, in which Type 304 or 316 stainless steels are resistant or 

susceptible to IASCC, in BWR-like oxidizing water. ................................................................... 29

25. Range of bulk S and C contents, in which Type 304, 316, 321, or 348 stainless steels are 

resistant or susceptible to IASCC, in PWR water.......................................................................... 30

26. IASCC susceptibility as measured by percent IG cracking as a function of stacking fault 

energy................................................................................................................................................ 30

27. Comparison of the measured yield strength and uniform elongation for several heats of 

Type 304 and 304L stainless steels irradiated in the Halden reactor at 290°C to 3 dpa and in 

BOR-60 fast breeder reactor at 320°C to 4.8 dpa. ......................................................................... 32

28. Plot of CGR in water vs. CGR in air showing environmental enhancement of growth rates in 

high-purity water at 289°C. ............................................................................................................. 34
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29. Plot of crack length vs. time for irradiated Type 304 stainless steel in high-purity water at 

289°C. ............................................................................................................................................... 35

30. Engineering stress vs. strain curve for Type 304 stainless steel irradiated to 3.0 dpa and 

tested in air at 289 and 325°C.......................................................................................................... 37

31. Strain distribution of a moving crack in a strain-softening and a strain-hardening material. ..... 37

32. Crack growth rates for irradiated stainless steels in BWR environments under loading 

conditions that meet or exceed the proposed K/size criterion....................................................... 38

33. Crack length and K vs. time for Type 304L SS in high-purity water at 289°C. .......................... 38

34. Side view of the first slice cut from Type 304L SS Specimen C3-C............................................ 39

35. Experimental SCC growth rates in NWC BWR environment of austenitic stainless steels 

irradiated to 0.75-3.0 dpa, 0.7-4.25 dpa and 5.5-37.5 dpa............................................................ 41

36. SCC growth rates in HWC BWR environment of austenitic stainless steels irradiated to 0.75-

3.0 dpa, 0.7-4.25 dpa, and 5.5-37.5 dpa.......................................................................................... 43

37. CGR under constant load for thermally aged and irradiated CF-8M cast stainless steel 

specimens in BWR environment at 289°C. .................................................................................... 45

38. Crack growth rates of irradiated austenitic stainless steels in NWC and HWC BWR 

environments at 288°C and K of 20 MPa m1/2 plotted as a function of neutron dose. ............... 47

39. Crack growth rates of irradiated austenitic stainless steels in NWC and HWC BWR 

environments at 288°C and K of 20 MPa m1/2 plotted as a function of yield stress. .................. 48

40. SCC growth rates of austenitic stainless steels irradiated to 3.0-37.8 dpa in PWR 

environment. ..................................................................................................................................... 49

41. Temperature dependence of SCC growth rates in irradiated CW Type 316+Ti stainless steel 

in PWR environment........................................................................................................................ 50

42. Normalized SCC growth rates in PWR environment for austenitic stainless steels irradiated 

to 3 dpa, 11-25 dpa, and 6-38 dpa. .................................................................................................. 50

43. Normalized SCC growth rates of austenitic stainless steels irradiated to 0.7-4.25 dpa in HWC 

BWR and PWR environments at 288°C. ........................................................................................ 52

44. Crack growth rates for irradiated austenitic stainless steels in HWC BWR water at 288°C or 

PWR environment at 320°C plotted as a function of neutron dose. ............................................. 52

45. CGR data under cyclic loading for irradiated SS weld HAZ materials in air at 289°C. ............. 54

46. CGR for irradiated specimens of austenitic SSs under continuous cycling at 289°C in highpurity water with ≈300 ppb and <30 ppb dissolved oxygen.......................................................... 55
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47. CGR data under cyclic loading for nonirradiated SS weld HAZ materials in high-purity water 

at 289°C. ........................................................................................................................................... 56

48. CGR for irradiated specimens of Type 304L SA weld HAZ from the Grand Gulf core shroud 

and laboratory-prepared Type 304 SS SMA weld HAZ under continuous cycling in highpurity water at 289°C. ...................................................................................................................... 56

49. CGR data under cyclic loading for two irradiated specimens of CF-8M cast austenitic SS in 

high-purity water at 289°C. ............................................................................................................. 57

50. Stress as percent of irradiated yield stress vs. time data for IASCC flaw initiation in 

austenitic stainless steels in a PWR environment. ......................................................................... 58

51. Stress as percent of irradiated yield stress vs. neutron dose for IASCC flaw initiation in 

austenitic stainless steels in a PWR environment. ......................................................................... 59

52. Percent of irradiated yield stress vs. time for IASCC flaw initiation in a PWR environment, 

in an austenitic stainless steel irradiated to 71-76 dpa. .................................................................. 59

53. Time for initiation of IASCC in irradiated austenitic stainless steels as a function of stress...... 60

54. Change in fracture toughness JIc as a function of neutron exposure for austenitic SSs 

irradiated in fast reactors.................................................................................................................. 66

55. Change in fracture toughness JIc as a function of neutron exposure for irradiated austenitic 

SSs..................................................................................................................................................... 67

56. Change in fracture toughness JIc as a function of neutron exposure for SSs. .............................. 68

57. Fracture toughness J-R curves for sensitized Type 304 SS, weld HAZ materials of Type 304 

and 304L SS, and CF-8M cast SS in high-purity water at 289°C................................................. 69

58. Fracture toughness J-R curves for irradiated specimens of Type 304 SS SMA weld HAZ and 

Type 304L SA weld HAZ in air and NWC BWR environment. .................................................. 70

59. Fracture toughness J-R curves for sensitized Type 304 SS in air and NWC BWR water at 

289°C. ............................................................................................................................................... 70

60. Fracture toughness J-R curves and load vs. loadline displacement curves for thermally aged 

and irradiated CF-8M cast SS.......................................................................................................... 71

61. Fracture toughness JIc of irradiated austenitic stainless steels and welds as a function of test 

temperature. ...................................................................................................................................... 71

62. The change in initiation toughness JIc of wrought austenitic SSs and cast austenitic SSs and 

weld metals as a function of neutron exposure. ............................................................................. 72

63. Change in coefficient C of the power-law J-R curve for wrought austenitic SSs and cast 

austenitic SSs and weld metals as a function of neutron exposure. .............................................. 75
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64. Experimental values of J-integral at a crack extension of 2.5 mm for wrought austenitic SSs 

and cast austenitic SSs and weld metals plotted as a function of neutron exposure.................... 78

65. Void swelling of PWR materials plotted as a function of neutron dose....................................... 84

66. Irradiation temperature and neutron dose for the void swelling data shown in Fig. 65. ............. 84

67. Void swelling of PWR materials plotted as a function of effective full power years.................. 86

68. Thermal creep and irradiation-enhanced creep of a 20% CW Type 316 stainless steel 

irradiated in EBR-II.......................................................................................................................... 90

69. Trend line for screening irradiation-induced stress relaxation based on available test data. ...... 91

Tables

1. Material property equations for irradiated CW Type 316 stainless steel. .................................... 14

2. Material property equations for irradiated solution-annealed Type 304 stainless steel............... 14

3. Material property equations for solution-annealed and nonirradiated Type 304 stainless steel. 21

4. Screening criteria for thermal-aging susceptibility of cast austenitic stainless steels.................. 77

5. Fatigue CGR correlations for wrought and cast austenitic stainless steels in BWR 

environments at 289°C..................................................................................................................... 99

6. Data gaps in the current understanding of irradiation effects on LWR core internal materials.. 108
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Executive Summary

Background

Austenitic stainless steels (SSs) are used extensively as structural alloys in the internal components 

of light water reactor (LWR) pressure vessels because of their relatively high strength, ductility, and 

fracture toughness. However, exposure to neutron irradiation for extended periods changes the 

microstructure (radiation hardening) and microchemistry (radiation-induced segregation or RIS) of these 

steels and degrades their fracture properties. This report addresses the concerns arising from irradiationinduced materials issues such as irradiation-assisted stress corrosion cracking (IASCC), neutron 

embrittlement, void swelling, and irradiation-induced stress relaxation. The existing data have been 

reviewed to identify the key material and environmental parameters that influence these processes and to 

determine their effects. The results are used to (a) define a threshold fluence above which irradiation 

effects on materials properties are significant, (b) develop disposition curves for cyclic and IASCC 

growth rates for reactor core internal materials, and (c) assess the significance of void swelling and 

irradiation creep relaxation on the structural and functional integrity of reactor internal components. This 

report also identifies potential deficiencies or knowledge gaps in the existing experimental data on 

degradation of LWR core internal materials due to neutron irradiation.

Susceptibility to IASCC 

Neutron irradiation of austenitic SSs can produce damage by displacing atoms from their lattice 

position, which creates point defects such as vacancies and interstitials. Most of these defects are 

annihilated by recombination. The surviving defects rearrange into more stable configurations such as 

dislocation loops, network dislocations, precipitates, and cavities (or voids). Changes in the 

microstructure of austenitic SSs due to neutron irradiation vary with the material composition, irradiation 

temperature, and neutron fluence, flux, and energy spectrum. At temperatures below 300°C, the material 

microstructure primarily consists of “black spot” defect clusters and faulted dislocation loops, whereas 

large faulted loops, network dislocations, cavities/voids (clusters of vacancies and/or gas bubbles), and 

precipitates are observed above 300°C. At LWR operating temperatures, the loop density saturates at 

about 1 dpa, and the average loop diameter saturates at 5 dpa. The loop size increases and loop density 

decreases with irradiation temperature. Cavities and voids form at high doses and high temperatures. 

High irradiation temperatures (above 350°C) also lead to the formation of second phase particles. Metal 

carbides are the primary stable precipitates in 300-series SSs under LWR conditions, although RIS of Ni 

and Si to sinks may lead to the formation of γ’ phase (Ni3Si) and G phase (M6Ni16Si7).

The microchemistry of the material is also changed due to RIS. Elements such as Si, P, and Ni are 

enriched at regions such as grain boundaries that act as sinks for point defects, while Cr, Mo, and Fe are 

depleted. The extent of RIS depends on the rate of generation and recombination of point defects (i.e., 

irradiation temperature and dose rate). At LWR temperatures, significant segregation is observed at an 

irradiation dose of 0.1 dpa, and the effect either saturates or changes very slowly beyond 5 dpa. 

The point defect clusters and precipitates act, to varying extent, as obstacles to a dislocation motion 

that leads to matrix strengthening, resulting in an increase in tensile strength and a reduction in ductility 

and fracture toughness of the material. The yield strength of irradiated SSs can increase up to five times 

that of the nonirradiated material after a neutron dose of about 5 dpa. In general, cavities (or voids) are 

strong barriers, large faulted Frank loops are intermediate barriers, and small loops and bubbles are weak 

barriers to dislocation motion. For austenitic SSs, the greatest increase in yield strength for a given 
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irradiation level occurs at irradiation temperatures near 300°C (572°F), which is in the temperature range 

of LWR operation. The change in yield strength correlates well with the microstructural changes. The 

yield strength also either saturates or increases very slowly above 5 dpa. Nearly all SSs exhibit strain 

softening (also termed “dislocation channeling”) and little or no uniform elongation at irradiation dose 

above 5 dpa. The enhanced planar slip leads to a pronounced degradation in the fracture toughness of 

austenitic SSs.

The significance of material yield strength on IASCC susceptibility, particularly in the LWR 

environments, is discussed. Data on stress corrosion cracking (SCC) susceptibility indicate that 

nonsensitized SSs with high tensile strength are susceptible to SCC in ultra-high purity water. The 

increases in yield strength can originate from surface or bulk cold work, weld shrinkage strain, and 

precipitation or irradiation hardening; all of these increase SCC susceptibility. At a given yield strength 

below 800 MPa, a similar cracking behavior is observed for different grades of SSs, although the crack 

growth rates (CGRs) in low-potential water [less than 50 ppb dissolved oxygen (DO)] are up to an order 

of magnitude lower. However, at yield strengths of 800 MPa or higher, the CGRs in low- and highpotential environments are also comparable. The significance of Si segregation at grain boundaries, S 

content in the steel, or stacking fault energy of the material is also discussed. 

Correlations, developed by the Materials Reliability Program (MRP), are presented for estimating tensile 

yield and ultimate strengths and uniform and total elongation of austenitic SSs as a function of neutron 

dose. The tensile properties reach saturation at 5-20 dpa and do not seem to change significantly at higher 

dose levels. However, the majority of the data have been obtained on materials irradiated in fast neutron 

reactors (e.g., EBR-II in the U.S. or BOR-60 in Russia), which use fast neutrons to burn more uranium238 than conventional LWRs. Data on LWR-irradiated materials are more limited, in particular above 

10 dpa. The available data indicate that the yield and ultimate strengths of the LWR-irradiated materials 

are typically higher, and uniform and total elongations are lower, than those of materials irradiated in fast 

reactors. 

IASCC Initiation

The existing IASCC initiation data have been reviewed to evaluate the adequacy of the database for 

developing screening criteria for IASCC susceptibility of core internal components in pressurized water 

reactors (PWRs). The screening criteria are defined for a given neutron fluence above a threshold value 

(i.e., 3 dpa), an apparent stress threshold below which IASCC will not occur in a PWR environment. The 

fluence threshold of 3 dpa that has been established by the nuclear industry is based on incidents of 

IASCC failures in operating PWRs and on results from slow strain rate tensile tests in the PWR 

environment. The stress thresholds are defined from IASCC initiation data. However, the database used 

for defining neutron dose/applied stress thresholds for austenitic SSs is inadequate to evaluate variability 

among different grades of SSs, or the differences among solution-annealed or cold-worked materials, or 

the effect of irradiation temperature. Most of the data are on cold-worked Type 316 SS irradiated at 

temperatures below 310°C. Also, although the fluence threshold for IASCC is defined as 3 dpa, there are 

no IASCC initiation data on materials irradiated below 10 dpa. Furthermore, in reactor core internal 

components, IASCC is likely to occur under creviced conditions (i.e., under off-normal water chemistry). 

IASCC Growth Rates

Irradiation-assisted stress corrosion cracking is a complex phenomenon that is dependent on several 

parameters, which interact with each other such that their effect on growth rates is not always the same. 

As a result of these complexities and uncertainties in experimental measurements, the SCC growth rate

data in LWR environments show significant variability, and specific effects of these parameters on 
xvii

growth rates cannot be accurately assessed under all conditions of interest. Crack growth rate data on 

irradiated wrought and cast austenitic SSs as well as SS weld metal and weld heat-affected zone (HAZ) 

materials in LWR environments have been compiled and evaluated to define the threshold fluence for 

IASCC and to develop disposition curves for cyclic and IASCC growth rates for reactor core internal 

materials. The importance of test procedures that closely reproduce the loading and environment 

conditions for reactor core internal components and provide accurate measurements and reproducible 

results is discussed. 

The results indicate that, in the normal water chemistry (NWC) of boiling water reactors (BWRs) 

environment, neutron irradiation up to 0.45 dpa (3 x 1020 n/cm2) has no effect on CGRs. Additionally, 

the rates are below the NUREG-0313 disposition curve for nonirradiated materials in high-purity water 

containing 8 ppm DO. The CGRs for materials irradiated to higher neutron dose levels can be up to a 

factor of 40 higher. The CGRs for some SSs irradiated to 5-8 dpa (corresponding to a 60-year end-of-life 

neutron dose for BWRs) are a factor of 20 higher than the NUREG-0313 curve and 2-3 higher than the 

Electric Power Research Institute (EPRI) disposition curve for BWR core internal components. The CGR 

data for highly irradiated materials show anomalous behavior. The SCC growth rates, in an NWC BWR 

environment, for SSs irradiated to 19 or 38 dpa are below the NUREG-0313 disposition curve for

nonirradiated SSs. The reason for the low CGRs observed for these highly irradiated materials in the 

high-potential BWR environment is currently unknown. However, both materials came from the same 

reactor - a decommissioned PWR. It is possible that this behavior is unique to the materials from this 

reactor, but their microstructural and microchemistry characterization is not available in the open 

literature. 

The SCC growth rates in the hydrogen water chemistry (HWC) BWR environment show a 

significant decrease relative to those in the NWC BWR environment for SSs irradiated to less than 3 dpa; 

little or no reduction for some SSs irradiated to dose levels as low as 3-4 dpa, at least at stress intensity 

factor (K) values above 18 MPa m1/2; and no decrease for all SSs irradiated to 12-14.5 dpa. The CGRs 

for some materials irradiated to 3-4 dpa are above the CGR disposition curve proposed by EPRI for the 

HWC BWR environment. However, these results are often screened out because they marginally 

exceeded the proposed K/size criterion for irradiated materials, in which the irradiation-induced increase 

in yield stress is discounted by a factor of 2 or 3. However, several investigations on the validity of the 

proposed K/size criterion do not support a criterion based on an effective yield stress rather than the 

measured yield stress. Additional CGR data on SSs irradiated to 3-8 dpa are needed to accurately 

establish the threshold for IASCC susceptibility in low-potential environments.

The SCC growth rates in PWR environment show significant variability. The CGRs for the same 

material and irradiation condition increase with increasing temperature. The temperature dependence of 

growth rates may be represented by an activation energy of 100 kJ/mol. In PWR water at 320°C, most of

the CGR data for SSs irradiated to 3 dpa are up to a factor of 5 above the NUREG-0313 curve for 

nonirradiated materials, and those for SSs irradiated to 11 dpa (and SSs irradiated in the BOR-60 reactor 

to 25 dpa) are up to two orders of magnitude above the same curve. Also, as observed earlier for HWC 

BWR water, the CGRs in the PWR environment of SSs irradiated to 19 or 38 dpa exhibit anomalous 

behavior as the rates are below the NUREG-0313 curve for nonirradiated SSs. As discussed above, while 

the reason for these low rates is unknown, both materials came from the same reactor. Additional data on 

PWR materials irradiated at different temperatures (e.g., between 300 and 350°C) are needed to better 

understand the IASCC susceptibility of austenitic SSs in PWR environments.
xviii

Fatigue Crack Growth Rates

Limited fatigue CGR data in air on austenitic SSs irradiated to 2 dpa under LWR irradiation 

conditions indicate little or no effect of neutron irradiation on growth rates. Additional results on SSs

irradiated in fast reactors at 427°C show only slightly higher CGRs (up to a factor of two higher) at low 

ΔK values (less than 44 MPa m1/2). Also, tests on Type 304 and 316 SS irradiated in a thermal reactor at 

288°C to 2.7 dpa and tested at 427°C showed superior resistance to crack growth; crack growth rates were 

25 to 50% lower than those for nonirradiated materials. 

In the NWC BWR environment, the cyclic CGRs of SSs irradiated to 0.45 dpa (3 x 1020 n/cm2) are 

the same as those for nonirradiated materials, whereas the CGRs of SSs irradiated to 0.75-3.0 dpa are 

higher. The cyclic CGRs at low frequencies (i.e., rise times ≥10 s) are lower by more than an order of 

magnitude when the DO level is decreased by changing from NWC to HWC. A superposition model has 

been used to represent the cyclic CGRs of austenitic SSs. The CGR in reactor coolant environment is 

expressed as the superposition of the rate in air (mechanical fatigue) and the rates due to corrosion fatigue 

and SCC. Studies at Argonne National Laboratory (NUREG/CR-6960) indicate that the superposition 

model generally provides the best estimate, or slightly conservative estimates, of fatigue CGRs in NWC 

and HWC BWR environments. The results indicate that for irradiated materials, the higher fatigue CGRs 

at long rise times (i.e., low frequencies) are primarily due to the SCC component of crack growth. 

Fatigue CGR data are not available for austenitic SSs irradiated to higher neutron dose levels in the 

BWR environments or on irradiated austenitic SSs in the PWR environment. Additional fatigue crack 

growth data under these material and environmental conditions are needed to develop fatigue CGR 

correlations for BWR and PWR core internal components. 

Neutron Embrittlement 

The fracture toughness of austenitic SSs has been divided into three broad categories. The fracture 

toughness (JIc) is above 150 kJ/cm2 for Category III materials, and 30-150 kJ/cm2 for Category II 

materials. These materials fracture after stable crack extension at stresses well above or close to the yield 

stress. Category I materials fracture at stress levels well below the yield stress by unstable crack 

extension, and their fracture toughness, as characterized with the critical stress intensity factor (KIc), is 

less than 75 MPa m1/2. Nonirradiated wrought and cast austenitic SSs and their welds typically fall in 

Category III. However, neutron irradiation degrades the fracture toughness of austenitic SSs to the level 

of Category II or even I at high dose levels.

Fracture toughness data on irradiated wrought and cast austenitic SSs and their welds have been 

compiled and evaluated to define the threshold neutron dose above which fracture toughness of these 

materials is reduced significantly. The validity of the fracture toughness J-R curve data and the different 

methods for determining JIc are discussed. Comparison of fracture toughness data obtained on different 

size specimens showed that small specimens yield valid J-R curve data, at least for materials with JIc

below 300 kJ/m1/2. 

The existing fracture toughness data on austenitic SSs irradiated in LWRs indicate little or no loss 

of fracture toughness below an exposure of about 0.5 dpa and a substantial and rapid decrease at 

exposures of 1-5 dpa. Also, fracture toughness appears to saturate at approximately 10 dpa. A similar 

trend was observed earlier for austenitic SSs irradiated in high-flux fast reactors at 350-427°C and tested 

at 300-427°C. However, for the LWR-irradiated materials, the saturation fracture toughness (KIc or KJc) 

values are in the range of 36.8-40.3 MPa m1/2 (33.5-36.6 ksi in1/2), and samples at this toughness 
xix

typically fail without stable ductile crack extension. These values are lower than the KJc of 75 MPa m1/2

(68.2 ksi in.1/2) observed for fast-reactor-irradiated materials.

Most of the fracture toughness data on LWR-irradiated austenitic SSs have been obtained on 

Type 304 and 304L SSs. Similar data on SS welds, cast SSs, or weld HAZ materials are limited. The 

available data indicate that for the same irradiation conditions, the fracture toughness of thermally aged 

cast SS and weld metal is lower than that of HAZ material, which, in turn, is lower than that of solutionannealed materials. Available data indicate a strong orientation effect on fracture toughness; fracture 

toughness in the T-L orientation* is significantly lower than that in the L-T orientation. Also, several 

fracture toughness J-R curve tests on Type 304 SS control-rod blade material irradiated to 4.5-5.5 dpa 

showed no ductile crack extension. Potential effects of the coolant environment and crack morphology on 

fracture toughness of irradiated SSs are discussed. The available data are inadequate to establish 

accurately the effects of the irradiation temperature on the fracture toughness of austenitic SSs. 

The existing fracture toughness data have been evaluated to develop a fracture toughness trend 

curve that includes (a) a threshold neutron exposure for radiation embrittlement of austenitic SSs and a 

minimum fracture toughness for these materials irradiated to less than the threshold value, (b) a saturation 

neutron exposure and a saturation fracture toughness for materials irradiated to greater than this value, 

and (c) a description of the change in fracture toughness between the threshold and saturation neutron 

exposures. However, a review of the existing data indicated very limited data on (a) materials irradiated 

in LWRs to neutron dose levels of 0.1-1.0 dpa or above 10 dpa and (b) LWR-irradiated cast austenitic SSs 

and welds. The contribution of additional precipitate phases, voids, and cavities on the fracture toughness 

of these materials is not accurately known. 

Synergistic Effects of Thermal and Neutron Embrittlement

Cast and welded austenitic SSs have a duplex structure consisting of austenite and ferrite phases, 

and are susceptible to thermal embrittlement during service at LWR operating temperatures. Formation 

of the Cr-rich α$ phase in the ferrite is the primary mechanism for thermal embrittlement of these 

materials; thermal aging has no effect on the austenite phase. Embrittlement of the austenite phase from 

neutron irradiation under LWR operating conditions occurs at dose levels above 0.5 dpa, and that of the 

ferrite phase occurs at lower dose levels (above 0.07 dpa). However, for reactor core internal 

components, concurrent exposure to neutron irradiation can result in a synergistic effect wherein the 

service-degraded fracture toughness can be less than that predicted for either thermal embrittlement or 

neutron irradiation embrittlement independently. The available fracture toughness data are inadequate to 

evaluate the synergistic effects of thermal and neutron embrittlement on the threshold dose for 

embrittlement. Additional data are needed to better establish the potential for synergistic loss of 

toughness in these materials in the transition dose range from 0.05 to 2 dpa. 

Void Swelling 

Void swelling refers to the volume change of materials under neutron irradiation. Volume changes 

can produce dimensional distortions and misfits in reactor internal components, particularly when coupled 

with irradiation creep, which can compromise the functional integrity of the components. Microscopic 

voids, developed from vacancy coalescence, give rise to this geometry instability, with important 

consequences on the strength and resistance to failure of these materials. The void swelling process is 

 

*The first letter represents the direction perpendicular to the plane of the crack and the second letter represents the direction of 

crack advance. L = longitudinal or rolling direction and T = transverse direction.
xx

divided into two regimes: a transient regime followed by a steady-state swelling rate. As demonstrated by 

Garner and coworkers, austenitic SSs irradiated in fast reactors at temperatures above 427°C eventually 

reach a steady-state swelling rate of ~1%/dpa. Additionally, available fast reactor data indicate that the 

steady-state swelling rate could eventually be attained in PWR internal materials during extended 

operation. The time to reach the steady-state rate (i.e., the transient regime) depends on the material 

composition and thermo-mechanical condition, stress state, neutron flux and spectrum, or irradiation 

temperature (above 300°C). 

The existing data have been reviewed and analyzed to assess void swelling in PWR core internal 

components and to evaluate the effects of key material and environmental parameters. The results 

indicate that material composition can influence void swelling. In the 300-series SSs, void swelling 

decreases with increasing Ni, P, or Si contents. However, radiation-induced precipitation of second phase 

particles can remove Ni, P, and Si from the alloy matrix and thereby increase void swelling. In addition, 

cold work prolongs the transient regime for void swelling in austenitic SSs. Void swelling in austenitic 

SSs is typically observed at irradiation temperatures above 340°C (644°F), and shows a strong 

dependence on temperature. The low temperatures, typical of LWRs, extend the transient regime to much 

longer times. Nonetheless, depending on the irradiation temperature, dose, and dose rate, voids can form 

under PWR operating conditions within the reactor lifetime. Available data obtained in fast reactors at 

373-444°C and displacement rates of 0.3-10.0 x 10-7 dpa/s suggest that under comparable irradiation 

temperatures and neutron dose levels, the lower displacement rates typical of PWRs would result in 

higher swelling. The effects of applied or swelling-induced stress on void swelling, including the 

interaction of swelling and irradiation creep relaxation, are discussed. 

This section also describes the stress-free swelling rate equations, including the incremental stressenhanced swelling, developed by the MRP for solution-annealed Type 304 and cold-worked Type 316 

SSs. The stress-free swelling rate in PWR core internal components is expressed in terms of the 

irradiation temperature, neutron dose, and dose rate. The incremental stress-enhanced swelling is 

expressed in terms of the stress-free rate and the Von Mises effective stress. The results indicate that void 

swelling for an annealed Type 304 SS former plate at an average temperature of 355°C could be above 

5% after 50 effective full power years.

Embrittlement of the material due to voids is also discussed. Typically, austenitic SSs with ≥10% 

void swelling suffer from severe embrittlement, particularly at room temperature. This behavior has been 

attributed to a second order process, which involves stress concentration between voids, Ni segregation at 

void surfaces, and a resultant tendency toward martensite formation when the steel is deformed at room 

temperature. However, there are few quantitative data correlating void swelling with fracture toughness 

of the material. The embrittlement of austenitic SSs due to void swelling under PWR irradiation 

conditions should be further investigated.

Irradiation-Induced Creep Relaxation 

Loss of preload for bolting and springs due to stress relaxation is another aging degradation effect 

that needs to be addressed to assure the functional integrity of the reactor core internal components. 

Stress relaxation represents plastic deformation that occurs under constant strain below the yield point of 

the material. It is caused by either primary creep (thermal stress relaxation) or a flux of high-energy 

neutrons (irradiation creep). Thermal stress relaxation data have been obtained from short-term tests (less 

than 1000 h). As discussed in Materials Reliability Program report MRP-175, the percentage of stress 

relaxation increases with increasing initial stress and saturates at about 207 MPa (30 ksi). At PWR 

operating temperatures, thermal stress relaxation saturates within 100 h with a maximum decrease of 10-
xxi

20% of the initial preload stress depending on the thermal-mechanical treatment of the material. For 

PWRs, plastic strains resulting from thermal stress relaxation were considered to be insignificant and not 

a concern for the dimensional change of core internal components. However, neutron-irradiationenhanced creep at PWR temperatures can greatly increase the plastic strain by increasing both the 

transient creep and steady-state creep rates. Empirical models have been developed for irradiation creep 

based on data obtained from fast reactors. The results have been used to develop expression for 

estimating stress relaxation as a function of initial stress and neutron dose.
xxii

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xxiii

Acknowledgments

The author thanks Torill Karlsen, OECD Halden Reactor Project, Halden, Norway; Anders Jenssen, 

Studsvik Nuclear AB, Sweden; and Raj Pathania, Electric Power Research Institute, for providing the 

experimental data and for their helpful comments. The author is grateful to Bill Shack, Rob Tregoning, 

and Sri Rao for their invaluable input and guidance in preparing this report. This work is sponsored by 

the Office of Nuclear Regulatory Research, U.S. Nuclear Regulatory Commission, under Job Code 

N6818; Program Manager: Appajosula S. Rao.
xxiv

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xxv

Acronyms and Abbreviations

ANL Argonne National Laboratory

ASTM American Society for Testing and Materials

ATEM Analytical Transmission Electron Microscopy

ATR Advanced Test Reactor

BWR Boiling Water Reactor

BWRVIP Boiling Water Reactor Vessels and Internals Project 

CGR Crack Growth Rate

CIR Cooperative IASCC Research

CNSR Chevron Notch Short Rod

CT Compact Tension

CTOD Crack Tip Opening Displacement

CW Cold Worked

DC Direct Current

DO Dissolved Oxygen

dpa Displacements per Atom

ECP Electrochemical Potential

EFPY Effective Full Power Year

EPFM Elastic-Plastic Fracture Mechanics

EPRI Electric Power Research Institute

FFTF Fast Flux Test Facility

GE General Electric

GG Grand Gulf

GTA Gas Tungsten Arc 

HAZ Heat-Affected Zone

HWC Hydrogen Water Chemistry

IASCC Irradiation-Assisted Stress Corrosion Cracking

IG Intergranular

IGSCC Intergranular Stress Corrosion Cracking

JAPEIC Japan Power Electric Engineering and Inspection Corp.

J-R J Integral Resistance

LEFM Linear-Elastic Fracture Mechanics

LWR Light Water Reactor

MRP Materials Reliability Program

NDT Nil-Ductility Transition

NRC Nuclear Regulatory Commission

NWC Normal Water Chemistry

PWR Pressurized Water Reactor 
xxvi

RIS Radiation-Induced Segregation

SA Submerged Arc

SCC Stress Corrosion Cracking

SE(B) Single Edge Bend

SFE Stacking Fault Energy

SFT Stacking Fault Tetrahedra

SHE Standard Hydrogen Electrode

SIA Self-Interstitial Atom

SMA Shielded Metal Arc

SR Swelling Rate

SS Stainless Steel

SSRT Slow Strain Rate Tensile

TG Transgranular

UE Uniform Elongation

UHP Ultra High Purity

YS Yield Strength
xxvii

Symbols

a Crack length

B Specimen thickness

Beff Effective specimen thickness 

C Power-law constant

D Neutron dose

da Increment in crack length

E Elastic modulus

F Neutron fluence

J J integral, a mathematical expression used to characterize the local stress-strain field at the 

crack tip region (parameter J represents the driving force for crack propagation)

JIc Value of J near the onset of crack extension

K Stress intensity factor

KIc Critical stress intensity factor

KJc Equivalent critical stress intensity factor

Kmax Maximum stress intensity factor

Kmin Minimum stress intensity factor

M Constraint factor

R load ratio

S Stress

t Test duration

T Tearing modulus or temperature

teff Effective time of loading

tr Rise time

W Specimen width

φ Neutron dose

!

! Dose rate

! Von Mises effective stress

σf Flow stress, defined as the average of yield and ultimate stress

σu Ultimate stress

σy Yield stress

υ Poisson ratio
xxviii

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1

1 Introduction

In light water reactors (LWRs), austenitic stainless steels (SSs) are used extensively as structural 

alloys in the internal components of reactor pressure vessels because of their relatively high strength, 

ductility, and fracture toughness. Fracture of these steels occurs by stable tearing at stresses well above 

the yield stress, and tearing instabilities require extensive plastic deformation. However, exposure to 

neutron irradiation for extended periods changes the microstructure (radiation hardening) and 

microchemistry (radiation-induced segregation or RIS) of these steels and degrades their fracture 

properties.1-11 Also, neutron irradiation changes the water chemistry (radiolysis). In boiling water 

reactors (BWRs), neutron irradiation results in an increase in corrosion potential, whereas it has no effect 

in pressurized water reactors (PWRs) because hydrogen scavenges the radiolysis products and the 

corrosion potential remains low.2

Loss of fracture toughness due to radiation embrittlement was not considered in the design of LWR 

core internal components constructed of austenitic SSs, but it has been considered in addressing nuclear 

plant aging and license renewal issues. In addition, irradiation-assisted stress corrosion cracking 

(IASCC),

1-3,9-16 change in dimensions due to void swelling,17-20 and stress relaxation due to radiation 

creep,21-25 are other aging degradation process that affect LWR core internal components exposed to fast 

neutron radiation, and need to be considered in addressing plant aging issues. 

The objective of this report is to address the concerns arising from irradiation-induced materials issues. 

The ultimate goal is to help assure the structural and functional integrity of reactor core internals. The 

report presents a critical assessment of the susceptibility of austenitic SSs to irradiation effects such as 

IASCC, neutron embrittlement, void swelling, and stress relaxation in PWR environments. The existing 

data, in the open literature as well as NRC and industry reports, have been evaluated to establish the 

effects of material parameters (such as composition, thermo-mechanical treatment, microstructure, 

microchemistry, yield strength, and stacking fault energy) and environmental parameters (such as water 

chemistry, irradiation temperature, dose, and dose rate) on these processes. Differences in the radiationinduced degradation of material properties between LWR and fast-reactor irradiations are also discussed. 

The results are used to:

(a) Identify the key material and environmental parameters that influence irradiation-induced 

degradation of reactor core internal materials. 

(b) Provide a better understanding of the threshold fluence above which irradiation effects on IASCC 

susceptibility and tensile and fracture properties are significant. 

(c) Develop disposition curves that describe the cyclic and stress corrosion cracking (SCC) growth 

rates as a function of stress intensity factor for austenitic SSs and Ni alloys used for reactor core 

internal components. 

(d) Evaluate the potential of radiation embrittlement under BWR and PWR operating conditions, 

including the synergistic effects of thermal and neutron embrittlement of cast austenitic SSs. 

(e) Assess the effects of void swelling, including its effect on fracture toughness. 

(f) Investigate the significance of irradiation creep relaxation on the functional integrity of reactor 

internal components.

 This report also identifies potential deficiencies or knowledge gaps in the existing information or 

technical bases. A research plan is proposed to address the issues related to these knowledge gaps and 

help the NRC to develop guidelines for the license renewal for plant life up to and beyond 60 years. 
2

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3

2 Irradiation Assisted Stress Corrosion Cracking (IASCC)

2.1 IASSCC Susceptibility

A major concern regarding the structural and functional integrity of core internal components is the 

IASCC of austenitic SSs. Several incidents of IASCC have occurred since the mid 1970s in BWR control 

blade handles and instrumentation tubes, and since the 1990s in BWR core shroud and PWR baffle bolts. 

As the name implies, IASCC is literally irradiation-assisted enhancement of SCC susceptibility of 

austenitic SSs. Neutron irradiation increases the susceptibility of austenitic SSs to IASCC by changing 

the material microstructure due to radiation hardening and also the material microchemistry due to RIS. 

The IASCC susceptibility of austenitic SSs has been investigated by conducting slow-strain-rate-tensile 

(SSRT) tests, crack growth rate (CGR) tests, and crack initiation tests on irradiated material in simulated 

LWR environments. The effects of fast neutron fluence on the IASCC of SSs have been investigated for 

BWR control blade sheaths26-28 and in laboratory tests on BWR-irradiated material.14,29-34 The results 

indicate that the extent of intergranular stress corrosion cracking (IGSCC) increases with fluence. The 

factors that influence the IASCC susceptibility of materials include neutron irradiation conditions (e.g., 

neutron fluence, flux, and energy spectrum), cold work, material composition, corrosion potential, water 

purity, temperature, and loading conditions. Both plant26,27 and laboratory data15 show increased IASCC 

susceptibility with neutron fluence, corrosion potential, and water conductivity. In BWRs, instrument dry 

tubes and control blade handles and sheaths, which are typically subject to very low stress, but under 

creviced conditions, have shown IASCC (Fig. 1a). Also, the frequency of IASCC increases with 

increasing conductivity of the reactor water (Fig. 1b). 

0

20

40

60

80

100

0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4

% of Spot Welds Inspected with IGSCC

Neutron Fluence (n/cm2 x 1021) (E > 1 MeV)

BWR creviced control blade sheath

Threshold fluence for IGSCC

!5 x 1020 n/cm2

(a)

0.0

0.2

0.4

0.6

0.8

1.0

1.2

1.4

0.1 0.2 0.3 0.4 0.5 0.6 0.7

% with IGSCC/On-line Month

Plant Average Conductivity (μmS/cm)

Frequency of SCC initiation

increases dramatically with

increasing conductivity

Upper bound

Best fit

(b)

Figure 1. The effect of (a) fast neutron fluence and (b) average water conductivity on susceptibility of 

austenitic stainless steels to IASCC shown in field correlations of the core component 

cracking (Refs. 27,26).

Furthermore, radiolysis of water leads to its dissociation into various molecular, ionic, and radical 

reaction products that interact to form H2O2, H2, and O2. In BWRs, these species increase the corrosion 

potential, which is known to increase the susceptibility of SSs to IASCC. The dependence of IASCC 

susceptibility of austenitic SSs on neutron fluence and corrosion potential is shown in Fig. 2. However, 

the addition of H2 to the reactor water greatly reduces the effect of radiolysis by scavenging the radiolysis 

products. For example, the presence of about 500 ppb H2 (5.6 cc/kg) in water can suppress the effect of 

radiolysis on the corrosion potential in BWRs.15 Because PWR coolants typically contain 2 ppm H2

(30 cc/kg), radiolysis has no effect on the corrosion potential in PWRs.
4

0

20

40

60

80

100

1019 1020 1021 1022

% Intergranular SCC Fracture

Neutron Fluence (n/cm2) (E > 1 MeV)

Data of Jacobs (!) and Kodama ("!)

Post-irrad. SSRT 2-3 x 10-7 s-1 288°C

Comm.-purity Type 304 SS (!") & Type 316 SS (!)

Data shifted right by init. GB Cr enrichment

!"! 42 ppm O2-saturated vs. !"! 0.02 ppm O2 %IGSCC

0

20

40

60

80

100

-0.4 -0.2 0 0.2

!c, VSHE

~42 ~2

~0.2

0.02

to 0.05

ppm O2

Figure 2.

Dependence of IASCC susceptibility of 

austenitic stainless steels on fast neutron 

fluence and corrosion potential as 

measured from SSRT tests in 288°C 

water. The inset shows the effect of 

corrosion potential on %IGSCC (Ref. 15). 

Laboratory SSRT data on irradiated austenitic SSs have been used to identify a threshold fluence 

above which IASCC is significant in austenitic SSs in normal water chemistry (NWC) BWR 

environments.14,35 Tests on Types 304 and 347 SS from manual shielded arc weld materials with a heataffected zone (HAZ) irradiated to 2 x 1020 n/cm2 (E >1 MeV)* (0.3 dpa) at 290°C show no IGSCC.36 

Also, data on commercial and high-purity austenitic SSs and Alloy 800 irradiated in a BWR between 

3 x 1020 and 9 x 1021 n/cm2 (0.45-13.5 dpa) show no IASCC below 5 x 1020 n/cm2 (0.75 dpa) and a rapid 

increase in IASCC susceptibility above 1 x 1021 n/cm2 (1.5 dpa).37 Figure 3 shows a plot of the percent 

IGSCC measured in irradiated SS specimens as a function of fast neutron fluence. Although a threshold 

fluence level of 5 x 1020 n/cm2 (0.75 dpa) has been proposed for austenitic SSs in NWC BWR 

environments, the results in Fig. 3 suggest a lower threshold for some materials (right triangles). The 

SSRT data indicate an increase in intergranular (IG) cracking susceptibility in some commercial-purity 

SSs at fluence levels of about 2 x 1020 n/cm2 (0.3 dpa) and in high-purity heats of SSs at even lower 

fluence levels. 

0

20

40

60

80

100

1020 1021 1022

Clark & Jacob

Kodama et al.

Jacobs et al.

Chung et al.

Chung et al.

%IGSCC

Neutron Fluence (n/cm2, E >1 MeV)

Austenitic Stainless Steels

High–Purity Water ≥8 ppm DO

Open Symbols: Type 304 SS

Closed Symbols: Type 316 SS

Commercial Heats

High-Purity Heats

Figure 3.

Susceptibility of irradiated austenitic SSs to 

IGSCC as a function of fluence in high-dissolvedoxygen water. From slow-strain-rate tensile tests 

(Refs. 29,31-33).

 

*All references to fluence levels are calculated for E ≥1 MeV.
5

The SSRT test data on Type 304 and 316 SS irradiated in commercial BWRs up to 4.0 x 1021 n/cm2

(6.0 dpa) indicate a beneficial effect of reducing the corrosion potential of the environment.16,37 A low 

corrosion potential does not provide immunity to IASCC if the fluence is high enough. For example, 

IGSCC has been observed in cold-worked (CW), irradiated SS baffle bolts in PWRs. However, threshold 

fluences for IASCC are higher under low-potential conditions such as BWR hydrogen water chemistry 

(HWC) or PWR primary water chemistry. The threshold fluences for low-potential environments are not 

well defined; they could be different for HWC and PWR environments. The SSRT data indicate that 

threshold fluences also vary with material composition and thermo-mechanical treatment. In general, 

Type 316 SS is less susceptible than Type 304 SS (i.e., the threshold fluence for IASCC is likely to be 

lower for Type 304 SS than for Type 316 SS).37 

It is well known that SCC of materials in high-temperature, high-pressure water depends on three 

factors: a susceptible material, relatively high stress, and aggressive environment. In the film rupture/slip 

oxidation model, the fundamental elements of SCC have been described in terms of the processes that 

occur at the crack tip14,38-40 and are considered to include: dynamic strain at the crack tip that disrupts the 

passive film and exposes fresh metal surface, followed by rapid metal oxidation and subsequent surface 

repassivation. This process can be quantitatively linked to crack advance.35 The kinetics of repassivation 

are primarily a function of the local chemistry and material composition, as well as the mass transport 

processes that establish the local crack tip chemistry.

40 The effect of water chemistry is particularly 

important because it can change the corrosion reactions at the crack tip. Metallurgical and microstructural 

parameters are also important because sensitization of the material by thermal treatment or irradiation 

hardening due to radiation damage can alter the localized deformation and creep processes at the crack 

tip, thereby changing the film rupture frequency. A schematic illustration of the significant metallurgical, 

mechanical, and environmental aspects of IASCC is presented in Fig. 4. It shows radiation-induced 

changes in reactor water chemistry, material microstructure, and grain boundary composition, and the 

localized deformation processes that impact the crack tip micromechanics that influence IASCC.

A review of the existing data on IASCC susceptibility has identified the following key material and 

environmental parameters that influence the susceptibility of LWR core internals materials to IASCC.

Figure 4.

Schematic illustrating mechanistic 

issues that influence IASCC in 

austenitic stainless steels (Ref. 4).
6

2.1.1 Microstructure

Neutron irradiation of austenitic SSs can produce damage by displacing atoms from their lattice 

position. Irradiation damage is characterized by either the neutron fluence [e.g., neutrons per square 

centimeter (n/cm2)] or the average number of displacements experienced by each atom, i.e., 

displacements per atom (dpa).* Each displaced atom creates a vacancy and self-interstitial atom (SIA) 

pair, known as a Frenkel pair. These defects are unstable, and most of them are annihilated by 

recombination. The surviving defects rearrange into more stable configurations such as dislocation loops, 

network dislocations, precipitates, and cavities (or voids), or migrate to sinks such as grain boundaries or 

surfaces of second phase particles. The production, annihilation, and migration of the point defects lead 

to changes in the microstructure and microchemistry of the material. Changes in the microstructure of 

austenitic SSs due to neutron irradiation vary with the material composition and thermo-mechanical 

treatment and irradiation temperature and dose rate. The effects of material and irradiation conditions on 

microstructural evolution vary with irradiation temperature.4 Also, because it affects the recombination 

rates, increasing temperature reduces embrittlement per dpa and increases swelling per dpa.

Under LWR conditions, the material microstructure produced by irradiation seems to change 

significantly for irradiation temperatures above 300°C. At LWR operating temperatures (i.e., 275-

300°C), the defect structure primarily consists of small “black spot” defect clusters and large dislocation 

loops. The small defect clusters are less than 4 nm in diameter and consist mostly of interstitial loops and 

vacancy clusters, which are created from the collapse of damage cascade during irradiation. At higher 

temperatures (> 300oC) the microstructure contains large faulted loops, network dislocations, and 

cavities/voids (clusters of vacancies and/or gas bubbles). The large dislocation loops are 4-20 nm in 

diameter and are primarily faulted interstitial Frank loops that are created from the clustering of selfinterstitial atoms.4 The size of the interstitial loops increases with increasing irradiation dose as more 

interstitials are absorbed than vacancies. These loops stop growing when vacancies and interstitials are 

absorbed at the same rate. The saturation size of the interstitial loops depends on the irradiation 

conditions and material characteristics. The data also indicate that alloying elements can affect the 

material microstructure during low-temperature irradiations (e.g., P, Ti, and Nb increase the loop density 

and decrease loop size).4 Also, precipitates form at higher temperatures and high doses (>20 dpa).4,41-47

The change in the density and size of the interstitial loops in austenitic SSs with irradiation dose 

under LWR conditions at 275-290°C is shown in Fig. 5. Under LWR conditions, the loop density 

saturates at a relatively low dose (about 1 dpa), and the average loop diameter saturates at about 5 dpa. 

Cavities or voids have not been reported for SS irradiated under LWR conditions at temperatures below 

300°C. The loop size increases and loop density decreases with irradiation temperature. At temperatures 

of 300-350°C, the microstructure primarily consists of large Frank loops and a network of tangled 

dislocations that evolve as the larger faults in the Frank loops disappear.4 Also, cavities and voids form at 

dose levels above about 20 dpa; they become more important at higher temperatures and higher doses. 

Typically cavities are associated with precipitates, dislocations, and grain boundaries.

The data on the microstructural evolution in reactor core internal materials irradiated to doses above 

10 dpa and at temperatures of 300-350°C are rather limited. The available data indicate that under LWR 

irradiation conditions at 275-300°C the Frank loop microstructures in Type 304 and 316 SS are similar, 

and both the Frank loop density and loop size reach a saturation value at 5 dpa. Also, the initial state of 

the material (i.e., solution annealed or cold worked) does not seem to be important. The initial dislocation 

 

*Unless otherwise noted, when neutron dose in dpa was not available, the values of neutron fluence (n/cm2) were converted to 

dpa as follows: for LWRs, E>1 MeV and 1022 n/cm2 ≈15 dpa; and for fast reactors, E>0.1 MeV and 1022 n/cm2 ≈5 dpa.
7

structure produced by the cold work is removed by the time a steady-state loop density is achieved and 

the irradiated microstructure of the CW SSs is similar to that of the solution-annealed SSs.48 However, 

although the dislocation microstructure does not change significantly at higher dose levels, cavities, voids, 

and precipitates are observed in materials irradiated to higher neutron doses. 

The Frank loop densities and loop sizes in CW Type 316 SS thimble tube materials irradiated to 

33 dpa at 290°C and to 70 dpa at 315°C in a PWR were comparable to the saturation values in Fig. 5.

48 

However, a high density of nanocavities (average size less than 3 nm) was observed in both materials, but

the density of cavities was slightly higher and their size larger in the higher dose material. There was also 

some evidence of precipitates in the material irradiated to 70 dpa. The cavities were believed to be He 

bubbles that form as He accumulates during irradiation from transmutation reactions. Helium can be 

produced by the 10B reaction to form 4He and 7Li, and also from the production of 59Ni from 58Ni by 

thermal neutrons. A fine dispersion of He bubbles has also been observed in SSs irradiated at 300°C in 

reactors with high thermal neutron fluxes.

As noted, some differences in irradiated microstructures are observed in materials irradiated at 

temperatures above 320°C. The microstructure of a CW Type 316 SS baffle bolt from the Tihange 1 

plant was characterized at three positions (1, 25, and 55 mm).18 Although the loop densities were similar 

at all three positions, the size distribution at higher temperature (330-343°C) and lower flux at the 25- and 

55-mm positions along the shank of the bolt shifted to larger sizes compared to that at the 1-mm position 

near the bolt head, which has the lowest temperature (320°C) and highest flux. The total dose at the three 

positions decreased from 19.2 dpa at the 1-mm position to 13.2 and 8.7 dpa at the 25- and 55-mm 

positions, respectively. Also, precipitates and cavities were found at all positions, but their density and 

size varied considerably. The cavities observed at the 1-mm position were less than 2 nm and not 

distributed homogenously throughout the matrix. The cavities observed at the 25- and 55-mm positions 

were faceted, indicating that they were voids and not He bubbles, and the average size was about 8 nm. 

The precipitates at the 25-mm position were identified as γ’phase. 

Irradiation at temperatures above 320°C also leads to the formation of second phase particles. Such 

radiation-induced precipitates are similar to those that form during thermal aging and depend on material 

composition, thermo-mechanical treatment, and radiation conditions (neutron dose, temperature, and 

spectrum). Radiation can either enhance or retard the formation of some phases (e.g., Laves, σ and χ

intermetallics, and MC, M23C6, and M6C carbides) and modify the composition of M6C and Laves 

phases.42,43 The radiation-induced phases include γ’ silicide (Ni3Si), phosphides (M2P and M3P), and G 

phase (M6Ni16Si7). However, most of these observations are based on fast reactor irradiations at 

1020

1021

1022

1023

4

5

6

7

8

9

10

11

12

0 1 2 3 4 5 6 7 8

Type 304 SS

Type 316 SS

Loop Density (m-3)

Average Loop Diameter (nm)

Irradiation Dose (dpa)

Loop Size

Loop Density

LWR Irradiation at 275-290°C

Figure 5.

The change in density and size of 

interstitial loops as a function of 

irradiation dose (dpa) during LWR 

irradiation at 275-290°C (Ref. 4).
8

temperatures above 350°C, and data obtained on materials irradiated under LWR irradiation conditions at 

320-350°C are limited. The existing data suggest that radiation-induced precipitation is not a concern for 

LWRs at temperatures below 350°C. Metal carbides are the primary stable precipitate phase in 300-series 

stainless steels under LWR operating conditions, although RIS of Ni and Si to sinks in the matrix and 

grain boundaries may lead to the formation of γ’ and G phases.

Because of the high displacement rates, fast reactors are an attractive alternative for irradiating test 

materials to dose ranges that cannot be achieved within a reasonable time in LWRs. However, the higher 

dose rates and differences in the thermal- and fast-neutron spectra may result in material microstructure 

and microchemistry that are not representative of LWR irradiations. Therefore, possible differences in 

material microstructure need to be examined. 

The microstructural evolution in austenitic SSs irradiated in LWRs at 275-300°C has been 

compared with samples irradiated in the BOR-60 fast breeder reactor at 320°C (Fig. 6).48 The results 

indicate that at 5-10 dpa the dislocation structure of commercial heats of austenitic SSs is comparable. 

Although the data show subtle variations in dislocation structure, the differences are minor. The loop 

density in these steels irradiated either in LWRs or the BOR-60 reactor is about 2 x 1023 m-3; the 

saturation value was reached after about 2 dpa. The loop density is the same even in the ultra high-purity 

(UHP) heats of Type 304L (with or without minor additions of C, N, and Si) that have been irradiated to 

20 dpa. However, there are some differences in the loop size of the irradiated steels. For the LWRirradiated SSs, the Frank loop diameter in two high-purity heats (304-E and 316-F) is larger than that in 

the less pure heats.44 Also, the loop size for the high-purity heats irradiated in BOR-60 reactor to 20 dpa

is much lower than for the high-purity heats (Heats 304-E and 316-F) irradiated in LWRs to about 8 dpa; 

the average loop diameter is about 6 nm and 14 nm, respectively. Typically, stacking fault tetrahedra 

(SFT) are also observed in high-purity materials. The LWR-irradiated Heat 304-E contained some SFT in 

addition to the Frank loops, and the UHP 304 SS heat irradiated in the BOR-60 contained a high density 

of Frank loops and SFT. The existing data on microstructural evolution in irradiated SSs also indicate 

that addition of C increases the density of Frank loops and alters the tensile response of the steel, whereas 

addition of Si decreases the loop density as well as tensile hardening.49 However, additions of C, Si, P 

BOR-60 5&6

1020

1021

1022

1023

1024

4

6

8

10

12

14

16

18

0 5 10 15 20 25

304-B

304-C

316-P

316-K

304-E

 316-F

Loop Density, m!3

Average Loop Diameter, nm

Irradiation Dose, dpa

Size Density Loop Density

Loop Size

BOR-60

LWR

304L Core Shroud

UHP 304L 

UHP 304L+C

UHP 304L+N

UHP 304L+Si

316CW Baffle Bolt 

Figure 6. Results from the BOR-60 irradiation experiment with the data for austenitic stainless 

steels irradiated in a BWR (Ref. 48).
9

together and with other impurities appear to alter the evolution and type of microstructure of irradiated 

SSs in a complex way. 

These results indicate that at irradiation temperatures of 275-320°C, the Frank loop microstructure 

appears to be relatively insensitive to the displacement rates between LWRs and fast reactors. The main 

difference between the two is the presence of He bubbles and cavities in materials irradiated to high dose 

levels under LWR irradiation conditions. Cavities were not observed in the BOR-60 materials. Also, He 

bubbles appear to form preferentially on the grain boundaries, which could lead to intergranular 

embrittlement of the material.

48 Additional data need to be obtained on reactor core internal materials 

irradiated to doses above 10 dpa and at temperatures of 300-350°C under LWR irradiation conditions. 

2.1.2 Microchemistry

Neutron irradiation also changes the microchemistry of the material due to RIS. The migration of 

vacancies and self-interstitial atoms to sinks such as grain boundaries, dislocations, or precipitate surfaces 

leads to local compositional changes. Elements such as Si, P, and Ni that are believed to migrate by 

interstitial mechanisms are enriched near regions that act as sinks for the point defects, while elements 

such as Cr, Mo, and Fe that exchange more rapidly with vacancies are depleted near sinks.2,4,14 The 

extent of segregation or depletion depends on the rate of generation and recombination of point defects 

(i.e., it depends on irradiation temperature and dose rate). 

Typically, RIS peaks at intermediate temperatures. It is reduced at low temperatures because of 

reduced mobility and at high temperatures due to back diffusion. Also, for a specific neutron dose, RIS is 

greater at lower dose rate. At LWR temperatures (275-300°C), significant RIS is observed at irradiation 

doses as low as 0.1 to 5 dpa.2-4,46 The grain boundary segregation/depletion profiles are extremely 

narrow at ≈300°C (≈572°F). The extent of segregation/depletion at grain boundaries is decreased, and the 

width of the profile is increased with an increase in irradiation temperature.2 In low-C Type 316 SS, presegregation of Cr and Mo to grain boundaries in mill-annealed materials (without sensitization) promotes 

the formation of W-shaped profiles. At low doses, a narrow depleted region forms next to a grain

boundary that is still higher than the matrix level. Dose levels above 3 dpa are needed to decrease the Cr 

level at the boundary below that in the matrix.46 Dose levels greater than 6 dpa are needed to completely 

remove the W-shaped profiles and establish the typical V-shaped depletion profile. In general, lower 

doses are required to remove the W-shaped Cr profile in Type 304 SS (less than 3 dpa) compared with

Type 316 SS. The depletion of Mo, particularly in Type 316 SS, is also controlled by pre-segregation of 

Mo to the grain boundary; it takes doses greater than 5 dpa to remove the W-shaped profile. Another 

minor alloying element Mn also shows significant depletion at the grain boundary after irradiation. 

However, quantitative measurements of the compositional profiles for Mn are not generally available

because of the measurement difficulties in analyzing neutron-irradiated SSs by analytical transmission 

electron microscopy (ATEM). 

The radiation-induced changes in Cr concentration at grain boundaries in several 300-series SSs 

irradiated in LWRs at 275-290°C are shown in Fig 7a as a function of neutron dose; additional data for Cr 

depletion from several studies at about 300°C are shown in Fig. 8. Most of the steels show a rapid 

decrease in Cr to about 13 wt.% at dose up to 5 dpa.46 The data at dose levels above 10 dpa are limited, 

but they indicate that Cr content can decrease to 8-10 wt.% in some materials. Under similar irradiation 

conditions, grain-boundary Cr and Ni concentrations in most 300-series SSs irradiated in LWRs at 300-

320°C up to 65 dpa are shown in Figs. 9a and b, respectively. The Cr concentrations measured by 

Fukuya et al.50 in CW Type 316 SS PWR thimble tube materials are somewhat lower than those obtained 
10

by other investigators (e.g., Fujimoto et al.).51 The irradiation temperature was comparable in the two 

studies (295-320°C), and the initial cold work was 10-12% and 15%, respectively.

The Si concentration in LWR-irradiated SSs is shown in Fig 7b as a function of neutron dose;4 Si 

contents in SSs irradiated in LWRs at 300-320°C up to 65 dpa are shown in Fig. 9c. The results show a 

rapid increase in Si to about 4 wt.% at 5 dpa and a gradual increase to 5-6 wt.% at higher dose levels. 

Neutron irradiation should lead to enrichment of P; however, measurements of grain boundary P in 

neutron irradiated SSs have not yielded consistent results, but it appears that P does not segregate during 

LWR irradiations. The RIS behavior of other minor elements such as S, C, N, and B, all of which should 

segregate to the boundary, has not been established. These elements typically enrich grain boundaries but 

are difficult to measure. Existing data on LWR-irradiated SSs have not indicated significant segregation 

of these elements. However, these observations are for doses less than 15 dpa. Transmutation of Mn at 

higher doses may release S from MnS inclusions, which may lead to segregation.46 Also, if B segregates 

to the boundaries, He produced by the transmutation reaction (from 10B to form 4He and 7Li) would 

enrich at grain boundaries, which may lead to intergranular embrittlement. The redistribution of S, C, N, 

10

12

14

16

18

20

22

24

26

0 5 10 15 20

Type 348 SS

Type 304 SS

Type 316 SS

Grain Boundary Cr Concentration (wt.%)

Radiation Dose (dpa)

Fit to Type 316 SS

Commercial Heats

LWR Irradiations

at 275 - 290°C

0

1

2

3

4

5

6

7

8

0 2 4 6 8 10

Grain Boundary Si Concentration (wt.%)

Radiation Dose (dpa)

ATEM Measurements on

Commercial 300-Series SSs

LWR Irradiated at ~288°C

Bulk Si in SS

(a) (b)

Figure 7. Radiation-induced grain boundary (a) Cr and (b) Si concentration in 300-series stainless 

steels irradiated in LWRs at 275-290°C (Refs. 46,4).

10

12

14

16

18

20

22

24

26

0 5 10 15 20

Type 304 (Asano et al., 1992)

Type 304 (Jacobs et al., 1995)

Type 304 (Jacobs et al., 1993)

Type 304 (Kenik, 1992) 

Type 304 (Nakahiigashi et al.)

Type 316 (Asano et al., 1992)

Type 316 (Jacobs et al., 1990)

Type 348 (Jacobs et al., 1993)

HP304 (Walmsley et al., 1995)

CP304 (Walmsley et al., 1995) 

CP316 (Walmsley et al., 1995)

CP304 (Bruemmer et al., 1999)

CP316 (Bruemmer et al., 1999)

HP316 (Bruemmer et al., 1999)

CP304 Protons (Was et al., 2002)

CP316 Protons (Was et al., 2002)

Grain Boundary Cr Concentration (wt.%)

Dose (Ddpa)

Fast Neutron Fluence (E > 1 MeV) x 1021 n/cm2

0 2 4 6 8 10 12 14

Figure 8.

Dependence of grain boundary Cr 

concentration for several 300-series 

austenitic stainless steels irradiated at 

about 300°C on neutron dose (Ref. 3).
11

and B in austenitic SSs irradiated in LWR at dose levels above 15 dpa needs to be investigated to 

establish its significance on IASCC susceptibility and neutron embrittlement of SSs. 

Measurements of RIS in SSs irradiated in LWRs at temperatures above 320°C are limited. 

Microchemical characterization of CW Type 316 SS materials taken from the Tihange 1 baffle bolt18

showed that the Cr depletion and enrichment of Ni and Si in these materials were consistent with the data 

shown in Fig. 9 for SSs irradiated at 275-320°C. The Cr, Ni, and Si concentrations, respectively, were 

10-12 wt.%, 24-27 wt.%, and about 4.5 wt.% at the 25-mm position (13.2 dpa at 343°C) and 15 wt.%, 16-

21 wt.%, and 2 wt.% at the 55-mm position (8.7 dpa at 333°C). 

The microchemical changes in irradiated material can result in the formation of various 

precipitates; the enrichment of Ni and Si at the grain boundaries may lead to the formation of γ’ silicide or 

G phase. Although γ’ precipitates have been observed in the matrix in SSs irradiated under LWR 

conditions and may affect radiation hardening, they have not been observed at the grain boundaries. The 

γ’ precipitates were identified in the Tihange 1 baffle bolt at the 25-mm position.18 The absence of γ’ 

silicide (Ni3Si) at the grain boundaries suggests that RIS of Ni and Si probably saturates at higher doses 

(the actual dose may vary with material composition and irradiation temperature).

As mentioned earlier, a large fraction of the high neutron dose data was obtained on materials 

irradiated in fast reactors. Therefore, possible differences in material microchemistry due to the higher 

8

10

12

14

16

18

20

0 10 20 30 40 50 60 70

316CW-B Fast PNNL

316CW+Ti-C Fast PNNL

HP316CW Fast Fujimoto

CP316CW Fast Fujimoto

316CW PWR Fujimoto

316CW PWR Fukuya 

316SS LWR PNNL

Grain Boundary Cr Concentration (wt.%)

Radiation Dose (dpa)

Heats of 316SS

Fast Reactor

LWR

High-purity

Heat

(a)

0

5

10

15

20

25

30

35

0 10 20 30 40 50 60 70

316CW-B Fast PNNL

316CW+Ti-C Fast PNNL

HP316CW Fast Fujimoto

CP316CW Fast Fujimoto

316CW PWR Fujimoto

316CW PWR Fukuya 

316SS LWR PNNL

Grain Boundary Ni Concentration (wt.%)

Radiation Dose (dpa)

Heats of 316SS

Fast Reactor

LWR

High-purity

Heat

(b)

0

1

2

3

4

5

6

7

8

0 10 20 30 40 50 60 70

316CW-B Fast PNNL

316CW+Ti-C Fast PNNL

HP316CW Fast Fujimoto

CP316CW Fast Fujimoto

316CW PWR Fujimoto

316CW PWR Fukuya 

316SS LWR PNNL

Grain Boundary Si Concentration (wt.%)

Radiation Dose (dpa)

Fast Reactor Heats of 316SS

LWR

(c)

Figure 9.

Comparison of radiation-induced segregation of 

(a) Cr, (b) Ni, and (c) Si in LWR and BOR-60 fast 

reactor irradiated stainless steels (Ref. 52).
12

dose rates and differences in the thermal- and fast-neutron spectra need to be examined. The RIS of Cr, 

Ni, and Si in LWR- and BOR-60-irradiated austenitic SSs as a function of the neutron dose is shown in 

Fig. 9.

52 In all cases, the data from fast reactor irradiation fall below the lower limit of the data from 

LWR irradiation (i.e., Cr depletion and Ni and Si enrichment are less in the fast reactor irradiated 

materials). These differences may influence the IASCC susceptibility or radiation embrittlement of these 

materials. The results indicate that the RIS in the fast reactor irradiated materials may not be 

representative of that observed under LWR irradiation conditions.

2.1.3 Radiation Hardening 

The point defect clusters and precipitates produced by irradiation act, to varying extent, as obstacles 

to a dislocation motion resulting in an increase in tensile strength and a reduction in ductility and fracture 

toughness of the material. In general, cavities (or voids) are strong barriers, large faulted Frank loops are 

intermediate barriers, and small loops and bubbles are weak barriers.1 The yield strength of irradiated 

SSs can increase up to five times that of the nonirradiated material after a neutron dose of about 5 dpa.3 

The change in engineering stress vs. strain behavior of Type 304 and 316 SSs with neutron dose is shown 

in Fig. 10.

53 The yield and ultimate stresses increase and ductility decreases with irradiation. 

0

200

400

600

800

0 10 20 30 40 50

Unirradiated

0.45 dpa

1.35 dpa

3.00 dpa

Engineering Stress (MPa)

Engineering Strain (%)

Type 304 SS

Irradiated & Tested at 289°C

High Purity Water with 8 ppm DO

(a)

0

200

400

600

800

0 10 20 30 40 50

Unirradiated

0.45 dpa

1.35 dpa

3.00 dpa

Engineering Stress (MPa)

Engineering Strain (%)

Type 316 SS

Irradiated & Tested at 289°C

High Purity Water with 8 ppm DO

(b)

Figure 10. Engineering stress-strain curves at 289°C for (a) Type 304 SS and (b) Type 316 SS steel 

irradiated up to 3 dpa in the Halden reactor at 289°C. From SSRT tests in high-purity water 

containing 8 ppm dissolved oxygen (Ref. 53).

At high neutron doses, as the irradiated yield strength approaches the ultimate strength of the 

material, the deformation mode changes. Deformation by a planar slip mechanism is promoted and the 

material exhibits strain softening.54 This process is also termed “dislocation channeling,” whereby 

dislocation motion along a narrow band of slip planes clears the irradiation-induced defect structure, 

creating a defect-free channel that offers less resistance to subsequent dislocation motion or deformation. 

Nearly all SSs exhibit strain softening, and little or no uniform elongation, at neutron dose above 3-5 dpa. 

The enhanced planar slip leads to a pronounced degradation in the fracture toughness of austenitic SSs.6

Such effects of irradiation on the fracture toughness of austenitic SSs appear to be strongly influenced by 

minor differences in the chemical composition of the steels;1 the chemical composition can influence the

stacking fault energy (SFE) and/or irradiation-induced microstructure. In general, a higher SFE enhances, 

and cold work inhibits, dislocation channeling.1 

The change in yield strength of austenitic SSs correlates well with microstructural changes in the 

irradiated materials. As discussed in the previous section, under LWR operating conditions, the 
13

microstructural and microchemistry of the irradiated material change rapidly at small neutron doses. The 

microstructure reaches saturation at about 5 dpa, and the changes in microchemistry are small above 

5 dpa. The yield strength of the 300-series SSs irradiated under LWR conditions is plotted as a function 

of neutron dose in Fig. 11. The yield strength also appears to saturate at a neutron dose of 5 dpa.54

The extent of irradiation hardening and the increase in yield stress of austenitic SSs depend on the 

material composition and thermo-mechanical treatment, as well as the irradiation temperature. The effect 

of irradiation temperature on the yield strength and ductility of 20% CW Type 316 SS irradiated in fast 

reactors55 is shown in Fig. 12. The effect of irradiation hardening and the loss of ductility are 

significantly more pronounced at 290°C (554°F), at least at low dose levels. Also, for a given neutron 

dose and irradiation temperature, the yield strength decreases with increasing test temperature. A similar 

behavior is observed for SSs irradiated under LWR conditions. The greatest increase in yield strength for 

a given irradiation dose occurs at irradiation temperatures near 300°C (572°F), which is in the 

temperature range of LWR operation. 

200

400

600

800

1000

0 10 20 30 40 50

290°C

371°C

427°C

483°C

500°C

Yield Strength (MPa)

Neutron Dose (dpa)

Type 316CW SS

Irradiated in Fast Reactors

Irradiation 

& Test Temp. (°C)

(a)

0

5

10

15

0 10 20 30 40 50

290°C

371°C

427°C

483°C

500°C

Uniform Elongation (%)

Neutron Dose (dpa)

Type 316CW SS

Irradiated in Fast Reactors

Irradiation 

& Test Temp. (°C)

290°C

371°C

427°C

483°C

500°C

(b)

Figure 12. Effect of irradiation temperature on the (a) yield strength and (b) uniform elongation of CW 

Type 316 stainless steel irradiated in fast reactors at temperatures of 290-500°C and tested 

at the irradiation temperature (Ref. 55).

0

200

400

600

800

1000

1200

0 2 4 6 8 10

Type 316 SS(Bergenlid et al., 1990)

Type 3xx SS (Odette and Lucas. 1991)

Type 3xx SS (Kodama et al., 1997)

Type 304 SS (Jenssen, 1998)

Type 316 SS (Jenssen, 1998)

Type 3xx SS (Fukuya et al., 1993)

Type 304 SS (Allen et al., 2002)

Type 316 SS (Furutani et al., 2001)

Yield Stress (MPa)

Neutron Dose (dpa)

Figure 11.

The effects of neutron dose on measured 

tensile yield strength of 300-series stainless 

steels, irradiated and tested at about 300°C 

(Ref. 3).
14

2.1.4 Tensile Properties 

Tensile properties data have been obtained on solution-annealed and CW Type 304, 304L, 316, 

316L, and 347 SSs, including weld HAZ material, Type 308 and 309 material, and CF-8 cast austenitic 

SSs, irradiated at temperatures of 300-400°C (572-752°F) in fast reactors and LWRs.57-62 The 0.2% 

yield strength, ultimate tensile strength, uniform elongation, and total elongation at elevated temperatures 

are plotted in Figs. 13-15 as a function of neutron dose. Most data in these figures were obtained on 

materials irradiated in the BOR-60 fast reactor. Data on LWR-irradiated materials are limited, in 

particular, at high neutron dose. 

The curves in these figures represent the Materials Reliability Program (MRP)-developed 

correlations63 for estimating the tensile properties as a function of neutron dose. The 0.2% yield strength, 

ultimate tensile strength, uniform elongation, and total elongation data at 330°C were fitted to an 

exponential equation of the form:

Property = A0 +A1 (1-exp(-d/d0)), (1)

where d is the neutron dose in dpa and the coefficients A0, A1, and d0 for the irradiated material property 

equations are listed in Tables 1 and 2. The baseline reference tensile properties of nonirradiated materials 

are represented by a fourth-order polynomial63 of the form:

Property = C0 + C1T + C2T2 + C3T3 + C4T4, (2)

where T is the temperature (°C) and the coefficients for the different tensile property equations are listed 

in Table 3. 

Table 1. Material property equations for irradiated CW Type 316 stainless steel.

Property = A0 +A1 (1-exp(-d/d0))

Property Units A0 A1 d0

0.2% yield strength MPa 500 470 3

Ultimate tensile strength MPa 650 330 3

Uniform elongation % 10 -9.7 2

Total elongation % 18 -11 5

Table 2. Material property equations for irradiated solution-annealed Type 304 stainless steel.

Property = A0 +A1 (1-exp(-d/d0))

Property Units A0 A1 d0

0.2% yield strength MPa 200 600 3

Ultimate tensile strength MPa 450 350 3

Uniform elongation % 40 -39.5 1

Total elongation % 45 -37 2.5
15

200

400

600

800

1000

0 20 40 60 80 100

304L-1 (A) SA, Bor-60 at 320°C, 330°C

304L-1 (A) SA, Samara at 300°C, 330°C

304L-1 (A) SA, Samara at 300°C, 380°C

304L-1 (A) SA, Bor-60 +Samara, 330°C

304L-1 (CB) SA, Bor-60 at 320°C, 330°C

304-3 (FD) SA, Bor-60 at 320°C, 330°C

304 (EH or J7) SA, Bor-60 at 320°C, 330°C

304-J3 (EH) 10%CW, Bor-60 at 320°C, 330°C

347-2 (EC) SA, Bor-60 at 320°C, 330°C

304 Decom. PWR at 288-315°C, 320°C

304 Oskarshamn 1/2 at 280°C, 270-290°C

304L Oskarshamn 1/2 at 280°C, 288-295°C

304L Barseback 1 at 288°C, 288°C

347 Point Beach 2, 320°C 

304 BWR at 288°C, 288°C

304 Chooz A at 288°C, 320°C

304 BWR at 288°C, 300°C

304 BWR at 288°C, 300°C

304L BWR at 288°C, 300°C

Yield Strength (MPa)

Neutron Dose (dpa)

Type 304, 304L, & 347 SS

MRP Curve

(a)

200

400

600

800

1000

0 20 40 60 80 100

304L-1 (A) SA, Bor-60 at 320°C, 330°C

304L-1 (A) SA, Samara at 300°C, 330°C

304L-1 (A) SA, Samara at 300°C, 380°C

304L-1 (A) SA, Bor-60 +Samara, 330°C

304L-1 (CB) SA, Bor-60 at 320°C, 330°C

304-3 (FD) SA, Bor-60 at 320°C, 330°C

304 (EH or J7) SA, Bor-60 at 320°C, 330°C

304-J3 (EH) 10%CW, Bor-60 at 320°C, 330°C

347-2 (EC) SA, Bor-60 at 320°C, 330°C

304 Decom. PWR at 288-315°C, 320°C

304 Oskarshamn 1/2 at 280°C, 270-290°C

304L Oskarshamn 1/2 at 280°C, 288-295°C

304L Barseback 1 at 288°C, 288°C

347 Point Beach 2, 320°C 

304 BWR at 288°C, 288°C

304 Chooz A at 288°C, 320°C

304 BWR at 288°C, 300°C

304 BWR at 288°C, 300°C

304L BWR at 288°C, 300°C

Ultimate Tensile Strength (MPa)

Neutron Dose (dpa)

Type 304, 304L, & 347 SS

MRP Curve

(b)

Figure 13. Change in (a) yield strength, (b) ultimate tensile strength, (c) uniform elongation, and (d) total 

elongation as a function of neutron dose for solution-annealed Type 304, 304L, and 347

stainless steels at elevated temperature (270-380°C) (Refs. 57-62).
16

0.1

1

10

100

0 20 40 60 80 100

304L-1 (A) SA, Bor-60 at 320°C, 330°C

304L-1 (A) SA, Samara at 300°C, 330°C

304L-1 (A) SA, Samara at 300°C, 380°C

304L-1 (A) SA, Bor-60 +Samara, 330°C

304L-1 (CB) SA, Bor-60 at 320°C, 330°C

304-3 (FD) SA, Bor-60 at 320°C, 330°C

304 (EH or J7) SA, Bor-60 at 320°C, 330°C

304-J3 (EH) 10%CW, Bor-60 at 320°C, 330°C

347-2 (EC) SA, Bor-60 at 320°C, 330°C

304, Decom. PWR at 288-315°C, 320°C

304 Oskarshamn 1/2 at 280°C, 270-290°C

304L Oskarshamn 1/2 at 280°C, 288-295°C

304L Barseback 1 at 288°C, 288°C

347 Point Beach 2, 320°C 

304 BWR at 288°C, 288°C

304 BWR at 288°C, 300°C

304 BWR at 288°C, 300°C

304L BWR at 288°C, 300°C

Uniform Elongation (%)

Neutron Dose (dpa)

Type 304, 304L, & 347 SS

MRP Curve

(c)

10

100

0 20 40 60 80 100

304L-1 (A) SA, Bor-60 at 320°C, 330°C

304L-1 (A) SA, Samara at 300°C, 330°C

304L-1 (A) SA, Samara at 300°C, 380°C

304L-1 (A) SA, Bor-60 +Samara, 330°C

304L-1 (CB) SA, Bor-60 at 320°C, 330°C

304-3 (FD) SA, Bor-60 at 320°C, 330°C

304 (EH or J7) SA, Bor-60 at 320°C, 330°C

304-J3 (EH) 10%CW, Bor-60 at 320°C, 330°C

347-2 (EC) SA, Bor-60 at 320°C, 330°C

304, Decom. PWR at 288-315°C, 320°C

304 Oskarshamn 1/2 at 280°C, 270-290°C

304L Oskarshamn 1/2 at 280°C, 288-295°C

304L Barseback 1 at 288°C, 288°C

347 Point Beach 2, 320°C 

304 BWR at 288°C, 288°C

304 BWR at 288°C, 300°C

304 BWR at 288°C, 300°C

304L BWR at 288°C, 300°C

316L Barseback 1, 288°C 

Total Elongation (%)

Neutron Dose (dpa)

Type 304, 304L, & 347 SS

MRP Curve

(d)

Figure 13. (Contd.)
17

300

400

500

600

700

800

900

1000

1100

0 20 40 60 80 100

316-1 (B) 15%CW, Bor-60 at 320°C, 330°C

316-1 (B) 15%CW, Samara at 300°C, 330°C

316-1 (B) 15%CW, Samara at 300°C, 380°C

316-1 (B) 15%CW, BOR-60 + Samara, 330°C

316-2 (DB=J6) 15%CW, BOR-60 at 320°C, 330°C

316-3 (DA) 11%CW, BOR-60 at 320°C, 330°C

316-4 (DC) CW, BOR-60 at 320°C, 330°C

316-L1 (EA) CW, BOR-60 at 320°C, 330°C

316-L2 (EB) CW, BOR-60 at 320°C, 330°C

316-H1 (J5) CW, BOR-60 at 320°C, 330°C

316-1 (C) SA, BOR-60 at 320°C, 330°C

316 CW Farley 1 at 303-357°C, 320°C

316 CW Ringhals 2 at 290°C, 320*C

316CW PWR at 310-335°C, 290-340°C

316CW PWR at 290-310°C, 290-340°C

Yield Strength (MPa)

Neutron Dose (dpa)

Type 316 Cold Worked SS

MRP Curve

(a)

300

400

500

600

700

800

900

1000

1100

0 20 40 60 80 100

316-1 (B) 15%CW, Bor-60 at 320°C, 330°C

316-1 (B) 15%CW, Samara at 300°C, 330°C

316-1 (B) 15%CW, Samara at 300°C, 380°C

316-1 (B) 15%CW, BOR-60 + Samara, 330°C

316-2 (DB=J6) 15%CW, BOR-60 at 320°C, 330°C

316-3 (DA) 11%CW, BOR-60 at 320°C, 330°C

316-4 (DC) CW, BOR-60 at 320°C, 330°C

316-L1 (EA) CW, BOR-60 at 320°C, 330°C

316-L2 (EB) CW, BOR-60 at 320°C, 330°C

316-H1 (J5) CW, BOR-60 at 320°C, 330°C

316-1 (C) SA, BOR-60 at 320°C, 330°C

316 CW Farley 1 at 303-357°C, 320°C

316 CW Ringhals 2 at 290°C, 320°C

316CW PWR at 310-335°C, 290-340°C

316CW PWR at 290-310°C, 290-340°C

Ultimate Tensile Strength (MPa)

Neutron Dose (dpa)

Type 316 Cold Worked SS

MRP Curve

(b)

Figure 14. Change in (a) yield strength, (b) ultimate tensile strength, (c) uniform elongation, and (d) total 

elongation as a function of neutron dose for CW Type 316 stainless steel at elevated 

temperature (288-380°C) (Refs. 57-62).
18

0.1

1

10

100

0 20 40 60 80 100

316-1 (B) 15%CW, Bor-60 at 320°C, 330°C

316-1 (B) 15%CW, Samara at 300°C, 330°C

316-1 (B) 15%CW, Samara at 300°C, 380°C

316-1 (B) 15%CW, BOR-60 + Samara, 330°C

316-2 (DB=J6) 15%CW, BOR-60 at 320°C, 330°C

316-3 (DA) 11%CW, BOR-60 at 320°C, 330°C

316-4 (DC) CW, BOR-60 at 320°C, 330°C

316-L1 (EA) CW, BOR-60 at 320°C, 330°C

316-L2 (EB) CW, BOR-60 at 320°C, 330°C

316-H1 (J5) CW, BOR-60 at 320°C, 330°C

316-1 (C) SA, BOR-60 at 320°C, 330°C

316 CW Farley 1 at 303-357°C, 320°C

316 CW Ringhals 2 at 290°C, 320°C

316CW PWR at 310-335°C, 290-340°C

316CW PWR at 290-310°C, 290-340°C

Uniform Elongation (%)

Neutron Dose (dpa)

Type 316 Cold Worked SS

MRP Curve

(c)

10

100

0 20 40 60 80 100

316-1 (B) 15%CW, Bor-60 at 320°C, 330°C

316-1 (B) 15%CW, Samara at 300°C, 330°C

316-1 (B) 15%CW, Samara at 300°C, 380°C

316-1 (B) 15%CW, BOR-60 + Samara, 330°C

316-2 (DB=J6) 15%CW, BOR-60 at 320°C, 330°C

316-3 (DA) 11%CW, BOR-60 at 320°C, 330°C

316-4 (DC) CW, BOR-60 at 320°C, 330°C

316-L1 (EA) CW, BOR-60 at 320°C, 330°C

316-L2 (EB) CW, BOR-60 at 320°C, 330°C

316-H1 (J5) CW, BOR-60 at 320°C, 330°C

316-1 (C) SA, BOR-60 at 320°C, 330°C

316 CW Farley 1 at 303-357°C, 320°C

316 CW Ringhals 2 at 290°C, 320°C

316CW PWR at 310-335°C, 290-340°C

316CW PWR at 290-310°C, 290-340°C

Total Elongation (%)

Neutron Dose (dpa)

Type 316 Cold Worked SS

MRP Curve

(d)

Figure 14. (Contd.)
19

200

400

600

800

1000

0 20 40 60 80 100

308-1 (SA), BOR-60 at 330°C, 330°C

308-2 (SB), BOR-60 at 330°C, 330°C

308-3 (FE), BOR-60 at 330°C, 330°C

308-4 (J1), BOR-60 at 330°C, 330°C

HAZ (SB5-SB6), 304-SA BOR-60 at 330°C, 330°C

HAZ Base (J2), 304-CW BOR-60 at 330°C, 330°C

CF-8 Unaged (J4), BOR-60 at 330°C, 330°C

CF-8 Aged 400°C/100 h (J8), BOR-60 at 330°C, 330°C

CF-8 Aged 400°C/950 h (J9), BOR-60 at 330°C, 330°C

E308L Oskarshamn 1/2 at 280°C, 270-297°C

308 Cladding UBR at 288°C, 288°C

309 Cladding UBR at 288°C, 288°C

Yield Strength (MPa)

Neutron Dose (dpa)

Type 308, 308L Weld, 

304 HAZ, & Cast SS

MRP Curve

(a)

200

400

600

800

1000

0 20 40 60 80 100

308-1 (SA), BOR-60 at 330°C, 330°C

308-2 (SB), BOR-60 at 330°C, 330°C

308-3 (FE), BOR-60 at 330°C, 330°C

308-4 (J1), BOR-60 at 330°C, 330°C

HAZ (SB5-SB6), 304-SA BOR-60 at 330°C, 330°C

HAZ Base (J2), 304-CW BOR-60 at 330°C, 330°C

CF-8 Unaged (J4), BOR-60 at 330°C, 330°C

CF-8 Aged 400°C/100 h (J8), BOR-60 at 330°C, 330°C

CF-8 Aged 400°C/950 h (J9), BOR-60 at 330°C, 330°C

E308L Oskarshamn 1/2 at 280°C, 270-297°C

308 Cladding UBR at 288°C, 288°C

309 Cladding UBR at 288°C, 288°C

Ultimate Tensile Strength (MPa)

Neutron Dose (dpa)

Type 308, 308L Weld, 

304 HAZ, & Cast SS

MRP Curve

(b)

Figure 15. Change in (a) yield strength, (b) ultimate tensile strength, (c) uniform elongation, and (d) total 

elongation as a function of neutron dose for Type 308 and 309 weld metal, Type 304 HAZ, 

and CF-8 cast stainless steel at elevated temperature (270-330°C) (Refs. 57-62).
20

0.1

1

10

100

0 20 40 60 80 100

308-1 (SA), BOR-60 at 330°C, 330°C

308-2 (SB), BOR-60 at 330°C, 330°C

308-3 (FE), BOR-60 at 330°C, 330°C

308-4 (J1), BOR-60 at 330°C, 330°C

HAZ (SB5-SB6), 304-SA BOR-60 at 330°C, 330°C

HAZ Base (J2), 304-CW BOR-60 at 330°C, 330°C

CF-8 Unaged (J4), BOR-60 at 330°C, 330°C

CF-8 Aged 400°C/100 h (J8), BOR-60 at 330°C, 330°C

CF-8 Aged 400°C/950 h (J9), BOR-60 at 330°C, 330°C

Uniform Elongation (%)

Neutron Dose (dpa)

Type 308, 308L Weld, 

304 HAZ, & Cast SS

MRP Curve

(c)

1

10

100

0 20 40 60 80 100

308-1 (SA), BOR-60 at 330°C, 330°C

308-2 (SB), BOR-60 at 330°C, 330°C

308-3 (FE), BOR-60 at 330°C, 330°C

308-4 (J1), BOR-60 at 330°C, 330°C

HAZ (SB5-SB6), 304-SA BOR-60 at 330°C, 330°C

HAZ Base (J2), 304-CW BOR-60 at 330°C, 330°C

CF-8 Unaged (J4), BOR-60 at 330°C, 330°C

CF-8 Aged 400°C/100 h (J8), BOR-60 at 330°C, 330°C

CF-8 Aged 400°C/950 h (J9), BOR-60 at 330°C, 330°C

E308L Oskarshamn 1/2 at 280°C, 270-297°C

308 Cladding UBR at 288°C, 288°C

309 Cladding UBR at 288°C, 288°C

Total Elongation (%)

Neutron Dose (dpa)

Type 308, 308L Weld, 

304 HAZ, & Cast SS

MRP Curve

(d)

Figure 15. (Contd.)
21

Table 3. Material property equations for solution-annealed and nonirradiated Type 304 stainless steel.

Property = C0 + C1T + C2T2 + C3T3 + C4T4

Property Units C0 C1 C2 C3 C4

0.2% yield strength MPa 270.09 -0.5702 9.1162E-4 -5.6198E-7

Ultimate tensile strength MPa 617.275 -1.7750 7.0659E-3 -1.0769E-5 4.8941E-9

Uniform elongation % 55.8688 -0.1893 5.3656E-4 -4.5779E-7

Total elongation % 71.6321 -0.1956 5.7562E-4 -7.1266E-7 3.2172E-10

The results indicate that the tensile properties reach saturation at 5-20 dpa and do not change 

significantly at higher dose levels. Also, although the data on materials irradiated in LWRs are limited, 

the yield and ultimate strengths of the LWR-irradiated materials (solid symbols in Figs. 13 and 15) are 

consistently higher, and the uniform and total elongations are consistently lower than those for materials 

irradiated in fast reactors (open symbols in Figs. 13 and 15). The uniform elongation of some SSs can 

decrease to 0.1% at neutron dose levels of 5-10 dpa. The MRP correlations are primarily based on the 

fast reactor data and predict a saturated yield and ultimate strength value of 800 MPa for solutionannealed Type 304 SS and of 970 and 980 MPa, respectively, for CW Type 316 SS. However, most of 

the data for yield and ultimate strength of LWR-irradiated materials are above the MRP curve. Similarly, 

most of the data for uniform and total elongation of LWR-irradiated materials are below the MRP curve. 

Consequently, calculations of the tensile properties of irradiated austenitic SSs based on the MRP 

correlations underestimate irradiation hardening and loss of ductility in LWR conditions. 

Additional data on austenitic SSs irradiated in LWRs to dose levels above 20 dpa are needed to 

better define the correlations for estimating tensile properties of these materials for LWR operating 

conditions. Also, the contribution of voids in these materials to irradiation hardening and loss of ductility 

should be investigated. The presence of voids is likely to degrade the ductility of irradiated SSs to the 

extent that a brittle fracture at room temperature may be a significant concern. The effect of voids on the 

tensile properties of SSs should be examined at LWR operating temperatures as well as at room 

temperature.

Other correlations have also been proposed to estimate the tensile properties of irradiated austenitic 

SSs. The most widely used equation for the yield stress of 300-series SSs irradiated in LWRs at about 

300°C is that developed by Odette and Lucas.64 The increase in yield stress (MPa) is expressed in terms 

of the neutron dose (dpa) by the relationship 

Δσy = 670 [1 - exp(-dpa/2)]0.5. (3)

To help assess irradiation-induced degradation of fracture properties of BWR core internal components 

Electric Power Research Institute (EPRI) has proposed a correlation for estimating the flow stress of 

irradiated SSs as a function of fluence.65 The correlation is based on BWR-irradiated base metal and 

welds tested at 150-300°C. The flow stress σf (MPa), defined as the average of yield stress (σy) and 

ultimate stress (σu) [i.e., (σy+σu)/2], is expressed in terms of neutron fluence (n/cm2) by a power law 

relationship:

σf = 113.67 ln(fluence) - 4952.8. (4)
22

2.1.5 Yield Strength 

The three elements that, in combination, influence the IGSCC of austenitic SSs in BWR 

environments are a susceptible (sensitized) material, a significant tensile stress, and an aggressive 

environment. The NRC staff technical recommendations to reduce the susceptibility to IGSCC included 

the use of low-carbon wrought austenitic SSs and weld metal, which are considered adequately resistant 

to sensitization by welding.66 However, Andresen has shown that nonsensitized SSs are not immune to 

SCC; SSs with high yield strength (produced by cold work) are susceptible to IGSCC in ultra-high purity 

water.15 The effect of yield strength on SCC growth rates in high-purity water at high and low potential 

is shown in Fig. 16. The CGRs increase with increasing yield strength. Increases in yield strength can 

originate from surface or bulk cold work, weld shrinkage strain, precipitation hardening, irradiation 

hardening, or oxide dispersion hardening; all of these increase SCC growth rates. At a given yield 

strength, a similar susceptibility to cracking is observed in high- and low-potential environments, 

although the CGRs in low-potential water are an order of magnitude lower. There is no significant 

difference among the different grades of SSs, as well as no effect of deformation-induced martensite on 

CGRs. However, SSs that are cold worked by rolling favor transgranular (TG) rather than intergranular 

(IG) crack morphology.67

10-8

10-7

10-6

0 100 200 300 400 500 600 700 800

Crack Growth Rate (mm/s)

Yield Strength (MPa)

Unsensitized 304, 304L, & 316L SS & A600

288°C High Purity Water, 2000 ppb DO

CT Tests at 27.5 - 30 MPa√m

Circles = high martensite

Triangles = Alloy 600

Annealed Cold Worked

2 Sensitized Points

for Comparison

Very high

martensite

Predicted

Response

Effect of Yield Strength on Crack Growth Rate

Very low or no

martensite

(a)

10-9

10-8

10-7

0 100 200 300 400 500 600 700 800

Crack Growth Rate (mm/s)

Yield Strength (MPa)

Unsensitized 304, 304L, & 316L SS & A600

288°C High Purity Water, 95 or 1580 ppb DH

CT Tests at 27.5 - 30 MPa√m

Circles = high martensite

Triangles = Alloy 600

Annealed Cold Worked at -55°C or +140°C

2 Sensitized Points

for Comparison

Very high

martensite

Predicted

Response

Effect of Yield Strength on Crack Growth Rate

Very low or no

martensite

(b)

Figure 16. Effect of yield strength and martensite content on the SCC growth rate of stainless steels and 

Alloy 600 at 288°C in high-purity water (<0.10 mS/cm outlet) containing (a) 2000 ppb O2 and 

(b) 95 or 1580 ppb H2 (Ref. 15). 

The increase in CGR with increasing yield strength has been explained by Andresen15,39,68 and 

Shoji69 on the basis of the redistribution of strain field ahead of an advancing crack. Under constant 

stress intensity factor K, the redistribution of stress and strain fields as the crack advances provides the 

dynamic deformation or strain rate at the crack tip, which leads to SCC. In high strength materials 

because the plastic zone at a given K is smaller, the strain gradient at the crack tip is higher.

The dependence of SCC growth rates on the yield strength of nonirradiated, cold-worked, 300-

series austenitic SSs in an NWC BWR environment is shown in Fig. 17. In this figure, the curve was 

empirically determined to bound the available data. The results indicate a large increase in CGRs as the 

yield strength increased from 150 to 800 MPa and little or no increase at higher values. The available 

data also indicate that the SCC growth rates in the NWC BWR environment for sensitized SSs are higher 

than those for nonsensitized steels. Figure 18 shows CGR vs. yield strength for sensitized and 

nonsensitized SSs in the NWC BWR environment at 289°C and in the PWR environment at 340°C.70,71 

At yield strengths less than 800 MPa, the CGRs of sensitized materials are greater than those of 
23

nonsensitized materials, but for yield strength values of 800 MPa or higher, the CGRs in the low- and 

high-potential environments may be comparable. 

10-11

10-10

10-9

10-8

0 200 400 600 800 1000

Type 304, 1.4301

Type 347, 1.4550

Type 312, 1.4541

Type 316Ti, 1.4571

Crack Growth Rate (mm/s)

Yield Strength (MPa)

Stress corrosion cracking,

nonsensitized austenitic stainless steels

in simulated BWR water 288°C

Figure 17.

Crack growth rates of nonirradiated, cold 

worked, 300-series austenitic SSs in NWC 

BWR environment as a function of yield 

strength (Ref. 3).

10-11

10-10

10-9

10-8

0 200 400 600 800 1000

Sen. 304 (Castano et al., 2003) 

347 (Castano Marin et al. 2003)

304L (Shoji et al., 2003)

348 (Shoji et al., 2003)

Sen. 304L (Shoji et al., 2003)

Sen. 316L (Shoji et al., 2003)

Crack Growth Rate (mm/s)

Yield Strength (MPa)

Nonirradiated Austenitic SSs

289°C NWC BWR

30 MPa√m 

10-11

10-10

10-9

10-8

0 200 400 600 800 1000

Sen. 304 (Castano et al., 2003)

347 (Castano et al., 2003)

Sen. 304L (Shoji et al., 2003)

Sen. 316L (Shoji et al., 2003)

Crack Growth Rate (mm/s)

Yield Strength (MPa)

Nonirradiated Austenitic SSs

340°C PWR

Normalized to K=30 MPa√m 

(a) (b)

Figure 18. Change in SCC growth rates of nonirradiated, sensitized and nonsensitized, stainless steels 

as a function of yield strength in (a) NWC BWR environment at 289°C and (b) PWR 

environment at 340°C (Refs. 70,71). The curve is replicated from Fig. 17.

10-11

10-10

10-9

10-8

0 200 400 600 800 1000

304L (Ht. C3)

316L (Ht. C16)

316 (Ht. C21)

304 Shroud

316 Top Guide

Crack Growth Rate (mm/s)

Yield Strength (MPa)

Irradiated Austenitic SSs

CT Tests Normalized to 25 MPa√m 

289°C High Purity Water 

Open Symbols: NWC BWR 

Closed Symbols: HWC BWR 

Ooki et al., 2005

Figure 19.

Change in SCC growth rates at 289°C of 

LWR-irradiated austenitic stainless steels as 

a function of yield strength in NWC and HWC 

BWR environments (Refs. 11,72).

Similar behavior is observed in irradiated SSs, for which the increase in yield strength is due to 

irradiation hardening (Fig. 19). In the NWC BWR environment, CGRs in irradiated SSs increase rapidly 
24

for yield strength values in the range of about 200 to 800 MPa. The rates are at least an order of 

magnitude lower in the HWC BWR environment. However, for yield strengths above 700 MPa (i.e., SSs 

irradiated above 3 dpa) the CGRs are comparable in both high- and low-potential (i.e., NWC BWR and 

HWC BWR or PWR) environments. The effect of yield strength on the SCC growth rates of irradiated 

austenitic SSs is discussed further in Section 2.2.2.1 of this report.

The SSRT tests on irradiated SSs show a different dependence on yield strength. The percent 

IGSCC in SSRT tests on austenitic SSs in high-DO water at 289°C is plotted as a function of yield 

strength in Fig. 20. As in the case of the CGR data, IASCC susceptibility, measured as the fraction of 

intergranular cracking, increases with yield strength, but the increase in percent IG fracture occurs at 

higher yield strength values of 400-800 MPa. These results suggest that irradiated SSs that have yield 

strengths of 150-400 MPa and are subject to SCC may not show an IG fracture mode. 

0

20

40

60

80

100

0 200 400 600 800 1000

HP 316 SS with solute 

 additions (Busby 2003)

3xx SS (Bruemmer et al. 1999)

304 & 316 SS (Chung 2006)

HP 304 & 316 SS (Chung 2006)

%IGSCC

Yield Stress (MPa)

Irradiated Austenitic SSs

SSRT Tests in High DO Water

Figure 20.

Effect of yield strength on percent IGSCC in 

SSRT tests on irradiated 300-series stainless 

steels in high-DO water (Refs. 3,4,53).

2.1.6 Corrosion Potential 

The increase in SCC susceptibility of SSs with an increase in corrosion potential is due to the effect 

of the corrosion potential on the crack-tip water chemistry. The CGR data demonstrate that decreasing 

the corrosion potential from that of an NWC BWR environment to that of HWC BWR or PWR 

environments decreases the CGRs for all nonirradiated SSs. Although SSs are not immune to SCC in 

low-potential environments, the SCC growth rates of nonirradiated SSs and SSs irradiated to moderate 

levels of neutron dose (less than 3 dpa) are more than an order of magnitude lower in such environments 

than those in high-potential environments. 

The SCC growth rates of austenitic SSs irradiated to 1.0-2.5, 3.0-4.0, and more than 5 dpa are 

plotted as a function of the steel electrochemical potential (ECP) in Fig. 21.

11,15,60,73-79 The CGRs 

(da/dt) are expressed by a power law relationship with the stress intensity factor, K. The growth rates 

(m/s) for nonirradiated austenitic SSs are represented by the correlation given in the NRC report NUREG0313, Rev. 2:66

da/dt = A1 (K)2.161, (5)

where K is in MPa m1/2, and the magnitude of the constant A1 depends on the water chemistry and 

composition and structure of the steel. A value for the constant of 2.1 x 10-13 was proposed in NUREG0313 for sensitized SS in water with 8 ppm DO. For water with 0.2 ppm DO, the CGR is taken as onethird that of the value given in NUREG-0313; in this case A1 is 7.0 x 10-14. The value of constant A1 is 
25

smaller in low-DO environments, such as HWC BWR or PWR environments. The CGRs plotted in 

Fig. 21 were corrected to a K value of 20 MPa m1/2 by using the experimental value of the exponent of K, 

and if not known, using a value of 2.5. For nonirradiated SSs, the CGRs are typically corrected to 

30 MPa m1/2 to determine the effect of corrosion potential on SCC growth rates.40,68 

10-12

10-11

10-10

10-9

10-8

10-7

-800 -600 -400 -200 0 200

ANL (1.3-2.3 dpa)

Halden (1.4-2.0 dpa)

JAEA (0.8-1.8 dpa) 

Experimental CGR (m/s)

ECP [mV (SHE)]

Austenitic SSs

288°C

All data normalized 

to K value of 20 MPa m1/2 10-12

10-11

10-10

10-9

10-8

10-7

-800 -600 -400 -200 0 200

ANL (3 dpa)

Halden (3 dpa)

Andresen (4 dpa)

Jenssen (3.8 dpa)

Halden (3.0 dpa)

Experimental CGR (m/s)

ECP [mV (SHE)]

Austenitic SSs

288°C

All data normalized 

to K value of 20 MPa m1/2

(a) (b)

10-12

10-11

10-10

10-9

10-8

10-7

-800 -600 -400 -200 0 200

Jenssen (13 dpa)

Jenssen (5-11 dpa Fast)

Jenssen (25 dpa Fast) PWR

Halden (13-18 dpa)

Halden (37.5 dpa)

Halden (11-18 dpa) PWR

Halden (37.5 dpa ) PWR

Experimental CGR (m/s)

ECP [mV (SHE)]

Austenitic SSs 288°C

PWR data at 320°C

Normalized to K=20 MPa m1/2

Figure 21.

Experimental CGR under constant load in NWC 

and HWC BWR environments at 288°C for 

austenitic stainless steels irradiated to (a) 1.0-

2.5 dpa, (b) 3.0-4.0 dpa, and (c) 13 dpa, plotted 

as a function of the steel ECP versus a standard 

hydrogen electrode (SHE) 

(Refs. 11,15,60, 73-79).

(c)

The effect of reduced corrosion potential on the CGRs of SSs that were irradiated to less than 3 dpa 

is seen clearly in these figures. Decreasing the corrosion potential has a beneficial effect on growth rates 

for all steels irradiated to 1.0-2.5 dpa (Fig. 21a), and for some steels irradiated to 3.0-4.0 dpa (Fig. 21b). 

Although the materials irradiated to about 4 dpa show some decrease in CGRs in the low-potential 

environment, the decrease is rather modest (about a factor of 5 lower growth rates). For these steels, the 

CGRs in the HWC BWR environment are greater than 3 x 10-10 m/s. The results also show that, except 

for the materials irradiated in fast reactor (BOR-60) to 5-11 dpa at 320°C,79 none of the other materials 

irradiated in BWRs (at 289°C) showed a benefit of decreased corrosion potential on SCC growth rates. 

For SSs irradiated to 11-38 dpa in BWRs, the CGRs in NWC and HWC are comparable. However, the 

rates vary by more than an order of magnitude with neutron dose. The CGRs are, in fact, lower for SS 
26

irradiated to 38 dpa compared to those irradiated to 11-13 dpa. The SSs irradiated in the BOR-60 reactor 

to 5-11 dpa at 320°C show more than an order of magnitude decrease in growth rate in the HWC BWR 

environment, compared to that in the NWC environment. The different behavior for the materials 

irradiated in BWRs and fast reactors may be attributed to differences in the material microstructure and 

microchemistry, as discussed in Sections 2.1.1 and 2.1.2, respectively. However, it is not clear whether 

these differences are due to differences in the irradiation temperature (320 versus 289°C) or neutron 

spectrum and flux. Data on SCC growth rates and microstructural characterization of SSs irradiated in 

PWRs (at 320°C or higher temperatures) are needed to establish the effect of irradiation temperature on 

IASCC susceptibility of austenitic SSs. 

The SCC growth rates for SSs irradiated above 11 dpa in the PWR environment at 320°C show 

significant variation between different grades and heats of materials. The CGRs range from 5 x 10-11 m/s 

for some steels to 5 x 10-9 m/s for others. The differences have been explained on the basis of differences 

in the material microchemistry, in particular, the Si concentration at the grain boundaries.60,76,79 

The CGR data also indicate that for some SSs a benefit of reduced potential is observed at low K 

values (<20 MPa m1/2) but not at a higher K value.11 The behavior could be explained on the basis of a 

loss of specimen constraint by exceeding the K/size criterion or may indicate some kind of threshold 

phenomenon. For example, loss of the beneficial effect of decreased corrosion potential seems to occur at 

a threshold growth rate of about 1 x 10-9 m/s. At this threshold growth rate, the crack tip strain rate may 

be influenced by some mechanism that helps sustain the high growth rates, which are independent of 

corrosion potential and to some extent K. Under constant K, crack advance occurs by dynamic 

deformation or strain rate at the crack tip due to the redistribution of stress and strain fields as the crack 

advances. Any mechanism that enhances crack tip strain rate would facilitate SCC. A possible 

mechanism is the effect of hydrogen, generated from the oxidation/dissolution of the freshly exposed 

metal surface, on the crack tip strain rate. Hydrogen lowers the SFE, which promotes localized 

deformation and thereby influences the crack tip strain rate. The effect of SFE on IASCC susceptibility is 

discussed in Section 2.1.9.

The lack of an effect of reduced corrosion potential on CGRs for some SSs irradiated to 3-4 dpa11

and nearly all SSs irradiated above 5 dpa60,74-79 has often been explained on the basis of the loss of 

specimen constraint. The argument is that the specimen K/size criterion was exceeded at the high K 

values used in these tests. However, the basis for the proposed K/size criterion for irradiated materials is 

not well documented. Some recent studies have shown that the ASTM E399 size criterion is not 

exceeded in 1⁄2-T or 1⁄4-T compact tension (CT) specimens of irradiated SSs that have yield stress above 

600 MPa.80,81 The effects of K/size criterion on SCC growth rates are discussed further in Section 

2.2.1.3. Post-test metallographic evaluation of such test specimens does not suggest a loss of specimen 

constraint; the fracture morphology in these specimens is completely IG, and the fracture surface is 

normal to the stress axis.11 The existing data suggest that some SSs may show IASCC susceptibility in 

low-potential environments at fluence levels as low as 3 dpa. Caution should be exercised in using a 

K/size criterion for irradiated SSs that screens out such experimental data simply because they do not 

agree with the current expectations for the IASCC behavior of austenitic SSs. Additional data on SSs 

irradiated to 3-8 dpa are needed to better determine the effect of low-potential environments on the CGRs 

of highly irradiated materials.

2.1.7 Silicon Segregation

The increased susceptibility to IASCC and the loss of benefit of reducing potential in highly 

irradiated SSs have been attributed to the segregation of Si at the grain boundaries. The Si segregation is 
27

detrimental because at all relevant pH/potential conditions, Si oxidizes to SiO2, which is highly soluble in 

hot water. 

Unusually high CGRs have been observed in crack growth tests on CW SSs, containing 1.5 to 

5 wt.% Si, in high-purity water with 2 ppm dissolved oxygen (DO), and the growth rates show no effect 

of corrosion potential and little effect of applied K.82 Crack growth tests were conducted in high-purity 

water with 2000 ppb O2 on Type 304L SS containing 5% Si (no Mo, Nb, or Ti additions) that had been 

cold worked to 15% at 220°C (yield strength about 500 MPa). The results yielded a CGR of 2.2 x 10-

9 m/s at K of about 30 MPa m1/2. The observed rate is a factor of 10 higher than the rates in most 

commercial SSs with similar levels of cold work/yield strength. More significant, the rate did not 

decrease when the environment was changed from high potential to low potential (95 ppb H2). In 

addition, decreasing K at a rate of about 22 MPa m1/2 per mm resulted in an initial decrease in CGR from 

2.2 to 1 x 10-9 m/s but thereafter it remained constant as the K decreased from 30 to 15 MPa m1/2 (Figs. 8 

and 9 in Ref. 82). 

A similar detrimental effect of Si, although not to the same extent, has been observed in custom 

heats of Type 304L SS with 3 or 1.5% Si, and elevated Ni and reduced Cr consistent with the segregation 

behavior expected in an irradiated material. The CGRs in these heats were nearly identical to those 

observed in the heat with 5% Si, but in this case a decrease in corrosion potential resulted in a modest 

decrease in growth rates. The CGR decreased from 2.1 to 1.2 x 10-9 cm/s in the heat with 3% Si to about 

7 x 10-10 m/s in the heat with 1.5% Si. Also, for the 1.5%-Si heat, decreasing K from 30 to 16 MPa m1/2

decreased the CGR by a factor of 9. The effect of Si concentration on SCC growth rates in these custom 

heats of Type 304L SSs in various water chemistries and at various stress intensity factors is summarized 

in Fig. 22. 

The effect of bulk Si content on the SCC behavior of Type 316L SS has also been evaluated. Crack 

growth tests were performed on heats of Type 316L with 2 and 4% Si.82 None of these heats showed 

elevated growth rates. The reason for the different behavior in Type 316L compared with Type 304L is 

not clear. It is possible that molybdenum silicides form in Type 316L SS thereby removing the effect of 

Si on SCC susceptibility. 

10-10

10-9

10-8

0 1 2 3 4 5 6

Crack Growth Rates (m/s)

Percent Silicon

Effect of Si on SCC in Type 304L SS

in Pure Water at 288°C

30 MPa m0.5

High corrosion potential (2 ppm O2)

30 MPa m0.5

Low corrosion potential (95 ppb H2)

22 MPa m0.5

Low corrosion potential (95 ppb H2)

17.5 MPa m0.5

Low corrosion potential (95 ppb H2)

Figure 22.

Crack growth rates versus Si concentration 

in custom heats of Type 304L stainless 

steel simulating expected grain boundary 

microchemistries for high and low corrosion 

potentials and at various stress intensity 

factors (Ref. 82).

Tests with irradiated austenitic SSs have shown behavior similar to that observed in the tests on the 

model alloys, i.e., a relatively low K dependence and no benefit of decreased corrosion potential on SCC 

growth rates. For example, CGRs in BWR control-rod blade material (Type 304L SS) irradiated to about 

12 dpa were one to two orders of magnitude higher than those in the same material irradiated to 3 dpa.76
28

The material irradiated to 12 dpa contained a high concentration of Si (~3 wt.%) at the grain boundary. 

Also, a cold-worked Type 316+Ti SS baffle bolt material, irradiated to 25 dpa in the BOR-60 reactor at 

320°C, showed growth rates that were two orders of magnitude higher than those observed in coldworked Type 316 SS baffle bolt material irradiated under the same conditions. The Ni and Si 

concentrations at the grain boundary were higher in the Type 316+Ti SS.79 However, the validity of 

these tests is considered questionable because either the stress intensity factors exceeded the allowed 

values based on the proposed K/size criterion for irradiated materials or the duration of the test period was 

inadequate. The effects of K/size criterion are discussed in Section 2.2.1.3.

A common feature among the tests on high-Si SSs and irradiated SSs is that they all seem to have a 

relatively high growth rate (i.e., about 1 x 10-9 m/s or higher). As discussed in the previous section, this 

rate could be associated with some kind of threshold phenomenon that influences the crack tip strain rate. 

In contrast to the increased susceptibility of high-Si materials indicated by crack growth tests on CT 

specimens, SSRT test data on several laboratory heats of irradiated Type 304 and 316 SSs with 0.5 to 

1.8 wt.% Si show the opposite behavior. For the same fluence level, steels with 1.5 to 1.8 wt.% Si 

showed less irradiation hardening and greater elongation than steels containing 0.5 wt.% Si.53 The 

percent IG cracking was not significantly different in the low- or high-Si heats of SSs. The addition of Si 

to pure austenitic SSs without Mo has been shown to promote dislocation loop nucleation. The presence 

of Mo negates the effect of Si on loop nucleation by trapping Si atoms and preventing them from binding 

with SIAs; Mo-Si clusters form instead.83 

Because most commercial SSs contain about 0.5% Si, its effect on irradiation hardening is not 

likely to be significant. The tests on model alloys certainly suggest that high concentrations of Si at grain 

boundaries due to RIS could greatly increase susceptibility to IASCC. However, the absence of γ’ silicide 

(Ni3Si) at the grain boundaries suggests that RIS of Ni and Si probably saturates at higher doses; the 

actual dose and saturation concentrations may vary with material composition and irradiation 

temperature. The significance of Si segregation at grain boundaries on the SCC behavior of irradiated 

SSs is not clear, but the available results suggest that this segregation could be important. Crack growth 

tests should be conducted on irradiated material from commercial heats of SSs with similar compositions 

but different Si contents to establish the role of Si segregation on the IASCC susceptibility of irradiated 

SSs. 

2.1.8 Sulfur Content

Based on SSRT tests, for several commercial heats of Type 304 and 316 SS that contain S 

concentrations of ≤0.002 wt.%, excellent resistance to IASCC is observed for neutron fluence up to 

≈3 dpa.53 As the S content is increased above 0.003 wt.%, IASCC susceptibility (expressed as percentage 

IGSCC) increased drastically as shown in Fig. 23 for irradiated Types 304, 304L, 316, and 316L SS in 

high-DO water at 289°C.53 

The IASCC susceptibility of Type 304 or 316 SSs has been represented by a two-dimensional map 

of bulk S and C contents in the steel to show the range in which these steels are either resistant or 

susceptible to IASCC in the NWC BWR environment. An S-C content map for Type 304 and 316 SS in 

BWR-like oxidizing environments is shown in Fig. 24. At sufficiently low concentrations of S 

(≤0.002 wt.%), a high concentration of C suppresses the deleterious effect of S. However, when S 

concentration is high (≥0.003 wt.% in Type 304 and 316 SS, and ≥0.01 wt.% for Type 348 SS), the 

deleterious effect of S is so dominant that the effect of C is negligible or insignificant, even at high 

concentrations. Typically, in-service cracking of BWR core internal components occurs at >2 dpa and is 
29

characterized by a virtually full intergranular fracture surface, sometimes accompanied by a small amount 

of TG fracture surface. In constructing Fig. 24, Chung and Shack determined the resistance to IASCC 

based on the negligible IGSCC observed in SSRT tests conducted in high-DO (8-32 ppm) water with 

fluence levels of 3 dpa or greater. Limited data indicate a similar behavior of the effect of S and C 

contents on IASCC susceptibility in the PWR environment. The S-C content map for austenitic SSs in 

the PWR environment is shown in Fig. 25.

0

0.01

0.02

0.03

0.04

0.05

0.06

0.07

0.08

0.09

0.1

0.11

0.12

0 0.002 0.004 0.006 0.008 0.01

R304/316 , He-ion, SSRT, Fukuya '93

S304, proton, Busby '03

S304, BWR, CGR, Jenssen '03

R316, BWR, SSRT, Jacobs '93

R316, BWR, SSRT, Tanaka '97

S1.4541/316, BWR, tube expand, Garzarolli '87

SHP304, He-ion, SSRT, Fukuya '92

S304/316, ATR, SSRT, Jacobs '87

S304/316, BWR, SSRT, Fukuya '91 '93

SHP304, BWR, SSRT, Chung '93 '95

SHP304, proton, SSRT, Cookson '93

S304/316, BWR, SSRT, Kodama '91 '93 '95 '99

S304/316, BWR, SSRT, Katsura '93

S304/316, JRR, SSRT, Tsukada '95 '97

S304/316, BWR, SSRT, Jenssen '95

S304/316, BWR, SSRT, Tanaka '97

R316, BWR, SSRT, Kodama '99

S304L, Halden, CGR, Karlsen '99

R304/316, Halden, SSRT, this study

S304/316, Halden, SSRT, this study

Bulk Carbon Content (wt.%)

Bulk Sulfur Content (wt.%)

slashed region

resistant to IASCC

this side

susceptible

to IASCC

304- or 316-type SSs

BWR-like water

DO 8-32 ppm

open symbols, resistant to IASCC

filled symbols, susceptible to IASCC

L14 = lab heat #14

C1 = commercial heat #1

32% = 32 percent IGSCC

R304 = resistant 304 SS

S316 = susceptible 316 SS

HP = high purity

proton

proton

proton

He ion proton CGR

CGR

CGR

Figure 24. Range of bulk S and C contents, in which Type 304 or 316 stainless steels are resistant or 

susceptible to IASCC, in BWR-like oxidizing water (Ref. 53). 

2.1.9 Stacking Fault Energy

As discussed earlier, irradiation hardening promotes dislocation channeling and localized 

deformation. The SFE of the material is an important parameter that determines the deformation mode. 

In SSs, the SFE increases with increased Ni content. Slow strain rate tensile test results have shown that 

high-Ni alloys have good resistance to IASCC. Alloys with Ni concentration >18 wt.% are highly 

0

10

20

30

40

50

60

70

80

90

100

110

120

0.0 0.0040 0.0080 0.012 0.016 0.020 0.024 0.028 0.032 0.036 0.040

Jacobs 1987, 304, 316L, 316NG, ATR 260-330°C

Fukuya 1993, 316L, BWR

Tanaka, 1997, 316L, BWR

this study, 304, 304L, 316, 316L, Halden HBWR, 289°C

Percent IGSCC from SSRT Test

Bulk Sulfur Content (wt.%)

SSRT at 289°C, DO !8-32 ppm

304, 304L, HP304L, 316, 316L, 

HP316L, 316NG (HP=high-purity)

Fluence 2.9 – 3.5 dpa

Tsukada 1995, HP 304+S, JRR 240°C

1.0 dpa, SSRT at 300°C, DO 32 ppm.

Higher IG% predicted at 3.0 dpa.

Figure 23.

Effect of bulk S content on IASCC 

susceptibility (percent IGSCC) of 

commercial heats of Type 304, 304L, 

316, and 316L tested in high DO water 

at 289°C (Ref. 53).
30

resistant to IASCC compared to Type 304 SS with 8 wt.% Ni.3 Alloys with low SFE, such as SSs with 

8 wt.% Ni, deform entirely by planar slip, whereas there is no evidence of planar slip in alloys with high

SFE (i.e., alloys with >20 wt.% Ni). The deformation microstructure of low-SFE materials consists of 

well-defined slip bands or dislocation channels. As the Ni content increases, the cross slip increases. The 

deformation microstructure of high-SFE materials consists of dislocation tangles. 

The effect of SFE on IASCC susceptibility (measured by percent IG cracking) of austenitic SSs is 

shown in Fig. 26. The results indicate a high susceptibility to IGSCC for materials with low SFE. 

Typically, Type 3xx SSs have relatively low SFE. In such materials, partial dislocations are widely 

separated and cannot readily combine; this condition reduces cross slip and promotes planar slip and 

localized deformation. The higher resistance of Type 316 SS compared to Type 304 SS may be explained 

on the basis of differences in the deformation mode because of the higher SFE in Type 316 SS.

A model for IASCC, based on the deformation mode, has been proposed.3 In the model, IASCC 

susceptibility of SSs is explained in terms of two processes. In the first, the pileup of dislocations in the 

0

0.01

0.02

0.03

0.04

0.05

0.06

0.07

0.08

0 0.004 0.008 0.012 0.016 0.02 0.024 0.028

348 resistant, Garzarolli 1995

3 48 susceptible, Garzarolli 1987, 1995

304 resistant, Garzarolli 1995

304 susceptible, Garzarolli 1995

316 susceptible, Garzarolli 1987

1.4541 susceptible, Garzarolli 1987

1.4981 susceptible, Garzarolli 1993 

Ti-321 susceptible, Ehrensten 1999, VVWR

CW316 susceptible, Fukuya 2001

CW316 susceptible, Shogan 2001

Bulk Carbon Content (wt.%)

Bulk Sulfur Content (wt.%)

this region

resistant to

IASCC

this side

susceptible

to IASCC

304, 316, 321, or 348

type SSs in PWR water

Ti-containing

321-typeSS

in VVER

open symbols resistant to IASCC

closed symbols susceptible to IASCC

Figure 25. Range of bulk S and C contents, in which Type 304, 316, 321, or 348 stainless 

steels are resistant or susceptible to IASCC, in PWR water. 

0

20

40

60

80

100

0 50 100 150

HP 316 p to 5.5 dpa (Busby 2003)

HP 3xx, p to 0.6 dpa (Cookson 1996)

3xx SS, >7 dpa (Jenssen et al. 1996)

Alloy 800, >7 dpa (Jenssen et al. 1996)

3xx SS, ~1 dpa (Kodama et al. 1994)

3xx, ~1 dpa (Chung et al. 1999)

3xx SS, ~3 dpa (Tanaka et al. 1997)

IGSCC (%)

Stacking Fault Energy (mJ/m2)

SFE = 1.2 + 1.4%Ni + 0.6%Cr 

 +17.7%Mn - 44.7%Si

 (Rhodes & Thompson 1977)

Figure 26.

IASCC susceptibility as measured by 

percent IG cracking as a function of 

stacking fault energy (Ref. 3).
31

slip bands or channels intersecting the grain boundaries creates progressively higher stresses at the grain 

boundary at the head of the pileup. When the stress exceeds a critical value, separation of the boundary 

may occur by a wedge cracking mechanism. This process does not depend on the environment and may 

be the mechanism for IG cracking observed in highly irradiated SSs in an inert environment. 

Susceptibility at lower fluences results from a second process. In this case, stress at the grain boundary 

can change the crack tip strain rate by rupturing the oxide film, thereby exposing bare metal surface and 

leading to oxidation/dissolution and crack extension. Alternatively, absorption of dislocations at the grain 

boundaries may cause grain boundary sliding ahead of the crack tip, resulting in crack extension and 

IASCC. The role of the localized deformation mode on IASCC of austenitic SSs and the possible 

contribution of SFE should be further investigated. 

2.1.10 LWR vs. Fast Reactor Irradiations

Comparison of data on materials irradiated in fast reactors and LWRs indicates that the fast reactor 

irradiations may not be prototypic of LWRs, particularly at fluence levels above 5 dpa. The differences in 

the neutron spectrums could lead to higher defect survival rates in LWRs compared to fast reactors, 

whereas the differences in flux and temperature could lead to lower defect survival rates in LWRs 

compared to fast reactors. In addition to possible differences in RIS and radiation hardening, differences 

in helium generation, because helium can migrate to grain boundaries, could play an important role in 

IASCC of PWR core internals. As discussed in Section 2.1.2, microchemistry characterization indicated 

that the RIS is greater in SSs irradiated in LWRs than in fast reactors. Also, for similar irradiation 

conditions, small cavities (probably He bubbles) are observed in LWR irradiated materials. 

The effect of accelerated irradiation and helium generation on the IASCC susceptibility of SSs in 

PWR environment has been evaluated by comparing the results for materials irradiated in a fast reactor to 

40 dpa with those irradiated in actual PWRs to 65 dpa.51 Although the tensile properties were not 

significantly different for the materials irradiated in a fast reactor or a PWR, the fast-reactor-irradiated 

materials showed less helium and lower IASCC susceptibility compared to the PWR-irradiated materials. 

Also, IG cracking was observed in highly irradiated materials tested in an inert environment.

Slow strain rate tensile test results from another study also indicate that fast reactor irradiations may 

not be representative of LWR irradiations.* In Fig. 27, the yield strength (YS) and uniform elongation 

(UE) measured from SSRT tests on several heats of Types 304 and 304L SS irradiated in the Halden 

reactor at ~290°C to 3 dpa are compared with those measured on the same heats of SSs irradiated in the 

BOR-60 fast breeder reactor at 320°C to 4.8 dpa. The results indicate that although the total dose for the 

Halden materials is lower than that of the BOR-60 materials, their YS is higher, and their UE is lower. 

However, it cannot be ascertained whether these differences are due to the differences in irradiation 

temperature or flux and energy spectrum for the two irradiations. Additional studies are needed to 

validate the applicability of the fast reactor data to LWRs.

2.2 Crack Growth Rates

Experimental data have been obtained from studies sponsored by the NRC at Argonne, the Halden 

Reactor Project in Norway, the Cooperative IASCC Research (CIR) program, industry sponsored 

programs such as the BWR Vessel and Internals Project (BWRVIP) and Materials Reliability Program 

(MRP), and studies conducted at various institutions in Japan, EdF/Framatome in France, and Studsvik in

 

*Yiren Chen, Argonne National Laboratory, private communications, “Evaluation of the Causes and Mechanisms of IASCC in 

PWRs – Preliminary Results of IASCC Study on BOR-60 Materials,” Sept. 2007.
32

Sweden. The data were obtained from specimens either machined from materials from LWR core 

internal components (mostly BWR components) or irradiated, after machining, under LWR irradiation 

conditions. However, some of the high neutron dose data have been obtained on materials irradiated in 

fast reactors (e.g., nearly all of the data obtained under the CIR program have been on specimens 

irradiated in the BOR-60 reactor in Russia). The applicability of the fast-reactor data for estimating 

irradiation effects under LWR operating conditions needs to be evaluated. 

As discussed previously, SCC is a complex phenomenon that is dependent on several parameters 

such as material composition and processing, irradiation effects, water chemistry and electrochemistry, 

loading condition, and temperature. In addition, these parameters interact with each other such that their 

effect on growth rates is not always the same. For example, low-potential HWC BWR or PWR 

environments reduce susceptibility in SSs irradiated up to about 5 dpa, but not in SSs irradiated to higher 

dose levels. Sensitization (Cr depletion) due to thermal treatment or irradiation, as well as the presence of 

chlorides or sulfates in water, is known to have a strong effect on SCC susceptibility of austenitic SSs in 

the high-potential NWC BWR environment but little effect on SCC in low-potential HWC BWR or PWR 

environments.82 As a result of these complexities and uncertainties in experimental measurements, the 

CGR vs. K data in LWR environments show significant variability (up to 1000X), and it is not yet 

possible to determine specific effects of these parameters on SCC growth rates with high confidence 

under all conditions of interest. Therefore, most investigators have emphasized the importance of high 

quality data and used caution in interpreting the results. 

Crack growth rates have been obtained on irradiated wrought and cast austenitic SSs as well as SS 

weld metal and weld HAZ materials. There is little or no CGR data on irradiated Ni alloys, such as 

Alloy 600 or X750, that are used for some core internal components.

The quality of the CGR data has been evaluated or screened according to the following factors: 

(a) test procedures that closely reproduce the loading and environment conditions for reactor core internal 

components, (b) data acquisition and analysis system that provides accurate measurements and 

reproducible results, and (c) well-characterized material. These concerns and problems regarding the 

quality of experimental data for CGR are discussed further in the following section. The potential sources 

of uncertainty in experimental measurements and data scatter are also discussed.

600

700

800

900

1000

600 700 800 900 1000

YS of Halden specimens (MPa)

YS of BOR-60 specimens (MPa)

~ 50 MPa

HP304L SS

HP304L SS

HP 304-like alloy

304 SS

304L SS 304 SS

(a)

0

5

10

15

0 5 10 15

UE of Halden specimens (%)

UE of BOR-60 specimens (%)

Estimated

HP304L SS

HP 304-like alloy

304L SS

304 SS

304 SS

HP304L

SS

304 SS

(b)

Figure 27. Measured (a) yield strength and (b) uniform elongation for several heats of Type 304 and 

304L stainless steels irradiated in the Halden reactor at 290°C to 3 dpa and in BOR-60 fast 

breeder reactor at 320°C to 4.8 dpa. 
33

2.2.1 Quality of Crack Growth Rate Data

2.2.1.1 Test Procedures

Nearly all of the recent CGR data on irradiated materials in simulated BWR and PWR 

environments have been obtained on CT specimens (1/4-T or 1/2-T CT) or four-point bend specimens. 

The tests are performed in accordance with American Society for Testing and Materials (ASTM) E-647, 

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

Test Method for Determining a Threshold Stress Intensity Factor for Environment-Assisted Cracking of

Metallic Materials under Constant Load.” 

The specimens are fatigue pre-cracked (typically in air) to create a starter crack and are actively 

loaded in the environment of interest at constant K or constant load to determine the SCC growth rates 

and under cyclic loading to measure fatigue CGRs. The fatigue CGR tests are conducted with a triangular 

or slow/fast saw-tooth waveform with rise times up to 10,000 s and load ratio R between 0.2 and 0.9, and 

the SCC growth rate tests are conducted under constant K with or without periodic partial unloading. The 

stress intensity factor for SCC tests (and maximum stress intensity factor Kmax for fatigue tests) is in the 

range 3-40 MPa m1/2. Also, the reversing direct-current (DC) potential difference method is used to 

monitor crack extension. The crack extensions determined from the DC potential drop method are 

corrected to match the final optically measured crack length. In most studies, crack lengths obtained from 

the DC potential drop method have shown good agreement with post-test fractography measurements; 

typically crack lengths measured by the potential drop method are 5-20% smaller.

During each test period (i.e., a specific loading and environmental condition), the crack extension is 

at least 10 times the resolution of the DC potential drop method (typically ≈5 μm). Thus, crack 

extensions are at least 50 μm and range up to 500 μm. The CGRs are determined from the slope of the 

corrected crack length (a) vs. time (t) plots. For cyclic loading, the CGR can be expressed in terms of a 

time-based growth rate (da/dt) or a growth rate per cycle (da/dN). However, the method for calculating 

da/dt varies among different investigators. Crack growth studies on austenitic SSs by Shack and 

Kassner84 in the simulated BWR environment, and by Tice et al.85,86 in the PWR environment have used 

a saw-tooth waveform (where the rise time was more than 85% of the period of the cycle) and da/dt was 

determined by dividing the crack advance per cycle by the rise time for the cycle. 

Other investigators have used triangular or sinusoidal waveforms and determined CGRs by dividing 

the da/dN by the total time period for the cycle. In some studies, to account for the effects of surface 

oxide, crack closure, and roughness between the edges of the crack, an effective time of loading, teff, is

used instead of the total time period t.87 Thus, da/dt is determined from the relationship

da

dt = 1t

eff

da

dN

!

"

# $

%

& , (6)

where the teff for loading cycle ranges from 0.6t, at R = 0.1, to t, at R = 0.7 or greater.

2.2.1.2 Intergranular Starter Crack

A major source for variability in measured SCC growth rates is crack morphology. Typically, SCC 

of austenitic SSs in high-temperature water is IG, whereas the starter crack for SCC tests is produced by 

fatigue cycling the specimen in air, which creates a TG crack. Depending on the susceptibility of the 

material to SCC, the TG crack may not transition to an IG crack during the test or may transition only 
34

along a portion of the specimen width. This behavior results in an irregular or finger-like crack front. 

Even if precracking is carried out in simulated LWR environments, an IG crack is not guaranteed. Under 

the more rapid cyclic loading typically used for precracking, the crack growth is dominated by 

mechanical fatigue, which nearly always creates a TG crack. 

10-12

10-11

10-10

10-9

10-8

10-7

10-12 10-11 10-10 10-9 10-8 10-7

Irradiated to 0.75 dpa

Irradiated to 2.16 dpa

CGRenv (m/s)

CGRair (m/s)

Austenitic Stainless Steels

289°C, !500 ppb DO

Precracking

Cycling at 

low frequencies

& high R 

Figure 28.

Plot of CGR in water vs. CGR in air showing 

environmental enhancement of growth rates 

in high-purity water at 289°C (Refs. 10,11). 

The specimens are, therefore, subsequently subjected to gentle cycling (to reduce the contribution 

of mechanical fatigue compared to corrosion fatigue or SCC) and facilitate the transition of a TG crack to 

an IG crack. To achieve this, the loading frequency is gradually decreased, and the load ratio R is 

increased above 0.3. This behavior is illustrated in Fig. 28, where CGRs in the environment are plotted as 

a function of the estimated CGRs in air under the same loading conditions. In this plot, since the CGRs 

of austenitic SSs in air are independent of frequency, the data points that lie along the diagonal represent 

predominantly mechanical fatigue, and those that lie above the diagonal indicate environmentally 

enhanced crack growth. Under gentle cycling as the frequency is decreased and R is increased, the CGRs 

first decrease along the diagonal, as shown by the curve denoted “Precracking” in Fig. 28, then jump to 

new, higher environmentally enhanced growth rates. Once environmental enhancement is achieved, the 

crack morphology for austenitic SSs is IG, and the crack front in test specimens is relatively straight. 

Environmental enhancement of growth rates is readily achieved for materials that are susceptible to SCC, 

as seen for the material irradiated to 2.16 dpa, whereas materials less susceptible to SCC (e.g., SS 

irradiated to 0.75 dpa) take much longer and often require some coaxing, i.e., very gradual reductions in 

the cyclic loading. For austenitic SSs, at Kmax = 15–18 MPa m1/2, environmental enhancement typically 

occurs at R ≥ 0.5 and rise times ≥ 30 s.10 

An example of enhanced growth rates for irradiated Type 304L SS in the simulated NWC BWR 

environment at 289°C is shown in Fig. 29.

9 The change in crack length is plotted as a function of time. 

For this specimen, environmental enhancement occurred after ≈170 h, when the load ratio and rise time, 

respectively, were changed from 0.5 and 60 s to 0.7 and 300 s. For the new loading condition, although 

the predicted CGR in air decreased by a factor of ≈15, the measured rate in the environment increased by a 

factor of ≈3. 

2.2.1.3 K/Size Criteria

Fracture mechanics is a correlative technology and does not attempt to describe the mechanisms 

that are occurring at the crack tip. It correlates the behavior of components with that of specimens 
35

through the use of the K parameter. If two cracks have the same K, then they have the same strains and 

stresses in a region near the crack tip. For this correlation between specimen and component to work, K 

has to characterize the stresses and strains at the crack tip in the process zone. Mathematically, it can be 

shown that this observation is true if the plastic zone size is “small enough.” The K/size criteria are 

combined theoretical and empirical results that have been found to ensure the plastic zone is small enough 

and K is controlling. The ASTM specifications for specimen K/size criteria are intended to ensure the 

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. For constant load tests, ASTM E 1681 requires that 

Beff and (W - a) ≥2.5 (K/σy)2, (7)

and for cyclic loading, ASTM 647 requires that 

Beff and (W - a) ≥(4/π) (K/σy)2, (8)

where K is the applied stress intensity factor, σy is the yield stress of the material, a is crack length, and 

the Beff is the specimen effective thickness, defined as (B BN)0.5. Thus, the specimen thickness or 

remaining ligament ahead of the advancing crack is at least a factor of 8 greater than the plastic zone size 

for tests conducted in accordance with the K/size criterion of Eq. 7. For high strain-hardening materials 

[i.e., ratio of ultimate stress to yield stress (σu/σy) ≥1.3], both criteria allow the use of the flow stress 

defined as σf = (σu + σy)/2 rather than the yield stress. 

At the present time, the database for defining the K/size criteria for irradiated materials is 

inadequate. The K/size criteria were developed for materials that show work hardening and, therefore, 

may not be applicable for materials irradiated to fluence levels where, on a local level, they do not strain 

harden. This lack of strain hardening, termed “strain softening,” is most dramatic when dislocation 

channeling occurs at high fluences. For moderate to highly irradiated material, Andresen15 has suggested 

an effective yield stress, defined as the average of the nonirradiated and irradiated yield stresses [σeff =

(σyirr + σynonirr)/2]; this discounts the irradiation-induced increase in yield stress by a factor of 2. 

Jenssen et al.78 obtained crack growth data in simulated BWR environments on Type 304L SS 

irradiated to ≈13 dpa and investigated the specimen K/size criterion for CGR testing of irradiated 

6.6

6.8

7.0

7.2

7.4

15

20

25

30

35

100 150 200 250 300 350 400

Crack Length (mm)

Kmax (MPa m0.5)

Time (h)

Type 304 SS (Heat C3)

Test CGRI–07 (Spec. C3-B)

Fluence 0.9 x 1021 n/cm2

289°C

High–Purity Water

CGR = 1.75 x 10–10 m/s

Kmax = 20.1 MPa m0.5

R = 0.5, Rise Time 60 s

Kmax Crack Length DO !250 ppb

Steel ECP 190 mV (SHE)

DO <30 ppb

1.06 x 10–9 m/s

21.4 MPa m0.5

R = 1.0

1.04 x 10–9 m/s

23.5 MPa m0.5

R = 1.0

CGR = 6.38 x 10–10 m/s

Kmax = 21.0 MPa m0.5

R = 0.7, Rise Time 300 s

Constant Load

Unload to R = 0.7 every 2 h

Figure 29. Plot of crack length vs. time for irradiated Type 304 stainless steel in 

high-purity water at 289°C (Ref. 9).
36

austenitic SSs. They concluded the following from a finite element study: if the strain softening found in 

highly irradiated materials is taken into account, and if the K/size criterion is defined as σeff = (σyirr + 

σynonirr)/2, there is a significant amount of plastic deformation in the plane of the growing crack. The 

authors argue that as a result of an increased tendency for “highly irradiated material” to deform by 

dislocation channeling, a K/size criterion based on the sum of irradiated and nonirradiated yield stress 

divided by 3 [i.e., σeff = (σyirr + σynonirr)/3] fits the crack growth behavior better.78 

Thus, in these proposed K/size criteria, the irradiation-induced increase in yield stress is discounted 

by a factor of 2 for moderately irradiated materials and by a factor of 3 for highly irradiated materials.

These proposed criteria were not developed based on analyses of crack tip plastic strain or 

experimental comparisons of specimens with different degrees of constraint. Rather they seem to be 

based on an implicit assumption that if the benefit of reduced corrosion potential on CGRs is effective at 

one K level and is not effective at a higher K level, then the difference must be due to a violation of the 

K/size criterion. The argument being that since austenitic SSs irradiated above approximately 3 dpa 

exhibit strain softening, the K/size criterion based on the measured yield strength of the irradiated 

material is inadequate for such materials. It is not clear whether due consideration has been given to the 

possibility of other effects. As discussed later in this section, strain softening in highly irradiated 

austenitic SSs is rarely more than 10-15%.

Nakamura et al.80 investigated the validity of a K/size criterion based on the effective yield stress 

instead of the measured yield stress of the material. They conducted crack growth tests with 1⁄2-T CT 

specimens of irradiated Type 304L SS in simulated BWR environments. The material was from a BWR 

control rod irradiated to 2.5 x 1021 n/cm2 (3.5 dpa), and the yield stress of the material at 288°C was 

600 MPa. The CGRs were obtained in BWR environments at K values of 23-35 MPa m1/2. The authors 

concluded that their results did not support a K/size criterion based on an effective yield stress rather than 

the measured yield stress. The crack growth behavior was relatively stable even after K exceeded the 

upper limit based on a small-scale yielding condition with effective yield stress for the irradiated 

material.80 The plastic zone, evaluated on the basis of the microstructure and hardness distribution 

around the crack tip, was estimated to be 0.2-0.4 mm at K = 30 MPa m1/2. For the same loading 

condition, the plastic zone size was estimated to be ≈0.8 mm for a material with 600-MPa yield stress, 

and ≈2.0 mm if the effective yield stress proposed by Andresen is considered.

Another study,81 evaluated the validity of the K/size criterion for irradiated materials by comparing 

the plastic strain distribution in a 1⁄2-T CT specimen estimated from finite element method calculations 

with experimentally observed plastic deformation area. The results indicate that for a Type 316L SS 

material irradiated to 3 x 1021 n/cm2 in the Japan Materials Test Reactor, the appropriate K for SCC 

growth rate tests is at least 30 MPa m1/2 for a 5.8-mm-thick 1⁄2-T CT specimen. 

Pettersson* has presented three arguments against the proposed reduced effective stresses. First, he 

suggests that the strain softening in irradiated austenitic SSs is rarely more than 10-15%. This behavior is 

clearly demonstrated in the engineering stress vs. strain curves shown in Fig. 30 for Type 304 SS 

irradiated to 3.0 dpa in the Halden reactor at 288°C and tested in air at 289 and 325°C.53 Secondly, in 

most of the plastic zone the plastic strains are so low that the material never passes the maximum tensile 

stress, so that it is effectively not a strain-softening material. Third, finite element analyses indicate that 

the difference is marginal between the strain distributions ahead of an advancing crack in a strain-

 

*Kjell Pettersson, Matsafe AB, private communication, “Some Aspects of Specimen Size Validity and Crack Tip Strain Rate,” 

Nov. 2006. 
37

hardening versus a strain-softening material (Fig. 31). These calculations do not support the suggestion 

that the nonirradiated yield strength should be involved in any calculations of specimen sizes for 

obtaining valid data.

0

200

400

600

800

1000

0 1 2 3 4 5

325°C air

289°C air

Stress (MPa)

Strain (%)

Type 304 SS (Heat C19)

Irradiated to 3.0 dpa

Strain rate 1.65 x 10-7 s-1

Rapid stress 

reduction is 

due to necking Figure 30.

Engineering stress vs. strain curve for 

Type 304 stainless steel irradiated to 

3.0 dpa and tested in air at 289 and 325°C 

(Ref. 53).

0.000

0.005

0.010

0.015

0.020

0 20 40 60 80 100 120

Strain softening

Strain

Strain hardening

Distance, μm

Figure 31.

Strain distribution of a moving crack in a 

strain-softening and a strain-hardening 

material (Ref. 11).

The adequacy of the proposed K/size criterion has also been evaluated by examining the effect of 

loss of specimen constraint on fracture morphology and crack growth behavior.11 The experimental data 

for CGRs obtained under loading conditions that either meet or exceed the K/size criterion proposed by 

Andresen are shown in Fig. 32; the numbers next to the data points in Fig. 32b represent the value 

(in percentages) by which applied K exceeded the allowed value. 

In Fig. 32b, although the applied K exceeds the value allowed by the K/size criterion proposed by 

Andresen by up to ≈40%, the CGRs measured from these tests are consistent with the results from tests 

shown in Fig. 32a that meet the criterion. Also, the K dependence for these tests is consistent with that 

observed for valid tests (e.g., the data yield an exponent of ≈2.1). Furthermore, in high-DO water, the 

CGR did not increase during the test period (for up to 200 h). Typically, the CGR increases rapidly when 

the applied load exceeds the specimen K/size criterion. For example, for a 1/2-T CT specimen of Type 

304 SS irradiated to 4.0 dpa, the CGR increased by a factor 5 over a period of 40 h in high-DO water at 

288°C and Kmax of 29-34 MPa m1/2.

15 

The two data points obtained in a low-DO environment on Type 304L and 316 SS irradiated to 

3.0 dpa and tested at 25-30 MPa m1/2 (solid triangle and solid right-angle triangle in Fig. 32b) did not 

show the expected decrease in CGR when the DO level in the environment was decreased. The change in 

crack length and Kmax with time for Type 304L SS irradiated to 2 x 1021 n/cm2 is shown in Fig. 33. The 
38

results indicate no change in CGR for more than 300 hours when the DO level was decreased from 

400 ppb to about 20 ppb. A similar behavior was observed for Type 316 SS. It can be argued that 

because the expected decrease in growth rate is not observed for these tests, the loading conditions must 

have exceeded the specimen K/size criterion, but it is not clear whether this behavior should be attributed 

to loss of specimen constraint. If specimen constraint had been lost, the growth rate should have rapidly 

increased in high-DO water.

7.20

7.40

7.60

7.80

8.00

8.20

20

25

30

35

40

45

200 240 280 320 360 400 440 480 520

Crack Length (mm)

Kmax (MPa m0.5)

Time (h)

Type 304 SS (Heat C3)

Test CGRI–08 (Spec. C3-C)

Fluence 2.0 x 1021 n/cm2

289°C

Constant Load, periodic unloading Kmax to R = 0.7 every 1 h

Crack Length

CGR = 6.91 x 10–10 m/s

Kmax = 27.5 MPa m0.5

Steel ECP -294 mV (SHE)

5.07 x 10–10 m/s

23.7 MPa m0.5

Steel ECP 165 – 10 mV (SHE)

!400 ppb DO !80 ppb DO !20 ppb DO

Figure 33. Crack length and K vs. time for Type 304L SS in high-purity water at 289°C (Ref. 11).

The specimen constraints based on Eq. 7 were lost for the two tests on irradiated Type 304L SS in 

low-DO water at K levels above 35 MPa m1/2 (solid triangles in Fig. 32b), as evidenced by the unusually 

high growth rates. This behavior was verified by fractographic examination of the specimen; under these 

loading conditions, the crack propagation was away from the normal plane. 

10-12

10-11

10-10

10-9

10-8

10-7

5 10 15 20 25 30 35 40

304L 0.3x10^21

304L 0.9x10^21

304L 2.0x10^21

316L 2.0x10^21

316 0.3x10^21

316 2.0x10^21

Experimental CGR (m/s)

Stress Intensity K (MPa m1/2)

BWR Environment 

289°C

NUREG–0313

Curve

6 x NUREG-0313

Curve

Steel & Fluence n/cm2

Open Symbols: high DO

Closed Symbols: low DO 10-12

10-11

10-10

10-9

10-8

10-7

5 10 15 20 25 30 35 40

304L 0.3x10^21

304L 0.9x10^21

304L 2.0x10^21

316L 2.0x10^21

316 0.3x10^21

316 2.0x10^21

Experimental CGR (m/s)

Stress Intensity K (MPa m1/2)

BWR Environment 

289°C

NUREG–0313

Curve

6 x NUREG-0313

Curve

Steel & Fluence n/cm2

5

27 16

49

2

44

30

32

9 4

31 29

4

Open Symbols: high DO

Closed Symbols: low DO

14 111

134

(a) (b)

Figure 32. Crack growth rates for irradiated stainless steels in BWR environments under loading 

conditions that (a) meet or (b) exceed the proposed K/size criterion. The numbers next to the 

data points represent the percentage difference between the applied and allowed K (Ref. 11).
39

A loss of specimen constraint should influence the fracture mode and morphology. For example, if 

the thickness criterion is exceeded, the crack plane is typically out-of-normal near the edges of the 

specimen, and if the specimen ligament criterion is exceeded, the crack propagates away from the normal 

plane at an angle of 45°. The fracture surface of Specimen C3-C was examined to investigate any change 

in fracture morphology and/or change in the fracture plane. Figure 34 shows a side view of a 1-mm-wide 

slice of the fracture surface (along the entire crack advance) cut from Specimen C3-C. The fracture 

surface is toward the top, in a plane perpendicular to the picture. (Although precautions were taken to 

ensure that the specimen was square to the movement of the electrical discharge machine wire, the cuts 

were not always straight; the bottom surface of this slice had an uneven cut.) The profile of the fracture 

surface indicates that the fracture plane is relatively straight and normal to the stress axis for the initial 

≈3.5-mm crack extension. The DO level was decreased from ≈400 to 20 ppb at ≈1.7-mm crack 

extension, which is equivalent to the middle of the relatively straight crack extension. The fracture plane 

is out-of-normal for crack extensions greater than 3.5 mm. A secondary crack propagated at an angle of 

≈45° to the original fracture plane. This region corresponds to the crack advance for the two tests for 

Type 304L SS at K levels of about 35 MPa m1/2, which are represented as solid triangles in Fig. 32b; the 

CGRs during these periods were unusually high. 

In addition, the fracture morphology did not change when the DO level was decreased from 400 to 

20 ppb. After the initial TG fracture during fatigue precracking, the fracture morphology during all other 

test periods was completely IG. These results do not show any fractographic indication of a loss in 

constraint in the specimen (i.e., the fracture morphology did not change, and the fracture plane was 

straight and normal to the stress axis). Furthermore, the growth rate was constant for the entire duration 

when the DO level was decreased. 

The results in Fig 32 indicate that some SSs may show IASCC susceptibility in low-potential 

environments at fluence levels as low as 3 dpa. Caution should be exercised to not propose a K/size 

criterion for irradiated SSs that screen out such experimental data simply because they do not agree with 

the current understanding of the IASCC behavior of austenitic SSs. The specimen size criteria for 

nonirradiated materials could be validated experimentally by using different size specimens. Such an 

approach is extremely difficult for irradiated materials, but more rigorous finite element studies of crack 

tip behavior may be helpful in developing appropriate validity criteria. In any case, the potentially 

conflicting results currently available indicate that additional data on SSs irradiated to 3-8 dpa are needed 

to accurately establish the threshold for IASCC susceptibility in low-potential PWR environments.

2.2.1.4 Rising K or dK/da Effects

Another factor that can influence measurements of CGRs, under conditions that are relevant for 

LWRs, is a variation in K as the crack advances. In LWR components, K can increase as the crack depth, 

a, increases with crack advance, and can increase and then decrease due to the through-wall weld residual 

stress profile, which is usually U-shaped. Consequently, as the crack advances through the wall, K can 

Figure 34. Side view of the first slice cut from Type 304L SS Specimen C3-C (Ref. 11).
40

increase, remain relatively constant, or decrease depending on the residual stress profile. Initially, when 

the crack depth is small, K could increase very rapidly because it is proportional to the square root of 

crack depth (K α σ√a). Studies on austenitic SSs and Ni alloys in high-purity water with 2000 ppb O2 at 

288°C have shown that rising K (+dK/da) at K values that are relevant for LWR components can increase 

the CGRs significantly, whereas decreasing K (-dK/da) has little or no effect.88 Under rising K 

conditions, growth rates can be more than two orders of magnitude higher than those under constant K 

conditions. 

As discussed by Andresen and Morra,88 dynamic strain at the crack tip is an important element of 

crack advance under constant K loading conditions. As the crack advances, the process of stress and 

strain redistribution sustains crack growth. It leads to slip offset, which ruptures the surface oxide film, 

and the crack advances due to oxidation/dissolution of the metal. Thus, there is an inherent synergy 

between crack advance and growth rate itself (or dK/da as the crack advances). The data presented by 

Andresen and Morra88 indicate that when changes in K are controlled by dK/da, very different responses 

can be seen for increasing versus decreasing K conditions with crack propagation. For example, 

increasing K results in an increased growth rate, which causes faster increase in K, which is a positive 

feedback. Decreasing K, on the other hand, yields negative feedback. These results indicate that CGR 

data obtained from tests where K was allowed to increase during the test or was decreased too fast may 

give erroneous results for the K dependence of SCC growth rates. Caution should be exercised while 

interpreting such results.

2.2.1.5 Reloading Effects

For some irradiated austenitic SS materials, small increases in the stress intensity factor K due to 

either load perturbations or the reloading portion of the partial unload/reload cycle during an SCC growth 

rate test result in a rapid or step-like crack extension.60 The growth rates during such periods are as high 

as 10-7 m/s. Such reloading effects produce a classic staircase appearance in the crack length vs. time 

plots. Similar behavior has been observed in nonirradiated CW SSs with high yield strength.39 This 

behavior has often been attributed to the breaking of uncracked ligaments in the material, and is believed 

to occur in materials with an uneven crack front. The significance and cause for this behavior should be 

investigated.

2.2.2 Stress Corrosion Cracking Growth Rates

Most of the CGR data on irradiated materials have been obtained under BWR operating conditions 

(i.e., relatively low neutron dose in NWC or HWC BWR environments). Consequently, the data in BWR 

environments are discussed first.

2.2.2.1 BWR Environments

Dependence of Stress Intensity Factor K

The majority of the SCC growth rate tests in BWR environments have been conducted at constant 

K with or without periodic partial unloading to R = 0.5-0.7. A few tests were performed using a 

trapezoidal waveform with rise and return times up to 5000 s. For these tests, the experimental data can 

be adjusted to account for the contribution of corrosion fatigue by using the cyclic CGR data; the details 

are discussed in Section 2.2.3. The SCC growth rates reported in the literature for various grades and 

heats of austenitic SSs irradiated from 5 x 1020 n/cm2 to 2.5 x 1022 n/cm2 (0.75 to 37.5 dpa) in NWC 

BWR and HWC BWR environments are shown in Figs. 35a-c and Figs. 36a-c, respectively. Crack 
41

growth rates in submerged arc (SA) weld and shielded metal arc (SMA) weld HAZ materials as well as 

sensitized SSs are included in the figures. Also included are the K vs. CGR disposition curves proposed 

in NUREG-0313 for nonirradiated sensitized austenitic SSs in high-purity water containing 8 ppm DO 

(Eq. 5) and the curve proposed by EPRI for austenitic SS BWR core internal components.89 The EPRI 

disposition curve for use in BWR core environments is expressed as 

da/dt (m/s) = A2 (K)2.5, (9)

where K is in MPa m1/2, and the constant A2 is 4.564 x 10-13 in NWC and 1.512 x 10-13 in HWC BWR 

environments. The EPRI correlations are based on two datasets: General Electric (GE)/Japan Power 

Engineering and Inspection Corp. (JAPEIC) data for SSs irradiated to 2.7 x 1021 n/cm2 (4.0-4.5 dpa),89

10-12

10-11

10-10

10-9

10-8

10-7

0 5 10 15 20 25 30 35 40

304 Sensi 0.75 dpa

304 Sensi 2.16 dpa

316 3.0 dpa

316 0.45 dpa

316 1.35 dpa

316L 3.0 dpa

304L 3.0 dpa

304L 0.45 dpa

304L 1.35 dpa

SAW HAZ 0.75 dpa

SAW HAZ 2.16 dpa

SMAW HAZ 0.75 dpa

SMAW HAZ 2.16 dpa

SMAW HAZ +TT 0.75 dpa

Experimental CGR (m/s)

Stress Intensity K (MPa m1/2)

Material & Dose

(ANL data)

NUREG-0313

Curve

Irradiated Stainless Steels 

BWR Environment

EPRI

NWC BWR

(a)

10 x NUREG-0313

Curve

10-12

10-11

10-10

10-9

10-8

10-7

0 5 10 15 20 25 30 35 40

316NG 1.5-2.3 dpa

347 SS 2.3-3.3 dpa Spec. CT-2

347 SS 2.3-3.3 dpa Spec. CT-3

347 SS 0.7 dpa

347 SS 0.95 dpa

304 SS 4.25 dpa

Experimental CGR (m/s)

Stress Intensity K (MPa m1/2)

Material & Dose

(Halden data)

NUREG-0313

Curve

Irradiated Stainless Steels 

BWR Environment

EPRI

NWC BWR

(b)

10 x NUREG-0313

Curve

(GE/JAPEIC data)

Figure 35. Experimental SCC growth rates in NWC BWR environment of austenitic stainless steels 

irradiated to (a) 0.75-3.0 dpa (Refs. 9,11), (b) 0.7-4.25 dpa (Refs. 74,75,77,89) and 

(c) 5.5-37.5 dpa (Refs. 74,75,77,79,78).
42

10-12

10-11

10-10

10-9

10-8

10-7

0 5 10 15 20 25 30 35 40

304 SS 18.0 dpa 277°C

304 SS 37.5 dpa 277°C

304L 12.3 dpa 277°C

304L 13.6-14.4 dpa 280°C

304L 12.9 dpa 288°C

304L 5.5 dpa Fast 288°C

304L 10.2 dpa Fast 288°C

304L 11.2 dpa Fast 288°C

Experimental CGR (m/s)

Stress Intensity K (MPa m1/2)

Material, Dose, Test Temp.

(Halden data)

NUREG-0313

Curve

Irradiated Stainless Steels 

BWR Environment

20 x NUREG-0313

Curve

Materilas Irradiated in

Open symbols: BWR 

Closed symbols: BOR-60

(c)

EPRI

NWC BWR

(Studsvik data)

Figure 35. (Contd.)

and the Halden IFA data (639 test series) for Type 304, 347, and 316NG SSs that were irradiated to 9.0, 

1.5, and 0.9 x 1021 n/cm2 (13.5, 2.25, and 1.35 dpa).74 The significant results in NWC and HWC BWR 

environments are summarized as follows.

1. In the NWC BWR environment, the SCC growth rates of austenitic SSs irradiated to 5 dpa could be a 

factor of 10 higher than the NUREG-0313 disposition curve for nonirradiated SSs in 8 ppm water, 

and those irradiated to 10-12 dpa could be a factor of 25 higher than the NUREG-0313 curve. 

Furthermore, as seen in Figs. 35a and b, for the end-of-60-year-life fluence level (3 x 1021 n/cm2 or 

4.5 dpa), the CGRs of some SSs materials could be significantly higher than the CGR disposition 

curve proposed by EPRI. In Fig. 35a, the data for SA weld HAZ materials irradiated to 0.75 or 

2.16 dpa11 are a factor of 2 above the EPRI curve. The results also indicate that the SCC growth rates 

for SSs irradiated in BWRs to 18 dpa or higher are below the NUREG-0313 disposition curve for 

nonirradiated SSs. The reasons for the low CGRs obtained for these highly irradiated BWR materials 

in the high-potential BWR environment are not clear.

2. A CGR disposition curve a factor of 10 higher than the NUREG-0313 curve bounds nearly 80% of 

the data obtained in the NWC BWR environment for SSs irradiated up to 4.25 dpa (Fig. 35b). The 

EPRI curve is about a factor of 6 higher than the NUREG-0313 curve, although the K dependence is 

higher (i.e., the exponent is 2.5 for the EPRI curve and 2.161 for the NUREG-0313 curve). 

3. In general, CGR data obtained using test procedures in which the specimen was precracked in air and 

then transferred to the autoclave for the SCC test show much more variability at K values of 8-

15 MPa m1/2 (Halden data in Figs 35b and c) than data obtained using test procedures in which the 

specimen was precracked in the NWC environment and the TG fatigue crack was transitioned to an 

IG crack using slow fatigue cycling in the same environment before starting the SCC test [Argonne 

National Laboratory (ANL) data in Fig. 35a and Studsvik data in Figs. 35b and c]. This difference 

does not appear to be due to differences in material or irradiation conditions because data obtained at 

Halden and Studsvik on the same heat of Type 304L SS BWR control-rod blade material irradiated to 

comparable neutron dose [e.g., the two horizontal triangles ( and ) in Fig. 35c for Type 304L 
43

irradiated to 12.3 and 12.9 dpa, respectively] show more than a factor 10 higher growth rates for the 

material tested at Studsvik than for the material tested at Halden.

The difference is also not related to failure mode differences, since post-test fractography of the 

Halden specimens showed that the fracture morphology was always IG as in the ANL and Studsvik 

testing. Finally, for the loading sequence used in the tests at ANL and Studsvik, the CGRs at K 

values of 3-11 MPa m1/2 were obtained by decreasing K from 14-20 MPa m1/2 to the low values. 

Thus, the cracks were growing at a rate ≥ 7 x 10-10 m/s before K was decreased, whereas, in the 

Halden data, tests were started at low K and the CGRs were <1 x 10-10 m/s as K was allowed to 

increase during the test. As discussed in Sections 2.1.6 and 2.1.7, the relatively weak K dependence 

10-12

10-11

10-10

10-9

10-8

10-7

0 5 10 15 20 25 30 35 40

316 3.0 dpa

316 1.35 dpa

316L 3.0 dpa

304L 3.0 dpa

304L 1.35 dpa

SAW HAZ 0.75 dpa

SMAW HAZ 0.75 dpa

SMAW HAZ +TT 0.75 dpa

Experimental CGR (m/s)

Stress Intensity K (MPa m1/2)

Material & Dose

(ANL data)

NUREG-0313 Curve

Irradiated Stainless Steels 

BWR Environment

EPRI

HWC BWR

(a)

10-12

10-11

10-10

10-9

10-8

10-7

0 5 10 15 20 25 30 35 40

316NG 1.5-2.3 dpa

347 SS 3.3 dpa Spec. CT-2

347 SS 3.3 dpa Spec. CT-3

347 SS 0.7 dpa

347 SS 0.95 dpa

304 SS 4.25 dpa

Experimental CGR (m/s)

Stress Intensity K (MPa m1/2)

Material & Dose

(Halden data)

NUREG-0313

Curve

Irradiated Stainless Steels 

BWR Environment

EPRI HWC BWR

(b)

(GE/JAPEIC data)

Figure 36. SCC growth rates in HWC BWR environment of austenitic stainless steels 

irradiated to (a) 0.75-3.0 dpa (Refs. 9,11), (b) 0.7-4.25 dpa (Refs. 74,75,77,89), 

and (c) 5.5-37.5 dpa (Refs. 74,75,77,79,78).
44

of CGRs may be associated with a threshold growth rate (e.g., about 1 x 10-9 m/s) that influences the 

crack tip strain rate. Therefore, differences between the CGR appear to be due to the differences in 

the loading sequence and not to differences in materials, failure mode, or irradiation fluence. 

4. For the tests where the fatigue TG crack was transitioned to an IG crack by slow fatigue cycling in the 

environment, periodic unloading during the SCC test at constant K had little or no effect on growth 

rates; the CGRs obtained with periodic unloading are comparable to those obtained without 

unloading. However, periodic unloading seems to have a significant effect on growth rates for some 

tests that did not use slow fatigue cycling to transition the fatigue TG crack to an IG crack. In these 

tests, for low dose materials (e.g., less than 5 dpa), the CGRs obtained with periodic unloading are 

comparable to those obtained without periodic unloading in the NWC BWR environment but are 

typically higher than those obtained without periodic unloading in the HWC BWR environment. This 

difference is associated with uncracked ligaments in the material. Also, for the Type 304 SS 

irradiated to 18.0 or 37.5 dpa (Fig. 35c), stable CGRs could be established only with periodic 

unloading and not under constant K. Both materials came from a decommissioned PWR. It is 

possible that this behavior is unique to the material microstructure and microchemistry.

5. The K dependence of growth rates shows some variability. The exponent of K, in Eq. 5, was 

typically between 1.7 and 2.7 for SSs irradiated up to about 5 dpa. For the Halden data, the exponent 

was about 2 for the IFA-639 test series and about 2.5 for the IFA-658 test series. Exponents between 

1.5 and 2.0 were observed for the ANL data sets. For SSs irradiated to higher dose levels, depending 

on the test procedure, the exponent could be as high as 7. As discussed above, the higher K 

dependence may be due to the effect of the loading sequence. However, as discussed above, for the 

SSs irradiated to 18.0 or 37.5 dpa, stable CGRs were observed only with periodic unloading, and the 

rates were relatively insensitive to changes in K.

6. The cracking behavior of various grades and heats of SSs irradiated to the same dose level seems to 

be the same. For example, in the Halden study, the CGRs for Type 316NG and 347 SSs irradiated to 

2.0-3.5 dpa are similar.74 Similarly, in the ANL study, the CGRs of Type 304L, 316L, or 316 SSs 

irradiated to 1.35-3.0 dpa are comparable.11 However, the rates for weld HAZ materials are slightly 

10-12

10-11

10-10

10-9

10-8

10-7

0 5 10 15 20 25 30 35 40

304 SS 18.0 dpa 277°C

304 SS 37.5 dpa 277°C

304L 13.6-14.4 dpa 280°C

304L 12.9 dpa 288°C

304L 5.5 dpa Fast 288°C

304L 10.2 dpa Fast 288°C

304L 11.2 dpa Fast 288°C

Experimental CGR (m/s)

Stress Intensity K (MPa m1/2)

Material, Dose, Test Temp.

(Halden data)

NUREG-0313

Curve

Irradiated Stainless Steels 

BWR Environment

Materilas Irradiated in

Open symbols: BWR 

Closed symbols: BOR-60

(c)

EPRI

HWC BWR

(Studsvik data)

Figure 36. (Contd.)
45

higher than those for solution-annealed materials (Fig. 35a). The higher rates most likely are due to 

weld residual strains, which could be as high as 30% of room-temperature tensile strain.40 Typically, 

residual strains are the highest at the root of the weld and near the weld fusion line. The significance 

of cold work in increasing SCC susceptibility is well known, and has been discussed in Section 2.1.5. 

Figure 37 gives the SCC growth rates for two irradiated specimens of CF-8M cast SS in the NWC 

BWR environment. The specimens were aged for 10,000 h at 400°C before irradiation. The large 

difference between the measured CGRs for the two specimens is because sustained crack growth was 

not achieved in Specimen 75-11TT under constant K loading (i.e., without periodic unloading). The 

results for the other specimen are comparable to the data obtained on solution-annealed SSs and weld 

HAZ materials irradiated to similar dose levels. The CGRs are a factor of 2-6 above the values 

predicted by the NUREG-0313 curve for nonirradiated austenitic SSs.66

10-12

10-11

10-10

10-9

10-8

10-7

0 5 10 15 20 25 30 35 40

75-11TT

75-11TM

Experimental CGR (m/s)

Stress Intensity K (MPa m1/2)

Specimen Number

NUREG-0313

Curve

6 x NUREG-0313

Curve

CF-8M Cast Stainless Steel

Aged 10,000 h at 400°C

Irradiated 1.63 x 1021 n/cm2 (2.46 dpa)

BWR Environment

289°C EPRI

NWC BWR

Figure 37.

CGR under constant load for thermally aged and 

irradiated CF-8M cast stainless steel specimens 

in BWR environment at 289°C (Ref. 11).

Effect of Corrosion Potential

The data in Fig. 36 for SCC growth rates in the HWC BWR environment show a significant 

decrease in growth rates relative to those in the NWC BWR environment for SSs irradiated to less than 

3 dpa; little or no reduction for some heats of SSs irradiated to 3-4 dpa, at least at K values above 

18 MPa m1/2; and no decrease for all SSs irradiated 12-14.5 dpa. The materials irradiated to 18 or 

37.5 dpa also did not show any benefit of decreased corrosion potential on CGRs, but a sustained crack 

advance was not observed in these materials under constant K. In these materials, growth rates were 

established with periodic unloading, and the rates even in the NWC BWR environment were below the 

NUREG-0313 curve for nonirradiated SSs. A significant feature of the cracking behavior of high dose 

materials was the “staircase” effect or reloading effect discussed in Section 2.2.1.5. In the HWC BWR 

environment, each unloading/reloading cycle was accompanied by a step or “jump” in crack length, 

creating a staircase or stepped crack growth. Although this behavior is often attributed to breaking of 

uncracked ligaments or straightening of the crack front, it is not clear why it should occur only in the 

HWC environment. Additional data in the HWC BWR or PWR environments are needed on high-dose 

materials to better understand the IASCC behavior of austenitic SSs.

The data for the 3-4 dpa materials (e.g., circles in Fig. 36a and diamonds and horizontal triangles in 

Fig. 36b) are often considered invalid because of the high K values that exceeded the proposed K/size 
46

criterion for irradiated materials (i.e., the irradiation-induced increase in yield stress is discounted by a 

factor of 2). As discussed in some detail in Section 2.2.1.3, the arguments against the validity of the data 

do not seem to have a well-developed technical basis. The results in Figs. 32 and 36 indicate that the 

benefit of low-potential environments on SCC growth rates may be lost for some heats of SSs irradiated 

to fluence levels as low as 3 dpa. Some internal consistency checks can be used to demonstrate the 

validity of the Halden data for Type 347 SS irradiated to 2.3-3.3 dpa (diamonds in Figs. 35b and 36b). 

The growth rates observed (not corrected) during the test time from 720 to 840 h for two identical 

specimens (CT-2 and CT-3) of this material tested in the NWC BWR environment were 7.32 x 10-10 m/s 

at 28.7 MPa m1/2 and 8.75 x 10-10 m/s at 23.6 MPa m1/2.

74 The environment was then changed to HWC 

and the rates for the two specimens during test time 840 to 950 h were 7.77 x 10-10 m/s at 29.3 MPa m1/2

and 8.9 x 10-10 m/s at 24.4 MPa m1/2. The growth rates for both specimens did not decrease after the 

corrosion potential was decreased. Also, note that the applied K for specimen CT-2 in NWC was higher 

than that for specimen CT-3 in HWC. If the measured CGR for specimen CT-3 is invalid in HWC, it 

should also be invalid for specimen CT-2 in NWC. The experimental data or the post-test fractography of 

these samples do not show any indication that the specimen constraints were exceeded for these 

specimens. 

Some investigators seem to dismiss such results simply because they do not agree with the current 

understanding of the IASCC behavior of austenitic SSs. The feature that is common between specimens 

CT-2 and CT-3 is that although the K values are different, the growth rates for the specimens are 

comparable and may represent a threshold rate above which CGRs are relatively insensitive to changes in 

stress intensity factor or corrosion potential. As discussed previously, additional data on SSs irradiated to 

3-8 dpa are needed to accurately establish the threshold for IASCC susceptibility in low-potential PWR 

environments.

IASCC Growth Rate Disposition Curve for BWR Core Internals

The IASCC growth rate vs. K disposition curves proposed by EPRI for BWR core internal 

components are based primarily on two datasets: GE/JAPEIC and Halden data for austenitic SSs 

irradiated up to 13 dpa. The database does not include results obtained at ANL on weld HAZ materials or 

at Studsvik on BWR-irradiated Type 304L SS, which show SCC growth rates that are a factor of 2-3 

higher than the EPRI disposition curve. Therefore, the EPRI disposition curves in NWC and HWC 

environments should be revaluated in light of recent data. The curve is clearly not bounding, and an 

updated estimate of the fraction of the population of CGR data that are encompassed by the curve is 

needed. Additional data should be obtained on austenitic SSs, including Type 304L SS and weld HAZ 

materials, in HWC to better define the threshold fluence for IASCC susceptibility in low-potential 

environments.

Effect of Neutron Fluence

For the various grades and heats of austenitic SSs shown in Figs. 35 and 36, the SCC growth rates 

in the NWC BWR and HWC BWR environments at 288°C and K = 20 MPa m1/2 are plotted as a function 

of neutron dose in Figs. 38a and b, respectively. The materials were irradiated in BWRs at 280-300°C 

and in a fast reactor at 320°C. For each data set, the CGR at 20 MPa m1/2 was determined using Eq. 5; 

the actual value of the exponent was used when known. In the NWC BWR environment, the CGR 

increases above the NUREG-0313 level for doses >0.5 dpa, and it also increases with dose to about 10 

dpa, although there is variability in the data. In the HWC BWR environment, the CGRs do not exceed the 

NUREG-0313 levels until the fluence is ≥3.0 dpa. The CGRs for materials irradiated >5 dpa are 

comparable to those observed in the NWC BWR environment. As noted previously, there are some 
47

anomalously low data at 18 and 38 dpa. The curves shown in Fig. 38 bound approximately the 75th

percentile of the data in the NWC environment and greater for the HWC environment. The curves 

assume a threshold dose of 0.45 dpa and 2.7 dpa (or 3 x 1020 and 1.8 x 1021 n/cm2) in the NWC and 

HWC environment, respectively. The curves assume that the value for the constant A1 in the NUREG0313 disposition curve (Eq. 5) varies with neutron dose. In the NWC environment A1 is taken as 

A1 = 1.21 x 10-12 + 2.65 x 10-12 log(dpa), (10)

and in the HWC environment as 

A1 = 2.76 x 10-12 + 6.82 x 10-12 log(dpa). (11)

Note that because Eqs. 10 and 11 define the constant A1 in Eq. 5, these values must be multiplied by 

(20)2.161 to obtain the expression for the CGR curves shown in Fig. 38.

10-11

10-10

10-9

10-8

10-1 100 101

ANL

Halden

Studsvik

Studsvik (Fast)

Crack Growth Rate (m/s)

Neutron Dose (dpa)

Irradiated Austenitic SSs

Normalized to 20 MPa m 1/2 ̃

NWC BWR Water at 288°C

CGR = 7.84x10-10 + 1.72x10-09log(dpa)

NUREG-0313

(a)

10-11

10-10

10-9

10-8

10-1 100 101

ANL

Halden

Studsvik

Studsvik (Fast)

Crack Growth Rate (m/s)

Neutron Dose (dpa)

Irradiated Austenitic SSs

Normalized to 20 MPa m1/2

HWC BWR Water at 288°C

CGR = -1.79x10-9 + 4.42x10-09 log(dpa)

NUREG-0313

(b)

Figure 38. Crack growth rates of irradiated austenitic stainless steels in (a) NWC and (b) HWC BWR 

environments at 288°C and K of 20 MPa m1/2 plotted as a function of neutron dose.
48

Effect of Yield Strength

The effect of neutron dose on SCC growth rates may also be determined from the change in CGRs 

as a function of material yield strength (measured at temperatures of 280-300°C). The data shown in 

Figs. 38a and 38b are plotted as a function of material yield stress in Figs. 39a and 39b for CGRs in NWC 

and HWC environments, respectively. The results show an increase in CGRs with an increase in yield 

stress due to irradiation hardening. The increase in growth rates occurs at about 300 MPa in the NWC 

environment and at about 550 MPa in the HWC environment. However, the correlation between yield 

stress and CGR is not as good as that for cold-worked, nonirradiated austenitic SSs (in Fig. 16 or 17). For 

example, the yield stress of SSs irradiated to 0.45 or 0.75 dpa is 500-550 MPa, but the growth rates vary 

by a factor of 7-10 for these materials. Although data scatter and material variability contribute to the 

large differences in the CGRs, the results for irradiated materials suggest that in addition to irradiation 

hardening, other factors such as the RIS of Cr and Si at the grain boundaries are also important. 

10-11

10-10

10-9

10-8

0 200 400 600 800 1000

ANL

Halden

Studsvik

Studsvik (Fast)

Crack Growth Rate (m/s)

Yield Strength (MPa)

Irradiated Austenitic SSs

Normalized to 20 MPa m1/2

NWC BWR Water at 288°C

NUREG-0313

CGR = -1.0x10-9 + 4.3x10-8log(MPa)

(a)

10-11

10-10

10-9

10-8

0 200 400 600 800 1000

ANL

Halden

Studsvik

Studsvik (Fast)

Crack Growth Rate (m/s)

Yield Strength (MPa)

Irradiated Austenitic SSs

Normalized to 20 MPa m1/2

NWC BWR Water at 288°C

NUREG-0313

CGR = -2.7x10-8 + 1.0x10-8 log(MPa)

(b)

Figure 39. Crack growth rates of irradiated austenitic stainless steels in (a) NWC and (b) HWC BWR 

environments at 288°C and K of 20 MPa m1/2 plotted as a function of yield stress.
49

2.2.2.2 PWR Environments

Data on the SCC growth rate for irradiated austenitic SSs in the PWR environment have been 

obtained at Halden60,76 on BWR irradiated materials, at Studsvik79 on BOR-60 irradiated materials, and 

from EPRI-sponsored studies on BWR and PWR irradiated materials. The total irradiation dose for these 

materials was in the range of 3.0-37.5 dpa, with test temperatures of 288-340°C. The measured growth 

rates in various grades and heats of austenitic SSs in the PWR environment are shown in Fig. 40. There 

is a great deal of variability in the reported data. For a given material and irradiation condition, the 

variability is on the order of a factor of 10. The variability between materials and irradiation conditions is 

on the order of 1000.

10-12

10-11

10-10

10-9

10-8

10-7

0 5 10 15 20 25 30 35 40

304, 37.8 dpa, 335°C

304, 18.7 dpa, 335°C

304L, 14.0 dpa, 335°C

304L, 11.4 dpa, 305°C

304L, 11.4 dpa, 305°C

304L, 3.0 dpa, 305-320°C

304L, 3.0 dpa, 305-320°C

304, 6.3 dpa, 288-340°C

316CW, 25 dpa (fast), 320-340°C

316+Ti CW, 25 dpa (fast), 288-340°C

304L SA, 25 dpa (fast), 288-320°C

Experimental CGR (m/s)

Stress Intensity K (MPa m1/2)

Material, Dose, Test Temp.

(Halden data)

NUREG-0313

Curve

Irradiated Stainless Steels 

PWR Environment 305-340°C

EPRI

HWC BWR

(GE data)

(Studsvik data)

Figure 40. SCC growth rates of austenitic stainless steels irradiated to 3.0-37.8 dpa in PWR 

environment (Refs. 60,76,79,90).

Tests conducted on the same material at different temperatures indicate that the CGRs increase with 

increasing temperature. The temperature dependence of the SCC growth rates in irradiated and 

nonirradiated austenitic SSs yields activation energies between 60 and 150 kJ/mol. The experimental data 

for CGRs79 of irradiated CW Type 316+Ti SS in PWR are plotted as a function of inverse temperature in 

Fig. 41. The data yield an activation energy of 105 kJ/mol. An activation energy of 100 kJ/mol was used 

to normalize the data shown in Fig. 40 to a temperature of 320°C; the normalized CGRs are shown in 

Fig. 42 for materials irradiated to (a) 3 dpa, (b) 11-25 dpa, and (c) 6-38 dpa. Note that normalization does 

not significantly decrease the overall data scatter in Fig. 42 compared to Fig. 40. The normalization 

process changes only the magnitude of CGRs for a specific data set. 

The following observations were made from the normalized CGRs:

1. In general, in the PWR environment at 320°C, the SCC growth rates of materials irradiated up to 

12 dpa are above the EPRI disposition curve for the HWC BWR environment at 288°C. More than 

half the data points for Type 304L SS irradiated to 3 dpa are above the EPRI curve (Fig. 42a). The 

CGRs for Type 304L SS irradiated to 11.4 dpa and for fast-reactor-irradiated materials (except coldworked Type 316 SS) irradiated to 25 dpa are about two order of magnitude higher than the NUREG-
50

10-9

10-8

10-7

10-6

0.00160 0.00165 0.00170 0.00175 0.00180

Crack Growth Rate (m/s)

Inverse Temperature (K-1)

Q = 105 kJ/mole

Irradiated CW Type 316+Ti SS

PWR Environment

Growth rates at K = 20 MPa m1/2

340°C

320°C

288°C

Figure 41.

Temperature dependence of SCC growth 

rates in irradiated CW Type 316+Ti stainless 

steel in PWR environment (Ref. 79).

10-12

10-11

10-10

10-9

10-8

10-7

0 5 10 15 20 25 30 35 40

304L, 3.0 dpa

304L, 3.0 dpa

Experimental CGR (m/s)

Stress Intensity K (MPa m1/2)

Material & Dose

(Halden data)

NUREG-0313

Curve

Irradiated Stainless Steels 

PWR Environment 

Normalized to 320°C

EPRI

HWC BWR

(a)

10-12

10-11

10-10

10-9

10-8

10-7

0 5 10 15 20 25 30 35 40

304L, 11.4 dpa

304L, 11.4 dpa

316CW, 25 dpa (fast)

316+Ti CW, 25 dpa (fast)

304L SA, 25 dpa (fast)

Experimental CGR (m/s)

Stress Intensity K (MPa m1/2)

Material & Dose

(Halden data)

NUREG-0313 Curve

Irradiated Stainless Steels 

PWR Environment 

Normalized to 320°C

EPRI

HWC BWR

(b)

(Studsvik data)

Figure 42. Normalized SCC growth rates in PWR environment for austenitic stainless steels irradiated to 

(a) 3 dpa, (b) 11-25 dpa, and (c) 6-38 dpa (Refs. 60,76,79).
51

10-12

10-11

10-10

10-9

10-8

10-7

0 5 10 15 20 25 30 35 40

304, 37.8 dpa

304, 18.7 dpa

304L, 14.0 dpa

304, 6.3 dpa

Experimental CGR (m/s)

Stress Intensity K (MPa m1/2)

Material & Dose

(Halden data)

NUREG-0313

Curve

Irradiated Stainless Steels 

PWR Environment 

Normalized to 320°C

EPRI

HWC BWR

(c)

(GE data)

Figure 42. (Contd.)

0313 curve (Fig. 42b). The high growth rates in the 11.4 dpa material have been explained on the 

basis of possible effects of short test duration, high K values, and high grain-boundary Si 

concentration. Figure 42c shows that, as observed earlier for the NWC BWR environment, CGRs in 

non-CW SSs irradiated to 14 dpa or higher are mostly below the NUREG-0313 curve for 

nonirradiated SSs.

2. The data presented in Fig. 42c show that the CGRs for a Type 304 SS irradiated to only 6.3 dpa90 are 

also below the NUREG-0313 curve. The low growth rates measured for this material might be due to 

the test procedure. These tests were conducted at constant K but, before initiation of the SCC growth 

rate test at constant K, the TG fatigue precrack was not transitioned to an IG crack by slow fatigue 

cycling in the environment. 

3. Figure 43 shows the SCC growth rates in the PWR and HWC BWR environments for austenitic SSs 

irradiated to 0.7-4.25 dpa. The rates in the PWR environment were normalized to 288°C by using an 

activation energy of 100 kJ/mol. Except for the CGRs for the 3.0 dpa materials the growth rates of 

most materials are bounded by the EPRI CGR curve for the HWC BWR environment.

Figure 44 shows the SCC growth rates as a function of neutron dose for various grades and heats of 

austenitic SSs that had been irradiated to 0.7-37.5 dpa in the HWC BWR water at 288°C or the PWR 

environment at 320°C, with K = 20 MPa m1/2. All materials except those irradiated in a fast reactor were 

irradiated in BWRs at temperatures of 280-300°C; the fast reactor irradiations were at 320°C. The data 

for the HWC BWR or PWR environments may be represented by the bounding curves shown in the 

figure. The curve represents a variation of A1 in the NUREG-0313 disposition curve (Eq. 5) with neutron 

dose that is expressed by Eq. 11. 
52

10-12

10-11

10-10

10-9

10-8

10-7

0 5 10 15 20 25 30 35 40

316NG 1.5-2.3 dpa

347 SS 2.3-3.3 dpa

347 SS 2.3-3.3 dpa

347 SS 0.7 dpa

347 SS 0.95 dpa

304 SS 4.25 dpa

304L SS 3.8 dpa

304L 3.0 dpa (PWR)

Experimental CGR (m/s)

304L 3.0 dpa (PWR)

Stress Intensity K (MPa m1/2)

Material & Dose

NUREG-0313

Curve

Irradiated Stainless Steels 

BWR Environment

EPRI

HWC BWR

PWR data normalized to 288°C

using Q = 100 kJ/mole

Figure 43. Normalized SCC growth rates of austenitic stainless steels irradiated to 0.7-4.25 dpa in HWC 

BWR and PWR environments at 288°C (Ref. 74-79).

10-11

10-10

10-9

10-8

10-1 100 101

ANL

Halden

Studsvik

Studsvik (Fast)

Halden

Studsvik (fast)

GE

Crack Growth Rate (m/s)

Neutron Dose (dpa)

Irradiated Austenitic SSs

Normalized to 20 MPa!m 

HWC BWR Water at 288°C

PWR Environment 320°C

CGR = -1.79x10-9 + 4.42x10-09 log(dpa)

NUREG-0313

Figure 44. Crack growth rates for irradiated austenitic stainless steels in HWC BWR water at 288°C or 

PWR environment at 320°C plotted as a function of neutron dose.

As noted previously, the available SCC growth rate data for irradiated austenitic SSs in PWR 

environments show significant variability. The reason for the high CGRs for some heats irradiated to 

11.4 dpa and the unusually low growth rates in SSs irradiated above 18 dpa is not known. Also, most of 

the data in the PWR environment were obtained either on BWR irradiated material or materials irradiated 

in the BOR-60 fast reactor. The possible effects of differences in neutron spectrum and flux (fast vs. 

LWR) or irradiation temperature (288°C for BWRs vs. >320°C for PWRs) cannot be determined from the 

existing data. Additional data on PWR irradiated materials and materials irradiated at different 

temperatures (e.g., between 300 and 350°C) are needed to better understand the IASCC susceptibility of 

austenitic SSs in PWR environments.
53

2.2.3 Fatigue Crack Growth Rates

Cyclic CGRs tests have been conducted with either a triangular or slow/fast sawtooth waveform. 

The values of load ratio, R, are ≤0.3 for the triangular waveform tests and 0.3-0.8 for the sawtooth 

waveform tests. The reported time-based CGRs (da/dt) were determined by using only the rise time for 

the fatigue cycle. Corresponding CGRs in air, under the same loading conditions, were determined from 

the correlations developed by James and Jones91 for solution-annealed SSs. 

Under cyclic loading in a test environment, 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, given as

CGRenv = CGRair + CGRcf +CGRscc. (12)

The CGRair (m/s) was determined from the James and Jones correlations:

91

CGRair = CSS S(R) ΔK3.3/tr , (13)

where R is the load ratio (Kmin/Kmax), ΔK is Kmax - Kmin in MPa m1/2, tr is the rise time (s) of the 

loading waveform, and the function S(R) is expressed in terms of the load ratio R as follows: 

S(R) = 1.0 R < 0

S(R) = 1.0 + 1.8R 0 < R <0.79

S(R) = -43.35 + 57.97R 0.79 < R <1.0. (14)

In Eq. 13, CSS is a function of temperature, which is expressed by James and Jones in terms of a thirdorder polynomial in temperature T (°C):

CSS = 1.9142 x 10-12 + 6.7911 x 10-15 T - 1.6638 x 10-17 T2 + 3.9616 x 10-20 T3. (15)

Shack and Kassner84 have investigated the effects of LWR coolant environments on the fatigue CGR of 

nonirradiated austenitic SSs. In the absence of any significant contribution of SCC to growth rate, the 

CGRs in water with ≈0.3 ppm DO were best represented by the expression 

CGRenv = CGRair + 4.5 x 10-5 (CGRair)0.5, (16)

and in water with ≈8 ppm DO by the expression, 

CGRenv= CGRair + 1.5 x 10-4 (CGRair)0.5. (17)

In Eq. 12, the SCC growth rates, CGRscc, are represented by the NUREG-0313 (Rev. 2)66 correlation 

shown in Eq. 5 for nonirradiated materials, and by Eq. 9 for irradiated materials. For fatigue loading, 

contributions from mechanical fatigue, and to some extent, from corrosion fatigue are always present for 

austenitic SSs and Ni alloys in LWR environments. The SCC contribution may not be significant for 

nonirradiated materials in low-potential HWC BWR or PWR environments. 

The only data on the effect of neutron irradiation on fatigue crack growth have been obtained under 

the fast breeder reactor program. The data were obtained on austenitic SSs irradiated in fast reactors 

(primarily EBR-II) at temperatures of 370-450°C and tested at 427-593°C. For Type 304 and 316 SS 

irradiated at 405-410°C to 1.2 x 1022 n/cm2 (E>0.1 MeV) (6.0 dpa), the CGRs at 427°C are less than a 
54

factor of 2 higher than those for nonirradiated material at ΔK values less than 44 MPa m1/2 (40 ksi in.1/2), 

and CGRs are comparable or lower at higher ΔK values.92 A similar behavior is observed for Type 316 

weld.93 Fatigue CGR data on Type 304 and 20% CW Type 316 SSs irradiated in the EBR-II reactor at 

slightly higher temperatures, 455-477°C, to 1.3 and 9.0 x 1021 n/cm2 (E>0.1 MeV) (0.65 and 4.5 dpa), 

showed no effects on growth rates at 427°C, and slightly lower growth rates at 538°C.94,95 Also, tests on 

Type 304 and 316 SS irradiated in a thermal reactor, the Advanced Test Reactor (ATR), at 288°C to 1.8 x 

1021 n/cm2 (E>0.1 MeV) and tested at 427°C had crack growth rates that were less than a factor of 2 

lower than those for nonirradiated materials.96 These results indicate no significant effect of irradiation 

on the fatigue CGRs for austenitic SSs at these temperatures and K values.

The measured values of fatigue CGRs in air of SS weld HAZ materials irradiated to ≈2.16 dpa are 

plotted in Fig. 45 as a function of the CGRs estimated from the James and Jones91 correlations for 

nonirradiated austenitic SSs under the same loading conditions. Once again, the results indicate no 

significant effect of irradiation on the fatigue CGRs for weld HAZ in air.

10-12

10-11

10-10

10-9

10-8

10-7

10-12 10-11 10-10 10-9 10-8 10-7

304 SS SMA Weld HAZ

304L SS SA Weld HAZ

304 SS SMA Weld HAZ 

 + 24 h at 500°C

Measured CGR (m/s)

Predicted CGR (m/s)

SS Weld HAZ

Irradiated to 2.16 dpa

Tested at 289°C in Air

Figure 45.

CGR data under cyclic loading for irradiated 

SS weld HAZ materials in air at 289°C 

(Ref. 11).

2.2.3.1 Solution-Annealed Materials

Under continuous cyclic loading, experimental data for CGR for solution-annealed Type 304 and 

316 SSs irradiated up to 3 dpa and tested in high- and low-DO environments are plotted in Fig. 46 as a 

function of CGRs predicted in air for the same loading conditions.9 The curves in the figures are based 

on the superposition model (Eq. 12). The cyclic CGRs in air (CGRair) were determined from Eq. 13, 

which was developed by James and Jones for nonirradiated materials.91 The corrosion fatigue

contribution (CGRcf) was determined from the Shack/Kassner model for nonirradiated SSs in high-purity 

water with either 8 or 0.2 ppm DO (Eqs. 16 and 17, respectively),84 and the SCC contribution (CGRscc) 

was determined from Eq. 5.

66 The SCC growth rate in SSs irradiated to >0.75 dpa was assumed to be a 

factor of six higher than that predicted by Eq. 5; as a result, the constant A in the equation was taken to be 

1.26 x 10-12 for irradiated SSs. For cyclic loading using either a triangular or a slow/fast sawtooth 

waveform, CGRscc is determined by considering the contribution of SCC during the slow rise time of the 

cycle; an equivalent K is computed to determine the contribution of SCC. The average values of K used 

in calculating the superposition curves are given in Fig. 46.
55

10-12

10-11

10-10

10-9

10-8

10-7

10-12 10-11 10-10 10-9 10-8 10-7

304L 0.45 dpa

304L 1.35 dpa

304L 3.00 dpa

316L 3.00 dpa

316 1.35 dpa

316 0.45 dpa

316 3.00 dpa

CGRenv (m/s)

CGRair (m/s)

Austenitic SSs, 289°C

!300 ppb DO

K = 14 MPa m1/2

Nonirradiated SS

Model 0.2 ppm DO

K = 16 MPa m1/2

Irradiated SS 

Model 8 ppm DO

10-12

10-11

10-10

10-9

10-8

10-7

10-12 10-11 10-10 10-9 10-8 10-7

304 SS 1.35 dpa

316L SS 3.00dpa

CGRenv (m/s)

CGRair (m/s)

Austenitic SSs, 289°C

<10 ppb Dissolved Oxygen

Nonirradiated SS

Model 0.2 ppm DO

CGRair +4.5x10-5CGRair0.5

No SCC in low–DO Water

(a) (b)

Figure 46. CGR for irradiated specimens of austenitic SSs under continuous cycling at 289°C in highpurity water with (a) ≈300 ppb and (b) <30 ppb dissolved oxygen (Refs. 9,11).

In Fig. 46, the data points that lie along the diagonal represent predominantly mechanical fatigue, 

and those that lie close to the model curve indicate environmentally enhanced crack growth. Austenitic 

SS irradiated to 0.45 dpa shows very little environmental enhancement of CGRs in high-DO water (open 

and closed diamonds in Fig. 46a). For austenitic SSs irradiated to less than 0.5 dpa, the fatigue CGRs in 

water with ≈ 300 ppb DO may be represented by superposition of the NUREG-0313 curve for 

nonirradiated SSs66 and by the Shack/Kassner model for nonirradiated austenitic SSs in high-purity water 

with 0.2 ppm DO.84 

The results for SSs irradiated to 1.35 or 3.0 dpa indicate significant enhancement of the CGRs in 

high-DO water under cyclic loading with long rise times. For austenitic SSs irradiated to 0.75-3.0 dpa, 

the fatigue CGRs in water with ≈ 300 ppb DO may be represented by superposition of the SCC curve for 

irradiated SSs (i.e., six times the NUREG-0313 curve) and by the Shack/Kassner model for nonirradiated 

SSs in high-purity water with 8 ppm DO.84 

For continuous cyclic loading, decreasing the DO level has a beneficial effect on the CGRs of 

irradiated SSs; for example, decreasing the DO from ≈300 ppb DO to <30 ppb DO lowers the CGR by a 

factor of 25. At 289°C, the fatigue CGRs for irradiated austenitic SSs in water with <30 ppb DO are 

lower than those predicted by the Shack/Kassner model for nonirradiated austenitic SSs in high-purity 

water with 0.2 ppm DO (Fig. 46b);84 there is no contribution of SCC in low-DO water. 

2.2.3.2 Stainless Steel Weld HAZ Materials

The experimental data for CGRs of nonirradiated SS weld HAZ materials in high-DO water10 are 

plotted in Fig. 47 as a function of those predicted in air for the same loading conditions. The curve 

(dotted and dashed line) is based on the superposition model. For the nonirradiated HAZ materials, the 

growth rate did not increase readily when the load ratio and rise time were increased. A large number of 

data points lie along or below the diagonal in Fig. 47. The applied Kmax had to be increased for 

environmental enhancement. The fatigue CGRs of nonirradiated SS weld HAZ materials may be 

conservatively represented by superposition of the SCC curve for nonirradiated SSs and the 
56

Shack/Kassner model for austenitic SSs in high-purity water with 8 ppm DO.84 The results also indicate 

that thermal treatment of the material for 24 h at 500°C has little or no effect on growth rates.

Experimental data for CGRs of irradiated GG Type 304L SA weld HAZ from the Grand Gulf (GG) 

core shroud and laboratory-prepared Type 304 SMA weld HAZ in high-DO water are plotted as a 

function of those predicted in air for the same loading conditions in Figs. 48a and 48b, respectively. The 

curves in the figures are based on the superposition model (Eq. 12). The results indicate significant 

environmental enhancement of CGRs for HAZ materials irradiated to 0.75 or 2.16 dpa. The CGRs of the 

GG Type 304L weld HAZ are slightly lower than those of the Type 304 SMA weld HAZ. The fatigue 

CGRs of SS weld HAZ materials irradiated to 0.75-2.16 dpa in water containing ≈ 500 ppb DO can be 

represented by superposition of the SCC curve for irradiated SSs (i.e., six times the NUREG-0313 curve) 

and the Shack/Kassner model for nonirradiated austenitic SSs in high-purity water with 8 ppm DO.84 

The estimates may be somewhat conservative for Type 304L weld HAZ materials. 

10-12

10-11

10-10

10-9

10-8

10-7

10-12 10-11 10-10 10-9 10-8 10-7

GG5T-A

GG5T-B

GG6T-A

CGRenv (m/s)

CGRair (m/s)

Grand Gulf Core Shroud

Type 304 L SA Weld HAZ

289°C, !500 ppb DO

K = 13 MPa m1/2

Irradiated SS

Model 8 ppm DO

Open Symbols: 0.75 dpa

Closed Symbols: 2.16 dpa

Specimen Number

10-12

10-11

10-10

10-9

10-8

10-7

10-12 10-11 10-10 10-9 10-8 10-7

85-1A-TT

85-7A

85-XA

85-3TT

CGRenv (m/s)

CGRair (m/s)

Type 304 SS (Heat 10285)

289°C, !500 ppb DO

K = 13 MPa m1/2

Irradiated SS

Model 8 ppm DO

Open Symbols: 0.75 dpa

Closed Symbols: 2.16 dpa

SMA Weld HAZ

Sensitized

(a) (b)

Figure 48. CGR for irradiated specimens of (a) Type 304L SA weld HAZ from the Grand Gulf core 

shroud and (b) laboratory-prepared Type 304 SS SMA weld HAZ under continuous cycling in 

high-purity water at 289°C (Ref. 11).

10-12

10-11

10-10

10-9

10-8

10-7

10-12 10-11 10-10 10-9 10-8 10-7

85-3A-TT

85-YA

GG3B-A-TT

GG5B-A

CGRenv (m/s)

CGRair (m/s)

Non Irradiated SS Weld HAZ

300–500 ppb Dissolved Oxygen

Symbols with +: K/size 

 criterion not satisfied

Specimen Number

Type 304 SMA Weld HAZ

Type 304L SA Weld HAZ

K = 17 MPa m1/2

Nonirradiated SS

Model 8 ppm DO

Figure 47.

CGR data under cyclic loading for nonirradiated 

SS weld HAZ materials in high-purity water at 

289°C (Ref. 11).
57

2.2.3.3 Cast Austenitic Stainless Steels

Experimental data for CGRs of CF-8M cast austenitic SS under continuous cycling in the NWC 

BWR environment are plotted in Fig. 49 as a function of those predicted for austenitic SSs under the same 

loading conditions in air. The two curves in the figure are based on the superposition model for different 

DO levels. The material was thermally aged for 10,000 h at 400°C and then irradiated to 2.46 dpa at 

≈300°C. As seen before (Fig. 47) for nonirradiated HAZ materials, environmental enhancement of CGRs 

did not occur readily for Specimen 75-11TT when the load ratio and rise time were increased; for this 

specimen, a large number of data points lie along the diagonal in Fig. 49. The applied Kmax had to be 

increased for environmental enhancement.

Under similar loading and environmental conditions, the fatigue CGRs of CF-8M cast austenitic SS 

appear to be lower than those of wrought SSs or SS weld HAZ materials. The limited data available 

indicate that the fatigue CGRs of SS weld HAZ materials irradiated to 0.75-2.46 dpa in water containing 

≈ 300 ppb DO can be represented by superposition of the SCC curve for irradiated SSs (i.e., six times 

the NUREG-0313 curve) and the Shack/Kassner model for nonirradiated austenitic SSs in high-purity 

water with 0.2 ppm DO.84 

10-13

10-12

10-11

10-10

10-9

10-8

10-7

10-13 10-12 10-11 10-10 10-9 10-8 10-7

75-11TT

75-11TM

CGRenv (m/s)

CGRair (m/s)

CF-8M Cast Austenitic SS

Irradiated to 2.46 dpa

289°C !300 ppb DO Water

K = 13 MPa m1/2

Irradiated SS

Model 8 ppm DO

Specimen Number

K = 13 MPa m1/2

Irradiated SS

Model 0.2 ppm DO

Figure 49.

CGR data under cyclic loading for two 

irradiated specimens of CF-8M cast austenitic 

SS in high-purity water at 289°C (Ref. 11).

2.3 Initiation of Irradiation-Assisted Stress Corrosion Cracking

As discussed previously, IASCC is an irradiation-induced increase in susceptibility of materials to 

SCC with increasing neutron fluence. The SCC of materials in high-temperature, high-pressure water 

depends on material susceptibility, high stresses, and an aggressive environment. However, the 

mechanism of IASCC in the PWR environment and the individual effect of various material and 

environmental parameters on IASCC are still not well understood. Although the degradation of tensile 

and fracture properties of austenitic SSs appears to saturate at 5-20 dpa, the susceptibility of these 

materials to IASCC continues to increase with neutron fluence. Both solution-annealed and cold-worked 

austenitic SSs are susceptible to cracking in the PWR environment, but there is significant heat-to-heat 

variability. In general, the irradiation-induced degradation of fracture properties of austenitic SSs is 

slower in cold-worked than solution-annealed materials because the cold-worked materials have a high 

dislocation density, which slows irradiation hardening and damage and suppresses void nucleation and 

swelling. However, at high neutron fluence levels (i.e., above 10-20 dpa or 6.7 x 1021 n/cm2 to 
58

1.3 x 1022 n/cm2, E>1.0 MeV), there is little or no difference in IASCC susceptibility of these materials 

in PWR environments. 

For all wrought and cast austenitic SSs and their welds there appears to be a threshold value of 

neutron fluence below which the materials can be considered not susceptible to IASCC in a PWR 

environment. Current estimates of a threshold value have been based on laboratory SSRT data. The 

following values have been proposed: a conservative threshold fluence of approximately 6.7 x 1020 n/cm2

(1 dpa),101 which represents the fluence at which IASCC can occur in a material after extremely high 

strains, and a more realistic value of 2 x 1021 n/cm2 (3 dpa),100,101 which represents the fluence at which 

IASCC can be initiated at above the yield stress of the material. For materials irradiated above this 

threshold, IASCC initiation data have also been used to define for a given neutron dose, an apparent stress 

threshold below which IASCC initiation will not occur in a PWR environment. The SCC initiation data 

were obtained by conducting constant-load SCC initiation tests in a simulated PWR environment on Oring, C-ring, or tensile specimens of irradiated materials.61,62,97-101 The specimens were tested as a 

function of fluence and applied stress. 

The constant-load IASCC initiation results plotted as stress (as percent of irradiated yield stress) 

versus time are shown in Fig. 50. The open symbols represent specimens that did not fail, and the closed 

symbols represent failed specimens. The same data are plotted as a function of neutron dose in Fig. 51. 

The results indicate that under a high enough stress, crack initiation in highly irradiated materials can 

occur quite rapidly (i.e., within 500 h). Furthermore, 80% of these failures occurred within 150 h. The 

results also indicate a stress threshold below which cracks did not initiate even after several thousands of 

hours. Data on CW Type 316 SS flux thimble tubes removed from commercial PWRs show that, at 

26 dpa, cracks initiated at stresses above approximately 62% of the irradiated yield stress.97 In Fig. 51, 

for SSs irradiated to very high fluence levels, the stress threshold has been defined as 50% of the 

irradiated yield stress, although two data points fall below this threshold. The IASCC initiation data also 

indicate significant scatter with respect to failure vs. non-failure, and the time to failure. For example, in 

a set of six identical specimens simultaneously tested, four failed and two did not. Also, of the four 

0

20

40

60

80

100

120

140

160

100 101 102 103 104

Beaver Valley

H.B.Robinson

Ringhals

Chooz A 304 

Barseback 304L

Nishioka 316CW

Takakura BFB

Takakura FTT

Bugey 2 316CW

Chooz A 304

Percent of Irradiated Yield Stress

Time (h)

Open Symbols: No failure

Closed Symbols: Failure

Figure 50. Stress as percent of irradiated yield stress vs. time data for IASCC flaw initiation in austenitic 

stainless steels in a PWR environment (Refs. 61,62,97-101).
59

specimens that did fail, time to failure varied from 29 to 483 h. Therefore, sufficiently large numbers of 

specimens have been tested to identify valid data trends. 

0

20

40

60

80

100

120

100 101 102 103 104 105 106

Percent of Irradiated Yield Stress

Time (h)

No IASCC

71-76 dpa

48 EFPY

Figure 52. Percent of irradiated yield stress vs. time for IASCC flaw initiation in a PWR 

environment, for an austenitic stainless steel irradiated to 71-76 dpa (Ref. 62,97,98).

The IASCC initiation data have been used to estimate, for a given neutron dose, a lower bound 

stress vs. time-to-failure curve for austenitic SSs irradiated to 71-76 dpa. The curve in Fig. 52 represents 

for a specific irradiated material, the shortest time to initiate IASCC at a given stress in the PWR 

environment. These curves are extrapolated to 40-y or 60-y reactor life [i.e., 32 or 48 effective full power 

year (EFPY)] to define an apparent stress threshold below which IASCC initiation will not occur for that 

material within a specified time, say, the reactor lifetime. Such results are used to develop the curves for 

stress vs. neutron fluence threshold that predict time for IASCC initiation as a function of stress and 

fluence. 

0

20

40

60

80

100

120

140

160

0 20 40 60 80 100

Beaver Valley

H.B.Robinson

Ringhals

Chooz A 304 CIR Halden

Barseback 304L CIR Halden

Nishioka 316CW

Takakura Baffle Former Bolt

Takakura Flux Thimble tube

Bugey 2 316CW EdF

Chooz A 304 JOBB

Percent of Irradiated Yield Stress

Neutron Dose (dpa)

Open Symbols: No failure

Closed Symbols: Failure

Figure 51. Stress as percent of irradiated yield stress vs. neutron dose for IASCC flaw initiation in 

austenitic stainless steels in a PWR environment (Refs. 61,62,97-101).
60

Figure 53 presents two curves that define the approximate stress and neutron fluence for IASCC 

initiation in a relatively short time (i.e., 100 h) and after a very long time. The latter may be used as the 

lower bound for stress and fluence below which IASCC is not likely to occur within the lifetime of the 

reactor. However, the data that were used to develop these curves do not account for the effect of 

material composition, material variability between heats, or the differences between solution-annealed or 

cold-worked material. To account for material variability and uncertainty due to data scatter, MRP has 

proposed a screening curve, shown in Fig. 53, for IASCC initiation of austenitic SSs. This curve is used 

to divide various PWR core internal components into different categories of aging management strategies 

during the initial screening process. 

0

200

400

600

800

0 10 20 30 40 50 60 70 80

Stress (MPa)

Neutron Exposure (dpa)

No IASCC

Short term

IASCC failure

possible

Long term

IASCC failure possible

Proposed Screening Criterion

Stress = 158.4 ln(dpa) + 789.65

Figure 53. Time for initiation of IASCC in irradiated austenitic stainless steels as a 

function of stress (Ref. 61,101).

Such an approach for defining neutron dose/applied stress thresholds for austenitic SSs is limited by 

the adequacy of the database used for developing the screening curve. First of all, insufficient data are 

currently available to adequately define an IASCC initiation curve correlating stress and neutron fluence. 

Considerably more data are needed to account for uncertainties due to data scatter, material variability, 

irradiation temperature, neutron spectrum effects, and water chemistry. As more information has been 

added to the IASCC initiation database, the threshold stress below which IASCC is considered not likely 

to occur during a 60-y reactor lifetime has been decreased from an initial value of 62% to 50% and now 

possibly 40% of the irradiated yield stress. Although a threshold fluence of 3 dpa has been defined for 

IASCC in the PWR environment, there are no IASCC initiation data on austenitic SSs irradiated between 

3 and 9.5 dpa. Thus, in Fig. 51, for materials irradiated to 3-10 dpa, the threshold stress for IASCC 

initiation is defined arbitrarily by a straight line between 3 and 10 dpa. In addition, the existing database 

for IASCC initiation does not consider the effects of the following parameters that are known to influence 

IASCC susceptibility:

1. Material Type, Composition, and Condition: The available IASCC initiation data have been obtained 

primarily on CW Type 316 SS and data on solution-annealed Type 304, 304L, or 316 SS are limited. 

Also, there are no data on cast austenitic SSs, SS welds, and weld HAZ materials. Additional IASCC 

initiation data should be obtained on these materials irradiated to neutron fluence levels relevant for 

PWR core internals. 

Also, the increased susceptibility to IASCC and the loss of benefit of reduced corrosion potential on 

IASCC susceptibility in highly irradiated SSs are often attributed to the segregation of Si at the grain 
61

boundaries. Initiation data should be obtained on irradiated material from commercial heats of SSs 

with similar compositions but different Si contents to establish the role of Si segregation on the 

IASCC susceptibility of irradiated SSs.

2. Irradiation Temperature: Nearly all of the high fluence (above 20 dpa) data and the majority of the 

low fluence data on IASCC initiation have been obtained on materials irradiated below 325°C. 

Consequently, the database does not include the potential effects of additional precipitate phases, 

voids, and cavities, which are observed in SSs irradiated to high fluence levels at temperatures above 

320°C. Also, He bubbles have been observed in SSs irradiated in reactors with high thermal neutron 

fluxes and could also influence IASCC initiation. 

3. Neutron Spectrum: The current IASCC initiation data include some tests on austenitic SS irradiated in 

the BOR-60 fast reactor. Several studies have shown that materials irradiated in fast reactor show 

lower susceptibility to IASCC than those irradiated in LWRs. Such data have not been included in 

Figs. 50 and 51. Caution should be exercised when using fast reactor data to assess the degradation 

of LWR core internal materials. 

4. Water Chemistry: The addition of H2 to the reactor water greatly reduces the effect of radiolysis by 

scavenging the radiolysis products. Because PWR coolants typically contain 2 ppm H2 (30 cc/kg), 

radiolysis has no effect on the corrosion potential in PWRs. However, in reactor core internal 

components, IASCC is likely to occur under creviced conditions (i.e., under off-normal water 

chemistry). Nearly all of the IASCC initiation tests have been conducted in the normal PWR 

environment, and the possible effects of impurities on initiation have not been investigated. 
62

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63

3 Neutron Embrittlement

Neutron irradiation can decrease the fracture toughness of austenitic SSs significantly, and failure 

may occur without general yielding. In such instances, a fracture mechanics methodology, such as 

elastic-plastic fracture mechanics (EPFM) or linear-elastic fracture mechanics (LEFM), is needed for 

analysis of structural integrity and development of inspection guidelines. The former involves the J 

integral-resistance (J-R) curve approach and is used where failure is caused by plastic deformation. The 

J integral is a mathematical expression used to characterize the local stress-strain field at the crack tip 

region (parameter J represents the driving force for crack propagation), and the J-R curve characterizes 

the resistance of the material to stable crack extension. The fracture toughness of such materials is 

represented by fracture mechanics parameters such as JIc, the value of J near the onset of crack extension, 

and the tearing modulus, T, which characterizes the slope of the J-R curve:

T = dJ

da

E

!f

2 , (18)

where E is the elastic modulus, a is the crack length, and σf is the flow stress defined as the average of the 

yield stress (σy) and ultimate stress (σu). The LEFM methodology is used where failure involves 

negligible plastic deformation. The fracture toughness of such materials is represented by the parameter 

KIc (i.e., plane strain fracture toughness), which characterizes the resistance of the material to unstable 

crack extension. Under EPFM conditions, the equivalent critical stress intensity factor, KJc, can be 

determined from the saturation JIc value using the relationship

KJc = E!J ( Ic )

1/2

, (19)

where E! = E / 1" #2 ( ), E is the elastic modulus, and υ is the Poisson ratio.

The fracture toughness of austenitic SSs has been divided into three broad categories.6 Category III 

corresponds to high toughness materials with JIc above 150 kJ/cm2 (857 in.-lb/in.2). In these materials, 

fracture occurs after stable crack extension at stresses well above the yield stress. Category II 

corresponds to materials with intermediate toughness with JIc in the range of 30-150 kJ/cm2 (171-857 in.-

lb/in.2). In these materials, fracture occurs by stable or unstable crack extension at stress levels close to 

the yield stress. Category I corresponds to low-toughness materials with KIc less than 75 MPa m1/2

(68.2 ksi in.1/2) [JIc < 30 kJ/cm2 (< 171 in.-lb/in.2)]. In these materials, fracture occurs by unstable crack 

extension at stress levels well below the yield stress. 

Nonirradiated wrought and cast austenitic SSs and their welds fall in Category III. The JIc values 

for Type 304 and 316 SS at temperatures up to 125°C (257°F) vary between 169 and 1660 kJ/cm2 (965 

and 9479 in.-lb/in.2), with a median value of 672 kJ/cm2 (3837 in.-lb/in.2).6 The JIc values at 400-550°C 

(752-1022°F) are ≈35% lower, with a median value of 421 kJ/cm2 (2404 in.-lb/in.2). Fracture in such 

high-toughness materials is by the nucleation and coalescence of microvoids and is characterized by a 

dimpled fracture morphology. 

Although cast austenitic SSs and SS welds also exhibit ductile fracture at temperatures up to 550°C 

(1022°F), their fracture toughness is lower than that of the wrought SSs. A dimpled fracture morphology 

is also observed in SS welds.102 Because of a high density of inclusions in the weld, the dimples are 
64

relatively small and shallow. Also, dimples are often associated with an inclusion and are initiated by 

decohesion of the particle/matrix interface. The overall fracture toughness of cast austenitic SSs and SS 

welds is controlled by the density and morphology of second-phase inclusions in these materials and 

varies with the cast or weld process. For example, static cast products have lower fracture toughness than 

centrifugally cast pipes. Gas tungsten arc (GTA) welds exhibit the highest toughness; SMA welds have 

intermediate toughness; and SA welds have the lowest toughness.6 The median value of JIc is 492 kJ/cm2

(2809 in.-lb/in.2) for GTA welds and 147 kJ/cm2 (839 in.-lb/in.2) for SA welds at temperatures up to 

125°C (257°F). 

Welding of austenitic SSs results in a HAZ adjacent to the fusion zone, where the material 

microstructure and microchemistry are greatly altered because of the precipitation of Cr-rich carbides at 

the grain boundaries. The formation of the carbides depletes Cr from the grain-boundary region, thereby 

creating a region that is susceptible to SCC. However, the fracture toughness of HAZ material is 

generally superior to that of the weld metal and may be comparable to that of the base metal. 

3.1 Fracture Toughness of Irradiated Austenitic Stainless Steels

Neutron irradiation can degrade fracture toughness of austenitic SSs to the level of Category II or I. 

Until recently, most of the published experimental data on neutron embrittlement of austenitic SSs had 

been obtained on materials irradiated in high-flux fast reactors.103-119 In these studies, the embrittlement 

of the materials has been characterized in terms of tensile properties, Charpy-impact properties, and 

fracture toughness. The tensile properties of Type 304, 316, and 347 SSs and their weld metals and HAZ 

materials, as well as CF-8 cast austenitic SSs irradiated in fast reactors or LWRs, have been discussed in 

Section 2.1.4 of this report. 

Fracture toughness is typically characterized by the initiation toughness JIc and tearing modulus T 

for materials that fail after substantial plastic deformation (for EPFM analysis) and by the critical stress 

intensity factor KIc for materials that fail after little or no deformation (for LEFM analysis). Typically, 

fracture toughness data for irradiated materials have been obtained from compact tension (CT) or single 

edge bend [SE(B)] specimens and, in some cases, from Chevron notch short rod (CNSR) specimens. To 

reduce activity and facilitate handling, small specimens (e.g., ≈8-mm thick 1⁄4-T CT) have been used in 

several studies. For these specimens, J values above 150 kJ/m1/2 and crack extensions beyond about 

1.2 mm are above the validity limits based on ASTM Specification E 1820-06. However, comparison of 

fracture toughness data obtained on 1-T CT and small-size CT or SE(B) specimens shows comparable J 

vs. Δa values even beyond the ASTM defined validity limits. The small specimens yield equivalent J-R 

curve data at least for materials with JIc values up to about 300 kJ/m1/2 and, maybe, even higher. 

Plots of JIc or KIc and KJc as a function of neutron dose are used for developing screening criteria 

for neutron embrittlement. In ASTM Specification E 1820-06, JIc is determined from the intersection of 

the best-fit power-law J-R curve with the 0.2-offset line parallel to the blunting line, provided the 

specimen size criterion of Eq. 7 is satisfied. The blunting line is defined as 

J = MσfΔa, (20)

where σf is the flow stress, Δa is crack extension, and the constraint factor M is 2 or a value determined 

from the best fit of the experimental data. However, the analysis procedures described in the ASTM 

specifications for JIc determination are not applicable to austenitic SSs because of their extremely high 

toughness, ductility, and strain hardening ability. The main difference concerns the expression for the 

crack-tip blunting line. For austenitic SSs, a value of 2 for M significantly overpredicts the crack 
65

extension due to crack tip blunting, and, therefore, it yields a non-conservative value of JIc.

6,120 For 

austenitic SSs, a value of 4 for M better defines the blunting line. 

The constraint factor M, which relates J to the crack tip opening displacement (CTOD) given by the 

expression 

J = Mσy(CTOD). (21)

The use of a higher value for M in Eq. 20 is consistent with the expected variation of M and σf with strain 

hardening. The factor M is 1 for materials with intermediate to high strength and low strain hardening, 

and 2 for materials with low strength and high strain hardening, such as austenitic SSs. For the latter, the 

yield strength is approximately two-thirds of the flow stress, and the crack extension associated with 

blunting is approximately one-third of CTOD.6 Thus, for such materials, the crack tip blunting line is 

given by

J = Mσy(CTOD) ≈ 2(2σf/3)(3Δa) = 4σfΔa, (22)

i.e., Eq. 20 with M = 4. This relationship has been used to determine JIc in most investigations on neutron 

embrittlement.11,59 A value of 2 for M has also been used by several investigators.

58 The latter typically 

yields a higher value of JIc for Category III materials (i.e., JIc above 150 kJ/cm). However, the difference 

in JIc values determined using values of M of 2 or 4 is insignificant for Category II materials (i.e., JIc

<100 kJ/cm2). Since it is primarily the cases in which the fracture toughness of irradiated austenitic SSs 

is reduced to Category II levels that are of interest, the effect of differences in the procedure to determine 

JIc is likely to be insignificant. 

Another factor that may influence the reported values of JIc is when an effective yield stress is used 

instead of the measured yield stress. As discussed in Section 2.2.1.3, the K/size criteria were developed 

for materials that show work hardening and, therefore, may not be applicable for materials irradiated to 

fluence levels where, on a local level, they do not strain harden. An effective yield stress, in which the 

irradiation-induced increase in yield stress is discounted by a factor of 2 for moderately irradiated 

materials and by a factor of 3 for highly irradiated materials, has been proposed to define K/size criteria 

for moderate to highly irradiated materials. Some studies have used such a yield stress to determine JIc.

11 

Because JIc is a measure of fracture toughness at instability without significant stable crack extension, the 

measured yield or flow stress of the irradiated materials seems more appropriate for JIc determinations. 

Nevertheless, the choice of measured or effective yield stress is likely to have an insignificant effect on 

the measured JIc of materials with poor fracture toughness. 

3.1.1 Fracture Toughness JIc

The effect of neutron exposure (in dpa) on the fracture toughness JIc at 25-427°C (77-842°F) of 

austenitic SSs irradiated at 90-450°C (194-842°F) up to 90 dpa in fast reactors is shown in Fig. 54.

103-119

The irradiation and test temperatures were 100-155°C and 125°C for the data obtained by Sindelar et 

al.,114 90-250°C and 25-250°C for the data obtained by Alexander et al.,113 and 325°C and 25°C for the 

data obtained by Kim et al.119 The rest of the data were obtained on materials irradiated at 350-427°C 

and tested at 300-427°C. The fast reactor data show a rapid decrease in fracture toughness at a neutron 

dose of 1-2 dpa; the neutron dose at the onset of the rapid decrease varies with the chemical composition 

and heat treatment of the steel. The effects of irradiation may be divided into three regimes: little or no 

loss of toughness below an exposure of ≈1 dpa, substantial decrease in toughness at exposures of 1-

10 dpa, and no further reduction in toughness above a saturation exposure of 10 dpa. The degradation in 
66

fracture properties saturates at a JIc value of ≈30 kJ/m2 (171 in.-lb/in.2) [i.e., KJc of 75 MPa m1/2

(68.2 ksi in.1/2)]. Also, the failure mode changes from dimple fracture to channel fracture. 

The fracture toughness trend for the LWR data5,11,57-59,65,121,122 is similar to that observed for fast 

reactor data (Fig. 55). Most of the fracture toughness JIc values for austenitic SSs irradiated in LWRs 

[288-316°C (550-601°F)] fall within the scatter band of the data obtained on materials irradiated in fast 

reactors, even though the LWR irradiations were at lower temperatures. There are only minor differences 

in the fracture toughness of the various wrought and cast austenitic SS materials. For the same irradiation 

conditions, the fracture toughness of thermally aged cast SS and weld metal is lower than that of HAZ 

material, which, in turn, is lower than that of solution-annealed materials. A similar behavior is also 

observed for the fast reactor data in Fig. 54. The JIc values of welds and HAZ materials are consistently 

lower than those for the solution-annealed and even CW materials. The data for CF-8 cast SS were 

obtained at room temperature and, therefore, are relatively high; the JIc values are expected to be lower at 

LWR operating temperatures. 

Some materials irradiated above 4 dpa at LWR temperatures show very poor fracture toughness; 

their JIc values are below the lower bound curve for the fast reactor data. For a Type 304 SS irradiated to 

4.5-5.3 dpa (shown as cross in Fig. 55), nine out of ten CT specimens showed no ductile crack extension, 

and the KIc values were 52.5-67.5 MPa m1/2 (47.7-61.4 ksi in1/2).57 The lowest fracture toughness, with 

KIc or KJc values in the range of 36.8-40.3 MPa m1/2 (33.5-36.6 ksi in1/2), was for a Type 347 SS 

irradiated to 16.5 dpa in a PWR57 and for a Type 304 SS irradiated to 7.4-8.4 dpa in a BWR.122 

Another significant result is a strong orientation effect on fracture toughness. Fracture toughness JR tests and microstructural and microchemistry characterization were performed on a Type 304 controlrod and Type 304L top guide materials irradiated to 4.7-12 dpa in a BWR, and on Type 304 control-rod 

0

100

200

300

400

500

600

0 20 40 60 80

316L Van Osch et al./Alexander et al.)

304 (Sindelar et al./Mills/De Vries/Michel & Gray)

316CW (Michel & Gray/Hamilton et al./Huang)

304 HAZ (Sindelar et al.)

321 (Little)

304L (Van Osch et al.)

316 (Picker et al./Mills/O'Donnell et al./Michel & Gray)

CF-8 (Burke et al.)

308 (Michel & Gray/Sindelar et al./Mills)

316 Weld (Picker et al./O'Donnell et al.)

316H Weld (Bernard & Verzeletti)

JIc (kJ/m2)

Neutron Exposure (dpa)

Austenitic SSs

Irradiated in Fast Reactors 90-427°C

Tested at 25-427°C

Figure 54. Change in fracture toughness JIc as a function of neutron exposure for austenitic SSs 

irradiated in fast reactors. Solid lines represent the scatter band for the fast reactor data on 

austenitic SSs (Refs. 104-119).
67

material irradiated to 7.4 and 8.4 dpa in another BWR. All materials consistently showed lower fracture 

toughness in the T-L orientation than in the L-T orientation,58 and these toughness values were lower 

than the limiting fracture toughness KIc of 55 MPa m1/2 (50 ksi in.1/2) that has been proposed by industry 

for flaw tolerance evaluation in austenitic SSs irradiated above 4.5 dpa (3 x 10-21 n/cm2).58,65 The lower 

fracture toughness along the T-L orientation has been attributed to the presence of stringers consisting of 

long, narrow particles oriented in the rolling direction, which result in a long and narrow quasi-cleavage 

structure parallel to the crack advance, thereby accelerating the crack advance.58 In addition, the 

Type 304 control-rod material (7.4-8.4 dpa) showed very poor fracture toughness (JIc of 40 kJ/m2 in L-T 

and 7.5 kJ/m2 in T-L orientation). The authors concluded that the low JIc of this material might be 

considered a special case of materials containing a high density of particles aligned in the rolling 

direction. Nonetheless, these results show that very low fracture toughness values are possible for 

irradiated austenitic SSs. Microstructural characterization of the Type 304 control-rod material showed a 

fine distribution of γ’ phase with size in the range of 2-10 nm and an average size of 4.4 nm.58 The 

density was 1-3 x 1022 m-3. This phase was not observed in the Type 304 top guide material, and may 

influence the fracture toughness of these materials. The γ’ phase has been observed at dose levels above 

4 dpa in CW Type 316 SS irradiated under the PWR condition.123 The contribution of additional 

precipitate phases, voids, and cavities on fracture toughness needs to be investigated. 

0

100

200

300

400

500

600

0 5 10 15 20 25

304 JAPEIC (CT)

304 JAPEIC (BB)

304 JAPEIC (SR)

304 MRP-160

316CW/347 MRP-79

304 MRP-79

304L MRP-79

E308L MRP-79

304L/316L (Ehrnsten 2006) 

304/316L NUREG/CR-6960

304/304L HAZ NUREG/CR-6960

CF-8M Aged NUREG/CR-6960

304 Sensi NUREG/CR-6960

304/304L L-T (Demma 2007)

304/304L T-L (Demma 2007)

304 (Fyfitch 2009)

JIc (kJ/m2)

Neutron Exposure (dpa)

Austenitic SSs

Irradiated in LWRs

Tested at 250-320°C

835 kJ/cm2

Closed Symbols: BWR Water

Open Symbols: Air

Figure 55. Change in fracture toughness JIc as a function of neutron exposure for irradiated austenitic 

SSs. Dashed lines represent the scatter band for the fast reactor data on austenitic SSs 

irradiated at 350-450°C (662-843°F) (Refs. 5,11,57-59,65,121,122).

3.1.2 Fracture Toughness J-R Curve

Fracture toughness J-R curve data have been obtained for the following materials: Type 304, 304L, 

and 316L SSs; SS weld materials and HAZs; and CF-3, CF-8, and CF-8M cast austenitic SSs. Data are 

available for materials irradiated in LWRs up to about 14 dpa,11,57-59,65,121,122 and in fast reactors to 

much higher dose levels. The change in the J-R curve with neutron dose is shown in Fig. 56. The 

decrease in fracture toughness is quite rapid up to about 6 dpa, and the toughness continues to decrease 
68

moderately at higher dose levels. The effects of various parameters such as material type and heat 

treatment, test and irradiation temperature, and neutron energy spectrum, flux, and dose are discussed 

below.

Irradiation Conditions: Fast reactor irradiations are at fluxes and temperatures higher than those 

typically observed in LWRs and have a different spectrum. Until recently, most of the high neutron 

exposure data were from fast reactor irradiations at temperatures above 350°C (662°F) (Fig. 54). An 

accurate determination of the effects of neutron spectrum, flux, and temperature on the fracture properties 

of these materials requires data on the same heat of material irradiated in a fast reactor and an LWR to 

comparable neutron dose. Such information is not available. Although the general data trends appear to 

be similar for fast reactor and LWR irradiations, the tensile property data discussed in Section 2.1.4 

indicate that tensile strength is higher and ductility is lower for the BWR-irradiated materials than 

materials irradiated in fast reactors. However, the existing data are inadequate to determine the individual 

contributions of irradiation temperature, flux, and energy spectrum to the degradation of fracture 

properties in irradiated austenitic SSs. Therefore, additional fracture toughness data should ideally be 

obtained on the same heat of material that has been irradiated in both fast and thermal reactors to 

comparable fluence levels at the same temperature. Only such carefully controlled data can be used to 

accurately assess the applicability of fast reactor data to LWR irradiation conditions. 

Material Type: Most of the J-R curve data on LWR-irradiated austenitic SSs have been obtained on 

Type 304 and 304L SS. Data on Type 316, 316L, 316CW, and 347 SSs are very limited. Also, the only 

data for SS welds are on Type 308L material irradiated to <1 dpa or 12 dpa. Similarly, there are only a 

few J-R curve tests on weld HAZ materials and CF-8M cast SS irradiated to 2.1-2.5 dpa. Some 

differences in the fracture toughness data trends appear for the various grades of wrought austenitic SSs, 

but these differences may be artifacts of the limited data. For example, the heat-to-heat variation for a 

particular grade may be comparable to the apparent differences between grades in the current data. The 

results for different heats of Type 304 SS indicate little or no effect of sensitization treatment (green open 

circles and green-filled right-angle triangles in Fig. 55, respectively). Also, for the same irradiation 

conditions, the fracture toughness of the weld HAZ materials is lower than that of the solution-annealed 

materials (olive-filled diamonds and green open circles in Fig. 55, respectively), and the toughness of the 

thermally aged cast SS is lower than that of the HAZ material (Fig. 57). 

Although the fracture toughness of nonirradiated CW steels is lower than that of nonirradiated 

solution-annealed steels, the decrease in toughness of CW steels with neutron exposure is lower and the 

0

100

200

300

400

500

600

700

0 0.5 1 1.5 2 2.5 3

0.45 dpa

1.35 dpa

3.00 dpa

J (kJ/m2)

Crack Extension (mm)

Type 304 SS

Irradiated at 289°C 

Tested at 289°C

J = 575!a0.17

JIC = 503 kJ/m2

J = 438!a0.33

JIC = 308 kJ/m2

J = 265!a0.29

JIC = 184 kJ/m2

(a)

0

100

200

300

400

500

600

700

0 0.5 1 1.5 2 2.5 3

#2C1 0.065 dpa

#1G1 5.9 dpa

#1D1 13.2 dpa

J (kJ/m2)

Crack Extension (mm)

Type 304 SS 

Irradiated at 295-305°C

Tested at 320°C

J = 85.7!a0.19

JIC = 65 kJ/m2

J = 651!a0.45

JIC = 744 kJ/m2

J = 143!a0.21

JIC = 109 kJ/m2

(b)

Figure 56. Change in fracture toughness JIc as a function of neutron exposure for SSs (Refs. 11,122).
69

JIc value at saturation is higher than those of irradiated solution-annealed steels (Fig. 54). However, the 

data for CW steels are from fast reactor irradiations at relatively high temperatures, 400-427°C (752-

800°F). The saturation JIc for CW SSs is likely to be lower for irradiations at LWR operating 

temperatures [i.e., 290-320°C (554-608°F)], and the differences may be small. 

0

100

200

300

400

500

600

0 0.5 1 1.5 2 2.5 3

J (kJ/m2)

Crack Extension (mm)

Austenitic Stainless Steels

289°C BWR Water

304 Sensitized 

10.5 h @ 600°C

CF-8M 2.46 dpa

304L SA Weld HAZ 

2.16 dpa

304 SMA Weld HAZ 

2.16 dpa

2.16 dpa

0.75 dpa

Figure 57.

Fracture toughness J-R curves for sensitized 

Type 304 SS, weld HAZ materials of Type 304 

and 304L SS, and CF-8M cast SS in highpurity water at 289°C (Ref. 8).

Nonirradiated weld metals and thermally aged cast SSs have lower fracture toughness than wrought 

austenitic SSs, and their fracture toughness generally decreases more rapidly with neutron exposure than 

that of solution-annealed material. However, the saturation toughness for the welds is not significantly 

different from that of solution-annealed SSs, and the same bounding curve for JIc appears to be applicable 

to both wrought materials and welds and cast austenitic SSs. Although LWR core internals are typically 

constructed of CF-8 or CF-3 steels, the only data for LWR irradiation of cast SS are for CF-8M steel. 

The data for thermally aged CF-8 cast SS shown in Fig. 54 are for materials that were irradiated in the 

BOR-60 fast reactor, and may be non-conservative for LWR irradiation conditions. Furthermore, the data 

were obtained at room temperature; as discussed below in this section, fracture toughness at higher 

temperatures is expected to be lower. For thermal embrittlement of cast SSs the fracture toughness of CF8M steel represents the worst-case scenario.102,124 It thus might also represent a bounding case also for 

the synergistic effects of irradiation and thermal aging. 

Test Environment: Nearly all of the existing fracture toughness data have been obtained from tests 

in air and on specimens that were fatigue precracked at relatively low load ratios (typically 0.1-0.2) in 

room-temperature air. However, in reactor core components, cracks are initiated primarily by SCC and 

have IG morphology, whereas the fatigue precracks in fracture toughness tests are always TG. Also, the

corrosion/oxidation reaction could influence fracture toughness. For example, hydrogen generated from 

the oxidation reaction could diffuse into the material and change the deformation behavior by changing 

the stacking-fault energy of the material. 

To investigate the possible effects of the BWR coolant environment on fracture toughness (e.g., the 

effect of the corrosion/oxidation reaction during crack extension or use of specimens with an IG rather 

than the TG fatigue crack generally employed in nearly all fracture toughness tests), J-R curve tests have 

also been conducted in a BWR NWC environment.

11 The J-R curve data on irradiated SS weld HAZ 

materials (Fig. 58) indicate that an NWC BWR environment has little or no effect on the fracture 

toughness. The J-R curves for irradiated Type 304L SA weld HAZ in air and water environments are 

essentially identical (Fig. 58b), and, although the complete J-R curve could not be obtained for Type 304 

SMA weld HAZ in air, ductile crack extension occurred at approximately the same value of J in air and 

water environments (Fig. 58a).
70

The J-R curves for a sensitized Type 304 SS in air and water environments are shown in Fig. 59. 

The results indicate slightly lower fracture toughness in water. Also, the material tested in water was

sensitized for a shorter time than the material tested in air. Therefore, for materials with identical 

sensitization treatment, the difference between the J-R curves in air and water environments may be 

greater than that indicated by Fig. 59. 

0

100

200

300

400

500

600

0 0.5 1 1.5 2 2.5 3

10.5 h at 600°C

24.0 h at 600°C

J (kJ/m2)

Crack Extension (mm)

Type 304 SS (Heat 10285 )

Fluence 1.44 x 1021 n/cm2

290°C 

J = 316!a0.45

JIC = 176 kJ/m2

Open Symbols: Water

Closed Symbols: Air

Heat Treatment

J = 376!a0.38

JIC = 238 kJ/m2

Figure 59.

Fracture toughness J-R curves for 

sensitized Type 304 SS in air and NWC 

BWR water at 289°C (Ref. 8).

The J-R curves and the load vs. load-line displacement curves for the two tests on thermally aged 

and irradiated CF-8M cast SS in NWC BWR water are shown in Fig. 60. Companion tests in air have not 

been conducted on this material. In both tests, large load drops, accompanied by crack extensions up to 

0.5 mm in Specimen 75-11TM and 1.0 mm in Specimen 75-11TT, were observed at the onset of crack 

extension. Such load drops are not typically observed during tests in air.124 The fracture surfaces of 

these specimens have not been examined to establish the fracture morphology. Additional tests on 

irradiated cast SSs or SS welds in air and water environments should be conducted to determine the 

possible effect of LWR coolant environments on their fracture toughness. 

Irradiation Temperature: The available data are inadequate to establish accurately the effects of the 

irradiation temperature on the fracture toughness of austenitic SSs. However, tensile data for austenitic 

SSs indicate that irradiation hardening is highest, and ductility loss is maximum, at an irradiation 

0

50

100

150

200

250

300

350

400

0 0.5 1 1.5 2 2.5 3 3.5

Spec. 85-XB in Air

Spec. 85-XA in BWR Water

J (kJ/m2)

Crack Extension (mm)

Type 304 SS, SMAW HAZ 

Fluence 1.44 x 1021 n/cm2

289°C

Estimated Effective 

Flow Stress: 528 MPa

J = 219!a0.43

JIC = 128 kJ/m2

(a)

0

50

100

150

200

250

300

350

400

0 0.5 1 1.5 2 2.5 3 3.5

GG6T-B in Air

GG6T-A in BWR Water

J (kJ/m2)

Crack Extension (mm)

Type 304L SS, SAW HAZ 

Fluence 1.44 x 1021 n/cm2

289°C Air

Estimated Effective 

Flow Stress: 502 MPa

(b)

Figure 58. Fracture toughness J-R curves for irradiated specimens of (a) Type 304 SS SMA weld HAZ 

and (b) Type 304L SA weld HAZ in air and NWC BWR environment (Ref. 8).
71

temperature of ≈300°C (≈572°F).54 Thus, the JIc values for all of the data at neutron exposures greater 

than 20 dpa shown in Fig. 62 may overestimate the toughness of materials irradiated at temperatures of 

290-320°C (554-608°F) because the irradiation temperatures for these data were above 350°C (662°F). 

0

100

200

300

400

500

600

700

0 0.5 1 1.5 2 2.5 3

J (kJ/m2)

Crack Extension (mm)

CF-8M Cast SS (28% Ferrite)

Aged 10,000 h at 400°C

289°C Air

Open Symbols: 1/4-T CT

Closed Symbols: 1-T CT

Irradiated to 2.46 dpa

Tested in BWR Water

0.0

1.0

2.0

3.0

4.0

5.0

6.0

0 0.5 1 1.5 2 2.5 3

75-11TM

75-11TT

Load (kN)

Displacement (mm)

CF-8M Cast SS (Heat 75, 28% ferrite) 

Aged 10,000 h at 400°C & Irradiated

Fluence 1.63 x 1021 n/cm2

289°C BWR Water

Specimen No.

Figure 60. Fracture toughness J-R curves and load vs. loadline displacement curves for thermally aged 

and irradiated CF-8M cast SS (Ref. 8). 

Test Temperature: The fracture toughness of nonirradiated austenitic SSs is known to decrease as 

the test temperature is increased. The change in the JIc of irradiated SSs as a function of test temperature 

is plotted in Fig. 61 for several grades of SSs and welds irradiated in LWRs and fast reactors. The 

fracture toughness of steels irradiated to relatively low dose (less than 5 dpa) decreases with increasing 

test temperature in most cases. However, for steels irradiated to more than 12 dpa, test temperature has 

little effect on fracture toughness. The data on materials irradiated in LWRs or fast reactors exhibit 

similar trends. It should be noted that, at this fluence level, the toughness value is already low, which 

makes it difficult to discern definitive trends.

0

100

200

300

400

500

600

700

0 100 200 300 400 500

316 SS, 56 dpa, 377–400°C

316CW, 82 dpa, 395–425°C

316L Weld, 3 dpa, 90°C

316L Weld, 3 dpa, 250°C

304L, 12 dpa, 280°C

308L Weld, 12 dpa, 280°C

304, 0.9 dpa, 280°C

308L, 0.7 dpa, 280°C

Fracture Toughness JIc (kJ/m2)

Test Temperature (°C)

Austenitic Stainless Steels

Open Symbols: Fast Reactor

Closed Symbols: LWRs Figure 61.

Fracture toughness JIc of irradiated 

austenitic stainless steels and welds 

as a function of test temperature 

(Ref. 11).

The effect of test temperature is also reflected in the fracture morphology of highly irradiated 

materials. At temperatures above 230°C (446°F) the failure mode is predominantly channel fracture 
72

characterized by a faceted fracture surface. It is associated with highly localized deformation along a 

narrow band of slip planes whereby the initial dislocation motion along the narrow band clears away the 

irradiation-induced defect structure, creating a defect-free channel that offers less resistance to subsequent 

dislocation motion. The localization of the deformation ultimately leads to channel failure. 

3.2 Fracture Toughness Trend Curve

The change in initiation toughness JIc of wrought austenitic SSs (including weld HAZ materials and 

sensitized SSs) and cast SSs and weld metals is shown in Fig. 62 as a function of neutron dose. The 

fracture toughness data from both fast reactor and LWR irradiations are included. The irradiation 

temperatures range from 90 to 427°C (194-800°F) and test temperatures generally from 250 to 427°C 

(212-800°F). Only the data for CF-8 cast SS irradiated in a fast reactor to 10-11 dpa (inverted triangles in 

the figure) were obtained at room temperature. The data in Fig. 61 indicate little or no effect of test 

temperature for materials irradiated to 12 dpa or higher, although the toughness values are already quite 

low above 12 dpa. Also, as discussed previously in Section 3.1, the procedures for determining JIc vary 

0

50

100

150

200

250

300

350

400

0.01 0.1 1 10 100

304

304 HAZ

321

304L

316

316L

316CW

304

304 HAZ

304L

304L HAZ

316L

316CW

304 Sensi.

347

BWR-irradiated 304

 (CNSR specimens)

JIc (kJ/m2)

Neutron Exposure (dpa)

Austenitic Stainless Steels

Irradiation Temp: 90–427°C

Test Temp: 250–427°C

JIc = 7.5 + 110 exp[–0.35(dpa)1.4]

JIc = 7.5

Open Symbols: Fast reactors

Closed Symbols: LWRs

EPRI Curve 

(Demma et al. 2007)

(a)

0

50

100

150

200

250

300

350

400

0.01 0.1 1 10 100

308

316 Weld

316H Weld

CF-8

308L

CF-3

CF-8M

JIc (kJ/m2)

Neutron Exposure (dpa)

Cast Austenitic SSs & Welds

Irradiation Temp: 90–427°C

Test Temp: 250–427°C

JIc = 7.5 + 110 exp[–0.35(dpa)1.4]

JIc = 7.5

EPRI Curve 

(Demma et al. 2007)

(b)

Figure 62. The change in initiation toughness JIc of (a) wrought austenitic SSs and (b) cast austenitic 

SSs and weld metals as a function of neutron exposure. The data points plotted at 0.005 dpa 

are for nonirradiated materials.
73

among these studies. For example, in earlier studies a bilinear J-R curve was used to fit the data, whereas 

a power-law curve was used in the more recent studies. Different expressions have also been used for the 

blunting line. Overall, the results indicate little or no change in toughness below 0.5 dpa, a rapid decrease 

between 1 and 5 dpa, and no further change (saturation) beyond 10 dpa. 

There appear to be some differences in behavior between subsets of the data in Fig. 62a. The 

threshold dose and the dependence of the decrease in fracture toughness JIc on neutron dose seem to vary 

for different grades of materials. The average JIc of the Type 304 SS drops from above 150 kJ/m2

(857 in.-lb/in.2) at 1 dpa to 12-24 kJ/m2 (69-137 in.-lb/in.2) at about 5 dpa. For Type 316L SS the 

decrease appears to occur at a somewhat higher fluence range (3 dpa to 10 dpa), and for Type 304L SS it 

appears to occur at a somewhat lower fluence. Therefore, the decrease in toughness with increasing 

fluence appears earliest in Type 304L SS, followed by Type 304 SS and then Type 316 SS. 

The fracture toughness data in Fig. 62b for cast SSs and welds are lower than those of the wrought 

SSs for all dose levels less than the 10-dpa saturation level. However, the available fracture toughness 

data for irradiated SS welds and cast austenitic SSs are extremely limited. There are no data on any of 

these materials for fluences above 20 dpa, and little or no data on cast austenitic SSs for fluences below 

about 2 dpa. The existing data for welds suggest that ≈0.3 dpa may be considered a threshold neutron 

dose below which irradiation has little or no effect on fracture toughness of SS welds. However, this 

threshold does not consider the possible synergistic effects of thermal and neutron embrittlement of 

welds. 

The existing data for cast austenitic SSs are inadequate to define a threshold dose for irradiation 

effects on fracture toughness. First, there are no fracture toughness data for dose levels of 0.1-2.0 dpa;

second, the potential synergistic interactions of thermal aging and neutron irradiation embrittlement of 

cast austenitic SSs and SS welds124-128 have yet to be addressed. 

Although wrought SSs are typically completely austenitic, welded and cast SSs have a duplex 

microstructure consisting of austenite and ferrite phases. The ferrite phase increases the tensile strength 

and improves resistance to SCC, but it is susceptible to thermal embrittlement after extended service at 

reactor operating temperatures. Thermal aging of cast SSs at 250-400°C (482-752°F) leads to 

precipitation of additional phases in the ferrite (e.g., formation of Cr-rich α$ phase by spinodal 

decomposition; nucleation and growth of α$; precipitation of a Ni- and Si-rich G phase, M23C6 carbide, 

and γ2 austenite; and additional precipitation and/or growth of existing carbides at the ferrite/austenite 

phase boundaries).129-132 The formation of the Cr-rich α$ phase by spinodal decomposition of ferrite is 

the primary mechanism for thermal embrittlement; it strengthens the ferrite phase by increasing strain 

hardening and the local tensile stress. Thermal aging has little or no effect on the austenite phase. Thus, 

thermal aging of cast SSs and SS welds leads to the development of a material with a brittle phase 

dispersed in a ductile matrix. 

Embrittlement of the ferrite phase due to neutron irradiation occurs much faster than for austenitic 

SSs; at reactor operating temperatures of 288-343°C (550-650°F), a shift in the nil-ductility transition 

(ΔNDT) temperature of up to 150°C (302°F) has been observed in pressure vessel steels after neutron 

exposures of 0.07-0.15 dpa (0.5-1.0 x 1020 n/cm2).133 The irradiation temperature is an important factor 

in establishing the extent of embrittlement of ferritic steels. Although both the thermal aging 

embrittlement of ferrite and the neutron irradiation embrittlement of ferrite are well characterized, the 

synergistic effect of thermal aging and neutron irradiation on the embrittlement of SS welds and cast SSs 

has not been investigated yet. The concurrent exposure to high neutron fluence levels could result in a 

synergistic effect wherein the service-degraded fracture toughness is reduced from the levels predicted 
74

independently for either of the two mechanisms. The possible synergistic effects of thermal aging and 

neutron irradiation embrittlement of SS welds and cast SSs are discussed further in Section 3.3. 

The two curves shown in Fig. 62 represent a disposition curve proposed by EPRI58 and a fracture 

toughness trend curve that bounds the existing data. The trend curve takes into consideration: 

(a) a threshold neutron exposure for radiation embrittlement of austenitic SSs and a minimum fracture 

toughness for these materials irradiated to less than the threshold value, (b) a saturation neutron exposure 

and a saturation fracture toughness for materials irradiated to greater than this value, and (c) a description 

of the change in fracture toughness between the threshold and saturation neutron exposures. As shown in 

Fig. 62, the fracture toughness JIc curve that bounds the existing data for JIc as a function of neutron dose 

(in dpa) may be represented by 

JIc = 7.5 + 110 exp[-0.35(dpa)1.4]. (23)

This lower bound curve represents a threshold dose of 0.3-0.5 dpa for neutron embrittlement, a minimum 

fracture toughness JIc of ≈118 kJ/m2 below the threshold dose, a saturation threshold of 5 dpa beyond 

which the fracture toughness of these materials appears to saturate, a saturation fracture toughness JIc of 

7.5 kJ/m2 (or KIc or KJc of 38 MPa m1/2), and a description of the change in toughness between 0.3 and 

5 dpa. The JIc value of ≈118 kJ/m2 below the threshold dose is appropriate for thermally aged and unaged 

cast SSs and SS flux welds. A value higher than 118 kJ/m2 maybe considered for the minimum fracture 

toughness JIc for wrought austenitic SSs irradiated below the threshold dose for neutron embrittlement. 

The description of the change in fracture toughness below 1.5 dpa will change accordingly. The lower 

bound trend curve given by Eq. 23 is consistent with the MRP lower bound model proposed for PWRs.121 

The MRP model is expressed in terms of a lower bound KJc (MPa m1/2) curve. It bounds all the fracture 

toughness data from fast reactors, BWRs, and PWRs as a function of the neutron dose (in dpa) and is 

given by the expression, 

KJc = 180 – 142[1-exp(-dpa)]. (24) 

Both Eqs. 23 and 24 predict a saturation fracture toughness KIc of 38 MPa m1/2. For materials irradiated 

below the threshold dose for irradiation embrittlement, Eq. 23 predicts a minimum KJc of about 

151 MPa m1/2, but the MRP expression predicts fracture toughness values that for some materials, such as 

SS welds or weld HAZ, may be higher than the minimum toughness of the materials in the nonirradiated 

condition. The disposition curve proposed by EPRI for BWRs is not bounding for the existing data for 

BWR-irradiated austenitic SSs. For example, at neutron doses <0.7 dpa, the JIc values based on the EPRI 

curve are higher than the minimum JIc of nonirradiated SS welds (particularly flux welds), some heats of 

wrought SSs, and most thermally aged cast austenitic SSs with >15% ferrite.124 The saturation KIc of 

55 MPa m1/2 at 4.5 dpa for the EPRI curve is also higher than the value of 38 MPa m1/2 previously 

proposed by MRP for PWRs.121 The saturation KIc for the EPRI curve was based on data for which the 

specimen orientation was unknown. Recent data indicate that fracture toughness in the transverse 

orientation is nearly half of that in the longitudinal orientation.58 Therefore, the bounding KIc values 

above 4.5 dpa are likely to be lower than 55 MPa m1/2. Also, as seen in Fig. 62b, some of the data for SS 

welds irradiated to 2-4 dpa are also below the EPRI curve. 

A fracture toughness J-R curve may be used to analyze material behavior for loading beyond JIc. 

The J-R curve is expressed in terms of the J integral and crack extension (Δa) by the power law 

J = C(Δa)n. At dose levels below the threshold dose for saturation (i.e., at dose levels less than ≈5 dpa), 

the effect of neutron irradiation on the fracture toughness of austenitic SSs can also be represented by a 

decrease in the coefficient C of the power-law correlation for the J-R curve with neutron dose. The 
75

variation of C for wrought and cast SSs and welds as a function of neutron dose is shown in Fig. 63. 

Except for the results for CF-3 (closed diamond in Fig. 63b) and CF-8 irradiated in BOR-60 reactor (open 

inverted triangle in Fig. 63b), the remaining data were obtained at temperatures above 250°C. Based on 

the data trends in Fig. 61, test temperature should have little or no effect on the data for CF-8 cast SS. 

However, the constant C for CF-3 cast SS may be more than a factor of two lower at LWR operating 

temperatures. The two curves in Fig. 63 represent the disposition curve proposed by EPRI for BWRs,65

and a trend curve for coefficient C that bounds the existing data.

0

100

200

300

400

500

600

700

0.01 0.1 1 10 100

304

348

316H

316

316H HAZ

316CW

304

304 HAZ

304L

304L HAZ

316L

316CW

304 Sensi.

Power-law Constant C (kJ/m2)

Neutron Exposure (dpa)

Austenitic Stainless Steels

Irradiation Temp: 90–427°C

Test Temp: 250–427°C

C = 25 + 175 exp[–0.35(dpa)1.4] C = 25

Open Symbols: Fast reactors

Closed Symbols: LWRs

EPRI Curve 

(Carter & Gamble 2002)

(a)

0

100

200

300

400

500

600

700

0.01 0.1 1 10 100

308

316 Weld

316L Weld

CF-8

308L

CF-3

CF-8M

CF-8 BOR-60

Power-law Constant C (kJ/m2)

Neutron Exposure (dpa)

Cast Austenitic SSs & Welds

Irradiation Temp: 90–427°C

Test Temp: 250–427°C

C = 25 + 175 exp[–0.35(dpa)1.4] C = 25

(b)

EPRI Curve 

(Carter & Gamble 2002)

Figure 63. Change in coefficient C of the power-law J-R curve for (a) wrought austenitic SSs and 

(b) cast austenitic SSs and weld metals as a function of neutron exposure.

Even for fluence levels above 10 dpa, most heats of wrought austenitic SSs show ductile crack 

extension in the toughness tests. Under similar irradiation conditions, the coefficient C for cast SSs and

welds is lower than that for wrought SSs. However, since most of the data are from irradiations in fast 

reactors and at temperatures of 370-427°C (698-800°F), the values for the power-law coefficient C are 

likely to be lower for irradiations at LWR operating temperatures. As mentioned previously, fracture 

toughness data are limited on materials irradiated in LWRs to neutron dose levels of 0.1-1.0 dpa or above 

10 dpa. Therefore, it is not possible to define accurately the lower bound trend curve for the power-law 

coefficient C as a function of neutron dose. For fluences less than 5 dpa, as shown in Fig. 63, the existing 

fracture toughness data can be bounded by the following expression for C:
76

C = 25 + 175 exp[-0.35(dpa)1.4], (25)

and an exponent n equal to 0.37 (the median value of the experimental data). The exponent n of the 

power-law curve typically ranges from 0.35 to 0.70 for nonirradiated materials and 0.16 to 0.65 for 

irradiated materials. Unlike the behavior for thermally aged cast austenitic SSs, where exponent n 

typically decreases with a decrease in fracture toughness,124 no obvious trend of n with fluence is evident. 

For irradiated materials, the median value of n is 0.37. Equation 25 yields a C value of ≈200 kJ/m2

(1285 in.-lb/in.2) for materials irradiated to less than 0.1 dpa and ≈31 kJ/m2 (≈160 in.-lb/in.2) for 

materials irradiated to ≈5 dpa. These values yield JIc values of 125 and 17 kJ/m2, respectively, for 

materials irradiated to <0.1 and 5 dpa. These values are consistent with the JIc trend curve of Fig. 62. 

The JIc at 5 dpa is also consistent with the data for the CT specimens of Type 304 SS irradiated to 

≈4.5 dpa in a BWR (closed triangles in Fig. 62a).

As noted previously, ductile crack extension was also not observed for some specimens of a 20% 

CW Type 316 SS irradiated to 74-88 dpa in a fast reactor at 410-425°C (770-797°F). The specimens 

failed by a quasi-cleavage fracture believed to be an indirect consequence of the onset of void swelling in 

the material. The KIc values were 74-90 MPa m1/2 (67-82 ksi in.1/2).

3.3 Synergistic Effect of Thermal and Neutron Irradiation

An issue that has been a concern for reactor core internal components is the possibility of a 

synergistic interaction between irradiation embrittlement and thermal embrittlement of cast austenitic SSs 

and SS weld metals. The effect of neutron irradiation on the fracture toughness of wrought and cast SSs 

and welds has been discussed in Section 3.1, and the change in fracture toughness JIc and coefficient C of 

the power-law J-R curve with neutron dose is shown in Figs. 62 and 63, respectively. Thermal aging of 

cast austenitic SSs and SS welds at reactor operating temperatures of 280-350°C (536-662°F) can lead to 

degradation of their fracture properties.102,124-126 Welded and cast SSs have a duplex structure 

consisting of austenite and ferrite phases. Formation of Cr-rich α$ phase in the ferrite is the primary 

mechanism for thermal embrittlement of these materials;124-132 thermal aging has no effect on the 

austenite phase. 

Thermal aging increases the tensile strength, hardness, and Charpy-impact transition temperature, 

and it decreases the ductility, fracture toughness, and impact strength.124-127 For cast austenitic SSs, the 

extent of mechanical-property degradation is essentially determined by the chemical composition of the 

steel, the casting process used to construct the component, the ferrite content and ferrite morphology of 

the steel, and the time and temperature of service for the component.124 Cast SSs with high levels of Mo 

(e.g., CF-8M) show greater susceptibility to thermal embrittlement than steels with low Mo content (e.g., 

CF-3 or CF-8). Static cast steels are more susceptible to thermal embrittlement than centrifugally cast 

components.124 The screening criteria to determine the susceptibility of cast SS components to thermal 

aging embrittlement are outlined in Table 4. 

For cast austenitic SSs, the minimum fracture toughness that can occur due to thermal 

embrittlement depends strongly on the ferrite content and morphology. A globular ferrite morphology in 

which the brittle ferrite phase is isolated in an austenitic matrix will have a higher fracture toughness than 

a lacy morphology in a material with greater than 9% ferrite, where a continuous fracture path through the 

brittle ferrite is possible. The minimum toughness due to thermal aging occurs when the ferrite is fully 

embrittled, and the remaining toughness depends on the toughness provided by the ductile matrix 

surrounding the embrittled phase. Based on an ANL study,124 the predicted saturation J-R curves for the 

various cast SSs in the thermally aged condition (i.e., the lowest fracture toughness that could be achieved 
77

for the steel after thermal aging) are expressed as J ≈ 264 Δa0.35, ≈251 Δa0.34, and ≈167 Δa0.31, 

respectively, for CF-3, CF-8, and CF-8M steels at 290°C (554°F). 

The results for SS welds indicate that the decrease in fracture toughness due to aging depends on 

the ferrite content and initial toughness of the weld.102 Differences in the fracture toughness of SS welds 

arise from differences in the density and size of inclusions in the material. Failure occurs by the 

formation and growth of microvoids near hard inclusions. Welds with relatively high fracture toughness 

(e.g., GTA or tungsten inert gas weld) show a significant decrease due to thermal aging whereas welds 

with poor fracture toughness (e.g., SA, SMA, or manual metal arc welds) show minimal change. In the 

latter, failure primarily occurs by the formation and growth of microvoids. Such processes are relatively 

insensitive to thermal aging. The existing data indicate that at 280-350°C, the fracture toughness JIc of 

thermally aged welds can be as low as 40 kJ/m2. A conservative estimate of the J-R curve for aged SS 

welds102 is given by J = 40 + 83.5 Δa0.643.

Reactor core internal components are subject to thermal and concurrent exposure to neutron 

irradiation. This condition could result in a synergistic effect wherein the service-degraded fracture 

toughness can be less than that predicted for either thermal embrittlement or neutron irradiation 

embrittlement independently.

For license renewal, to account for the effects of thermal aging and neutron irradiation 

embrittlement on the fracture toughness of reactor core internal components, the NRC staff has proposed 

that for cast austenitic SS components that have a fluence of greater than 1 x 1017 n/cm2 (0.00015 dpa) or 

are determined to be susceptible to thermal aging embrittlement, an aging management program should 

be implemented. This program should consist of either a supplemental examination of the affected 

components as part of the applicant’s 10-year inservice inspection program during the license renewal 

term, or a component-specific evaluation to determine the susceptibility to loss of fracture toughness.134 

Furthermore, the program should consider the synergistic loss of fracture toughness due to neutron 

irradiation embrittlement and thermal aging embrittlement. 

An EPRI report on thermal aging embrittlement of cast SS components proposed the use of the J 

value at a crack extension of 2.5 mm (0.1 in.), J2.5, to differentiate between nonsignificant and potentially 

significant reductions in fracture toughness of cast austenitic SSs.135 Flaw tolerance evaluations were 

presented in Appendices A and B of the EPRI report to support the choice of a threshold value of 

J2.5 = 255 kJ/m2 (1456 in.-lb/in.2). The NRC staff has found that using J2.5 = 255 kJ/m2 is an acceptable 

screening approach for fracture toughness of cast SSs.134 This concept can be extended to irradiated 

materials. However, the applicability of the flaw tolerance evaluations in Appendices A and B of the 

EPRI report would have to be demonstrated to support the use of the J2.5 parameter for evaluating the 

toughness of irradiated SSs.

Table 4. Screening criteria for thermal-aging susceptibility of cast austenitic stainless steels (Ref. 134).

Mo Content (wt.%) Casting Method Ferrite Content Susceptibility Determination

High (2.0-3.0) Static ≤14% Not susceptible

>14% Potentially susceptible

Centrifugal ≤20% Not susceptible

>20% Potentially susceptible

Low (0.5 max.) Static ≤20% Not susceptible

>20% Potentially susceptible

Centrifugal All Not susceptible
78

For the wrought and cast austenitic SSs and welds listed in Fig. 63, the experimental J-integral 

values at a crack extension of 2.5 mm are plotted as a function of neutron dose in Fig. 64. The solid curve 

in Fig. 64 represents the predicted values of J at 2.5-mm crack extension that are expected to bound the 

existing experimental data shown in Fig. 63. The curve was obtained using the power-law J-R curve 

relationship, with coefficient C determined from Eq. 25 and an exponent n of 0.37. The lower bound 

curve indicates that for cast SSs and welds irradiated up to 1.0 dpa, the predicted J at 2.5 mm is above the 

screening value of 255 kJ/m2 (1456 in.-lb/in.2). However, additional fracture toughness data on irradiated 

SS welds and cast SSs, particularly at 0.1-2.0 dpa, are needed to better define the threshold dose when the 

fracture toughness of austenitic SSs begins to significantly decrease.

This evaluation does not consider the synergistic interaction of neutron irradiation embrittlement 

and thermal aging embrittlement. Embrittlement of ferrite phase from neutron irradiation occurs at lower 

dose levels than does embrittlement of the austenite phase. A shift in the NDT temperature of up to 

150°C (302°F) has been observed in pressure vessel steels irradiated to 0.07-0.15 dpa.133 Thus, 

embrittlement of ferrite is expected to occur at 0.05-1.0 dpa, whereas, as discussed in Section 3.1, any 

0

200

400

600

800

1000

0.01 0.1 1 10 100

304

304 HAZ

321

304L

316

316L

316CW

304

304 HAZ

304L

304L HAZ

316L

304 Sensi.

J at 2.5 mm Crack Extension (kJ/m2)

Neutron Exposure (dpa)

Austenitic Stainless Steels

Irradiation Temp: 90–427°C

Test Temp: 100–427°C

Open Symbols: Fast reactors

Closed Symbols: LWRs

(a)

J = 255 kJ/m2 Lower bound of the

existing data

0

200

400

600

800

1000

0.01 0.1 1 10 100

308

316 Weld

316H Weld

CF-8

308L

CF-3

CF-8M

Neutron Exposure (dpa)

Cast Austenitic SSs & Welds

Irradiation Temp: 90–427°C

Test Temp: 25–427°C

(b)

Open Symbols: Fast reactors

Closed Symbols: LWRs

J = 255 kJ/m2

Lower bound of the

existing data

J at 2.5 mm Crack Extension (kJ/m2)

Figure 64. Experimental values of J-integral at a crack extension of 2.5 mm for (a) wrought austenitic 

SSs and (b) cast austenitic SSs and weld metals plotted as a function of neutron exposure. 
79

significant effect of neutron irradiation on embrittlement of the austenite phase occurs only above 

≈0.5 dpa (Fig. 62). 

In addition to possibly altering the threshold dose for neutron embrittlement, synergistic effects of 

neutron irradiation and thermal aging embrittlement could decrease the saturation fracture toughness of 

irradiated welded and cast SSs and accelerate the change in fracture toughness between the threshold and 

saturation neutron exposures. Unfortunately, limited data are available to assess such effects. Figure 64

shows the results of two tests on a CF-8M steel that was thermally aged for 10,000 h at 400°C and then 

irradiated to well above the threshold dose for neutron embrittlement (inverted triangles in Fig. 64). The 

resulting toughness is bounded by the curve for other SSs irradiated to a similar level, i.e., thermal aging 

does not seem to lower the toughness below that expected for irradiation alone at these neutron dose 

levels. 

Additional fracture toughness data are needed to better establish the potential for synergistic loss of 

toughness in these materials in the transition dose range from 0.05 to 2 dpa. A program is being 

conducted at ANL on fracture toughness characterization of CF-3, CF-8, and CF-8M cast SS and SS 

welds irradiated in the Halden reactor in Norway at 315°C and doses of 5 x 1019 n/cm2 (0.075 dpa) and 

2 x 1021 n/cm2 (3.0 dpa). Although cast CF-8M SSs are not used in LWR core internal components 

because of the difficulty of testing irradiated materials, it may be helpful to study this material as a 

“worst-case” in lieu of testing a number of heats of CF-3 and CF-8. 
80

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81

4. Void Swelling 

4.1 Background Information

Void swelling refers to the volume change of materials under neutron irradiations. Void swelling 

was first observed in the late 1960s in austenitic SSs irradiated in fast reactors to neutron exposures above 

1 x 1023 n/cm2 at temperatures of 370-650°C (698-1202°F).17 As discussed in Section 2.1.1, neutron 

irradiation damages materials by displacing atoms from their lattice position. Such displacements create 

vacancies and interstitials, most of which are annihilated by recombination. The surviving defects lead to 

microstructural changes as they rearrange into more stable configurations. Because of the relatively large 

strain field that surrounds interstitials, there is a strong interaction between interstitials and dislocations, 

which results in a preferential flux of interstitials toward dislocations. The remaining vacancies cannot 

annihilate by recombining with interstitials, and this condition leads to the nucleation of cavities or 

microvoids. Under certain conditions of temperature and dose rate, these cavities or microvoids can 

eventually grow to larger sizes. Fission products such as He and H also play an important role in void 

formation. By combining with these gas atoms, void nucleation is facilitated through reduction of the 

surface energy of vacancy clusters. The fundamental driving force of void formation, however, remains 

the excess vacancy flux toward voids. 

The formation of a large number of voids results in an increase in the volume of the material; this 

process is referred to as void swelling. The volume changes from void swelling are normally isotropic 

and occur in all directions. However, there are always constraints on swelling, both internally due to 

gradients of temperature and dose rate within the component, and externally from neighboring 

components. Such constraints result in anisotropic swelling and stress fields within the components that 

activate irradiation creep in unconstrained directions, producing dimensional distortions and misfits that 

can compromise the functional integrity of the reactor core internal components.

Most void swelling data have been obtained from materials irradiated in fast breeder reactors at 

temperatures above 385°C (725°F) and at dose rates that are orders of magnitude higher than those in 

PWRs, and extrapolation of these results to estimate the void swelling behavior for PWR end-of-life or 

extended life conditions introduces substantial uncertainties. 

The void swelling process is divided into two regimes: a transient regime followed by a steady-state 

swelling rate. All materials regardless of material composition and thermo-mechanical condition, stress 

state, neutron flux and spectrum, or irradiation temperature (above 300°C) are believed to reach a steadystate swelling rate of ~1%/dpa.17 However, the duration of the transient regime varies with the various 

material, stress, and irradiation parameters. In general, higher neutron flux and stress accelerate, and cold 

work in the material delays, the onset of the steady-state swelling rate. To achieve steady-state swelling 

rate within the life time of PWRs may require irradiation tempertures above 400°C. Significant effects of 

various material and environmental parameters on void swelling are as follows:

Irradiation Temperature: In the temperature range for fast reactors, void swelling shows a strong 

dependence on irradiation temperature. For 20% CW Type 316 SSs a steady-state swelling rate of 

1%/dpa is observed in fast reactor data at 427°C (800°F) or higher temperatures.18 As discussed in 

Section 2.1.1, voids typically appear in austenitic SSs at temperatures above 340°C. The PWR coolant 

temperatures do not exceed 345°C (653°F), although gamma heating of thick section components near the 

reactor core could increase the local temperature within the component to 370°C (698°F) or even higher. 

The lower irradiation temperatures, typical of LWRs, extend the transient regime to much longer times 
82

and result in the formation of a higher density of smaller voids. However, depending on the irradiation 

temperature, dose, and dose rate, voids can form under PWR operating conditions within the reactor 

lifetime. Voids have been observed in austenitic SSs irradiated at temperatures as low as 300°C 

(572°F).136 

Material Composition: Reactor core internal components are primarily constructed of austenitic 

SSs. These materials have a face-centered cubic crystal structure and are more susceptible to void 

swelling (i.e., swelling occurs earlier and at a higher rate) than ferritic steels, which have a body-centered 

cubic crystal structure. The material composition can influence the transient regime for void swelling. 

The most important factor is the Ni content in the steel. In the 300-series SSs, swelling decreases rapidly 

with increasing Ni content up to about 30-40%.137 In contrast, Cr has the opposite effect; it decreases the 

transient regime and increases swelling.137 However, at low temperatures typical of LWRs, swelling is 

not expected to change significantly over the range of Cr found in the 300-series SSs used for nuclear 

service. Because of compositional differences, Type 304 SS typically swells more than Type 316 SS 

under similar thermal-mechanical conditions. 

The effect of minor trace elements on void swelling is sensitive to material and environmental 

conditions and, therefore, is less predictable. Similar to the effect of Ni, P contents above 0.01 wt.% and 

Si contents above 0.2 wt.% decrease swelling in the 300-series SSs.138,139 The effect of these solutes at 

small concentrations may be different at lower temperatures typical of PWRs. However, since most SSs 

used for the construction of LWR core internals typically contain 0.03-0.04 wt.% P and ~0.5 wt.% Si, 

these elements are expected to decrease void swelling. Both change the effective diffusivity of vacancies 

and thereby the vacancy supersaturation. Minor differences in the concentration of these minor elements 

and their distribution due to differences in thermal-mechanical history may explain the large differences 

in the swelling behavior sometimes observed in essentially identical heats of SSs.101

Phase Changes during Irradiation: As discussed in Section 2.1.1, irradiation at higher 

temperatures, particularly above 340°C, leads to the formation of second phase particles. The radiationinduced formation of γ’ silicide (Ni3Si), phosphides (M2P and M3P), and G phase (M6Ni16Si7) can 

remove Ni, P, and Si from the alloy matrix, and thereby increase void swelling. However, radiationinduced precipitation of second phases is strongly dependent on both temperature and displacement 

rate.17

Material Condition: Cold work in the material prolongs the transient regime for void swelling in 

austenitic SSs. The high dislocation density produced by the cold work provides additional 

recombination sites for vacancies and interstitials, thereby decreasing supersaturation of vacancies in the 

material and delaying void nucleation. The high dislocation density also interferes with diffusion of 

minor elements and, therefore, delays or prevents the formation of second phase particles, which remove 

elements (such as Ni, Si, or P) that are known to suppress void swelling. However, the benefit of cold 

work is short lived because absorption of vacancies increases the mobility of the dislocations, permitting 

them to interact and decline in density to levels similar to those achieved in solution-annealed materials 

with comparable irradiation doses. In solution-annealed Type 304 SS, voids nucleate easily, and the 

transient regime is characterized by a slowly increasing rate of swelling with increasing neutron dose. In 

CW Type 316 SS, void nucleation is difficult or does not occur for some time, but the transient regime 

ends, and the steady-state swelling rate starts abruptly.18 

Stress and Stress History: In general, stress accelerates void swelling and decreases the transient 

regime. The sign of the stress state, tensile or compressive, does not matter; it is the shear component and 

not the hydrostatic component that accelerates swelling.17 As discussed above, constraints due to 
83

gradients in temperature and/or dose rate and from adjoining components lead to anisotropic swelling and 

buildup of stresses within the component. These stresses, however, are limited by irradiation-induced 

creep and will relax in time; the rate of relaxation is proportional to the swelling rate. Therefore, the 

swelling-induced stresses often vary with time, increasing due to swelling and component constraints, and 

decreasing with creep relaxation. The interaction of swelling and irradiation creep relaxation plays an 

important role in evaluation of swelling in PWR core internal components. The deformation due to 

irradiation creep is in the direction that relieves the shear component of the applied stress. In a bolt, the 

primary stress is tensile along the bolt axis. Thus, the bolt would creep to increase its length and decrease 

the diameter. However, if the plate in which the bolt is embedded swells at a greater rate, the stress in the 

bolt would be reestablished. In the case of the hot spot in the former plate behind the reentrant baffle 

plate corner, the local stresses that are generated because of external and internal constraints would 

redirect swelling in the vertical direction.101 Another effect of stress is that a high swelling region within 

a component would produce stresses in an adjacent lower swelling region, thereby increasing their 

swelling rate. Therefore, calculations of swelling based on stress-free material would tend to 

overestimate swelling gradients in a component. 

Displacement Rate: Void swelling data on annealed Type 304L SS irradiated to ~30 dpa at 390°C 

in fast reactors indicate that swelling increases as the displacement rate is decreased.140 This finding 

suggests that estimates of void swelling based on high flux fast reactor data could yield nonconservative 

results for PWR core internal components. In other words, under comparable irradiation temperatures 

and neutron dose levels, the lower displacement rates typical of PWRs could result in higher swelling. 

However, the exact mechanism by which displacement rate affects swelling is not well understood. 

Initially, it was believed to be due to an upward shift in the temperature range for void swelling with 

increases in displacement rate. Such a temperature shift is observed under ion bombardment in pure 

metals.141 At low temperatures, the displacement rate effect is believed to arise from displacement-ratedependent changes in point defect concentration and their effect on vacancy-interstitial recombination. At 

high temperatures, the higher production rate of vacancies and interstitials at the high dose rates is 

balanced by the enhanced defect migration rate so that the point defects that survive recombination are 

comparable in number to those surviving from the low displacement rate neutrons. Such an argument 

explains the change in swelling in terms of a change in void growth rate rather than a change in void 

nucleation. However, several studies on austenitic SSs irradiated in EBR-II at 373-444°C (703-831°F) 

indicate that the effect of displacement rate on swelling is due to its effect on the transient regime and not 

void growth.142 The data also indicate that the displacement rate effect in solution-annealed materials is 

similar to that in cold-worked materials.

Helium Production: Under neutron irradiation, He can be produced by the 10B reaction to form 4He 

and 7Li, and also from the two-step production of 59Ni from 58Ni by thermal neutrons [i.e., 

58Ni(n,γ)59Ni(n,α) 56Fe]. In the latter, the He generation rate is proportional to the 59Ni content and the 

thermal neutron flux. In PWRs, transmutations can produce both He and H. Hydrogen is produced from 

the two-step reaction 58Ni(n,γ)59Ni(n,p) 59Co with thermal neutrons as well as from other nontransmutation sources. The nucleation of microvoids due to migration and condensation of vacancies 

produced by neutron displacements can be stabilized by the transmuted gas atoms (He or H), thereby 

increasing void swelling. A fine dispersion of He bubbles has been observed in SSs irradiated in LWRs 

at 300-340°C, especially in reactors with high thermal flux.48,51 However, the effect of He on swelling is 

not as strong as that on Ni or P contents in the steel. Most of the earlier studies have focused on the 

effects of He on swelling. Studies on the effects of H are limited.
84

4.2 Assessment of Void Swelling in PWR Core Internal Components 

The void swelling measured in PWR internal components, such as the flux thimble tubes and baffle 

bolts,20,143,144 is plotted as a function of neutron dose in Fig. 65. Except for two data points that show 

0.2-0.25% swelling, void swelling in these SSs irradiated up to 80 dpa is insignificant. However, the 

irradiation temperature for the low-swelling materials was less than 330°C (626°F). A review of the 

existing void swelling data on austenitic SSs irradiated in fast reactors and the limited data available on 

PWR-irradiated materials indicated that they all can reach the steady-state swelling rate of ~1%/s. The 

duration of the transient regime to reach this steady-state rate depends on material composition and 

thermal-mechanical condition, stress state, irradiation temperature, and neutron displacement rate. The 

state of the knowledge on the effects of these variables is discussed next. 

1. Temperature: Although PWR coolant temperatures do not exceed 345°C (653°F), gamma heating can 

increase the local temperature in thick sections of reactor core components to values as high as 420°C 

(788°F). However, as shown in Fig. 66, most of the void swelling data on PWR-irradiated materials 

0.00

0.05

0.10

0.15

0.20

0.25

0.30

0 20 40 60 80 100

Ringhals Flux thimble tube, CW316

Farley/Point Beach Lock bar, SA304

Tihange 1 Baffle bolt, CW316

Point Beach Baffle bolt, SA347

Japan PWR Flux thimble tube, CW316

Farley Baffle bolt, CW316

Void Swelling (%)

Neutron Dose (dpa)

Figure 65. Void swelling of PWR materials plotted as a function of neutron dose (Refs. 20,143,144).

280

300

320

340

360

380

0 20 40 60 80 100

Ringhals Flux thimble tube, CW316

Farley/Point Beach Lock bar, SA304

Tihange 1 Baffle bolt, CW316

Point Beach Baffle bolt, SA347

Japan PWR Flux thimble tube, CW316

Farley Baffle bolt, CW316

Temperature (°C)

Neutron Dose (dpa)

Figure 66. Irradiation temperature and neutron dose for the void swelling data shown in Fig. 65.
85

are at temperatures below 330°C, where swelling is minimal. There are no swelling data at 370-

420°C (698-788°F), where swelling is expected to be the highest. Estimates of void swelling in PWR 

components are based on materials irradiated in fast reactors at temperatures above 370°C (698°F). 

2. Displacement Rate: Typical displacement rates in PWRs are 0.3-9.0 x 10-8 dpa/s, whereas rates used 

in the void swelling studies are 0.3-10.0 x10-7 dpa/s. Void swelling data on annealed and CW 

austenitic SSs irradiated in fast reactors at 373-444°C (703-831°F) indicate that swelling increases as 

the displacement rate is decreased. The displacement rate effect is primarily due to its effect on the 

transient regime. In other words, the dose required to achieve the 1%/dpa steady-state swelling rate is 

lower for austenitic SSs irradiated at the lower dose rate. These results suggest that under comparable 

irradiation temperatures and neutron dose levels, the lower displacement rates typical of PWRs would 

result in higher swelling. Limited void swelling data on PWR-irradiated materials are consistent with 

this prediction, which is based on fast reactor data at temperatures relevant for PWR internals and at 

displacement rates below 5.0 x 10-8 dpa/s. 

3. Stress State: In general, irrespective of whether it is tensile or compressive, applied stress or swellinginduced stress accelerates void swelling. Secondary stresses due to thermal gradients and geometric 

constraints, however, are limited by irradiation creep and relax rather quickly. Therefore, the 

interaction of swelling and irradiation creep relaxation plays an important role in evaluation of 

swelling in PWR core internal components. Also, another stress-related effect is that stresses 

produced by a high swelling region in an adjacent lower swelling region force the lower swelling 

region to try to catch up with the faster growing regions. Thus, estimates of swelling based on stressfree material would tend to overestimate swelling gradients in a component. 

4. Material Composition and Condition: The data on LWR-irradiated materials are inadequate to 

establish the differences in the swelling behavior of austenitic SSs under LWR and fast reactor 

irradiation conditions. The material composition can influence the transient regime for void swelling. 

In general, an increase in Ni, P, or Si content decreases swelling by increasing the transient regime. 

In contrast Cr has an opposite effect. Under PWR irradiation conditions, void swelling is not 

expected to change significantly with minor differences in the Cr content. Cold work also increases 

the transient regime for swelling but does not affect the steady-state swelling rate. Based on the fast 

reactor data on CW Type 316 SS, which indicate that the steady-state swelling rate of 1%/s is not 

reached after very high fluence and a temperature of 400°C (750°F), Byrne et al. concluded that it is 

unlikely that the 1%/dpa swelling rate would ever be attained in commercial PWR core internal 

materials.18 However, their assessment is for 60-year operation; a steady-state swelling rate could 

eventually be attained in PWR internal materials during extended operation beyond 60 years. The

differences between LWR and fast reactor irradiations need to be investigated.

5. Phase Change: As discussed above, radiation-induced formation of second phases can remove 

elements such as Ni, P, and Si, which are known to decrease swelling, from the alloy matrix. This 

effect increases void swelling. The precipitation of second phases is strongly dependent on material 

composition and thermal-mechanical condition as well as temperature, dose, and dose rate. However, 

another aspect of precipitation of second phase particles may tend to decrease swelling. Typically, 

irradiation-induced precipitation at PWR operating temperatures is extremely small (i.e., a few to 

several tens of nanometers), and the precipitate density is large. Therefore, at low temperatures, 

precipitation creates very large internal surfaces (i.e., interfaces between the precipitates and the 

metal matrix), which would act as sinks for vacancies, thereby suppressing the accumulation and 

condensation of vacancies to form voids. This behavior has been well known to occur in vanadium 
86

alloys.20 The role of precipitation on void swelling in austenitic SSs under PWR irradiation 

conditions needs to be investigated. 

Most of the void swelling data on LWR-irradiated materials have been obtained on CW Type 316 

SS, and data on solution-annealed Type 304 SS are limited. As discussed above, because of the higher Ni 

content, void swelling in Type 316 SS is less than that in Type 304 SS. Also, cold work further reduces 

swelling in CW Type 316 SS. Therefore, additional data are needed on PWR-irradiated annealed 

Type 304 SS. Also, it is important to develop a technical basis for extrapolating the fast reactor data and 

the limited PWR data to predict void swelling in PWR core internal components. Although extrapolation 

of these results suggests that void swelling will not be a significant problem during the first license 

extension period, investigations are still needed on the effects of irradiation temperature (particularly 

above 370°C), displacement rate, and stress on void swelling, including the interaction of swelling and 

irradiation creep relaxation.

Stress-free swelling equations have been developed by the MRP for annealed Type 304 and CW 

Type 316 SSs used in PWR core internal components.63 The swelling rate SR (%/dpa) is expressed in 

terms of the temperature T (°C), neutron dose φ (dpa), and dose rate !

! (10-7 dpa/s). The swelling rate for 

annealed Type 304 SS is given by 

0.731

SA304

18558 SR 2 exp 22.106

T 273.15

! " # = $$ % ! & ' + (

! , (26)

and the swelling rate for CW Type 316 SS by 

( ( )) 0.731

CW316

18558 SR 0.9 10 2 1 exp 0.01 exp 22.106

T 273.15

! " # = $

) + & ! ! & %*& ' ! ( + + ,

! . (27)

0

1

2

3

4

5

0 10 20 30 40 50 60

Void Swelling (%)

90% EFPY

Core Barrel-toFormer Bolt

CW Type 316 SS

355°C, 8.4E-09 dpa/s

Core Barrel 60 EFPY 14.3 dpa

Annealed Type 304 SS

340°C, 7.3E-09 dpa/s

60 EFPY 12.4 dpa

Former Plate

Annealed Type 304 SS

355°C, 5.6E-08 dpa/s

60 EFPY 95.4 dpa

20 dpa

40 dpa

60 dpa

Figure 67. Void swelling of PWR materials plotted as a function of effective full 

power years (Refs. 20,143,144).
87

Equation 27 predicts a slightly higher swelling rate for Type 316 SS than that for annealed Type 304 SS 

up to 10 dpa and a lower swelling rate for higher dose levels. The predicted void swelling for Type 304 

SS core barrel and former plate and for CW Type 316 SS core barrel-to-former bolt for a 60-y life at 90% 

EFPY is shown in Fig. 67. Note that the effect of temperature is much stronger than that of dose rate. 

For the core barrel, although the dose rate is low, the 60-y end-of-life void swelling is quite low because 

the total dose is only 12.4 dpa. 

An expression to correlate stress-enhanced swelling has also been developed. The incremental stressenhanced swelling, ΔS (%), is given by 

!S = S"!#(1+ 0.005$), (28)

where S! is stress free swelling rate SRSA304 or SRCW316, !" is the increment of neutron dose, and ! is 

the von Mises effective stress (MPa). As discussed previously, any stress state, negative or positive, will 

accelerate the swelling rate. 

Equations 26-28 are based on isothermal irradiation, and the maximum swelling rate is limited to 

1%/s. 

4.3 Embrittlement due to Void Swelling 

As the volume of voids in the material continues to increase beyond 3-5%, the size and distribution 

of voids essentially control the fracture properties of the material. Several studies have shown that 

austenitic SSs irradiated at ≤400°C to neutron dose levels that produce ≥10% void swelling suffer from 

severe embrittlement.144-147 Furthermore, the critical swelling level required to induce an extremely 

brittle state decreases with decreasing irradiation temperature.145 A Type 316 SS irradiated in the EBR-II 

fast reactor at 400°C to 130 dpa, producing 14% void swelling, fractured during handling at room 

temperature.146 Also, a Russian SS, EI-847, irradiated in the BN-350 fast reactor in the annealed 

condition to 73 dpa at 335°C and in the CW condition to 82 dpa at 365°C had more than 10% void 

swelling, and was found to be exceptionally brittle.147 

In austenitic SSs with ≥10% void swelling, the high embrittlement has been attributed to a process 

that involves stress concentration between voids, Ni segregation at void surfaces, and a tendency toward 

martensite formation when the steel is deformed at room temperature.145 As discussed previously, the 

removal of Ni and Si from the metal matrix to form γ’ or G phase promotes void nucleation and swelling. 

Once voids form, Ni segregates to void surfaces due to the inverse Kirkendall effect, in which the slowest 

diffusing elements segregate. This further depletes Ni from the matrix. The loss of Ni and relative 

increase of Cr in the matrix in the region between the voids decrease the SFE substantially, which 

promotes planar slip and formation of martensite during room-temperature deformation. The 

austenite/martensite boundaries provide a low-energy crack propagation path, resulting in a brittle, quasicleavage fracture at room temperature.146,147 At higher temperatures, the effect of Ni segregation on SFE 

is much less pronounced, and large levels of flow localization are required for failure, which occurs along 

the flow localization path, and is referred to as channel fracture.

115

There are few quantitative data correlating the void swelling with fracture toughness of the 

material. A Fe-14Cr-16Ni-1.5Mo SS (D9) irradiated in the Fast Flux Test Facility (FFTF) at 390°C 

(734°F) to 114 dpa showed a KIc of about 30 MPa m1/2 measured at 410°C and zero tearing modulus.148

Although the void swelling was not measured for this material, it is expected to be above 10%. Swelling 
88

in the same material irradiated at 410°C to 175 dpa showed about 30% void swelling, and 20% CW 

Type 316 SS irradiated at 410°C to 170 dpa showed 18-23% swelling. 

The fracture toughness of highly irradiated SSs with >10% swelling is lower than that observed in 

Type 304 or 316 SS irradiated to 4.5-8.5 dpa in BWRs (at 290°C). As discussed in Section 3.1.2, a KIc of 

55 MPa m1/2 has been reported for BWR control blade material irradiated to 4.5 dpa, and 37 MPa m1/2 in 

material irradiated to about 8 dpa. The latter value is currently defined as the saturation fracture 

toughness for irradiated austenitic SSs. The results for high swelling suggest that the fracture toughness 

of PWR core internal components of austenitic SS could be lower than the saturation value 

(KIc=38 MPa m1/2) proposed by MRP.121 The embrittlement of austenitic SSs due to void swelling 

under PWR irradiation conditions should be further investigated. 
89

5. Stress Relaxation and Creep

Loss of preload for bolted joints and redistribution of stresses in components due to stress 

relaxation are other aging degradation concerns that need to be addressed to assure the functional integrity 

of the reactor core internal components. Stress relaxation represents plastic deformation that occurs with 

time under constant strain below the yield point of the material. In other words, stress relaxation may be 

considered as creep that occurs under constant strain instead of constant load or stress. It can occur either 

by thermal creep21 or neutron-irradiation-assisted creep.23,24 

Thermal creep is represented by primary (or transient), secondary (or steady-state), and tertiary 

creep stages. For most metals and alloys, at temperatures below 0.4 or 0.5 of their melting temperature 

(in degrees kelvin), creep under constant load is characterized by only primary creep with decreasing 

creep rates, and steady-state creep is insignificant. For the 300-series austenitic SSs and Ni alloys used in 

LWR core internal components, this temperature is about 394°C (740°F), which is higher than the typical 

temperature range for PWR core components. Only a few locations close to the core may be above this 

temperature due to gamma heating. Thus, for PWR internals, the contribution to total strain due to 

steady-state thermal creep under constant load is likely to be relatively small. 

Stress relaxation data are generally obtained from tests performed in accordance with ASTM 

Specification E 328-86 under tension, compression, torsion, or bending to reflect the stress state imposed 

on the component. Most of the available data are for stress relaxation tests that have been conducted 

under tension (for bolted joints) and bending (for leaf springs). The stress state behavior of coil springs is 

more complex, and there are no stress relaxation data for coil springs in the open literature. In general, 

the thermal stress relaxation of coil springs appears to be much larger and depends on the coil design 

geometry.101 

The existing data on thermal stress relaxation have been obtained from short-term tests (less than 

1000 h) at 260-705°C (500-1300°F) on Type 304 SS as a function of cold work, thermo-mechanical 

condition, and stress level.22 The results indicate that the percentage of stress relaxation increases with 

increasing initial stress and levels off at about 207 MPa (30 ksi). At PWR operating temperatures, stress 

relaxation saturates within 100 h with a maximum decrease of 10-20% of the initial preload stress

depending on the thermal-mechanical treatment of the material.101. At temperatures below 315°C 

(600°F), for the same initial stress, the extent of stress relaxation decreases with increasing amount of 

cold work in the material. The effect of cold work is less significant at temperatures above 371°C 

(700°F). Stress relaxation reaches saturation in the short-term tests because only two components, 

anelastic strain and microplastic strain, contribute to the total time-dependent strain (or creep strain). 

Secondary or steady-state creep is insignificant in these tests at relatively low temperatures (i.e., less than 

394°C). However, over the reactor lifetime, even a rather low steady-state creep rate could lead to greater 

stress relaxation than anticipated from the short-term tests. For example, at 350°C (662°F), the steadystate creep rate for Type 304 SS is estimated to be 5.2 x 10-13 /s, which corresponds to 0.049% strain over 

30 y or a total stress relaxation of 76.5 MPa (11.1 ksi).21 

Neutron-irradiation-enhanced creep can greatly increase the plastic strain by increasing both the 

primary creep and steady-state creep rates. Unlike thermal creep, neutron irradiation can result in 

significant creep strain at PWR temperatures. Irradiation creep can be divided into three stages: transient, 

steady-state, and void-swelling-induced creep. The magnitude of both transient creep and steady-state 

creep is essentially proportional to the applied stress level. Transient creep is typically short and, 

although its magnitude depends on material condition and structure and irradiation temperature, the 
90

contribution to total creep strain is relatively small (e.g., PWR fuel rods show <0.05% transient creep);17

the steady-state creep rate is independent of these parameters. Figure 68 plots the thermal creep and 

irradiation-enhanced creep of 20% CW Type 316 SS in the EBR-II fast reactor at 454°C (850°F) and 

under an uniaxial stress of 138 MPa (20 ksi).

0

2

4

6

8

10

12

0 500 1000 1500 2000 2500

Strain (10-4)

Time (h)

20% CW Type 316 SS

138 MPa (20ksi)

454°C (850°F)

Irradiation-enhanced creep

Thermal exposure only

Figure 68. Thermal creep and irradiation-enhanced creep of a 20% CW Type 316 stainless steel 

irradiated in EBR-II (Ref. 101).

Empirical models have been developed for irradiation creep using data obtained from fast reactors. 

Because of the relatively weak dependence of steady-state creep on dose rate or temperature, the fast 

reactor data have been used to estimate irradiation creep at PWR operating temperatures.101 The existing 

data are inadequate to determine the effects of cold work or material composition on irradiation-enhanced 

steady-state creep. 

Initially, void-swelling-induced creep was not considered to be significant for PWR core internal 

components. However, as discussed in Section 4.2, void swelling could be significant for PWR core 

internals because of the low dose rates, which could lead to considerable creep strain. A detailed 

evaluation based on typical temperature/dose rate information for core components is needed to estimate 

the contribution of void-swelling-induced creep. 

As discussed previously, stress relaxation is creep that occurs at constant strain. It follows the 

trends observed in thermal and irradiation creep. Initially, there is a sharp decrease in stress representing 

transient creep, followed by a gradual decrease in stress due to essentially steady-state irradiation creep. 

Empirical equations have been developed from the irradiation creep data to estimate stress relaxation as a 

function of initial stress and neutron dose, and the constants in the equations were derived by regression 

fitting to the available stress relaxation data. One such curve representing stress relaxation (%) as a 

function of neutron dose is given by,101

!

!o

= exp{"0.774(1" exp("23f)) " 0.688f} , (29)
91

where f is the accumulated neutron fluence in dpa, and ! and !o are the remaining and initial stress. 

The first term [i.e., 0.774(1 - exp(-23f))] in the equation represents stress relaxation due to irradiationenhanced transient creep, and the second term [i.e., 0.688f] represents stress relaxation due to steady-state 

creep. The screening trend curve obtained with the available stress relaxation data is shown in Fig. 69. 

The curve bounds all available test data for Type 304 SS, 20% CW Type 316 SS, and Alloy X-750.23-25 

The results indicate approximately 60% stress relaxation at a neutron dose of 0.2 dpa (or 1.3 x 1020 n/cm2

E >1 MeV). Based on the screening trend curve in Fig. 69, all bolted PWR internals locations or spring 

component items that require preload for functionality and reach 0.2 dpa or higher are identified for 

further evaluation.101 

0

20

40

60

80

100

0.0 0.5 1.0 1.5 2.0 2.5 3.0

SA Type 304 uniaxial tension <38°C

Alloy X-750 bent-beam 297°C

Alloy X-750 HTH bent-beam 57°C

SA Type 304 bent-beam 57°C

SA Type 304 bent-beam 297°C

20% CW Type 316 bent-beam 370°C

Remaining Stress (%)

Neutron Dose (dpa)

Figure 69. Trend line for screening irradiation-induced stress relaxation based on available test data

(Refs. 23-25,101).

However, most of the existing data on irradiation-enhanced stress relaxation/irradiation creep have 

been from fast reactor irradiations, and data obtained under PWR irradiation conditions is limited. 

Therefore, the effects of neutron spectrum and He production rate are not known. 

Contributions from other mechanical factors related to washers, gaskets, O-ring, or cycling loading, 

including vibrations, need to be considered in evaluating stress relaxation of bolted joints.
92

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93

6. Summary 

The existing data on the effect of neutron irradiation on PWR core internal materials have been 

reviewed to evaluate the susceptibility of structural materials to irradiation effects such as IASCC, 

neutron embrittlement, void swelling, and irradiation-induced stress relaxation. The results are used to 

(a) define threshold fluence for irradiation effects, (b) develop disposition curves for crack growth rates 

for reactor core internal materials, and (c) assess the significance of void swelling and irradiation creep 

relaxation on the structural and functional integrity of reactor internal components. 

6.1 Susceptibility to IASCC 

A review of the literature has identified the following key material and environmental parameters 

that influence the susceptibility of LWR core internals materials to IASCC susceptibility.

Microstructure: Neutron irradiation of austenitic SSs creates point defects that rearrange into more 

stable configurations such as dislocation loops, network dislocations, precipitates, and cavities (or voids). 

Such changes in the microstructure of austenitic SSs vary with material composition, irradiation 

temperature, neutron fluence and flux, and energy spectrum. At temperatures below 300°C, the material 

microstructure primarily consists of “black spot” defect clusters and faulted dislocation loops, whereas 

large faulted loops, network dislocations, cavities/voids (clusters of vacancies and/or gas bubbles), and 

precipitates are observed above 300°C. At LWR operating temperatures (i.e., 275-300°C), the loop 

density saturates at a relatively low dose (about 1 dpa), and the average loop diameter saturates at 5 dpa. 

Alloying elements can affect the material microstructure (e.g., P, Ti, and Nb increase the loop density and 

decrease loop size). Cavities or voids have not been observed in SSs irradiated below 300°C. 

The loop size increases and loop density decreases with increasing irradiation temperature. At 

temperatures of 300-350°C, the microstructure primarily consists of large Frank loops and a network of 

tangled dislocations. Cavities and voids form at high doses and high temperatures. High irradiation 

temperatures also lead to the formation of second phase particles. However, radiation-induced 

precipitation is not a concern at temperatures below 350°C. Metal carbides are the primary stable 

precipitates in 300-series SSs under LWR conditions, although RIS of Ni and Si to sinks may lead to the 

formation of γ’ phase (Ni3Si) and G phase (M6Ni16Si7).

Most of the studies under LWR conditions have been conducted on BWR components such as 

control-rod blades or top guides irradiated at 290°C to up to 13 dpa. Studies on PWR materials such as 

flux thimble tubes or baffle bolts are limited, particularly on materials irradiated at 320-350°C. 

Additional microstructural data, including the distribution of voids/cavities and second phase particles, 

should be obtained on austenitic SSs irradiated in PWRs above 10 dpa and at temperatures above 300°C.

Microchemistry: The microchemistry of the material is also changed due to RIS. Elements such as 

Si, P, and possibly Ni are enriched at regions such as grain boundaries that act as sinks for point defects, 

while Cr, Mo, and Fe are depleted. The extent of segregation or depletion depends on the rate of 

generation and recombination of point defects (i.e., irradiation temperature and dose rate). Typically, RIS 

peaks at intermediate temperatures. It is reduced at low temperatures because of reduced mobility and at 

high temperatures due to back diffusion. For a given neutron dose, RIS increases at lower dose rate. At 

LWR temperatures, significant segregation is observed at irradiation dose of 0.1 dpa, and the effect either 

saturates or changes very slowly beyond ≈5 dpa. Although measurements of RIS in austenitic SSs 

irradiated in LWRs at temperatures above 320°C are limited, the available data indicate that Cr content at 
94

grain boundaries can decrease to 8 wt.%, and Si and Ni contents can increase to 6 and 30 wt.%, 

respectively, at high dose levels. Such enrichment of Ni and Si at the grain boundaries may lead to the 

formation of γ’ silicide or G phase. 

As discussed above for material microstructure, additional studies are needed to characterize the 

microchemistry of austenitic SSs irradiated in PWRs above 10 dpa and at temperatures above 300°C. 

Also, the redistribution of S, P, C, N, and B in austenitic SSs irradiated in LWRs at dose levels above 

15 dpa needs to be investigated to establish its significance on IASCC susceptibility and neutron 

embrittlement of SSs.

Radiation Hardening: The point defect clusters and precipitates act, to varying extent, as obstacles 

to a dislocation motion that leads to matrix strengthening, resulting in an increase in tensile strength and a 

reduction in ductility and fracture toughness of the material. The yield strength of irradiated SSs can 

increase up to 5 times that of the non-irradiated material after a neutron dose of about 5 dpa. In general, 

cavities (or voids) are strong barriers, large faulted Frank loops are intermediate barriers, and small loops 

and bubbles are weak barriers to dislocation motion. For austenitic SSs, the greatest increase in yield 

strength for a given irradiation level occurs at LWR operating temperatures. The change in yield strength 

correlates well with microstructural changes in the material. The yield strength of solution-annealed SSs 

saturates at 3-5 dpa. At higher neutron dose (above 3 dpa) as the yield strength approaches the ultimate 

strength of the material, deformation by planar slip is promoted, and the material exhibits strain softening, 

also termed “dislocation channeling.” The enhanced planar slip leads to a pronounced degradation in the 

fracture toughness of austenitic SSs. 

Tensile Properties: Irradiation at LWR operating temperatures increases the tensile strength and 

decreases the ductility of wrought and cast austenitic SSs and their welds. The tensile properties reach 

saturation at 5-20 dpa and do not seem to change significantly at higher dose levels. Uniform elongation 

of some LWR-irradiated SSs can decrease to 0.1% at neutron dose levels of 5-10 dpa. Correlations have 

been developed by MRP for estimating tensile yield and ultimate strengths and uniform and total 

elongation of austenitic SSs as a function of neutron dose. However, the majority of the data are for

materials irradiated in a fast reactor, and data on LWR-irradiated materials are limited, in particular, 

above 10 dpa. The available data for the LWR-irradiated materials indicate that the yield and ultimate 

strengths are typically higher, and the uniform and total elongations lower, than those for materials 

irradiated in fast reactors. Consequently, estimates of the tensile properties of irradiated austenitic SSs 

based on the MRP correlations are likely to be non-conservative for LWR operating conditions. 

Additional data on austenitic SSs irradiated in LWRs to dose levels above 20 dpa are needed to 

better define the correlations for estimating the tensile properties of these materials for LWR operating 

conditions. Also, the contribution of voids and second phase particles in these materials to irradiation 

hardening and loss of ductility should be investigated. Voids and second phase particles may not affect 

the tensile strength, but depending on whether they form preferentially at the grain boundaries, fracture 

toughness and ductility and IASCC susceptibility could be affected significantly. 

Yield Strength: It is well known that a susceptible material, significant tensile stress, and an 

aggressive environment, in combination, can increase the IGSCC of austenitic SSs in BWR 

environments. The susceptibility to IGSCC is decreased for low-C SSs and weld metal, which are 

considered resistant to sensitization by welding. However, nonirradiated, nonsensitized, austenitic SSs 

with high tensile strength are susceptible to IGSCC in ultra-high purity water. Increases in yield strength 

can originate from surface or bulk cold work, weld shrinkage strain, and precipitation or irradiation 

hardening; all of these increase SCC susceptibility. 
95

Corrosion Potential: For nonirradiated SSs, decreasing the corrosion potential from that of an 

NWC BWR environment to that of HWC BWR or PWR environments decreases the CGRs. Although 

SSs are not immune to SCC in low-potential environments, the CGRs of nonirradiated SSs and SSs 

irradiated to a neutron dose of less than 3 dpa are more than an order of magnitude lower in such 

environments than those in high-potential environments. The results also indicate that the beneficial 

effect of decreased corrosion potential on CGRs is lost for most SSs irradiated above 10 dpa and some 

heats of SSs irradiated to dose levels as low as 3-5 dpa. Furthermore, for some SSs, the benefit of 

reduced potential is observed at K values less than 20 MPa m1/2 but not at higher K values. 

Silicon Segregation: The increased susceptibility to IASCC and the loss of benefit of reducing 

potential in highly irradiated SSs have often been attributed to the segregation of Si at the grain 

boundaries. The Si segregation is detrimental because at all relevant pH/potential conditions, Si oxidizes 

to SiO2, which is highly soluble in hot water. The possible effect of Si segregation is based on crack 

growth studies on CW Type 304L SS containing 3-5 wt.% Si. Unusually high CGRs were observed in 

these materials in high-purity water with 2000 ppb DO, and the rates did not decrease when the 

environment was changed from high to low potential, or K was decreased from 30 to 15 MPa m1/2. 

However, in contrast to the increased susceptibility of high Si materials indicated by crack growth tests on 

CT specimens, SSRT test data on irradiated Type 304 and 316 SSs with 0.5 to 1.8 wt.% Si show the 

opposite behavior. For the same fluence level, steels with 1.5 to 1.8 wt.% Si showed less irradiation 

hardening and greater elongation than steels containing 0.5 wt.% Si. 

The significance of Si segregation at grain boundaries on the SCC behavior of irradiated SSs is not 

clear, but it could be important. Crack growth tests should be conducted on irradiated material from 

commercial heats of SSs with similar compositions but different Si contents to establish the role of Si 

segregation on the IASCC susceptibility of irradiated SSs.

Sulfur Content: Several commercial heats of Types 304 and 316 SS with ≤0.002 wt.% S show 

excellent resistance to IASCC in the NWC BWR environment at neutron dose up to ≈3 dpa. At 

sufficiently low concentrations of S (≤0.002 wt.%), a sufficiently high concentration of C suppresses the 

deleterious effect of S. However, when S concentration is high (≥0.003 wt.%), the deleterious effect of S 

is so dominant that the effect of C is negligible or insignificant, even at high concentrations. Limited data 

indicate a similar behavior of the effect of S and C contents on IASCC susceptibility in the PWR 

environment. 

Stacking Fault Energy: The SFE is an important parameter that determines the deformation mode. 

Alloys with high SFE (e.g., Ni concentration >18 wt.%) are highly resistant to IASCC compared to 

Type 304 SS with 8 wt.% Ni. Alloys with low SFE, such as SSs with 8 wt.% Ni, deform entirely by 

planar slip, whereas there is no evidence of planar slip in alloys with high SFE (with >20 wt.% Ni). The 

increased susceptibility to IASCC is attributed to absorption of dislocations at the grain boundaries that 

may cause grain boundary sliding ahead of the crack tip, resulting in crack extension and IASCC. 

Alternatively, progressively higher stresses at the grain boundary can change the film rupture frequency, 

thereby exposing bare metal surface and leading to oxidation/dissolution and crack extension. However, 

the role of the localized deformation mode on IASCC of austenitic SSs, and the possible contribution of 

SFE, should be further investigated. 

LWR vs. Fast Reactor Irradiations: Fast reactors are an attractive alternative for irradiating 

materials to high dose levels because of the higher displacement rates. However, differences in the 

thermal- and fast-neutron spectra result in differences in microstructure and microchemistry that may not 

be representative of LWR irradiations. At 275-320°C, although the loop microstructure appears to be 
96

relatively insensitive to displacement rates, the main difference in the microstructure is the presence of He 

bubbles and cavities in materials irradiated to high dose levels in LWRs. Also, Cr depletion and Si and 

Ni enrichment are less in fast reactor irradiated materials. These differences are likely to influence 

neutron embrittlement and IASCC susceptibility of PWR core internal materials. Caution should be 

exercised when using fast reactor data to determine the effects of neutron irradiation on the fracture 

properties of LWR core internal components.

6.2 IASCC Growth Rates

Irradiation-assisted SCC is a complex phenomenon. It is dependent on several parameters that

interact with each other such that their effect on growth rates is not always the same. As a result of these 

complexities and uncertainties in experimental measurements, the SCC crack growth rate data in LWR 

environments show significant variability (up to 1000X), and specific effects of these parameters on 

growth rates cannot be accurately determined under all conditions of interest. Therefore, the SCC crack 

growth data have been typically screened to ensure: (a) well-characterized material, (b) test procedures 

that closely reproduce the loading and environmental conditions for reactor core internal components, and 

(c) data acquisition and analysis system that provides accurate measurements and reproducible results. 

Crack growth rate data on irradiated wrought and cast austenitic SSs, as well as SS weld metal and 

weld HAZ materials in LWR environments, have been compiled and evaluated to define the threshold 

fluence for IASCC and to develop disposition curves for cyclic and IASCC growth rates for reactor core 

internal materials. The significance of test procedures, importance of IG starter crack and specimen 

K/size criterion, and effects of rising K and reloading are discussed. 

Data on the SCC growth rate have been obtained for NWC and HWC BWR environments as well 

as the PWR environment on wrought and cast austenitic SSs and weld HAZ materials irradiated up to 

38 dpa in LWRs or fast reactors. In the NWC BWR environment, neutron irradiation up to 0.45 dpa 

(3 x 1020 n/cm2) has no effect on CGRs; the SCC growth rates are below the NUREG-0313 CGR 

disposition curve for nonirradiated materials in high-purity water containing 8 ppm DO. The CGRs in the 

NWC BWR environment for materials irradiated to higher neutron dose levels can be up to factor of 40 

higher than the NUREG-0313 CGR disposition curve. The growth rates for some materials irradiated to 

5-8 dpa (corresponding to a 60-y end-of-life neutron dose for BWRs) are a factor of 20 higher than the 

NUREG-0313 curve and 2-3 higher than the EPRI disposition curve for BWR core internal components. 

In NWC, the cracking behavior of various grades and heats of SSs irradiated to the same dose level seems 

to be the same. However, the rates for weld HAZ materials are slightly higher than those for solutionannealed SSs, most likely due to weld residual strains. 

The CGR data for highly irradiated materials (greater than 18 dpa) show unusual behavior. For SSs 

irradiated in a PWR to 18 or 37.5 dpa, SCC growth rates in the NWC BWR environment are below the 

NUREG-0313 disposition curve for nonirradiated SSs. The reasons for the low CGRs observed for these 

highly irradiated BWR materials in the high-potential BWR environment are not clear. The low growth 

rates may be an artifact associated with the test procedure used in obtaining these results. In general, the 

SCC growth rate data obtained using test procedures in which the specimen was precracked in air and 

then transferred to the autoclave for the SCC test in the test environment show much more variability at K 

values of 8-15 MPa m1/2 than the data obtained using test procedures in which the specimen was 

precracked in the test environment and the TG fatigue crack was transitioned to an IG crack by using slow 

fatigue cycling before starting the SCC test. The data on the high irradiation dose materials were obtained 

by a procedure that did not involve slow cycling in the environment. Also, CGRs are lower by a factor of 
97

10 for the same heat of BWR control-rod blade material tested using a procedure that did not include slow 

cycling to transition the TG crack into an IG crack compared with a procedure that included slow cycling. 

The SCC growth rates in the HWC BWR environment consistently show the following effects: a 

significant decrease in growth rates relative to those in the NWC BWR environment for SSs irradiated to 

less than 3 dpa. However, some heats of SSs irradiated to dose levels as low as 3-4 dpa show little or no 

reduction, at least at K values above 18 MPa m1/2. No decrease has been observed for any SS irradiated 

to 12-14.5 dpa. The CGRs for some of the materials irradiated to 3-4 dpa are above the CGR disposition 

curve proposed by EPRI for the HWC BWR environment. Some investigators seem to screen out the 3-

4 dpa data that do not show the benefit of the reduced corrosion potential on growth rates on the premise 

that they exceeded a proposed K/size criterion for irradiated materials (i.e., an irradiation-induced 

increase in yield stress is discounted by a factor of 2 or 3). However, the arguments against the validity of 

the data do not seem to have a well-developed technical basis. Additional CGR data on SSs irradiated to 

3-8 dpa are needed to accurately establish the threshold for IASCC susceptibility in low-potential 

environments.

The results also indicate a threshold dose of 3 x 1020 and 1.8 x 1021 n/cm2 (or 0.45 and 2.7 dpa) in 

the NWC and HWC environments, respectively, above which the CGRs of austenitic SSs are greater than 

the NUREG-0313 curve for nonirradiated materials. The CGR data on irradiated austenitic SSs have 

been used to develop correlations for determining the constant A1 in the CGR disposition curve as a 

function of neutron dose. These correlations could be used to determine the CGR disposition curve for 

irradiated austenitic SSs in NWC and HWC BWR environments. However, the existing database is quite 

limited, and additional CGR data should be obtained on SA and SMA SS welds and weld HAZ materials 

in BWR environments to better define the CGR disposition curve for BWRs. 

For the 3-5 dpa materials, the lack of an effect of reduced corrosion potential on growth rates has 

often been explained on the basis of the loss of specimen constraint by exceeding the K/size criterion. 

However, the basis for the proposed K/size criterion for irradiated materials is not well documented. 

Some recent studies have shown that the ASTM E399 size criterion is not exceeded in 1⁄2-T or 1⁄4-T CT 

specimens of irradiated SSs that have yield stress above 600 MPa. The loss of the beneficial effect of 

decreased corrosion potential in some cases may be an indication of a threshold phenomenon. For 

example, the loss seems to occur under conditions that result in a growth rate of about 1 x 10-9 m/s. At 

this threshold growth rate, the crack tip strain rate may be influenced by an alternative mechanism that 

helps sustain the high growth rates, which are independent of corrosion potential. Additional data on SSs 

irradiated to 3-8 dpa are needed to better determine the effect of low-potential environments on the CGRs 

of highly irradiated materials.

The SCC growth rates have been obtained in the PWR environment at 288-340°C on austenitic SSs 

irradiated to 3.0-37.5 dpa in LWRs or fast reactors. The CGRs for the same material and irradiation 

conditions increase with increasing temperature. The temperature dependence of growth rates may be 

represented by an activation energy of 100 kJ/mol. In PWR water at 320°C, most of the CGR data for 

SSs irradiated to 3 dpa are up to a factor of 5 above, and those for SSs irradiated to 11-25 dpa are nearly 

two orders of magnitude above, the NUREG-0313 curve for nonirradiated materials. Also, as observed 

earlier in HWC BWR water tests, the CGRs in the PWR environment of SSs irradiated above 18 dpa are 

below the NUREG-0313 curve for nonirradiated SSs. The results also indicate that in the PWR 

environment at 320°C, the SCC growth rates of materials irradiated up to 12 dpa are above the EPRI 

disposition curve for the HWC BWR environment at 288°C. 
98

The SCC growth rates in the PWR environment show significant variability. The reason for the 

high CGRs for some SSs irradiated to 11.4 dpa, as well as the unusually low CGRs in SSs irradiated to 18 

and 39 dpa, is not known. The available CGR data in the PWR environment at 320°C could also be 

represented by the same expressions used to represent CGRs in the HWC BWR environment. Additional 

data on materials irradiated at different temperatures (e.g., between 300 and 350°C) are needed to better 

understand the IASCC susceptibility of austenitic SSs in PWR environments.

The SCC growth rate data show rapid increase in CGRs as the yield strength of the material 

increases from 150 to 600 MPa, and only a moderate increase at higher yield strengths. At yield strength 

below 800 MPa, a similar growth rate behavior is observed for different grades of SSs, although the 

CGRs in low-potential water are up to an order of magnitude lower. Also, in a high-potential 

environment, the CGRs of sensitized materials are higher than those for nonsensitized materials. 

However, at yield strengths of 800 MPa or higher, the CGRs in low- and high-potential environments are 

comparable. The significance of material yield strength on IASCC susceptibility is not well established 

and should be further investigated, particularly in the PWR environment.

6.3 Fatigue Crack Growth Rates

The limited fatigue CGR data in air on austenitic SSs irradiated to 2 dpa indicate little or no effect 

of neutron irradiation on growth rates. Additional data on austenitic SSs irradiated to higher neutron 

doses have been obtained on materials irradiated in fast reactors at 400-482°C. The results at 427°C show 

only slightly higher CGRs (up to a factor of 2 higher) for ΔK less than 44 MPa m1/2 (40 ksi in.1/2). Also, 

tests on Type 304 and 316 SS irradiated in a thermal reactor at 288°C to 2.7 dpa and tested at 427°C 

exhibited no significant adverse effects of irradiation on crack growth rate. In fact, the crack growth rates 

of the irradiated materials were actually 25 to 50% lower than those for the same heats of nonirradiated 

304 and 316 SS materials.

Fatigue CGRs have been obtained in NWC and HWC BWR environments on wrought and cast 

austenitic SSs and weld HAZ materials irradiated up to 3 dpa. In the NWC BWR environment, the cyclic 

CGRs of SSs irradiated to 0.45 dpa (3 x 1020 n/cm2) are the same as those for nonirradiated materials, 

whereas the CGRs of SSs irradiated to 0.75-3.0 dpa are higher. Limited data indicate that the growth 

rates of irradiated CF-8M cast SS are lower than those of wrought materials irradiated to the same neutron 

dose. The cyclic CGRs at low frequencies are decreased by more than an order of magnitude when the 

DO level is decreased by changing from NWC to HWC. A superposition model has been used to 

represent the cyclic CGRs of austenitic SSs. The CGR in the test environment is expressed as the 

superposition of the rate in air (mechanical fatigue) and the rates due to corrosion fatigue and SCC. The 

calculated results indicate that the higher fatigue CGRs for the irradiated materials are primarily due to 

the SCC component of crack growth. The correlations for the various material and environmental 

conditions are listed in Table 5. The fatigue CGRs for irradiated SSs in the NWC BWR environment can 

be described by superposition of the Shack/Kassner model for nonirradiated SSs in high-purity water with 

8 ppm DO and by the SCC curve for irradiated SSs. Note that although irradiation had no significant 

effect on fatigue CGRs in air, both fatigue and SCC growth rates for austenitic SSs irradiated above 

0.45 dpa are higher than those for nonirradiated SSs or SSs irradiated below 0.45 dpa.

Fatigue CGR data on wrought and cast austenitic SSs and their welds irradiated in LWRs to high 

neutron dose levels and tested in BWR or PWR environments are not available. Additional fatigue crack 

growth data under these material and environmental conditions are needed to develop fatigue CGR 

correlations for BWR and PWR core internal components. The new reactor designs minimize use of cast 

austenitic SSs for reactor core internal components to avoid loss of toughness due to thermal aging effects 
99

and difficulties encountered in performing nondestructive examination of cast SSs and their welds. 

However, such information would be used to address the license renewal issues for operating reactors. 

Table 5. Fatigue CGR correlations for wrought and cast austenitic stainless steels in BWR 

environments at 289°C. 

CGR (m/s) Correlations Material & Environmental Conditions

CGRair CGRair = da dtair = CSS S(R)!K3.3 t

rise

CSS =1.9142x10!12 +6.7911x10!15T !1.6638x10!17 T2 +3.9616x10!20T3

S(R) = 1.0 R < 0

S(R) = 1.0 + 1.8R 0 < R < 0.79

S(R) = -43.35 + 57.97R 0.79 < R < 1.0

Nonirradiated or irradiated

CGRCF

( ) 5 0.5 CGRCF air 4.5x10 CGR ! =

( ) 4 0.5 CGRCF air 1.5x10 CGR ! =

( ) 4 0.5 CGRCF air 1.5x10 CGR ! =

Nonirradiated and irradiated <0.45 dpa

0.3 ppm DO

8.0 ppm DO

Irradiated >0.5 & ≤3.0 dpa

0.2 - 0.5 ppm DOa

CGRSCC

( ) 13 2.161 CGRSCC 2.1x10 K ! =

( ) 14 2.161 CGRSCC 7.0x10 K ! =

( ) ( ) 12 12 2.161 CGRSCC 1.21x10 2.65x10 log dpa K ! ! = " + # $ %

( ) ( ) 12 12 2.161 CGRSCC 2.76x10 6.82x10 log dpa K ! ! = " + # $ %

Nonirradiated and irradiated <0.45 dpa

8.0 ppm DO

0.2 ppm DO

Irradiated >0.45 dpa & ≤12 dpa

Normal water chemistry BWR

Hydrogen water chemistry BWR

aCorrelation may yield conservative estimates of CGR for cast austenitic SSs and low-C Type 304L SS weld HAZ materials. 

6.4 IASCC Initiation

All wrought and cast austenitic SSs and their welds, irradiated below a threshold neutron fluence 

are considered not susceptible to IASCC in a PWR environment. Based on laboratory SSRT data and 

PWR operating experience, a threshold fluence of 2 x 1021 n/cm2 (3 dpa) has been proposed. For 

materials irradiated above this threshold, IASCC initiation data have been used to define for a given 

neutron dose, an apparent stress threshold below which IASCC initiation will not occur in a PWR 

environment. The existing IASCC initiation data have been reviewed to evaluate the adequacy of the 

database for developing screening criteria for the IASCC susceptibility of PWR core internal components. 

The following conclusions are reached: 

1. Most of the IASCC data have been obtained on CW Type 316 SS, while data on solution-annealed 

Type 304, 304L, or 316 SS are limited. Also, there are no data on cast austenitic SSs, SS welds, and 

weld HAZ materials. 

2. Although a threshold fluence of 3 dpa has been defined for IASCC in the PWR environment, no 

IASCC initiation data are available on austenitic SSs irradiated between 3 and 9.5 dpa.
100

3. Nearly all of the high fluence (>20 dpa) data and the majority of the low fluence data are for materials 

irradiated below 325°C and, therefore, do not include the potential effects of additional precipitate 

phases, voids/cavities, and He generation rate on IASCC.

4. In reactor core internal components, IASCC is likely to occur under creviced conditions. Nearly all of 

the IASCC initiation tests have been conducted in the normal PWR environment

5. Several studies indicate that materials irradiated in fast reactors show lower susceptibility to IASCC 

than those irradiated in LWRs. The current IASCC initiation database, however, includes some tests 

on austenitic SS irradiated in the BOR-60 fast reactor. 

6.5 Neutron Embrittlement

The fracture toughness of austenitic SSs has been divided into three broad categories. Category III 

materials with JIc >150 kJ/cm2 fracture after stable crack extension at stresses well above the yield stress. 

Category II materials with JIc of 30-150 kJ/cm2 fracture at stress levels close to the yield stress by stable 

or unstable crack extension. Category I materials with KIc <75 MPa m1/2 fracture at stress levels well 

below the yield stress by unstable crack extension. Nonirradiated wrought and cast austenitic SSs and 

their welds fall in Category III. The fracture toughness of cast austenitic SSs and SA and SMA welds is 

generally lower than that of the wrought SSs. Neutron irradiation can degrade fracture toughness of 

austenitic SSs to the level of Category II or I. 

Fracture toughness data on irradiated wrought and cast austenitic SSs and their welds have been 

compiled and evaluated to define the threshold neutron dose above which the fracture toughness of these 

materials is reduced significantly. Fracture toughness is typically characterized by the initiation 

toughness JIc and tearing modulus for materials that fail after substantial plastic deformation (EPFM 

analysis) and by the critical stress intensity factor KIc for materials that fail after little or no deformation 

(LEFM analysis). To reduce activity and facilitate handling, small specimens have been used in most 

studies. The validity of the J-R curve data and the different methods for determining JIc is discussed. The 

fracture toughness data obtained on different size specimens show that small specimens yield valid 

J-R curve data, at least for materials with JIc below 300 kJ/m1/2. 

Until recently, most published experimental data on the neutron embrittlement of austenitic SSs 

were obtained on materials irradiated at 350-427°C in high-flux fast reactors and tested at 300-427°C. 

The effects of fast reactor irradiation are divided into three regimes: little or no loss of toughness below 

an exposure of ≈1 dpa, substantial and rapid decrease in toughness at exposures of 1-10 dpa, and no 

further reduction in toughness above a saturation exposure of 10 dpa. The degradation in fracture 

properties saturates at a JIc value of ≈30 kJ/m2 or an equivalent KJc of 75 MPa m1/2 (68.2 ksi in.1/2). The 

fracture toughness trend for the LWR data is similar to that observed for fast reactor data. However, the 

tensile property data discussed in Section 2.1.4 indicate that tensile strength is higher and ductility is 

lower for the LWR-irradiated materials than those for materials irradiated in fast reactors. Additional 

fracture toughness data obtained on austenitic SS, preferably the same heat of material, irradiated under 

comparable conditions in fast and thermal reactors are needed to determine the applicability of fast 

reactor data to LWR irradiation conditions. 

Most of the J-R curve data on LWR-irradiated austenitic SSs have been obtained on Type 304 and 

304L SSs, and similar data on SS welds, cast SSs, or weld HAZ materials are limited. The available data 

indicate minor differences in the fracture toughness of the various wrought and cast austenitic SS 

materials. For the same irradiation conditions, the fracture toughness of thermally aged cast SS and weld 
101

metal is lower than that of HAZ material, which, in turn, is lower than that of solution-annealed materials. 

Available data indicate a strong orientation effect on fracture toughness. Tests performed on Type 304 or 

304L SS irradiated to 4-12 dpa in BWRs consistently show lower fracture toughness in the T-L than in 

the L-T orientation. Some SSs irradiated above 4 dpa in BWRs show KIc or KJc values in the range of 

36.8-40.3 MPa m1/2 (33.5-36.6 ksi in1/2). Several J-R curve tests on Type 304 SS control-rod blade 

material irradiated to 4.5-5.5 dpa showed no ductile crack extension. 

Fracture toughness J-R curve tests have also been conducted in a BWR NWC environment to 

investigate the possible effects of the coolant environment and crack morphology on fracture toughness. 

Nearly all of the existing fracture toughness data have been obtained from tests in air and on specimens 

that were fatigued to produce a crack that has TG morphology. In reactor core components, cracks 

initiate primarily by SCC and have IG morphology. Limited data suggest that the BWR environment has 

little or no effect on materials with poor fracture toughness, whereas materials with relatively high 

fracture toughness may be decreased in the test environment. In particular, the effect of environment on 

the fracture toughness of cast SS and welds should be investigated.

The available data are inadequate to establish accurately the effects of the irradiation temperature 

on the fracture toughness of austenitic SSs. However, tensile data for austenitic SSs indicate that 

irradiation hardening is the highest, and ductility loss is maximum, at an irradiation temperature of 

≈300°C (≈572°F). Additional data should be obtained on austenitic SSs irradiated to similar dose levels 

at temperatures between 290 and 340°C to establish the effects of the irradiation temperature.

The material fracture toughness can vary with test temperature. In general, the fracture toughness

of nonirradiated austenitic SSs is known to decrease as the test temperature is increased. The fracture 

toughness of steels irradiated to relatively low dose (less than 5 dpa) shows the same trend. However, for 

steels irradiated to more than 12 dpa, test temperature has little effect on fracture toughness. 

The existing fracture toughness data have been used to develop a fracture toughness trend curve 

that includes (a) a threshold neutron exposure for radiation embrittlement of austenitic SSs and a 

minimum fracture toughness for these materials irradiated to less than the threshold value, (b) a saturation 

neutron exposure and a saturation fracture toughness for materials irradiated to greater than this value, 

and (c) a description of the change in fracture toughness between the threshold and saturation neutron 

exposures. These trend curves describe the change in either JIc or the coefficient C of the power-law J-R 

curve as a function of neutron dose (Eqs. 23 and 25, respectively). The results indicate that the saturation 

fracture toughness proposed by MRP for PWRs is consistent with the existing data. However, the 

saturation fracture toughness proposed by EPRI for BWR internal materials irradiated above 4.5 dpa is 

higher than that observed for several materials irradiated to 4.5-8.5 dpa.

A review of the existing data indicated the following: (a) limited data on materials irradiated in 

LWRs to neutron dose levels of 0.1-1.0 dpa or above 10 dpa, and (b) little or no data on LWR-irradiated 

cast austenitic SSs and welds. As discussed in Section 6.3, the new reactor designs minimize use of cast 

austenitic SSs for reactor core internal components. However, this information would be used to address 

license renewal issues for operating reactors. Also, a fine distribution of γ’ phase is often associated with 

BWR-irradiated materials with poor fracture toughness. The contribution of additional precipitate phases, 

voids, and cavities on the fracture toughness of these materials needs to be investigated. 
102

6.6 Synergistic Effects of Thermal and Neutron Embrittlement

Cast and welded austenitic SSs have a duplex structure consisting of austenite and ferrite phases. 

Formation of Cr-rich α$ phase in the ferrite is the primary mechanism for thermal embrittlement of these 

materials; thermal aging has no effect on the austenite phase. Embrittlement of the austenite phase from 

neutron irradiation under LWR operating conditions occurs at dose levels above 0.5 dpa, and that of the 

ferrite phase occurs at lower dose levels (above 0.07 dpa). However, for reactor core internal 

components, concurrent exposure to neutron irradiation could result in a synergistic effect wherein the 

service-degraded fracture toughness can be less than that predicted for either thermal embrittlement or 

neutron irradiation embrittlement independently.

The screening criteria are discussed for thermal-aging susceptibility of cast austenitic SSs and 

estimates of the saturation J-R curve for specific grades of cast SSs as well as for aged SS welds. A 

methodology is also presented for determining the threshold neutron dose for irradiation embrittlement of 

wrought and cast austenitic SSs and welds. The available fracture toughness data indicate a threshold 

dose of about 0.3 dpa for these materials. However, the data are inadequate to evaluate the synergistic 

effects of thermal and neutron embrittlement on the threshold dose for embrittlement. Additional fracture 

toughness data are needed to better establish the potential for synergistic loss of toughness in these 

materials in the transition dose range from 0.05 to 2 dpa. 

6.7 Void Swelling

Void swelling refers to the volume change of materials under neutron irradiations. Voids form in 

the material when vacancies produced by the displacements migrate and accumulate to form small 

cavities, which eventually grow to large sizes. The fundamental driving force of void formation, 

however, remains the excess vacancy flux toward voids. Most of the void swelling data have been 

obtained from materials irradiated in fast breeder reactors at temperature above 385°C (725°F) and at 

dose rates that are orders of magnitude higher than those in PWRs. Therefore, caution should be 

exercised in extrapolation of the fast reactor data to estimate the void swelling behavior for PWR end-oflife or extended life conditions. 

Void swelling is divided into two regimes: a transient regime followed by a steady-state swelling 

rate. All materials are believed to reach a steady-state swelling rate of ~1%/dpa. The time to reach the 

steady-state rate (i.e., the transient regime) depends on the material composition and thermo-mechanical 

condition, stress state, neutron flux and spectrum, and irradiation temperature (above 300°C). The 

existing data were reviewed to assess void swelling in PWR core internal components and to evaluate the 

effects of key material and environmental parameters. 

Reactor core internal components are primarily constructed of austenitic SSs and some of Ni alloys, 

both of which have a face-centered cubic crystal structure. These materials are more susceptible to void 

swelling than ferritic steels, which have a body-centered cubic crystal structure. The material 

composition can influence void swelling, in particular, the Ni content. In the 300-series SSs, void 

swelling decreases with increasing Ni content. Increases in the Si and P contents also decrease swelling. 

Thus, under similar thermal-mechanical conditions, Type 304 SS typically swells more than Type 316 SS. 

However, radiation-induced precipitation of second phase particles can remove Ni, P, and Si from the 

alloy matrix, and thereby increase void swelling. In addition, cold work prolongs the transient regime for 

void swelling in austenitic SSs. The high dislocation density produced by the cold work provides 

additional recombination sites for vacancies and interstitials, thereby decreasing supersaturation of 

vacancies in the material, which retards void nucleation. 
103

Void swelling in the 300-series austenitic SSs is typically observed at irradiation temperatures 

above 340°C (644°F), and shows a strong dependence on temperature. However, the low temperatures 

typical of LWRs extend the transient regime to much longer times and result in the formation of a higher 

density of smaller voids. Nonetheless, depending on the irradiation temperature, dose, and dose rate, 

voids can form under PWR operating conditions within the reactor lifetime. The displacement rate also 

influences void swelling; swelling increases as the displacement rate is decreased. Available data 

obtained in fast reactors at 373-444°C and displacement rates of 0.3-10.0 x 10-7 dpa/s suggest that under 

comparable irradiation temperatures and neutron dose levels, the lower displacement rates typical of 

PWRs would result in higher swelling. 

Applied or swelling-induced stress, irrespective of whether it is tensile or compressive, accelerates 

void swelling. These stresses, however, are limited by irradiation creep relaxation. Therefore, the 

interaction of swelling and irradiation creep relaxation plays an important role in evaluation of swelling in 

PWR core internal components. Another stress-related effect is that a high swelling region produces 

stresses in the adjacent lower swelling region, forcing it to try to catch up. Thus, estimates of swelling 

based on stress-free material tend to overestimate the swelling gradients in a component. 

Most of the void swelling data have been obtained on materials irradiated in fast reactors at 

temperatures and dose rates that are higher than those typical for LWRs. Also, data on PWR-irradiated 

materials and on CW Type 316 SS are limited. It is important to develop a technical basis for 

extrapolating the fast reactor data and the limited PWR data to predict void swelling in PWR core internal 

components. Although a reasonable extrapolation of these results suggests that void swelling will not be 

a significant problem during the first license extension period, the effects of irradiation temperature 

(particularly above 370°C), displacement rate, and stress on void swelling, including the interaction of 

swelling and irradiation creep relaxation, should be investigated.

Stress-free swelling equations, including stress-enhanced swelling, have been developed by the 

MRP for solution-annealed Type 304 and CW Type 316 SSs. These equations correlate the swelling rate 

in PWR core internal components with irradiation temperature, neutron dose, and dose rate. 

Embrittlement of SS materials due to voids is also discussed. Typically, austenitic SSs irradiated at 

≤400°C to neutron dose levels to produce ≥10% void swelling suffer from severe embrittlement, 

particularly at room temperature. This behavior has been attributed to a second-order process, which 

involves stress concentration between voids, Ni segregation at void surfaces, and a resultant tendency 

toward martensite formation when the steel is deformed at room temperature. At room temperature, the 

austenite/martensite boundaries provide a low-energy crack propagation path, resulting in a brittle quasicleavage fracture. At higher temperatures, the effect of Ni segregation on SFE is much less pronounced, 

and failure occurs along flow localization path, referred to as channel fracture. However, few quantitative 

data correlate void swelling with fracture toughness of the material. A heat of SS irradiated in the FFTF 

at 390°C (734°F) to 114 dpa showed KIc of 30 MPa m1/2 at 410°C. This value is slightly lower than the 

saturation KIc of 38 MPa m1/2 for several SSs irradiated in BWRs to 5-8 dpa. However, these results 

suggest that the fracture toughness of PWR core internal SS components could be lower than the 

saturation value (KIc=38 MPa m1/2) proposed by MRP. The embrittlement of austenitic SSs due to void 

swelling under PWR irradiation conditions should be further investigated.

6.8 Irradiation Creep Relaxation

Loss of preload for bolting and springs due to stress relaxation is another aging degradation effect 

that needs to be addressed to assure the functional integrity of the reactor core internal components. 
104

Stress relaxation represents plastic deformation that occurs under constant strain below the yield point of 

the material. It is caused by either primary creep (thermal stress relaxation) or a flux of high-energy 

neutrons (irradiation creep). Thermal stress relaxation data have been obtained from short-term tests (less 

than 1000 h). At PWR operating temperatures, stress relaxation saturates within 100 h with a maximum 

decrease of 10-20% of the initial preload stress. The extent of relaxation depended on thermalmechanical treatment of the material; it increased with increasing initial stress and saturated at about 

207 MPa (30 ksi). For PWRs, plastic strains resulting from thermal stress relaxation were considered to 

be insignificant and not a concern for the dimensional change of core internal components. Although 

steady-state creep is insignificant at PWR operating temperatures, over the reactor lifetime it could lead to 

appreciable stress relaxation not anticipated in the short-term tests. 

However, neutron-irradiation-enhanced creep at PWR temperatures can greatly increase the plastic 

strain by increasing both the transient creep and steady-state creep rates. The magnitude of transient 

creep and steady-state creep is essentially proportional to the applied stress level. Transient creep is 

typically short and, although its magnitude depends on material condition and structure and irradiation 

temperature, the contribution to total creep strain is relatively small; the steady-state creep rate is 

independent of these parameters. Empirical models have been developed for irradiation creep using data 

obtained from fast reactors. The results have been used to develop an expression for estimating stress 

relaxation as a function of initial stress and neutron dose.
105

7. Gaps in Available Data

1. Applicability of Fast Reactor Data to LWRs

The following subtopics are related to the applicability of fast reactor data to LWRs. 

1a. Material Microstructure and Microchemistry

A large fraction of the high neutron dose data has been obtained on materials irradiated in fast 

reactors. Therefore, possible differences in material microstructure and microchemistry due to higher 

dose rates and differences in the thermal- and fast-neutron spectra need to be examined. A significant 

difference between the two is the presence of He bubbles and cavities in materials irradiated to high dose 

levels under LWR irradiation conditions. Also, data on material microchemistry indicate that RIS in the 

fast reactor irradiated materials may not be representative of LWR irradiation conditions; Cr depletion 

and Ni and Si enrichment are less in the fast-reactor-irradiated materials. These differences are likely to 

influence neutron embrittlement and IASCC susceptibility of LWR core internal materials. Additional 

studies are needed to validate the applicability of fast reactor data to LWRs. 

1b. Irradiation Hardening and Tensile Properties

Correlations have been developed by industry for estimating tensile yield and ultimate strengths 

and uniform and total elongation of austenitic SSs as a function of neutron dose. However, most data 

have been obtained on materials irradiated in a fast reactor and data on LWR-irradiated materials are 

limited, in particular above 10 dpa. Additional data on austenitic SSs irradiated in LWRs to dose levels 

above 20 dpa are needed to better define the correlations for estimating the tensile properties of these 

materials for LWR operating conditions. Also, the influence of voids and second phase particles on 

irradiation hardening and loss of ductility should be investigated.

1c. Fracture Toughness Database for Irradiated LWR Core Internal Materials

Although the general trends for the fracture toughness data for austenitic SSs appear to be similar 

for fast reactor and LWR irradiations, the significant differences in irradiation hardening and tensile 

properties indicate the need for fracture toughness data on materials irradiated in LWRs. A review of the 

existing fracture toughness database indicated the following: (a) limited data on materials irradiated in 

LWRs to neutron dose levels of 0.1-1.0 dpa or above 10 dpa, and (b) little or no data on LWR-irradiated 

cast austenitic SSs and welds. Also, the contribution of additional precipitate phases, voids, and cavities 

on the fracture toughness of these materials needs to be investigated. 

1d. Void Swelling 

Most of the void swelling data have been obtained on materials irradiated in fast reactors at 

temperatures and dose rates that are higher than those typical for LWRs. Data on PWR-irradiated 

materials are limited, and predominantly on CW Type 316 SS. To develop a technical basis for 

extrapolating the fast reactor data and the limited PWR data to predict void swelling in PWR core internal 

components, additional void swelling data should be obtained to establish the effects of irradiation 

temperature (particularly above 370°C), displacement rate, and stress on void swelling.
106

1e. Irradiation Creep Relaxation 

Most of the existing data on irradiation-enhanced stress relaxation/irradiation creep have been from 

fast reactor irradiations, and data under PWR irradiation conditions are limited. Therefore, the effects of 

neutron spectrum and He production rate, are not known. 

2. Microstructure and Microchemistry Characterization of PWR-Irradiated Materials

Most of the studies on characterization of the microstructure and microchemistry of irradiated 

austenitic SSs have been on BWR components such as control-rod blades or top guides irradiated at 

290°C up to 13 dpa. Studies on PWR materials such as flux thimble tube or baffle bolts are limited, 

particularly on materials irradiated at 320-350°C. Additional microstructural and microchemistry data, 

including the distribution of voids/cavities and second phase particles, should be obtained on austenitic 

SSs irradiated in PWRs to neutron dose levels above 10 dpa and at temperatures above 300°C. Also, the 

redistribution of S, P, C, N, and B in austenitic SSs irradiated in LWRs at dose levels above 15 dpa needs 

to be investigated to establish its significance on IASCC susceptibility and neutron embrittlement of SSs.

3. Effect of Si Segregation on IASCC Susceptibility

The increased susceptibility to IASCC and the loss of benefit of reduced corrosion potential on 

IASCC susceptibility in highly irradiated SSs have often been attributed to the segregation of Si at the 

grain boundaries. However, this proposition is based on data for nonirradiated cold-worked materials. 

The significance of Si segregation at grain boundaries on the SCC behavior of irradiated SSs is not clear, 

but it could be important. Crack growth tests should be conducted on irradiated material from 

commercial heats of SSs with similar compositions but different Si contents to establish the role of Si 

segregation on the IASCC susceptibility of irradiated SSs.

4. Validity of Proposed K/Size Criteria

The existing data suggest that some SSs may show IASCC susceptibility in low-potential 

environments at fluence levels as low as 3 dpa. However, such data are often screened out simply 

because they do not agree with the current expectations for the IASCC behavior of austenitic SSs. The 

justification is that the applied K for these tests may have exceeded the K/size criteria that have been 

proposed for irradiated austenitic SSs to account for strain softening in the material. These proposed 

K/size criteria could be validated experimentally by using different size specimens. Such an approach 

would be extremely difficult for irradiated materials, but more rigorous finite element studies of crack tip 

behavior may be helpful in developing appropriate validity criteria. In any case, the potentially 

conflicting results currently available indicate that additional data on SSs irradiated to 3-8 dpa are needed 

to accurately establish the threshold for IASCC susceptibility in low-potential PWR environments.

5. IASCC Crack Growth Rate Disposition Curve for PWR Core Internals

The available crack growth data on irradiated austenitic SSs in the PWR environment show 

significant variability. Also, crack growth data on austenitic SSs irradiated above 12 dpa have been 

obtained on materials that were either from a decommissioned PWR or were irradiated in the BOR-60 fast 

reactor. The SCC growth rate tests on the PWR-irradiated material show anomalous results (i.e., the 

CGRs on material irradiated to 18 or 38 dpa are lower than the growth rates for nonirradiated material). 

Additional data on austenitic SSs, including SS welds and weld HAZ materials, irradiated in LWRs at 

300-350°C, should be obtained in PWR environments to better define a CGR disposition curve for PWRs. 
107

6. Fatigue Crack Growth Rates

The available fatigue crack growth data for irradiated austenitic SSs are limited to materials 

irradiated under BWR operating conditions to neutron doses up to 3 dpa. Additional fatigue crack growth 

data on irradiated wrought and cast austenitic SSs and their welds in BWR and PWR environments are 

needed to develop fatigue CGR correlations for LWR core internal components. 

7. IASCC Initiation 

The screening criteria used to disposition various PWR core internal components into different 

categories for aging management strategies define for a given neutron fluence above a threshold value 

(i.e., 3 dpa) an apparent stress threshold below which IASCC will not occur in a PWR environment. The 

fluence threshold of 3 dpa is based on incidents of IASCC failures in operating PWRs and on SSRT test 

results in a PWR environment. The stress thresholds are based upon the IASCC initiation data. 

However, the database used for defining neutron dose/applied stress thresholds for austenitic SSs is 

inadequate to evaluate the variability between different grades of SSs, or the differences between 

solution-annealed or cold-worked material, or the effect of irradiation temperature. Most of the data are 

on cold-worked Type 316 SS irradiated at temperatures below 310°C. Also, although the fluence 

threshold for IASCC is defined as 3 dpa, there are no IASCC initiation data on materials irradiated below 

10 dpa. Furthermore, in reactor core internal components, IASCC is likely to occur under creviced 

conditions (i.e., under off-normal water chemistry). The variability and uncertainties due to these factors 

need to be addressed in developing threshold and screening criteria for IASCC in reactor core internal 

materials. Additional data should be obtained to determine the possible effects of material condition and 

composition, irradiation temperature, water chemistry, and neutron fluence below 10 dpa on IASCC 

initiation in a PWR environment. 

8. Effect of Irradiation Temperature on Fracture Toughness 

The tensile data for austenitic SSs indicate that irradiation hardening is the highest, and ductility 

loss is maximum, at an irradiation temperature of ≈300°C (≈572°F). However, the existing data are 

inadequate to establish accurately the effects of the irradiation temperature on the fracture toughness of 

austenitic SSs. Additional data should be obtained on austenitic SSs irradiated to similar dose levels at 

temperatures between 290 and 340°C to establish the effects of the irradiation temperature.

9. Lower Bound Fracture Toughness of Irradiated Austenitic Stainless Steels 

The existing data indicate a strong orientation effect on fracture toughness. Tests performed on 

Type 304 or 304L SS irradiated to 4-12 dpa in BWRs consistently show lower fracture toughness in the 

T-L than in the L-T orientation. In addition, some SSs irradiated above 4 dpa in BWRs show poor 

fracture toughness; the KIc is 36.8-40.3 MPa m1/2. Several J-R curve tests on Type 304 SS control-rod 

blade materials irradiated to 4.5-5.5 dpa showed no ductile crack extension. The lower bound fracture 

toughness curve proposed by industry for BWR and PWR core internal components does not consider this 

data. The lower bound curve should be updated to include these results. 

10. Void Swelling and its effect on Fracture Toughness

Most of the void swelling data have been obtained on materials irradiated in fast reactors, and data 

on PWR-irradiated matereials are quite limited. Additional void swelling data on PWR-irradiated 

materials should be obtained to establish the effects of irradiation temperature, diaplacement rate, and 
108

stress on void swelling. In addition, austenitic SSs with ≥10% void swelling suffer from severe 

embrittlement. There are few quantitative data correlating void swelling with fracture toughness of the 

material. Limited data indicate KIc values that are slightly lower than the saturation value of 38 MPa m1/2

proposed by the MRP for austenitic SSs irradiated to a very high neutron dose in PWRs. However, these 

results indicate the importance of void swelling on fracture toughness. The embrittlement of austenitic 

SSs due to void swelling under PWR irradiation conditions should be further investigated

11. Effect of Coolant Environment on Fracture Toughness 

Limited data suggest a possible effect of reactor coolant environment and crack morphology on the 

fracture of LWR structural materials. Nearly all fracture toughness tests are conducted on specimens 

containing a TG fatigue crack, whereas core internal components are likely to be IG SCC cracks. 

Fracture toughness J-R curve data should be obtained in LWR environments to evaluate the possible 

effects of coolant environment, in particular, for cast austenitic SSs and SS welds.

The priorities for addressing the various data gaps in the current understanding of irradiation effects 

on LWR core internal materials are assessed based on the effect of these knowledge gaps on the following 

criteria: 

1. The basic mechanism of the degradation process,

2. The application and use of these materials for LWR core internal components, and 

3. The performance of the component as the material ages.

Based on these criteria, the assessed priorities are given in Table 6.

Table 6. Data gaps in the current understanding of irradiation effects on LWR core internal materials. 

Item Description Priority

1 Applicability of fast reactor data to LWRs

a. Material microstructure and microchemistry M

b. Irradiation hardening and tensile properties M

c. Fracture toughness database for irradiated LWR core internal materials H

d. Void swelling L

e. Irradiation creep relaxation L

2 Microstructure and microchemistry characterization of PWR-irradiated material M

3 Effect of Si segregation on IASCC susceptibility L

4 Validity of proposed K/size criterion H

5 IASCC crack growth rate disposition curve for PWR core internals H 

6 Fatigue crack growth rates M

7 IASCC initiation H 

8 Effect of irradiation temperature on fracture toughness H

9 Lower bound fracture toughness of irradiated austenitic stainless steels H

10 Embrittlement due to void swelling M

11 Effect of coolant environment on fracture toughness L

Priority: H → high, M → medium, and L → low
109

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