Effect of Thermal Aging and 

Neutron Irradiation on Crack 

Growth Rate and Fracture 

Toughness of Cast Stainless 

Steels and Austenitic 

Stainless Steel Welds

Office of Nuclear Regulatory Research

NUREG/CR-7185

ANL-14/10
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infringe privately owned rights.
Effect of Thermal Aging and 

Neutron Irradiation on Crack 

Growth Rate and Fracture 

Toughness of Cast Stainless 

Steels and Austenitic 

Stainless Steel Welds

Manuscript Completed: September 2014

Date Published: July 2015

Prepared by: 

Omesh K. Chopra 

Argonne National Laboratory

Argonne, IL 60439

Appajosula S. Rao, NRC Technical Monitor and Project Manager

Office of Nuclear Regulatory Research

NUREG/CR-7185

ANL-14/10

 

iii

ABSTRACT 

Cast austenitic stainless steel (CASS) materials and austenitic stainless steel (SS) welds, both 

of which have a duplex structure, are susceptible to thermal aging embrittlement during reactor 

service. In addition, prolonged exposure of these materials used in core support structure and 

reactor internals to neutron irradiation changes their microstructure and microchemistry, which 

further degrades their fracture properties. The core support structures and reactor internals also 

contain austenitic SS welds. The purpose of this effort was to review and evaluate the existing 

thermal aging and neutron irradiation embrittlement data on these materials to update the 

methodology developed earlier at Argonne National Laboratory (Argonne) for estimating the 

degradation of mechanical properties of these materials during reactor service. The 

methodology developed earlier in NUREG/CR-4513 Rev. 1, 1994, to determine the kinetics and 

extent of thermal embrittlement from known material information as a function of time and 

temperature of service, has been reviewed and, where necessary, updated to incorporate 

recent data published after the NUREG/CR-4513 Rev.1 report was published. The updated 

methodology is used to estimate the fracture toughness J-R curve of thermally aged CASS 

materials as a function of time and temperature of service. The existing data on irradiated 

materials were also evaluated to define a threshold fluence level above which irradiation effects 

on materials properties are significant. This report also examines the combined effects of 

thermal and neutron embrittlement, which can degrade the fracture toughness of these 

materials more than predicted from either of the processes independently. Potential information 

gaps in the data needed to adequately address the effects of thermal aging and neutron 

irradiation on fracture properties of CASS materials and austenitic SS welds were identified. 

v

FOREWORD 

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

with austenitic stainless steels (SSs) because of their relatively high strength, ductility, and 

fracture toughness. In addition, cast austenitic stainless steel (CASS) materials are used for 

reactor pressure-boundary components, such as valve bodies, pump casings, primary coolant 

piping, and some reactor core internal components. Although austenitic SSs are completely 

austenitic in the wrought condition, welded SSs and CASS materials have a duplex structure 

consisting of austenite and ferrite phases. The ferrite phase increases the tensile strength and 

improves resistance to stress corrosion cracking (SCC). However, the duplex structure of ferrite 

and austenite phases is susceptible to thermal embrittlement after service at reactor operating 

temperatures. In addition, during normal reactor operational conditions, the reactor core internal 

components are exposed to high-energy neutron irradiation and high-temperature reactor 

coolant. Exposure to neutron irradiation for extended periods changes the microstructure and 

microchemistry (radiation-induced segregation) of wrought and CASS materials and their welds, 

thus degrading the component fracture properties and increasing their susceptibility to 

irradiation-assisted stress corrosion cracking (IASCC). These mechanical property changes are 

therefore caused by a complex interaction among changed or ongoing changes in material 

microstructure (material variable), irradiation, temperature, stress, and environment (viz. boiling 

water reactor [BWR] and pressurized water reactor [PWR] environments), which are field 

variables. 

As nuclear power plants age and neutron irradiation dose increases, the degradation of the 

reactor components becomes more likely. This represents an important technical issue to 

assess the long-term operation of nuclear power plants. The U.S. Nuclear Regulatory 

Commission (NRC) has been sponsoring research on reactor core internals, in particular the 

duplex steels and welds, to: (a) define the screening criteria for the maximum acceptable level 

of percent delta ferrite in these duplex steels, above which significant decreases in the fracture

toughness of these components begin due to both thermal and neutron fluence levels;

(b) develop curves for IASCC growth rates, which are used for flaw-acceptance criteria and the 

disposition of detected flaws during inspection; and (c) assess the significance of both thermal 

and neutron embrittlement on the structural and functional integrity of these CASS components 

and welds. This report provides technical data and microstructural analysis on the above 

mentioned complex variables on the fracture toughness of both irradiated and unirradiated aged 

and unaged CASS components.

This report provides technical data and microstructural analysis as function of ferrite 

composition on the fracture toughness of both irradiated and unirradiated thermally aged and 

unaged CASS components.

The results of this final report may be used to: (a) to determine when active aging management 

of reactor vessel internal (RVI) components manufactured from CASS materials is needed for 

license renewal (LR) of LWRs under 10 CFR Part 54; (b) support revision of the Generic Aging 

Lessons Learned (GALL) report; (c) determine appropriate inspection and flaw disposition 

procedures for reactor internals for use in American Society of Mechanical Engineers (ASME)

code development, and for developing appropriate staff positions for Title 10, Section 50.55a, 

“Codes and Standards,” of the Code of Federal Regulations (10 CFR 50.55a); and (d) identify 

technical issues related to screening criteria for the lower bound value of the delta ferrite limit at 

which significant loss of the fracture toughness of the CASS components and the welds could 

potentially occur during the normal operation of the nuclear power plants. This report may also 
vi

be used by the NRC to evaluate the sufficiency of planned industry and NRC-sponsored 

research to address gaps identified in this report.
vii

TABLE OF CONTENTS 

ABSTRACT ................................................................................................. iii

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

TABLE OF CONTENTS ............................................................................... vii

FIGURES .................................................................................................... xi

EXECUTIVE SUMMARY ............................................................................ xvii

ACKNOWLEDGMENTS ............................................................................ xxiii

ACRONYMS AND ABBREVIATIONS ........................................................... xxv

NOMENCLATURE .................................................................................... xxvii

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

2. THERMAL EMBRITTLEMENT OF DUPLEX STAINLESS STEELS .............. 5

2.1 Mechanism of Thermal Embrittlement ..................................................................... 5

2.2 Distribution and Morphology of Ferrite..................................................................... 9

2.2.1 Ferrite Morphology...................................................................................... 9

2.2.1.1 CASS Materials ............................................................................ 9

2.2.1.2 Austenitic Stainless Steel Welds ................................................ 12

2.2.2 Estimation of Ferrite Content .................................................................... 14

2.2.2.1 Hull’s Equivalent Factor.............................................................. 15

2.2.2.2 ASTM A800/800M Methodology................................................. 16

2.3 Kinetics of Thermal Embrittlement ......................................................................... 17

2.4 Extent of Thermal Embrittlement of CASS Materials ............................................. 22

2.4.1 Charpy-Impact Energy.............................................................................. 24

2.4.2 Fracture Toughness J-R Curve................................................................. 25

2.5 Assessment of Thermal Embrittlement of CASS Materials.................................... 34

2.5.1 Estimation of Thermal Embrittlement of CASS Materials of Known 

Composition and Service Condition – Service Time Values..................... 34

2.5.2 Estimation of Thermal Embrittlement of CASS Materials of Unknown 

Composition – Lower-bound values ......................................................... 36

2.5.3 Updated Fracture Toughness Data for CASS Materials........................... 56

2.5.4 Screening Criteria for Susceptibility of CASS Materials to Thermal 

Embrittlement............................................................................................ 58

2.5.5 Methodology for Estimating Thermal Embrittlement of CASS 

Materials ................................................................................................... 61
viii

2.6 Extent of Thermal Embrittlement of Austenitic Stainless Steel Welds ................... 64

2.6.1 Update of Thermal Embrittlement Data for Austenitic Stainless Steel 

Welds ........................................................................................................ 69

3. IASCC CRACK GROWTH RATES .......................................................... 73

3.1 Effects of Material and Environmental Parameters on IASCC Susceptibility......... 73

3.1.1 Microstructure ........................................................................................... 74

3.1.2 Microchemistry.......................................................................................... 74

3.1.3 Radiation Hardening ................................................................................. 75

3.1.4 Yield Strength ........................................................................................... 75

3.1.5 Silicon Segregation................................................................................... 76

3.1.6 Stacking Fault Energy............................................................................... 77

3.2 IASCC Crack Growth Rates................................................................................... 77

3.2.1 In BWR and PWR Environments .............................................................. 77

3.2.2 Update of the CGR Data........................................................................... 81

3.2.3 Effect of Neutron Fluence ......................................................................... 83

4. NEUTRON EMBRITTLEMENT ............................................................... 89

4.1 Fracture Toughness of Irradiated Austenitic Stainless Steels ............................... 89

4.1.1 Fracture Toughness JIc ............................................................................ 91

4.1.2 Fracture Toughness J-R Curve................................................................. 93

4.2 Fracture Toughness Trend Curve........................................................................ 100

5. COMBINED EFFECTS OF THERMAL AND NEUTRON 

EMBRITTLEMENT .............................................................................. 109

6. SUMMARY ......................................................................................... 115

6.1 Significant Results ............................................................................................... 115

6.1.1 Thermal Embrittlement of Duplex Stainless Steels................................. 115

6.1.2 IASCC Crack Growth Rates.................................................................... 118

6.1.3 Neutron Embrittlement............................................................................ 119

6.1.4 Combined Effects of Thermal and Neutron Embrittlement ..................... 121

6.2 Information Gaps ................................................................................................. 122

6.2.1 Thermal Embrittlement of Duplex Stainless Steels................................. 122

6.2.2 IASCC Crack Growth Rates.................................................................... 122

6.2.3 Neutron Embrittlement............................................................................ 122
ix

6.2.4 Combined Effects of Thermal and Neutron Embrittlement ..................... 122

7. REFERENCES .................................................................................... 123

xi

FIGURES 

1. Microstructure along the axial and circumferential sections of centrifugally-cast CF-8 

and CF-3 stainless steel pipes showing equiaxed grain structure......................................... 9

2. Microstructure along axial section of a check valve from the decommissioned 

Shippingport reactor showing columnar grains...................................................................... 10

3. Ferrite content and morphology along circumferential section of centrifugally-cast CF-8 

(Heat P1) and CF-3 (Heat P2) pipes from regions near the inside and outside diameter, 

and two vanes of a static cast CF-3 pump impeller (Heat I). ................................................. 11

4. Typical ferrite morphology of four different welds. ................................................................. 13

5. Plots of measured ferrite content and values calculated from Hull's equivalent factor for 

various CASS materials......................................................................................................... 16

6. Plots of measured ferrite content and values calculated from ASTM A800/A800M 

methodology for various CASS materials. ............................................................................. 17

7. Plots of ferrite content calculated from Hull's equivalent factor and those estimated from 

the ASTM A800/A800M methodology for various CASS materials. ...................................... 17

8a. Plots of measured ferrite content and values calculated from Hull's equivalent factor for 

various austenitic SS welds. .................................................................................................. 18

8b. Plots of measured ferrite content and values calculated from ASTM A800/A800M 

methodology for various austenitic SS welds. ....................................................................... 18

9. Arrhenius plots for the formation of G phase and reduction in impact energy....................... 20

10. Decrease in Charpy-impact energy for various heats of cast austenitic SSs aged at 

400°C..................................................................................................................................... 21

11. Correlation between RT Charpy-impact energy at saturation and the material 

parameter φ for CF-3, CF-8, and CF-8M steels. .................................................................... 26

12. Measured and estimated values of saturation RT Charpy-impact energy for CF-3, CF-8, 

and CF-8M CASS materials. ................................................................................................. 27

13. Correlation between RT Charpy-impact energy at saturation and the ferrite content for 

CF-3, CF-8, and CF-8M steels. ............................................................................................. 28

14. Correlation between RT Charpy-impact energy and coefficient C at 290–320°C and RT 

for CF-3, CF-8, and CF-8M CASS materials. ........................................................................ 29

15. Correlation between RT Charpy-impact energy and exponent n of the power-law J-R 

curve at RT and 290°C for CF-8M materials. ........................................................................ 31

16. Distribution of ferrite content in CASS CF-8M piping material in Westinghouse PWRs. ....... 34
xii

17. Comparison of the experimental values of coefficient C of the power-law J-R curve at 

(a) RT and (b) 290 or 320°C as a function of aging time and temperature for static-cast 

CF-3 plate with the values estimated from material composition and the initial RT 

Charpy-impact energy. .......................................................................................................... 37

18. Comparison of the experimental values of coefficient C of the power-law J-R curve at 

(a) RT or 54°C and (b) 290 or 315–320°C as a function of aging time and temperature for 

static-cast CF-8 plate with the values estimated from material composition and the initial 

RT Charpy-impact energy...................................................................................................... 38

19. Comparison of the experimental values of coefficient C of the power-law J-R curve at 

(a) RT and (b) 290 or 320°C as a function of aging time and temperature for static-cast 

CF-8M plate with the values estimated from material composition and the initial RT 

Charpy-impact energy. .......................................................................................................... 39

20. Comparison of the experimental values of coefficient C of the power-law J-R curve at 

(a) RT and (b) 290°C as a function of aging time and temperature for centrifugally-cast 

CF-8 pipe with the values estimated from material composition and the initial RT Charpyimpact energy. ....................................................................................................................... 40

21. Correlation between saturation RT Charpy-impact energy and ferrite content for CF-3, 

CF-8, and CF-8M CASS materials. ....................................................................................... 42

22a. Estimated lower-bound J-R curves at RT for static-cast CASS materials with ferrite 

contents of 25–40%, 15–25%, 10–15%, and <10%. ............................................................. 44

22b Estimated lower-bound J-R curves at 290–320°C for static-cast CASS materials with 

ferrite contents of 25–40%, 15–25%, 10–15%, <10%. .......................................................... 46

23a. Estimated lower-bound J-R curves at RT for centrifugally-cast CASS materials with ferrite 

contents of 25–40%, 15–25%, 10–15%, and <10%. ............................................................. 48

23b Estimated lower-bound J-R curves at 290–320°C for centrifugally-cast CASS materials 

with ferrite contents of 25–40%, 15–25%, 10–15%, and <10%............................................. 50

24. Comparison of the lower-bound fracture toughness J-R curve for thermally aged 

(a) static-cast and (b) centrifugally-cast CF-8M steels at 290°C estimated using the 

NUREG/CR-4513, Rev. 1 expressions, with the fracture toughness of SA welds used to 

develop ASME Section XI Subsection IWB-3640 evaluations............................................... 52

25. Comparison of the lower-bound fracture toughness J-R curve for thermally aged staticcast CF-8M steels containing (a) <10% Ni and (b) ≥10% Ni at 290°C estimated using the 

updated expressions presented in this report, with the fracture toughness of SA welds 

used to develop ASME Section XI Subsection IWB-3640 evaluations.................................. 54

26. Comparison of the lower-bound fracture toughness J-R curve for thermally aged 

centrifugally-cast CF-8M steels containing (a) <10% Ni and (b) ≥10% Ni at 290°C 

estimated using the updated expressions presented in this report, with the fracture 

toughness of SA welds used to develop ASME Section XI Subsection IWB-3640 

evaluations............................................................................................................................. 55
xiii

27. Fracture toughness J-R curve data at 320°C for static-cast and thermally aged 

plates (Heats CC and DI) and an elbow (Heat EL) of CF-8M material with ≥10% Ni 

and (a) >25% and (b) 15–25% ferrite. ................................................................................... 56

28. Fracture toughness J-R curve data at 320°C for static-cast and thermally aged elbow 

(Heat EK) of CF-8M material with <10% Ni and >25% ferrite................................................ 57

29. Fracture toughness J-R curve data for thermally aged Heat 68 of CF-8M plate at 54°C. ..... 58

30. The change in saturation J at 2.5 mm crack extension determined at (a) RT and (b) 290-

320°C as a function of ferrite content for static-cast CF-3, CF-8, and CF-8M CASS 

materials. ............................................................................................................................... 59

31. The change in saturation J at 2.5 mm crack extension determined at (a) RT and (b) 290-

320°C as a function of ferrite content for centrifugally-cast CF-3, CF-8, and CF-8M CASS 

materials. ............................................................................................................................... 60

32. Flow diagram for estimating mechanical properties of thermally aged CASS materials in 

LWR systems......................................................................................................................... 62

33. Fracture toughness J-R curves for stainless steel welds at (a) RT and (b) 288–427°C. 

Solid line represents lower-bound curve................................................................................ 65

34. Fracture toughness J-R curves for Linde 80 welds at 200 and 288°C and 

the lower-bound J-R curve for unaged austenitic SS welds. ................................................. 66

35. Fracture toughness J-R curves for aged stainless steel welds at (a) RT and (b) 288°C. 

Solid line represents lower-bound curve................................................................................ 67

36. Fracture toughness JIc for unaged and aged stainless steel welds. ...................................... 68

37. Fracture toughness lower-bound J-R curves and the data on unaged and aged 

304, 316L, and CF-3 welds used to develop the ASME Code IWB-3640 analysis. .............. 68

38. Fracture toughness J-R curve data for GTA, SMA, and SA welds in air at RT and 427°C. .. 69

39. Fracture toughness J-R curve data for unaged and aged Type 308 SS GTA and SMA 

welds in air at 427°C.............................................................................................................. 70

40. Fracture toughness J-R curve data for unaged and aged (a) Type 304/308 and (b) Type 

304/308L pipe welds in air at 25 and 290°C. ......................................................................... 71

41. Fracture toughness J-R curve data for as-welded Type 316L GTA weld at 288°C in 

air and BWR environment with 300 ppb DO.......................................................................... 72

42. Susceptibility of irradiated austenitic SSs to IGSCC as a function of fluence in high-DO 

water. ..................................................................................................................................... 74

43. Change in SCC growth rates at 289°C of LWR irradiated austenitic SSs as a function of 

yield strength in NWC and HWC BWR environments. .......................................................... 76
xiv

44. SCC growth rates in NWC BWR environment on austenitic SSs irradiated to 0.75–

37.5 dpa................................................................................................................................. 78

45. SCC growth rates in HWC BWR environment on austenitic SSs irradiated to 0.75–

37.5 dpa................................................................................................................................. 79

46. Normalized SCC growth rates in PWR environment for austenitic SSs irradiated to 0.75–

37.5 dpa................................................................................................................................. 81

47. SCC growth rates for nonirradiated and irradiated austenitic SS (a) SA and (b) SMA weld 

HAZ materials at 289°C in high-purity water with 600 or <50 ppb DO................................... 82

48. SCC growth rates for aged and unaged (a) Type 316L welds and (b) CASS materials 

in irradiated or nonirradiated conditions and tested at 289 or 320°C in high-purity water 

with 300–2,000 or <50 ppb DO.............................................................................................. 84

49. SCC growth rates of irradiated 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............ 85

50. SCC growth rates of nonirradiated and irradiated austenitic SS weld HAZs and welds 

and CASS materials 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................................................................ 87

51. Crack growth rates for irradiated (a) austenitic SSs and (b) CASS materials in HWC BWR 

water at 288°C or PWR environment at 320°C plotted as a function of neutron dose. ......... 88

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

irradiated in fast reactors. ...................................................................................................... 91

53. Change in fracture toughness JIc as a function of neutron exposure for LWR irradiated 

austenitic SSs. ....................................................................................................................... 92

54. Change in fracture toughness J-R curve and JIc as a function of neutron exposure for 

Type 304 SS at (a) 289 and (b) 320°C. ................................................................................ 94

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

Type 304 and 304L SS, and CF-8M steel in high-purity water at 289°C. .............................. 95

56a. Fracture toughness J-R curves for irradiated specimens of Type 304 SS SMA weld HAZ 

in air and NWC BWR environment. ....................................................................................... 97

56b. Fracture toughness J-R curves for irradiated specimens of Type 304L SA weld HAZ in air 

and NWC BWR environment. ................................................................................................ 97

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

at 289°C................................................................................................................................. 98

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

aged and irradiated CF-8M steel. .......................................................................................... 99
xv

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

test temperature..................................................................................................................... 100

60. The change in initiation toughness JIc of (a) wrought austenitic SSs and (b) CASS 

materials and weld metals as a function of neutron exposure............................................... 101

61. Plots of fracture toughness JIc values as a function of neutron dose for (a) austenitic SS 

welds, (b) CASS materials, and (c) austenitic SS HAZ. ........................................................ 103

62. Change in coefficient C of the power-law J-R curve for (a) wrought austenitic SSs and 

(b) CASS materials and weld metals as a function of neutron exposure............................... 105

63. Plots of constant C of the power-law J-R curve as a function of neutron dose for 

(a) austenitic SS welds, (b) CASS materials, and (c) austenitic SS HAZ. ............................. 107

64. Experimental values of J-integral at a crack extension of 2.5 mm for (a) wrought 

austenitic SSs and (b) cast materials and SS weld metals plotted as a function of neutron 

dose. ...................................................................................................................................... 111

65. Plots of fracture toughness J2.5 values as a function of neutron dose for (a) austenitic SS 

welds, (b) CASS materials, and (c) austenitic SS HAZ. ........................................................ 113

xvii

EXECUTIVE SUMMARY 

Background

Austenitic stainless steels (SSs) are used extensively as structural alloys in light water reactor 

(LWR) systems, including reactor core internal components, because of their excellent ductility, 

high notch toughness, corrosion resistance, and good formability. In addition, cast austenitic

stainless steel (CASS) materials are used for reactor pressure-boundary components, such as 

valve bodies, pump casings, primary coolant piping, and some reactor core internal 

components. Although austenitic SSs are completely austenitic in the wrought condition, 

welded SSs and CASS materials have a duplex structure consisting of austenite and ferrite 

phases. The ferrite phase increases the tensile strength and improves resistance to stress 

corrosion cracking (SCC). However, exposure to neutron irradiation for extended periods,

changes the microstructure (radiation hardening) and microchemistry (radiation-induced 

segregation) of wrought and CASS materials and their welds, degrades their fracture properties, 

and increases their susceptibility to irradiation-assisted stress corrosion cracking (IASCC). 

Furthermore, CASS materials and austenitic SS welds are susceptible to thermal embrittlement 

of the ferrite phase after extended operation at reactor operating temperatures.

The scope of this effort was to compile and evaluate the thermal and neutron embrittlement data 

on CASS materials and austenitic SS welds since the NUREG/CR-4513 Rev. 1 and 

NUREG/CR-6428 topical reports on these subjects were published, and (a) update the 

methodology for estimating the loss of fracture toughness of these materials during extended 

service at reactor operating temperatures, (b) define the threshold neutron fluence above which 

irradiation effects on mechanical properties are significant and develop disposition curves for 

IASCC growth rates for these materials used in the reactor core internals, and (c) identify 

information gaps in the data needed to address the effects of thermal and neutron embrittlement

on the structural and functional integrity of CASS components and austenitic SS welds in LWR 

systems. In addition, the effect of neutron irradiation on IASCC crack growth rates in these 

materials was evaluated. Furthermore, limited data suggest that the concurrent exposure to 

neutron irradiation during reactor service can result in a combined effect wherein the servicedegraded fracture toughness can be less than that predicted for either thermal or neutron 

irradiation embrittlement independently. The combined effects of thermal and neutron

embrittlement were also examined.

Thermal Embrittlement

Data on thermal embrittlement was compiled and evaluated as discussed below.

A. CASS Materials

Thermal aging of these materials increases their hardness and tensile strength and decreases 

ductility, impact strength, and fracture toughness. The low-carbon (low-C) American Society of 

Testing and Materials (ASTM) Specifications A351 Grade CF-3 CASS materials are the most 

resistant to thermal embrittlement, and the Mo-bearing, high-C Grade CF-8M CASS materials 

are the least resistant. The formation of chromium-rich α' regions in the ferrite is the primary 

mechanism for thermal embrittlement. The α' phase strengthens the ferrite matrix, which 

promotes brittle fracture even at higher temperatures. Furthermore, precipitation and/or growth 

of phase boundary carbides or nitrides results in brittle failure due to phase boundary 

separation, and may facilitate cleavage of the ferrite by particle cracking. The degree of
xviii

embrittlement and, hence, the toughness of the material, is controlled by the amount of brittle 

fracture. Consequently, the amount, size, and distribution of the ferrite phase in the duplex 

structure have a strong effect on the extent of embrittlement. 

Significant variations in ferrite content have been observed within CASS components. For 

example, the ferrite content ranged from 3.0 to 22.5% with an average of 13.3% ferrite for a hotleg elbow, and from 1.5 to 15.0% with an average of 9.8% ferrite for the crossover-leg elbow 

removed from the Ringhals reactor in Sweden. Similarly, the ferrite contents in centrifugally 

cast pipes were higher near the outer than the inner surface. In general, ferrite contents 

estimated from the material chemical composition using the Hull’s equivalent factors compared 

well with the values obtained using metallographic techniques. These observations indicated

that when measured ferrite content is used for design analyses, it is important to ensure that it is 

representative of the entire component. The distribution and morphology of ferrite in CASS 

materials is discussed. Parametric studies of the manufacturing process parameters have 

indicated that spinodal decomposition of the ferrite at reactor temperatures is very sensitive to 

the initial state of the ferrite. Process parameters that improve the homogeneity of the ferrite 

solid solution delay the beginning of the decomposition, which results in a more aging-resistant 

CASS material.

In this study, the degree of thermal embrittlement has been characterized in terms of the room 

temperature (RT) Charpy-impact energy. All CASS materials reach “saturation” RT impact 

energy (CVsat), the minimum impact energy that can be achieved for a material after long-term 

aging. The actual CVsat for a specific CASS material is independent of aging temperature but 

depends strongly on the chemical composition of the material. It is lower for ASTM A351 Grade 

CF-8M materials than for Grades CF-3 or CF-8, and decreases with an increase in ferrite 

content and the concentration of C or nitrogen (N) in the steel. The kinetics of thermal 

embrittlement are also discussed. The activation energy for the process of thermal 

embrittlement has been estimated from the chemical composition of the material; the values 

range between 65 and 105 kJ/mole (15 and 25 kcal/mole). 

The methodology developed earlier in NUREG/CR-4513, Rev. 1 for estimating the Charpyimpact energy and fracture toughness J-R curve of CASS components under LWR operating 

conditions from material information readily available in certified material test records was then

reviewed and, if necessary, updated to incorporate recent data published after the NUREG/CR4513 Rev. 1 report was published. The extent of thermal embrittlement as a function of time 

and temperature of reactor service was estimated from the extent of embrittlement at saturation 

and from the correlations describing the kinetics of embrittlement, which were expressed in 

terms of the chemical composition of the material. The fracture toughness J-R curve was then 

obtained from the correlation between the fracture toughness parameters and the RT Charpyimpact energy. Common lower-bound J-R curves at RT and 290–320°C for static and 

centrifugally cast Grades CF-3, CF-8, and CF-8M materials with <10%, 10–15%, 15–25%, and 

25–40% ferrite are presented. A significant modification of the earlier methodology based on 

the present evaluation is that separate expressions have been proposed for Grade CF-8M 

materials containing less than 10% Ni and those containing equal to or greater than 10% Ni. 

The latter are more susceptible to thermal embrittlement compared with the other grades.

The screening criteria established by the NRC to determine the susceptibility of CASS materials 

to thermal aging embrittlement, were then assessed to examine their validity. The results of the 

review indicated that for low-Mo CF-3 and CF-8 materials, the existing criterion for potentially 

susceptible material of greater than 20% ferrite for static-cast material and all centrifugally cast 

material are valid and somewhat conservative in some cases. However, for CF-8M materials, 
xix

the existing criterion for potentially susceptible material of greater than 14% ferrite for static-cast 

material and greater than 20% ferrite for centrifugally cast material needed to be updated. The 

updated criteria are presented. The criterion for static-cast and centrifugally-cast CF-8M 

materials has been decreased to 13% ferrite and 17% ferrite, respectively, for materials with 

less than 10% Ni, and to 10% ferrite and 13% ferrite, respectively, for materials with equal to or 

greater than 10% Ni. 

A flow diagram of the methodology for estimating mechanical properties of thermally aged 

CASS materials in LWR systems has also been updated. Three schemes are presented for 

estimation of mechanical properties of CASS material of (a) unknown composition, (b) known 

composition and unknown service history, and (e) known composition and service history. 

Lower bound fracture toughness J-R curves at RT and 290–320°C are presented for static-cast 

and centrifugally-cast CASS materials containing 25–40%, 15–25%, 10–15%, and <10% ferrite. 

However, the updated methodology does not consider the effect of reactor coolant environment 

on fracture toughness. Recent data indicate about a 40% reduction in fracture toughness J-R

curve data of long-term aged Grade CF-8 material at 54°C in PWR shutdown water chemistry 

relative to that in air. Additional fracture toughness tests on thermally aged CASS materials are 

needed to better understand the possible effects of LWR environments on fracture toughness.

B. Austenitic Stainless Steel Welds

Austenitic SS welds typically contain 5–15% ferrite. However, they do not have a single 

representative microstructure. Four distinct ferrite morphologies have been identified in 

austenitic SS welds: vermicular, acicular, lacy, and globular. The vermicular morphology is 

most commonly observed in welds with 5–15% ferrite, and a lacy or acicular morphology is 

observed in welds containing 13–15% ferrite. However, the mechanical properties of austenitic 

SS welds differ from those of CASS materials. For a given ferrite content, the tensile strength of 

welds is higher and fracture toughness is lower than that of CASS materials. The gas tungsten 

arc (GTA) welds exhibit the highest toughness and submerged arc (SA) welds the lowest. The 

toughness of shielded metal arc (SMA) welds is either comparable or slightly superior to that of 

SA welds. The JIc values for the SA welds are about one-third those for the GTA welds. In SS 

welds, failure is controlled by void formation near inclusions or other flaws in the material.

The mechanism of thermal embrittlement of austenitic SS welds is the same as that of CASS 

materials. However, since the ferrite content in welds is limited to about 15%, the degradation 

of mechanical properties occurs primarily because of the presence of grain/phase boundary 

carbides and second phase inclusions. The results presented earlier in NUREG/CR-6428, 

1996, indicate that the fracture toughness J-R curve for thermally aged austenitic SS welds 

might be decreased by up to 40%. The NUREG/CR-6428 report defined a lower-bound J-R 

curve for thermally aged SS welds. Furthermore, recent data indicate significant environmental 

effects on the fracture toughness of SS welds; fracture toughness in LWR environments can be 

up to 40% lower than in air. Additional data are needed to establish fracture toughness 

degradation of SS welds in LWR environments. 

IASCC Crack Growth Rates

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 IASCC 

susceptibility of austenitic SSs were discussed in an early study at Argonne, in NUREG/CR7027, 2010, and are summarized in this report. The existing data indicate that, in the normal 
xx

water chemistry (NWC) of boiling water reactors (BWRs), neutron irradiation up to 0.42 dpa has 

no effect on crack growth rates (CGRs). The CGRs for materials irradiated above 0.42 dpa are 

above the NUREG-0313 disposition curve for nonirradiated materials in high-purity water 

containing 8 ppm dissolved oxygen (DO). The CGRs can be up to a factor of 40 higher for 

materials irradiated to higher neutron dose levels. The growth rates for some irradiated SSs are 

a factor of 2–3 higher than the Electric Power Research Institute (EPRI) disposition curve for 

BWR core internal components. However, the data for highly irradiated materials show 

anomalous behavior. The CGRs in an NWC BWR environment, for SSs irradiated to 19 or 

38 dpa, are below the NUREG-0313 disposition curve for nonirradiated SSs. 

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 2.2 dpa. The CGRs increase rapidly at neutron dose levels above 2.2 dpa, and above 

about 4.0 dpa, the CGRs in HWC BWR environment are comparable to those in NWC BWR 

environment. This behavior is observed at least at stress intensity factor (K) values above 

18 MPa m1/2. Additional CGR data on SSs irradiated to 3–8 dpa are needed to better define the 

threshold neutron dose 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 CGRs obtained at 

different temperatures have been normalized to 320°C using an activation energy of 100 kJ/mol. 

In PWR water at 320°C, the CGRs for SSs irradiated above 3 dpa may be up to an order of 

magnitude above the NUREG-0313 curve for nonirradiated materials in high-purity water with 

0.2 ppm DO. The CGRs for SSs irradiated to 11–25 dpa may be up to two orders of magnitude 

above the same curve. As observed earlier for HWC BWR water, the CGRs in the PWR 

environment of SSs irradiated to 37.5 dpa exhibit anomalous behavior as the rates are only 

marginally above the NUREG-0313 curve. The reasons for the high CGRs for some SSs 

irradiated to 7–18 dpa and relatively low CGRs in SSs irradiated to 37.5 dpa are not known. 

The limited IASCC CGR data indicate that the CGRs for austenitic SS welds and weld heat 

affected zone (HAZ) materials, as well as CASS materials in NWC BWR environments, are 

comparable to those for wrought austenitic SSs. The CGRs increase rapidly for materials 

irradiated above 0.42 dpa. The growth rates for materials irradiated to 2.0–2.5 dpa may be up 

to an order of magnitude higher than the CGRs predicted by the NUREG-0313 curve for 

nonirradiated sensitized austenitic SSs in high-purity water with 8 ppm DO. The very limited 

CGR data for these materials in HWC BWR environments are mostly below the 0.2-ppm-DO 

NUREG-0313 curve. 

This study identified the following information gaps. There are no IASCC CGR data for neutron 

dose levels above about 2.3 dpa. Additional CGR data in LWR environments are needed for 

irradiated Type 316 or 316L welds. IASCC CGR data are also needed for CASS materials and 

austenitic SS welds and HAZ materials irradiated above 2.2 dpa, in HWC BWR environment at 

288°C or PWR primary water chemistry at 320°C. Furthermore, the effect of irradiation 

temperature should be established to better understand the IASCC susceptibility of austenitic 

SSs in PWR environments. 

Neutron Embrittlement 

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

fracture toughness (JIc) is above 150 kJ/m2 for Category III materials, and 30–150 kJ/m2 for 

Category II materials. These materials fracture after stable crack extension at stresses well 
xxi

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.

The fracture toughness data on irradiated wrought and cast austenitic SSs and their welds were 

compiled and evaluated in NUREG/CR-7027 (2010) to define the threshold neutron dose above 

which fracture toughness of these materials is reduced significantly. 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. In addition, 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 in.1/2), and samples at this toughness 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.

In NUREG/CR-7027, existing fracture toughness data were 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. 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 CASS materials and austenitic SS welds. 

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

of these materials is not accurately known. 

The limited data indicate that for the same irradiation conditions, the fracture toughness of 

thermally aged CASS materials and austenitic SS welds is lower than that of austenitic SS HAZ 

material, which, in turn, is lower than that of solution-annealed wrought SSs. However, the 

lower-bound fracture toughness trend curves developed in the Argonne studies may not be 

adequate for Type 316 welds. Furthermore, the available data also indicate that LWR 

environments can also decrease fracture toughness J-R curve of some of these materials. 

Typically, the effect of environment is larger for materials with higher fracture toughness. 

Furthermore, for some materials, the fracture toughness considering the effect of the reactor 

coolant environment is below the lower-bound curve. Additional data are needed to accurately 

establish the bounding fracture toughness of CASS materials and austenitic SS welds for LWR 

applications.

An assessment of the existing fracture toughness data on austenitic SS welds and associated 

HAZ materials, as well as CASS materials, indicate that some of the data for austenitic SS 

welds, particularly Type 316 welds, may not be bounded by the lower-bound fracture toughness 

versus neutron dose curve proposed in NUREG/CR-7027. The limited data on J-R curve tests 

in LWR coolant environments indicate that the reactor coolant environment can further decrease 

fracture toughness of these materials relative to that obtained in air. For some heats, fracture 

toughness in coolant environment is likely to be below the lower bound trend curve proposed in 

NUREG/CR-7027.
xxii

Combined Effects of Thermal and Neutron Embrittlement

The limited existing fracture toughness data indicate that for reactor core internal components, 

concurrent exposure to neutron irradiation can result in a combined effect wherein the servicedegraded 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 combined effects of thermal and neutron embrittlement on the 

threshold dose for embrittlement. Additional data are needed to better establish the potential 

loss of fracture toughness due to combined effects of thermal and neutron embrittlement in 

these materials in the transition dose range of 0.05 to 2.0 dpa. 

 
xxiii

ACKNOWLEDGMENTS 

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

Regulatory Commission, under Job Code V6455; Program Manager: Appajosula S. Rao.

xxv

ACRONYMS AND ABBREVIATIONS 

AMP aging management program

APFIM atom probe field ion microscopy

Argonne Argonne National Laboratory

ASME American Society of Mechanical Engineers

ASTM American Society for Testing and Materials

AWS American Welding Society

BWR boiling water reactor

C Carbon 

CASS cast austenitic stainless steels

CGR crack growth rate

CMTR certified material test record

CNSR Chevron notch short rod

Cr Chromium

CR composition ratio

CT compact tension

CTOD crack tip opening displacement

CW cold worked

DBTT ductile-brittle transition temperature

DO dissolved oxygen

dpa displacements per atom

EdF Electricité de France

EPFM elastic-plastic fracture mechanics

EPRI Electric Power Research Institute

Fe Iron

FN ferrite number

FRA Framatome

GE General Electric

GF George Fischer

GTA gas tungsten arc 

HAZ heat-affected zone

HWC hydrogen water chemistry

IASCC irradiation-assisted stress corrosion cracking

IG intergranular

IASCC irradiation-assisted stress corrosion cracking

IGSCC intergranular stress corrosion cracking

ISI in-service inspection

JAPEIC Japan Power Electric Engineering and Inspection Corp.

J-R J-integral resistance
xxvi

LEFM linear-elastic fracture mechanics

LWR light water reactor

MEA Materials Engineering Associates, Inc.

MMA manual metal arc 

Mn Manganese 

Mo Molybdenum

MRP Materials Reliability Program

MSIP mechanical stress improvement process

N Nitrogen

NDT nil-ductility transition

Ni Nickel 

NP National Power

NRC Nuclear Regulatory Commission

NWC normal water chemistry

PIFRAC pipe fracture (database)

PWR pressurized water reactor 

RIS radiation-induced segregation

RT room temperature

SA submerged arc

SANS small angle neutron scattering

SCC stress corrosion cracking

SE(B) single edge bend

SFE stacking fault energy

Si Silicon 

SMA shielded metal arc

SS stainless steel

SSRT slow strain rate tensile

TEM transmission electron microscopy

TG transgranular

Ti Titanium 

TWI The Welding Institute

USE upper-shelf energy

WH Westinghouse 
xxvii

NOMENCLATURE 

a Crack length

B Specimen thickness

Beff Effective specimen thickness 

C Coefficient of the power-law J-R curve

Creq Chromium equivalent for the material 

CV Room-temperature “normalized” Charpy-impact energy, i.e., Charpy-impact energy per 

unit fracture area, at any given service and aging time (J/cm2). The fracture area for a 

standard Charpy V-notch specimen (ASTM Specification E 23) is 0.8 cm2. The value 

of impact energy in J has been divided by 0.8 to obtain “normalized” impact energy in 

J/cm2.

CVint Initial room–temperature “normalized” Charpy-impact energy of a material, i.e., unaged 

material (J/cm2).

CVsat Room-temperature “normalized” Charpy-impact energy of a material at saturation, i.e., 

the minimum impact energy that would be achieved for the material after long–term 

service (J/cm2).

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)

Jd Deformation J per ASTM Specification E 813-85 or E 1152–87 (kJ/m2)

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

n Exponent of the power-law J-R curve

Nieq Nickel equivalent for the material

P Aging parameter, i.e., the log of the time of aging at 400°C

Q Activation energy for the process of thermal embrittlement (kJ/mole)

R load ratio

S Stress

t Service or rest or aging time 

teff Effective time of loading

tr Rise time

T Tearing modulus or temperature

W Specimen width

α Shape factor of the curve for the change in room-temperature Charpy-impact energy 

with time and temperature of aging
xxviii

β Half the maximum change in room-temperature Charpy-impact energy

δc Ferrite content calculated from the chemical composition of a material (%)

υ Poisson ratio

θ Aging behavior at 400°C, i.e., the log of the time to achieve β reduction in impact 

energy at 400°C.

σ Von Mises effective stress

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

σu Ultimate stress

σy Yield stress

φ Material parameter

In this report, all values of impact energy have been normalized with respect to the actual crosssectional area of the Charpy-impact specimen. Thus, for a standard Charpy-V-notch specimen 

per American Society for Testing and Materials (ASTM) Specification E 23 (i.e., 10 x 10-mm 

cross section and 2-mm V notch), impact energy value in J has been divided by 0.8 cm2 to 

obtain impact energy in J/cm2. Impact energies obtained on subsize specimens were 

normalized with respect to the actual cross-sectional area and appropriate correction factors 

were applied to account for size effects. Similarly, impact energies from other standards such 

as U-notch specimen were converted to a Charpy-V-notch value by appropriate correlations.

SI units of measurements have been used in this report. Conversion factors for measurements 

in British units are as follows:

To convert from to multiply by

in. mm 25.4

J* ft·lb 0.7376

kJ/ m2 in.-lb/in.2 5.71015

kJ/mole kcal/mole 0.239

* When impact energy is expressed in J/cm2, first multiply by 0.8 to obtain impact energy of 

a standard Charpy V-notch specimen in J.
1

1. INTRODUCTION 

Austenitic stainless steels (SSs) are used extensively as structural alloys in light water reactor 

(LWR) systems, including reactor core internal components, because of their excellent ductility, 

high notch toughness, corrosion resistance, and good formability. In addition, cast austenitic

stainless steel (CASS) materials are used in LWR systems for reactor pressure-boundary 

components, such as valve bodies, pump casings, primary coolant piping, and some reactor 

core internal components. Although these steels are completely austenitic in the wrought 

condition, welded SSs and CASS materials have a duplex structure consisting of austenite and 

ferrite phases. The ferrite phase provides additional benefits; it increases tensile strength and 

improves resistance to stress corrosion cracking (SCC).

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

(radiation hardening) and microchemistry (radiation-induced segregation, or RIS)1–5 of wrought 

and cast austenitic SSs and their welds, degrades their fracture properties6–21 and increases 

their susceptibility to irradiation-assisted stress corrosion cracking (IASCC).20–29 Furthermore, 

CASS materials and austenitic SS welds, which have a duplex structure, are susceptible to 

thermal embrittlement after extended operation at reactor operating temperatures,30–53 typically 

282°C (540°F) for boiling water reactors (BWRs), 288–327°C (550–621°F) for pressurized water 

reactor (PWR) primary coolant piping, and 343°C (650°F) for PWR pressurizers. 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, 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/σf

2), (1)

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. For small-scale yielding (such as deep 

cracks in bending in sufficiently large specimens), the fracture toughness can be characterized 

by JIc. Since JIc and KIc have different units, it is often convenient to represent JIc in terms of a 

parameter KJc, which has the units of the stress intensity factor and is determined from the JIc

value using the relationship

KJc = (E ́JIc)1/2, (2)

where the normalized elastic modulus is given by E ́ = E/(1 – υ2), E is the elastic modulus, and 

υ is the Poisson ratio. Only in cases where LEFM is applicable, is KJc equal to the critical stress 

intensity. 
2

The fracture toughness of austenitic SSs has been divided into three broad categories.8 

Category III corresponds to high toughness materials with JIc above 150 kJ/m2 (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 and with JIc in the 

range of 30–150 kJ/m2 (171–857 in.-lb/in.2). In Category II 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 JIc < 30 kJ/m2 (< 171 in.-lb/in.2), and fracture occurs below the 

yield stress with little or no controlled crack extension. 

Nonirradiated wrought and CASS materials 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/m2 (965 and 

9479 in.-lb/in.2), with a median value of 672 kJ/m2 (3837 in.-lb/in.2).

8 The JIc values at 400–

550°C (752–1022°F) are approximately 35% lower, with a median value of 421 kJ/m2 (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 CASS materials and SS welds exhibit ductile fracture at temperatures up to 550°C 

(1022°F), their fracture toughness is lower than that of the wrought SSs. In addition, a dimpled 

fracture-morphology is observed in SS welds.

53 Because of a high density of inclusions in the 

weld, the dimples are relatively small and shallow, and often associated with inclusions. The 

overall fracture toughness of CASS materials 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 centrifugallycast pipes. Gas tungsten arc (GTA) welds exhibit the highest toughness; shielded metal arc 

(SMA) welds have intermediate toughness; and submerged arc (SA) welds have the lowest 

toughness.

8 The median value of JIc is 492 kJ/m2 (2809 in.-lb/in.2) for GTA welds and 

147 kJ/m2 (839 in.-lb/in.2) for SA welds at temperatures up to 125°C (257°F). 

Welding of austenitic SSs results in a heat affected zone (HAZ) adjacent to the fusion zone, 

where the material microstructure and microchemistry are greatly altered because of the 

precipitation of chromium (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. 

A critical assessment of the susceptibility of austenitic SSs to IASCC was presented in U.S. 

Nuclear Regulatory Commission (NRC) topical report NUREG/CR-7027.21 The existing data 

were evaluated to establish the effects of material parameters (such as composition, thermomechanical treatment, microstructure, microchemistry, yield strength, and stacking fault energy) 

and environmental parameters (such as water chemistry, irradiation temperature, dose, and 

dose rate) on IASCC susceptibility. The results were used to (1) define a threshold-fluence 

above which irradiation effects on susceptibility to IASCC and fracture toughness of wrought 

and CASS materials and their welds are significant and (2) evaluate the potential of neutron 

embrittlement under LWR operating conditions. A superposition model is used to represent the 

fatigue crack growth rates (CGRs). The results indicate that for the same irradiation conditions, 

the fracture toughness of thermally aged CASS material and austenitic SS weld metal is lower 

than that of the HAZ of SS base materials, which, in turn, is lower than that of solution-annealed 

SS base materials. The potential deficiencies or knowledge gaps in the existing experimental 

data on degradation of LWR core internal materials due to neutron irradiation are also 

discussed. 
3

The thermal embrittlement of CASS materials and austenitic SS welds has also been 

investigated at Argonne National Laboratory (Argonne) and the results were published in 

several NRC topical reports.48–53 A procedure and correlations have been developed at 

Argonne for estimating fracture toughness, tensile, and Charpy-impact properties of CASS

components during service from known material information.50 The Argonne estimation scheme 

is applicable to compositions within the American Society for Testing and Materials (ASTM)

Specifications A351 for Grades CF-3, CF-3A, CF-8, CF-8A, and CF-8M. A correlation for 

Charpy-impact energy at saturation (i.e., the minimum impact energy that would be achieved for 

the material after long-term aging) is given in terms of chemical composition. Change in impact 

energy as a function of time and temperature of service is estimated from saturation impact 

energy and from the correlations that describe the kinetics of embrittlement, which are also 

given in terms of chemical composition. The fracture toughness J-R curve for the material is 

then obtained from the correlation between the fracture toughness parameters and the Charpyimpact energy. Tensile yield and flow stresses, and Ramberg/Osgood parameters are 

estimated from the flow stress of the unaged material and the kinetics of embrittlement.52

Although austenitic SS welds have a duplex structure, and their chemical compositions are 

similar to those of CASS materials, the Argonne correlations for CASS materials are not 

applicable to austenitic SS welds. The correlations account for mechanical-property 

degradation of typical heats of CASS, and do not consider the effects of compositional or 

structural differences that may arise from various welding processes, and that typically change 

the failure mechanism of the material. Degradation of fracture toughness and Charpy-impact 

energy of several SS pipe welds was evaluated in another NRC report, NUREG/CR-6428.53 

The welds were aged for 7,000 to 10,000 h at 400°C (752°F) to simulate saturation conditions; 

the lowest impact energy that would be achieved by the material after long-term aging. The 

results were compared with data from other studies. 

Furthermore, limited data suggest that the concurrent exposure to neutron irradiation during 

reactor service can result in a combined effect wherein the service-degraded fracture toughness 

can be less than that predicted for either thermal or neutron irradiation embrittlement 

independently. The combined effects of these two embrittlement processes are also evaluated. 

The scope of this report is to compile and evaluate the thermal and neutron embrittlement data 

on CASS materials and austenitic SS welds since the Argonne topical reports were published,

as well as the following: 

(a) Update the methodology for estimating the loss of fracture toughness of these materials 

during extended service at reactor operating temperatures,

(b) Define the threshold neutron fluence above which irradiation effects on mechanical 

properties are significant and develop disposition curves for IASCC growth rates for these 

materials used in the reactor core internals, and 

(c) Identify information gaps in the data needed to address the effects of thermal and neutron 

embrittlement on the structural and functional integrity of CASS components and austenitic 

SS welds in LWR systems.

Data have been obtained on thermally aged and unaged CASS materials (with ferrite content up 

to 40%) and austenitic SS welds with and without neutron irradiation. The methodology 

presented in NUREG/CR-4513, Rev. 1 for predicting the loss of fracture toughness of CASS 

components due to thermal aging during service at reactor operating temperatures has been 

updated to include CASS materials with up to 40% ferrite. The combined effects of thermal and 

neutron embrittlement on loss of fracture toughness of these materials have also been included. 
4

In addition, an assessment presented in NUREG/CR-6428 (1996) of aging degradation effects

on mechanical properties of austenitic SS welds, has been updated to establish the 

lower-bound fracture toughness J-R curve for SS welds during extended reactor operation. The 

effect of thermal and neutron irradiation on the bounding values of IASCC CGRs for these 

materials is evaluated as a function of neutron dose. The combined effects of thermal and 

neutron embrittlement as a function of ferrite content are also summarized.
5

2. THERMAL EMBRITTLEMENT OF DUPLEX STAINLESS 

STEELS 

It has been known that binary iron-chromium alloys and ferritic SSs are susceptible to severe 

embrittlement when exposed to temperatures in the range of 270 to 475°C (518 to 887°F).54–57 The 

potential for significant embrittlement of cast duplex SSs, consisting of both austenite and ferrite

phases, has been confirmed by studies at Argonne48–53 and elsewhere30–34,42–44 on materials that 

were aged at temperatures between 290–450°C (554–842°F) for times up to 70,000 h (~8 yr). The 

results indicate that thermal aging of CASS materials (ASTM Specification A-351 for Grades*

CF-3, CF-3A, CF-8, CF-8A, and CF-8M) at reactor operating temperatures of 270–350°C (518–

662°F) increases their hardness and tensile strength; decreases ductility, impact strength, and 

fracture toughness; and shifts the Charpy transition curve to higher temperatures. The effect of 

thermal aging is observed to decrease at temperatures above 400°C (752°F). For example, the 

extent of thermal embrittlement in CASS materials aged at 450°C (842°F) is less than that in 

materials aged for similar times at 400°C (752°F).20

Various grades and heats of CASS exhibit varying degrees of thermal embrittlement. The 

low-carbon (low-C) CF-3 steels are the most resistant to thermal embrittlement, and the 

Molybdenum (Mo)-bearing, high-C CF-8M steels are the least resistant. Ferrite morphology had 

a strong effect on the degree or extent of embrittlement, whereas small changes in material 

composition significantly altered the kinetics of embrittlement. In addition, the results of studies on 

the kinetics of thermal aging demonstrate that thermal embrittlement of cast duplex SSs can 

occur during the reactor design life of 40 yr.48 

In austenitic SS welds, the ferrite phase is desired for controlling the weld solidification behavior 

and inhibiting the formation of low-melting-point compounds such as sulfides and phosphides, 

which promote microfissuring.8 Since the ferrite phase is brittle at low temperatures, austenitic 

SS welds also exhibit a ductile-brittle transition temperature (DBTT) phenomenon. However, at 

ambient and elevated temperatures, the ferrite phase shows a ductile deformation behavior. 

The fracture toughness of Type 304/308 and 316/16-8-2 welds is dependent on the weld 

process but not composition.8 For a given weld process, both these weld metals exhibit similar 

fracture toughness.40 In general, the GTA welds exhibit the highest toughness and SA welds 

the lowest. The JIc values for the SA welds are about one-third those for the GTA welds. 

Type 308 SS welds generally contain 5–15% ferrite, but their mechanical properties typically 

differ from those of CASS materials. Earlier studies conducted at Argonne53 indicated that for a 

given ferrite content, the tensile strength of austenitic SS welds is higher and fracture toughness 

is lower than that of CASS materials. Experimental data53 indicate that CASS materials with 

very poor fracture toughness are relatively insensitive to thermal aging. In these steels, failure 

is controlled by void formation near inclusions or other flaws in the material (i.e., by processes 

that are not sensitive to thermal aging). These results suggest that austenitic SS welds with 

poor fracture toughness (e.g., SA and SMA welds) should be relatively less sensitive to thermal 

aging than GTA welds.

2.1 Mechanism of Thermal Embrittlement

Fracture of the 300 series austenitic SSs is essentially by nucleation, growth, and coalescence 

of microvoids resulting in a dimpled fracture morphology, irrespective of test temperature.8 The 

 *

The CF-3A and CF-8A grades represent high tensile strength material. The chemical composition of these grades is 

further restricted within the composition limits of CF-3 and CF-8, to obtain a ferrite/austenite ratio that results in 

higher ultimate and yield strengths. In this report, they are considered equivalent to CF-3 and CF-8 grades.
6

heat-to-heat variability is due to the differences in the density and morphology of inclusions such 

as carbides, calcium aluminates, and manganese sulfides, which serve as nucleation sites for 

void formation. Large inclusions or inclusion clusters aligned in the working direction fail early in 

the deformation process, thereby resulting in poor fracture toughness. In relatively tough 

materials, microvoids are nucleated away from the primary fracture plane, and significant plastic 

deformation is required for void coalescence. 

A similar fracture behavior is also observed in CASS materials and austenitic SS welds. 

Fracture occurs by microvoid coalescence. In CASS materials, voids nucleate preferentially 

within the ferrite phase, or at inclusions and phase boundaries, whereas in austenitic SS welds, 

microvoids forms at manganese silicates and silicides, oxides, and delta ferrite.8,41,43 In both 

materials, the overall fracture toughness is controlled by the density and morphology of the 

second phase particles, and for CASS materials, to some extent by the volume fraction of 

ferrite. The fracture toughness of austenitic SS welds is generally lower than that of cast or 

wrought austenitic SSs because of the higher density of inclusions. Among austenitic SS welds, 

the SA and SMA welds have poor fracture toughness relative to the GTA welds; they have a 

high density of manganese- and silicon-rich silicates and silicides. High silicon contents are 

generic to the SA and SMA welds because of silicon pickup from the flux. Typically SA welds 

have 0.6–1.0% Si, SMA welds have 0.5–0.8% Si, and GTA welds have less than 0.5% Si.8

Furthermore, in materials with a duplex structure, such as CASS and austenitic SS welds, the 

ferrite phase exhibits a ductile-to-brittle-transition temperature. Its plastic straining capacity is 

substantially decreased at low temperatures. However, the ferrite phase is ductile at room 

temperature and higher temperatures. Therefore, in the unaged condition, both CASS materials 

and austenitic SS welds exhibit a ductile dimpled fracture. The transition temperatures of 

unaged materials are relatively low. The differences in the transition temperature for the various 

unaged heats and grades of CASS materials are due to the amount of ferrite and the differences in 

the mechanism of brittle fracture. The high-carbon CF-8 or CF-8M steels have a higher transition 

temperature than CF-3 steels because of the presence of phase boundary carbides. The carbides 

weaken the boundaries and lead to premature phase boundary separation with little or no strain 

hardening. For austenitic SS welds, since the ferrite volume fraction is typically less than 15%, 

the ferrite content has little effect on the overall fracture toughness of welds. 

The available data indicate that the fracture toughness of wrought austenitic SSs and their 

welds is strongly influenced by specimen orientation. For austenitic SS pipe welds, specimens 

with a crack propagating in the axial direction (C-L orientation)* exhibit fracture toughness JIc

and tearing modulus values that are 60–70% lower than those propagating in the circumferential 

direction (L-C orientation).41 This behavior has been attributed to the alignment of stringers in 

the primary working direction (axial or longitudinal). When stringers are aligned parallel to the 

crack propagation direction C-L orientation), tear ridges are nucleated ahead of the crack front, 

the coalescence of which results in premature crack advance without extensive plastic 

deformation. When stringers are normal to the crack propagation direction (L-C orientation), 

microvoids are very deep and equiaxed, which blunts the advancing crack tip, and significant 

plastic deformation is needed for failure. Similarly, fracture toughness J-R tests on Type 304 

control-rod and Type 304L top guide materials irradiated to 4.7–12.0 dpa in a BWR, showed 

lower fracture toughness in the T-L orientation than in the L-T orientation.10 The lower fracture 

toughness along the T-L orientation is also due to the presence of stringers consisting of long, 

 *

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; C = circumferential direction for pipes; T = transverse direction for 

plates.
7

narrow particles oriented in the rolling direction, which result in a long and narrow quasicleavage structure parallel to the crack advance, thereby accelerating the crack advance.

The available data also indicate that thermal aging of CASS materials and austenitic SS welds 

at 290–450°C (554–842°F) results in thermal embrittlement of the ferrite, and the ductile-tobrittle-transition temperature for these materials is increased to higher temperatures.8,48–52

Thermal aging of CASS materials at 290–450°C leads to (a) spinodal decomposition of the 

ferrite into high-Cr α' and low-Cr α regions; (b) nucleation and growth of the high-Cr α' phase; 

(c) precipitation of a nickel (Ni)- and silicon (Si)-rich G phase, M23C6, and γ2 (austenite); and 

(d) additional precipitation and/or growth of existing carbides at ferrite/austenite phase 

boundaries.42,58–66 Thermal aging has little or no effect on the austenite phase. The formation 

of Cr-rich α' regions in the ferrite is the primary mechanism for thermal embrittlement. The α' 

phase strengthens the ferrite matrix and increases strain hardening, which raises the local 

tensile stresses above the critical value for brittle fracture at higher temperatures. The formation of 

G phase and precipitation and growth of carbides or nitrides at phase boundaries also 

contribute to some extent, particularly in high-C and Mo-bearing materials. Furthermore, the 

presence of these phases leads to a brittle failure by phase boundary separation, and may facilitate 

cleavage of the ferrite by particle cracking. 

Thermal aging of austenitic SS welds leads to spinodal decomposition of the ferrite to form the 

α' phase, formation of Ni- and titanium (Ti)-rich silicides (the G phase, Ti6Ni16Si7) in the ferrite, 

precipitation of M23C6 carbides on the phase boundaries, and limited M6C carbides in the 

matrix.8 However, the degradation of fracture toughness occurs primarily due to the presence 

of grain boundary or phase boundary carbides, which initiate localized intergranular cracking. 

Fracture occurs along the delta ferrite regions where the second phase particles initiate 

voids/cracks either by decohesion of the ferrite/austenite interphase or particle cracking.40 The 

dominant failure process is transgranular dimple fracture, and intergranular cracking is limited to 

a few isolated regions.8 

A study conducted at Electricité de France (EdF) on the evolution of thermally aged CASS 

materials, using Charpy U-notch impact test specimens, indicated that at room temperature, the 

brittle transgranular fracture of ferrite takes place primarily by cleavage; some twinning is also 

observed in highly embrittled material.43 Multiple-cleavage of ferrite islands is observed 

throughout the material. The cleavage facets are all parallel to each other, irrespective of the 

geometry of the individual islands. The percent of cleavage increases with increase in the 

degree of thermal embrittlement, or with decrease in test temperature. At higher temperature, 

although cleavage of the ferrite is reduced, twinning appears to plays an important role in brittle 

failure of ferrite. The ferrite/austenite phase boundary exhibits a jagged appearance with very 

fine lines and straight deformation bands within the ferrite islands (most likely slip bands and 

twins).43 The fracture behavior of the austenite also changes from a dimple fracture to shear 

fracture. The presence of completely embrittled ferrite skeleton promotes the low-energy failure 

mode of ductile shearing or tearing of the austenite ligaments between the islands of ferrite.

37

Typically, in duplex materials, cavities initiate from the cleavage cracks in the ferrite and not 

from matrix-inclusion decohesions. 

Another study on microstructural characterization and fracture behavior of unaged and aged CF-8M 

steel at room temperature using Charpy V-notch and fracture toughness compact tension (CT)

specimens also showed similar failure mechanisms.67 At room temperature, fully ductile dimple 

fracture was observed for unaged material and material aged up to 40,000 h at 300°C (572°F). 

Ductile dimple fracture and brittle cleavage facets were observed for material aged up to 10,000 h at 

350 and 400°C (662 and 752°F). In addition to these two fracture modes, ferrite/austenite phase 
8

boundary separation was also observed in materials aged at 450°C (842°F). Phase boundary 

separation is generally observed in high-C CASS materials aged at high temperatures (i.e., 400 or 

450°C). Fractographic evaluation of both Charpy V-notch and fracture toughness CT specimens 

showed similar fracture modes for aged CF-8M material. The only difference was that longer aging 

times were needed before brittle cleavage fracture was observed in the CT specimens. Although 

only room-temperature tests were conducted in this study, tests at higher temperatures are expected 

to show a larger fraction of ductile dimple fracture and less brittle fracture. 

Thus, severely embrittled CASS materials generally exhibit large areas of brittle transgranular 

fracture of ferrite, linked by ductile shearing or tearing of the austenite ligaments. At room 

temperature, brittle fracture of the ferrite is in the form of cleavage facets. The degree of 

embrittlement and, hence, the toughness of the material, is controlled by the amount of brittle 

fracture. CASS materials with poor toughness and impact strength exhibit greater fraction of brittle 

fracture. For some CASS materials, although a fraction of the material may fail in a brittle fashion, 

the surrounding austenite provides ductility and toughness. Such materials have adequate 

toughness even after long-term aging. A predominantly brittle failure occurs when either the 

ferrite phase is continuous (e.g., in cast material with a large ferrite content) or the 

ferrite/austenite phase boundary provides an easy path for crack propagation (e.g., in high-C 

grades of cast steel with large phase-boundary carbides). For CASS materials with lacy ferrite 

morphology, a ferrite content of about 10% or greater can provide a continuous ferrite phase. 

Consequently, the amount, size, and distribution of the ferrite phase in the duplex structure, and 

the presence of phase-boundary carbides are important parameters in controlling the degree or 

extent of thermal embrittlement.

Cleavage of ferrite occurs when the local tensile stress reaches the critical cleavage fracture stress. 

At low temperatures (i.e., high values of yield stress), cleavage cracks nucleate in the ferrite in the 

plastic zone near the notch tip at loads that are below general yielding. At high temperatures (i.e., 

low values of yield stress), strain hardening is needed to raise the local tensile stress to the cleavage 

fracture stress. Ductile fracture results when strain hardening is insufficient to raise the tensile 

stress to the critical value. The relationship between the degree of cleavage fracture and toughness, 

however, is complex since cleavage cracks can be initiated by several mechanisms (i.e., dislocation 

pile up, cracking of carbide or nitride particles, and cracking of twin intersections). Each of the 

mechanisms requires a unique stress level. Thus, for the same amount of cleavage fracture, the 

toughness may vary in different cast materials.

The time-temperature curves for the formation of various phases and the change in impact 

strength of thermally aged cast duplex stainless steel32 indicate that at temperatures above 

550°C (1022°F), the embrittlement is largely due to formation of sigma phase, and below 500°C 

(932°F), precipitation of α' phase leads to embrittlement. Furthermore, the time-temperature 

transformation curves for iron-chromium (Fe-Cr) alloys indicate that the α' phase is not stable at 

550°C (1022°F). At 550°C, the iron-chromium alloys are embrittled after aging for >10 h, owing 

to the formation of a sigma (σ) phase.33 Consequently, the fracture toughness of embrittled 

CASS materials can be recovered virtually completely by annealing for 1 h at 550°C (1022°F) 

and water quenching. This short heat treatment dissolves the α' phase but prevents formation 

of sigma phase;48 the dissolution of α' has been confirmed by microstructural studies.

42 
9

2.2 Distribution and Morphology of Ferrite

2.2.1 Ferrite Morphology

2.2.1.1 CASS Materials

The degree and kinetics of thermal embrittlement of duplex structures such as CASS materials 

are controlled primarily by the amount, size, and distribution of the ferrite phase, and to some 

extent the presence of carbides or nitrides at the phase boundaries. These material parameters 

in turn depend on the chemical composition and the manufacturing process. Differences in the 

thermal aging behavior have been observed is austenitic SS castings produced by different 

foundries, suggesting that manufacturing parameters may also be important for thermal aging. 

The effect of manufacturing process on the thermal embrittlement of CASS materials has been 

investigated by Massoud et al.68 The manufacturing process for castings consists of 

solidification of the material from a liquid state at around 1500°C (2732°F), followed by a 

homogenizing heat treatment of the solid at 1050–1150°C (1922–2102°F) and water quenching. 

Duplex SSs with high Cr contents solidify into a primary δ ferrite phase. The high-Cr SSs are 

also more sensitive to thermal aging. Typically, the core of the large sand-cast components 

consists of equiaxed grains, and the surface regions contain columnar grains, elongated in the 

direction of the temperature gradient; steeper temperature gradients result in smaller and more 

elongated columnar grains.68 Examples of equiaxed and columnar grain structures are shown 

in Figs. 1 and 2, respectively.

During cooling of the solidified material, if the concentrations of the γ-forming elements such as 

C, N, Mn, and Ni, are sufficient, a fraction of the δ-ferrite transforms to γ-austenite by solid-state 

δ to γ transformation. Furthermore, if the cooling rate between 900 and 600°C (1652 and 

1112°F) is slow, δ ferrite can transform to the brittle sigma (σ) phase, and M23C6 carbides can 

also precipitate at the δ/γ phase boundaries, both of which lead to degradation of the 

mechanical properties of the material.68 Therefore, the manufacturing processes require that 

Cast CF-8 Steel Cast CF-3 Steel

Axial Circum. Axial Circum.

Figure 1. Microstructure along the axial and circumferential sections of centrifugally-cast

CF-8 and CF-3 stainless steel pipes showing equiaxed grain structure.
10

the casting must be annealed at around 1100°C (2012°F) followed by a rapid quench to avoid 

the formation of these deleterious precipitates. 

The significant results of the parametric study of the manufacturing process parameters 

conducted by Massoud et al.68 are as follows: 

• The solidification speed affects the morphology of the ferrite-austenite microstucture and the 

characteristics of the toughness transition curve, but does not seem to affect the aging 

behavior of the steel.

• The homogenizing temperature especially affects the ferrite content of the material and the 

chemical composition of each phase. Thus, it affects the aging behavior of the steel.

• The homogenizing treatment holding time and the quenching rate affect the beginning of the 

decomposition of the ferrite and consequently the overall kinetics of aging.

The results indicate that the microstructural changes in the ferrite by spinodal decomposition 

during thermal aging at temperatures between 290 and 400°C (554 and 752°F), are very 

sensitive to the initial state of the ferrite in the as-quenched condition. Any manufacturing 

process parameter that improves the homogeneity of the ferrite solid solution, such as long 

homogenizing heat treatments and rapid quench, delay the beginning of the ferrite 

decomposition, which results in a more aging-resistant CASS material. 

The ferrite morphology of the various CASS materials varies with the ferrite content, chemical 

composition, and size of the casting. Studies conducted at Argonne on 30 heats of CF-3, CF-8, 

and CF-8M CASS materials showed that globular ferrite morphology was observed for materials 

containing <5% ferrite.48 Some differences in morphology were observed between the different 

grades of CASS materials containing >5% ferrite. The CF-8 and CF-8M steels had a lacy 

morphology, while the CF-3 cast steels showed a mixture of lacy and acicular ferrite. Significant 

variations in ferrite content within a CASS component have been observed in hot-leg and 

crossover-leg elbows removed from the Ringhals reactor.69 The material of the elbows was 

ASTM 351 CF-8M steel, with ferrite content determined from the modified Schaefler diagram, of 

20.1% for the hot-leg and 19.8% for the crossover-leg elbow. Test rings were cut from the two 

elbows and the ferrite content measured with a ferritemeter at the inside surface at three 

different heights of the ring, as well as through the thickness of the ring. The results from

around 700 measurements showed that the ferrite content varies significantly in the actual 

components. The ferrite content for the hot-leg elbow ranged from 3.0 to 22.5% with an 

average of 13.3% ±4.2% ferrite, and for the crossover-leg elbow ranged from 1.5 to 15.0% with 

an average of 9.8% ±3.4% ferrite. The results also indicate that for both elbows, the measured 

Figure 2. Microstructure along axial section of a check valve from the 

decommissioned Shippingport reactor showing columnar grains.
11

Centrifugally-cast CF-8 Stainless Steel (Heat P1) Pipe

Centrifugally-cast CF-3 Stainless Steel (Heat P2) Pipe

Static cast CF-3 Stainless Steel (Heat I) Pump Impeller

Figure 3. Ferrite content and morphology along circumferential section of centrifugally-cast

CF-8 (Heat P1) and CF-3 (Heat P2) pipes from regions near the inside and 

outside diameter, and two vanes of a static cast CF-3 pump impeller (Heat I).
12

average ferrite contents are comparable with the values estimated from the material chemical 

composition using the Hull’s equivalent factors or the ASTM A800/A800M methodology (see 

Section 2.2.2). 

Significant variations in the ferrite content were also observed in the Argonne studies on CASS 

materials, which were obtained from actual components. The morphology and measured ferrite 

content along circumferential sections of centrifugally-cast CF-8 (Heat P1) and CF-3 (Heat P2) 

pipes from regions near the inside and outside diameter, and two vanes of a static cast CF-3 

pump impeller (Heat I) are shown in Fig. 3. The results indicate that for both centrifugally-cast 

pipes, the ferrite contents are higher near the outer surface than the inner surface. The values 

are about 20% higher for the CF-3 pipe (Heat P2) and 40% higher for the CF-8 pipe (Heat P1). 

In addition, the ferrite content of pump impeller vane 3 was about 40% higher than that of 

Vane 1. The ferrite contents estimated from the chemical composition using Hull’s equivalent 

factors, were 17.7%, 15.6%, and 17.1 % for CF-8 pipe, CF-3 pipe, and CF-3 pump impeller, 

respectively. The estimated values are comparable to the measured ferrite contents for the 

CF-3 pipe and pump impeller, but are lower than the measured values for the CF-8 pipe. 

These observations indicate that if measured value of ferrite content is used for design 

calculations or analyses, several locations should be selected across the length and width of the 

component to ensure that it is representative of the entire component. The results also indicate 

that for centrifugally-cast SS pipes, the extent of embrittlement is likely to increase from the 

inner surface towards the surface because of the increasing ferrite content.

2.2.1.2 Austenitic Stainless Steel Welds

A duplex structure is also observed in austenitic SS welds, with ferrite being the minor phase 

distributed in various forms in the austenite matrix. Similar to the CASS materials, the 

solidification behavior and subsequent solid state transformation within the weld metal during 

cooling control the microstructural characteristics of the weld. The solidification sequences in 

Type 308 SS welds include primary crystallization of δ-ferrite with subsequent envelopment by 

austenite, followed by further transformations from liquid to γ and δ to γ.

70 As the sample cools 

to a temperature below that of the solidus, the transformation at the liquid-g interface is 

completed, leaving behind a skeletal network of untransformed d ferrite along the cores of the 

primary and secondary dendrite arms. This residual ferrite is rich in Cr, which makes it very 

stable.70 However, primary ferrite with a lower average Cr concentration of about 24–25 wt.%, 

may transforms into Widmanstatten austenite and ferrite during rapid cooling. These two types 

of transformations involve extensive solute redistribution by diffusion, and may be used to 

explain the various ferrite morphologies observed in austenitic SS welds.70 

Austenitic SS welds do not have a single representative microstructure. Examples of typical 

ferrite morphology of four different welds is shown in Fig. 4. PWWO is a 0.305-m (12-in.) 

schedule 100 pipe mockup weld with overlays, PWCE is a 0.71-m (28-in.) Type 304 pipe weld, 

PWDR is a 0.254-m (10-in.) Type 304 SS pipe weld after service in the Dresden reactor, and 

PWMS is a 0.71-m (28-in.) SS pipe weld treated by the Mechanical Stress Improvement 

Process (MSIP). In another study on the ferrite morphology and variations in ferrite content in 

two austenitic SS welds, four distinct ferrite morphologies were identified in Type 308 SS 

multipass welds: vermicular, lacy, acicular, and globular.70 These morphologies were observed 

on the same multipass weld. The welds were deposited by the GTA process on a 25-mm thick 

Type 304L SS plate (0.019 C, 1.75 Mn, 0.63 Si, 0.029 P, 0.006 S, 10.0 Ni, 18.55 Cr, and 

balance Fe [wt.%]) containing a single-V butt joint, using Type 308 SS filler metal (0.016 C, 1.95 

Mn, 0.35 Si, 0.029 P, 0.004 S, 9.76 Ni, 20.14 Cr, and balance Fe [wt.%], and Creq/Nieq = 1.66). 
13

One of the welds was made with the joint surfaces buttered with the weld metal and the other 

without buttering. Ferrite number (FN) was measured using a Magne-gage in accordance with 

American Welding Society (AWS) A4.2-74 (see Section 2.2.2).

71 

The average FNs of the root pass deposited in the buttered and not buttered weld were 13 and 

8, respectively. The lower ferrite content of the weld prepared without buttering was attributed 

to weld-metal dilution with the base metal; it is indicative of a change in weld metal composition. 

Furthermore, the FN value of the root pass decreased further after the second and third passes 

because of the dilution of ferrite from the thermal effects to which the root pass is subjected 

during a multipass weld. Variations in ferrite content were also observed in both welds within a 

cross section of the bead, along the length and width of the weld. The FN values* at various 

weld locations varied from 9 to 13 and 5 to 14 for the buttered and not-buttered welds, 

 * FN values determined from the modified Schaeffler diagram.

 

PWWO PWCE

 

PWDR PWMS

Figure 4. Typical ferrite morphology of four different welds.
14

respectively. Based on the composition of the weld metal, the calculated ferrite content for the 

weld, without dilution, is 8.1% from Hull’s equivalent factor,72 5.9 from the ASTM A800/A800M 

methodology (based on the Schoefer diagram),73,74 and about 13.8% from the modified 

Schaeffler (or Delong) diagram.75,76 

The value based on the ASTM A800 methodology is significantly lower than the average 

measured value of FN 11 for the buttered weld. However, since the ferrite measurements were 

made using a Magne-gage, such instruments are very sensitive to surface roughness or surface 

curvature. Furthermore, phases other than ferrite and austenite may form at higher 

temperatures during welding, which may alter the magnetic response of the material such that 

the indicated ferrite content is quite different than of the same material not subjected to the 

welding process. It should also be mentioned that the variations in ferrite content in part might 

be due to differences in the nitrogen (N) pickup during welding.70

The vermicular morphology is most commonly observed in austenitic SS welds with FN 5–15, 

and was predominantly observed in the weld root pass and the two subsequent passes. The 

ferrite content varied between FN 9 and 12. The vermicular (meaning the form, markings, 

motion, or tracks of worms) morphology, depending on the sectional cut viewed, appears as an 

aligned skeletal network of ferrite or as a curved skeletal form. The alignment is along the heat 

flow direction, which is also the primary dendrite growth direction. 

A lacy morphology was observed predominantly in the third pass of the weld; the ferrite content 

varied between FN 13 and 15. The lacy structure looked very regular and aligned. The lacy 

form of ferrite is characterized by long columns of interlaced ferrite oriented along the growth 

direction in an austenite matrix. Most likely, it forms by the transformation of primary δ-ferrite to 

Widmanstatten austenite and ferrite. 

An acicular morphology was present in the sixth and crown passes of the weld; the ferrite 

content was about FN 14. However, unlike the previous two morphologies, the acicular 

structure had no directionality and did not conform to the solidification substructure in any way. 

This morphology is typical of weld metals with Creq/Nieq = 2. It also forms by the lowtemperature transformation of primary ferrite to austenite and ferrite. 

The globular form is characterized by ferrite in the form of globules randomly distributed in a 

matrix of austenite. The structure has no directionality and is not related to the overall 

solidification substructure. It was commonly observed in weld passes 4, 5, and 6 of the multipass weld. The ferrite content was around FN 10. A globular morphology is formed because of 

thermal instability of any of the other forms of ferrite, particularly the acicular form. 

It should be noted that these ferrite morphologies may appear continuous or discontinuous 

depending on the section of the weld viewed. Therefore, caution must be exercised in such 

characterization of the ferrite since only vermicular and lacy morphologies are continuous.

2.2.2 Estimation of Ferrite Content

As discussed above, the ferrite content in the CASS structure is essentially a function of the 

chemical composition and the manufacturing process history of the material. Typically, the 

ferrite content of the CASS duplex structures is determined by from the (a) chemical 

composition, (b) magnetic response, or (c) metallographic examination of the material. Among 

the magnetic methods, the Magne-Gage and Ferritescope are the most commonly used 

instruments for measuring the ferrite content in CASS materials or austenitic SS welds. The 
15

Magne-Gage is a continuous-reading instrument that utilizes a spring to measure the attraction 

between a magnet and the material of unknown ferrite content, and the response is compared 

with that of a calibrated sample. The Ferritescope operates on the magneto-induction principle 

where the relative magnetic permeability of the specimen is measured. 

However, because the magnets or the probes of these instruments are small, the surface 

roughness or curvature of the sample is an important parameter that can vary the magnetic 

linkage with the material being measured. In addition, phases other than ferrite and austenite 

may form in the material during service, which may alter the magnetic response of the material 

such that the indicated ferrite content is quite different from that of the same chemical 

composition that has undergone a different heat treatment.

Until 1973, ferrite contents in duplex structures such as CASS materials and austenitic SS 

welds were determined by metallographic examination of the structure. A sample of the 

material was polished and etched to reveal the ferrite and austenite phases, and a grid was 

superimposed over the image of an optical microscope to determine by point counting the 

percentage of ferrite in the sample. The main drawback with this method is that the point count 

estimates of ferrite may vary with the etching technique used to reveal the ferrite phase, and 

with the number of grid points used in the measurements. Furthermore, as discussed above, 

the ferrite content in most CASS components varies significantly in different regions of the 

component, and obtaining metallographic samples from various regions may not be practical. 

Although a quantitative metallographic method provides the most accurate estimate of ferrite 

content, determination of ferrite percent from chemical composition of the material offers the 

most useful and most common method for ferrite control during solidification of the metal from a 

melt during casting or welding. However, the accuracy of these estimations depends on the 

accuracy of the chemical analysis procedure, and the degree of variability of composition within 

the casting or weld. The most commonly used methods are described below.

2.2.2.1 Hull’s Equivalent Factor

When a certified material test record (CMTR) is available, the ferrite content is calculated from 

chemical composition in terms of Hull's equivalent factors for nickel and chromium given by

Creq = Cr + 1.21(Mo) + 0.48(Si) – 4.99 (3)

and

Nieq = (Ni) + 0.11(Mn) – 0.0086(Mn)2 + 18.4(N) + 24.5(C) + 2.77, (4)

where the concentrations of the various alloying and interstitial elements is in wt.%. The 

concentration of N is often not available in a CMTR; if not known, it is assumed to be 0.04 wt.%. 

The ferrite content δc is given by

δc = 100.3(Creq/ Nieq)2 – 170.72(Creq/ Nieq) + 74.22. (5)

The measured ferrite content and values calculated from Hull's equivalent factor for the various 

CASS heats used in studies at Argonne,

48 the Georg Fischer Co. (GF),30 EdF,43 National Power 

(NP),44 Framatome (FRA),

34 and the Electric Power Research Institute (EPRI)35 are shown in 

Fig. 5. For most heats, the difference between the estimated and measured values is ±6% 

ferrite. The results also indicate that the calculated ferrite content was generally lower than the 
16

measured values for CF-8M heats that contained ≥10% Ni.

Figure 5.

Plots of measured ferrite content 

and values calculated from Hull's 

equivalent factor for various 

CASS materials.

2.2.2.2 ASTM A800/800M Methodology

The ferrite content of the casting shall be estimated from the central line of the of the diagram at 

the composition ratio of chromium equivalent, Creq, to nickel equivalent, Nieq, determined from 

the following formula: 

(Creq/Nieq) = (Cr + 1.5Si + 1.4Mo + Nb – 4.99)/(Ni +30C + 0.5Mn + 26(N – 0.02) + 2.77) (6)The 

values of the composition ratio (CR) for a given ferrite content (F), or vice versa, is then 

determined mathematically from the equation of the central line:

CR = 0.9 + 3.38883 × 10–2F – 5.58175 × 10–4F2 + 4.22861 × 10–6F3 (7)

The measured ferrite content and values calculated from ASTM A800/A800M methodology for 

the same heats of CASS materials plotted in Fig. 5 are shown in Fig. 6. The results indicate 

that for ferrite contents greater than 20%, the calculated ferrite content for several heats is lower 

than the measured values. Most of these heats with significantly lower calculated values 

contained 22.0–23.0 wt.% Cr and about 8.0–8.5 wt.% Ni. Compared to the ferrite content 

calculated from Hull’s equivalent factor, the ASTM A800/A800M methodology under predicts the 

ferrite content for CASS materials with greater than 15% ferrite. The difference between the two 

methods can be seen clearly in Fig. 7.

The measured ferrite content and values calculated from Hull’s equivalent factor or the ASTM 

A800/A800M methodology for several austenitic SS welds used in studies at Argonne and by 

Slama et al., Mills, and Lucas et al., are plotted in Figs. 8a and 8b, respectively. The results 

indicate that the measured ferrite contents in the welds are significantly higher than the 

calculated values. They are about 50% higher than the values calculated from Hull’s equivalent 

factor, and are 80% higher than those determined from the ASTM A800 methodology. 

Furthermore, as seen before for CASS materials, the ferrite content calculated from the ASTM 

A800 methodology are about 20% lower than those calculated from the Hull’s method.

0

10

20

30

40

50

0 10 20 30 40 50

ANL

Framatome

George Fischer

EdF

National Power

EPRI

Ferrite Content from Hull's Equivalent Factor (%)

Measured Ferrite using Ferritescope (%)

Cast CF-3, CF-8, & CF-8M SSs

Open Symbols: CF-3 & CF-8

Closed Symbols: CF-8M
17

Figure 6.

Plots of measured ferrite content and 

values calculated from ASTM 

A800/A800M methodology for various 

CASS materials.

Figure 7.

Plots of ferrite content calculated 

from Hull's equivalent factor and 

those estimated from the ASTM 

A800/A800M methodology for 

various CASS materials.

2.3 Kinetics of Thermal Embrittlement

The degree of embrittlement has generally been characterized in terms of Charpy-impact 

energy of notched toughness specimens. The current “best estimates” of the degree of 

embrittlement at reactor operating temperatures are obtained from Arrhenius extrapolations of 

laboratory data obtained at higher temperatures.30 The aging time to reach a given degree of 

embrittlement at different temperatures is determined from the equation: 

0

10

20

30

40

50

0 10 20 30 40 50

ANL

Framatome

George Fischer

EdF

National Power

EPRI

Ferrite Content from ASTM A800/A800M Methodology (%)

Measured Ferrite using Ferritescope (%)

Cast CF-3, CF-8, & CF-8M SSs

Open Symbols: CF-3 & CF-8

Closed Symbols: CF-8M

0

10

20

30

40

50

0 10 20 30 40 50

CF-3 and CF-8 CASS Materials

CF-8M CASS Materials

Ferrite Conent from ASTM A800/A800M Methodology (%)

Ferrite Content from Hull's Equivalent Factor (%)

Cast CF-3, CF-8, & CF-8M SSs
18

Figure 8a.

Plots of measured ferrite content 

and values calculated from Hull's 

equivalent factor for various 

austenitic SS welds.

Figure 8b.

Plots of measured ferrite content 

and values calculated from ASTM 

A800/A800M methodology for 

various austenitic SS welds. 

(8)

where Q is the activation energy, R the gas constant, T the absolute temperature, and P 

an aging parameter that describes the combined effect of time and temperature on aging 

represents the degree of aging reached after 10P h at 400°C (752°F). Thus, P = 1 for 

0

5

10

15

0 5 10 15

316L

308

308L

16-8-1

Ferrite Content from Hull's Equivalent Factor (%)

Measured Ferrite Content (%)

Austenitic SS Welds

Open Symbols: SA or SMA Welds

Closed Symbols: GTA Welds

0

5

10

15

0 5 10 15

316L

308

308L

16-8-1

Ferrite Content from ASTM A800/A800M (%)

Measured Ferrite Content (%)

Austenitic SS Welds

Open Symbols: SA or SMA Welds

Closed Symbols: GTA Welds

t = 10P exp

Q

R

1

T − 1673

⎧

⎨

⎩

⎫

⎬

⎭

⎡

⎣

⎢ ⎤

⎦

⎥
19

aging 10 h at 400°C. The aging parameter for any given aging condition is obtained by 

rewriting Eq. 8 such that, . (9)

The activation energy for the process of embrittlement has been described as a function of the 

chemical composition of the cast material.34 Thus,

Q(kJ/mole) = –182.6 + 19.9(% Si) + 11.08(% Cr) + 14.4(% Mo),

or

Q(kcal/mole) = –43.64 + 4.76(% Si) + 2.65(% Cr) + 3.44(% Mo). (10)

The activation energy calculated from Eq. 10 for the process of embrittlement, ranges between 

65 and 105 kJ/mole (15 and 25 kcal/mole) for CF-3, CF-8, and CF-8M cast SSs. These values 

are significantly lower than the activation energy of spinodal decomposition in CASS materials. 

The spinodal decomposition and G-phase precipitation in low-temperature-aged cast duplex SS

have been investigated by transmission electron microscopy (TEM), atom probe field ion 

microscopy (APFIM), small-angle neutron scattering (SANS), and extraction replica 

techniques.58–66 The activation energy of the spinodal reaction in CF-3 SS was 250 ± 

30 kJ/mole (~60 ± 7 kcal/mole).63 This value is comparable to that for Cr diffusion in Fe-Cr 

alloys. The lower values of the activation energy of thermal embrittlement of CASS materials 

are most likely due to other factors, such as the effect of the formation of carbides and nitrides

at the phase boundaries, or G-phase and/or γ2 precipitation in ferrite, all of which may change

the fracture mechanism of the aged material. 

For example, precipitation of large carbides or nitrides at phase boundaries can initiate phase 

boundary separation by particle cracking. Consequently, a lower degree of spinodal decomposition

(i.e., smaller amplitude of Cr fluctuation) is needed for a given change in mechanical properties. The 

material would show a faster reduction in impact strength relative to a material without phase 

boundary carbides. However, precipitation of carbides or nitrides occurs primarily at 400 or 450°C 

and is extremely slow at lower temperatures. Thus, the influence of phase boundary carbides would 

tend to increase the apparent activation energy of embrittlement measured from mechanicalproperty data. 

The other factor that can influence the overall activation energy for embrittlement is the precipitation 

of other second-phase particles in ferrite, in particular the G phase, a multicomponent phase 

consisting of Ni, Si, Mo, Cr, Fe, and some Mn, and C.

61,65 The kinetics of G-phase precipitation 

depend on the chemical composition of the cast material.42,58 For some heats, G phase is observed 

after times as short as 10,000 h at 400°C, while other heats require up to 70,000 h of aging at 400°C 

for G-phase formation. In general, precipitation of G phase is faster in the molybdenum-containing 

CF-8M steels.42,58,59,61 The aging conditions for which G phase has been detected by TEM or 

SANS techniques in various cast SSs, are shown in Fig. 9. The kinetics for the decrease in the 

Charpy-impact energy of the aged CASS material are also plotted in the figure. The actual aging 

times for a given decrease in impact energy varied significantly for the various heats (shown by the 

horizontal scatter bars in Fig. 9). Generally, the aging times for the CF-8M steels were lower from

those for the CF-3 or CF-8 steels. 

At 400°C, the reduction in impact energy appears to be essentially complete before G phase is 

detected in any of the heats. For example, G phase was not detected in any of the heats aged for 

P = log10 (t) − 1000Q

19.143

1

Ts + 273 − 1673

⎛

⎝

⎜

⎞

⎠

⎟
20

3,000 h, whereas an 80% reduction in impact energy occurred by 4,600 h in all heats. However, at 

temperatures <350°C, G-phase precipitation and the decrease in impact energy occurred 

simultaneously. At 320°C, only a 50% reduction in strength occurred in most heats by 30,000 h, and 

G phase was observed in the specimens. These results indicate that the influence of G-phase 

precipitation on embrittlement would be greater at low temperatures, since at 400°C the kinetics of 

spinodal decomposition are much faster than G-phase precipitation, which follows nucleation and 

growth.

The mechanism by which the G phase influences thermal embrittlement of CASS materials is not 

well understood. The precipitation of G phase can influence the kinetics of embrittlement by either 

directly altering the kinetics of spinodal decomposition or changing the deformation and fracture 

behavior of the ferrite matrix and thereby influencing the effectiveness of spinodal decomposition. 

The only experimental data on the kinetics of spinodal decomposition in cast duplex SS were 

obtained by modeling the amplitude of Cr fluctuations, measured by APFIM, in thermally aged CF-3 

steel.63 The results yield an activation-energy of 250 kJ/mole. No G phase was observed in the 

steel after 5,000 h at 400 or 350°C. As shown in Fig. 9, most heats require ≥10,000 h at 400°C and 

≥30,000 h at 350°C before G-phase precipitates are detected. 

The low values of activation energy obtained from mechanical property data most likely are due to

the effect of G-phase precipitation on the deformation behavior of the ferrite matrix. The concomitant 

precipitation of G phase may alter the frequency (spacing) of Cr fluctuations produced by spinodal 

decomposition, which would be more effective in strain hardening. Thus, a lower degree of spinodal 

decomposition (i.e., lower amplitude of the Cr fluctuations) would be needed for a given change in 

mechanical properties; embrittlement would be faster. 

For a specific composition of the CASS material, Eqs. 9 and 10 can be used to determine the 

aging conditions that are representative of end-of-life reactor service. For a material with an 

activation energy of 100 kJ/mole (24 kcal/mole), the end-of-life condition for cold-leg piping 

(40 yr at 290°C [555°F]), is equivalent to 10,000 h at 400°C. The end-of-life condition for hot-leg 

piping (40 y at 320°C [610°F]), is equivalent to 30,000 h at 400°C. Consequently, the laboratory 

Figure 9. Arrhenius plots for the formation of G phase and reduction in impact energy.

1.2

1.3

1.4

1.5

1.6

1.7

1.8

1.9

101 102 103 104 105 106

50% Decrease in Impact Energy

G Phase Detection

80% Decease in Impact Energy

1000/T(K)

Aging Time (h)

Thermal Aging of CASS

50% Dectease in 

Impact Energy

G Phase Detected

80% Dectease in 

Impact Energy

400°C

350°C

300°C
21

data obtained for materials aged at 400°C are currently used to predict the end-of-life impact 

strength at reactor temperatures.

The above methodology, however, assumes a unique aging behavior at 400°C (752°F), which is 

not observed for CASS materials produced at different foundries using different manufacturing 

processes. The decrease in room-temperature (RT) Charpy-impact energy during thermal 

aging at 400°C (752°F) of various heats of cast stainless steel30,34,35,48,77,78 is shown in Fig. 10. 

The results indicate that all materials reach “saturation” RT impact energy (i.e., a minimum 

value that would be achieved by the material after long-term aging). The actual saturation value 

of RT impact energy for a specific CASS material is independent of aging temperature but 

depends strongly on the chemical composition of the steel. It is lower for the Mo-bearing CF-8M 

steels than for the Mo-free CF-3 or CF-8 steels, and decreases with an increase in ferrite 

content and the concentration of C or N in the steel. 

Figure 10. Decrease in Charpy-impact energy for various heats of cast austenitic SSs 

aged at 400°C.

However, Fig. 10 indicates that for a given decrease in the RT Charpy-impact energy, the time 

for aging at 400°C varies by more than two orders of magnitude for the various heats. For 

example, the time required for the impact energy to start decreasing varies from less than 100 h 

(e.g., the EPRI heat) to more than 1,000 h (GF heats). The time for the start of thermal aging 

effects varies from 50 to 500 h for the Argonne and FRA heats. Furthermore, for some 

materials, the decrease in impact energy is very fast (i.e., low activation energy) and for others it 

is slow (i.e., high activation energy). Typically, CASS materials that take longer for thermal 

embrittlement to start have low activation energy and materials that take short times for

embrittlement to start have high activation energy. The results indicate that for most CASS 

materials, the time to reach saturation impact energy is the same for all grades and heats of 

CASS materials; the aging time for saturation is independent of material composition or 

manufacturing process. The logarithm of the aging time at 400°C (parameter θ) for a 50% 

reduction in RT Charpy-impact energy (parameter β) has been shown to be an important 

parameter for characterizing the kinetics of thermal embrittlement.50 Table 1 lists the values of 

0

50

100

150

200

250

300

350

400

101 102 103 104 105

ANL

George Fischer

Framatome

EPRI

Impact Energy, C

V (J/cm2)

Aging Time (h)

Cast Austenitic Stainless Steels

Aged at 400°C
22

parameters β and θ for the various heats of CASS materials shown in Fig. 10; parameter α is a 

shape factor. 

The activation energies of thermal aging embrittlement vary from 65 to 230 kJ/mole (15 to 

55 kcal/mole).30,34,43,48,50,58,59,77–79 Small changes in the composition of the material seem to 

cause the kinetics of thermal embrittlement to vary significantly. Microstructural examination of 

aged CASS materials suggests that steels that take longer for embrittlement to start at 400°C 

and have low activation energy are associated with clusters of Ni-Si, Mo-Si, and Ni-Si-Mo in the 

ferrite matrix.42,58,59 These clusters are considered precursors of G-phase nucleation and 

precipitation. CASS materials with low activation energy (i.e., fast embrittlement) show G-phase 

precipitation after aging but embrittlement at 400°C takes longer to start. CASS materials with 

high activation energy (i.e., slow embrittlement) do not contain a G phase but embrittlement at 

400°C takes relatively short time to start.

42,58–61 The presence of Ni-Si-Mo clusters in the ferrite 

matrix of an unaged material may be considered a signature of steels that are potentially 

sensitive to thermal embrittlement; steels with Ni-Si-Mo clusters in the ferrite matrix show low 

activation energy for thermal embrittlement but take longer to embrittle at 400°C.

Microstructural characterization and annealing studies on thermally aged cast stainless steel 

show that strengthening of ferrite is caused primarily by spinodal decomposition of ferrite to form 

Cr-rich α' phase.42,48,49 Consequently, the kinetics of thermal embrittlement should be 

controlled by the amplitude and frequency of Cr fluctuations produced by spinodal

decomposition (i.e., by the size and spacing of the α' phase). The low activation energies of

thermal embrittlement are most likely caused by variations in the spacing of Cr fluctuations;

atom probe field-ion microscopy studies indicate that the spacing between Cr fluctuations 

decreases with decreasing temperature.60,64 During thermal aging, production heat treatment

and possibly the casting process, both of which affect ferrite composition and microstructure of

unaged material, can influence microstructural evolution and, therefore, the kinetics of 

embrittlement.

2.4 Extent of Thermal Embrittlement of CASS Materials

As discussed in the previous section, all CASS materials reach “saturation” RT impact energy

(i.e., a minimum value that would be achieved by the material) after long-term aging. The actual 

value of saturation RT impact energy for a specific cast SS is independent of aging temperature 

but depends strongly on the chemical composition of the steel. It is lower for the Mo-bearing 

CF-8M steels than for the Mo-free CF-3 or CF-8 steels, and decreases with an increase in 

ferrite content or the concentration of C or N in the steel.50 Typically, the extent of thermal 

embrittlement has been characterized by the RT “normalized” Charpy-impact energy (Charpyimpact energy per unit fracture area). 

In the Argonne studies,48–51 correlations were developed for the extent of thermal embrittlement 

at “saturation” (i.e., the minimum Charpy-impact energy that would be achieved for the material 

after long-term aging) in terms of the chemical composition of the material. The extent of

thermal embrittlement as a function of time and temperature of reactor service was estimated 

from the extent of embrittlement at saturation and the correlations describing the kinetics of 

embrittlement, which are also given in terms of chemical composition. The fracture toughness 

J-R curve for the material was then obtained from the correlation between fracture toughness 

parameters and RT Charpy-impact energy used to characterize the extent of thermal 

embrittlement. A common lower-bound J-R curve for various grades of CASS materials of 

unknown chemical composition was defined for a given material specification, ferrite content, 

and temperature.
23

Table 1. Chemical composition, ferrite content, and kinetics of thermal embrittlement for 

various heats of CASS materials.

Chemical Composition (wt.%) Ferrite (%) CVsat Constants Q

Heat Cr Mo Si Ni Mn C N Calc. Mea

s.

(J/cm2) β θ α (kJ/mole)

Argonne National Laboratory

52 19.49 0.35 0.92 9.40 0.57 0.009 0.052 10.3 13.5 161.8 – – – –

51 20.13 0.32 0.86 9.06 0.63 0.010 0.058 14.3 18.0 115.9 0.139 3.53 1.15 204.7

47 19.81 0.59 1.06 10.63 0.60 0.018 0.028 8.4 16.3 163.7 0.069 2.29 1.20 195.7

P2 20.20 0.16 0.94 9.38 0.74 0.019 0.040 12.5 15.6 141.3 0.258 2.83 1.09 218.6

I 20.20 0.45 0.83 8.70 0.47 0.019 0.032 20.4 17.1 134.3 0.094 2.10 1.00 250.0

69 20.18 0.34 1.13 8.59 0.63 0.023 0.028 21.0 23.6 76.7 0.214 3.21 1.07 175.9

P1 20.49 0.04 1.12 8.10 0.59 0.036 0.057 17.6 24.1 53.7 0.305 2.57 0.75 252.7

61 20.65 0.32 1.01 8.86 0.65 0.054 0.080 10.0 13.1 93.3 0.214 3.48 1.20 197.8

59 20.33 0.32 1.08 9.34 0.60 0.062 0.045 8.8 13.5 89.1 0.197 3.14 1.20 249.4

68 20.64 0.31 1.07 8.08 0.64 0.063 0.062 14.9 23.4 47.1 0.301 2.88 0.68 161.1

60 21.05 0.31 0.95 8.34 0.67 0.064 0.058 15.4 21.1 44.8 0.291 2.89 0.88 210.9

56 19.65 0.34 1.05 9.28 0.57 0.066 0.030 7.3 10.1 117.6 – – – –

74 19.11 2.51 0.73 9.03 0.54 0.064 0.048 15.5 18.4 63.1 0.269 3.44 0.70 95.0

75 20.86 2.58 0.67 9.12 0.53 0.065 0.052 24.8 27.8 32.1 0.436 2.82 0.51 139.0

66 19.45 2.39 0.49 9.28 0.60 0.047 0.029 19.6 19.8 87.9 0.208 3.16 1.57 163.9

64 20.76 2.46 0.63 9.40 0.60 0.038 0.038 29.0 28.4 41.1 0.338 2.81 0.60 147.3

65 20.78 2.57 0.48 9.63 0.50 0.049 0.064 20.9 23.4 59.7 0.260 2.99 0.59 153.8

P4 19.64 2.05 1.02 10.00 1.07 0.040 0.151 5.9 10.0 62.7 0.289 2.70 0.62 158.7

63 19.37 2.57 0.58 11.85 0.61 0.055 0.031 6.4 10.4 126.5 0.119 2.83 1.11 155.5

Georg Fischer Co.

284 23.00 0.17 0.52 8.23 0.28 0.025 0.037 43.6 42.0 20.5 0.551 3.66 0.39 85.9

280 21.60 0.25 1.37 8.00 0.50 0.028 0.038 36.3 38.0 19.6 0.609 3.20 0.73 88.9

282 22.50 0.15 0.35 8.53 0.43 0.035 0.040 29.7 38.0 28.5 0.500 3.65 0.39 91.6

281 23.10 0.17 0.45 8.60 0.41 0.036 0.053 31.4 30.0 17.2 0.618 3.76 0.47 89.8

283 22.60 0.23 0.53 7.88 0.48 0.036 0.032 42.6 42.0 18.6 0.599 3.60 0.44 83.7

278 20.20 0.13 1.00 8.27 0.28 0.038 0.030 18.5 15.0 68.3 0.347 3.90 0.29 63.1

279 22.00 0.22 1.36 7.85 0.37 0.040 0.032 39.5 40.0 23.8 0.546 3.06 0.58 93.5

277 20.50 0.06 1.81 8.13 0.54 0.052 0.019 22.5 28.0 30.7 0.466 3.54 0.49 87.7

291 19.60 0.66 1.59 10.60 0.28 0.065 0.054 4.2 6.0 121.9 0.195 3.65 0.35 71.2

292 21.60 0.13 1.57 7.52 0.34 0.090 0.039 23.9 28.0 17.2 0.373 3.07 0.44 98.8

290 20.00 2.40 1.51 8.30 0.41 0.054 0.050 31.3 32.0 15.8 0.624 3.48 0.12 81.0

288 19.60 2.53 1.70 8.40 0.47 0.052 0.022 35.6 28.0 14.9 0.671 2.96 0.66 105.3

287 20.50 2.58 0.51 8.46 0.50 0.047 0.033 37.2 38.0 20.5 0.555 3.46 0.36 90.3

286 20.20 2.44 1.33 9.13 0.40 0.072 0.062 18.9 22.0 15.5 0.594 3.03 0.72 106.4

289 19.70 2.30 1.44 8.25 0.48 0.091 0.032 22.6 30.0 16.2 0.580 3.29 0.41 90.1

285 18.80 2.35 0.86 9.49 0.48 0.047 0.039 14.0 10.0 61.1 0.313 3.60 0.20 89.3

Framatome

A 18.90 0.10 0.99 8.90 1.14 0.021 0.074 6.0 6.3 166.0 0.090 3.44 0.20 111.7

E 21.04 0.08 0.54 8.47 0.80 0.035 0.051 17.6 16.5 45.7 0.334 2.63 0.65 132.9

F 19.72 0.34 1.16 8.33 0.26 0.038 0.026 17.7 12.0 83.2 0.282 2.45 1.23 176.2

C 20.73 0.13 1.09 8.19 0.91 0.042 0.035 20.9 20.1 51.1 0.393 3.30 0.45 83.1

G 20.65 0.02 1.03 8.08 0.74 0.040 0.073 15.3 17.0 62.5 – – – –

H 20.70 0.05 1.18 8.07 0.71 0.050 0.045 18.3 21.5 50.6 – – – –

D 19.15 2.50 0.94 10.32 1.12 0.026 0.063 12.2 13.9 33.0 0.439 3.30 0.40 89.7

I 19.36 2.40 0.98 10.69 0.70 0.020 0.039 14.1 15.5 150.7 – – – –

K 20.80 2.62 0.75 10.45 1.09 0.060 0.056 15.4 14.0 48.5 – – – –

L 20.76 2.48 0.81 10.56 0.79 0.040 0.042 18.6 19.0 30.4 – 3.00 – –

B 20.12 2.52 0.93 10.56 0.83 0.053 0.042 14.0 17.3 28.2 0.478 2.55 0.47 128.6

Westinghouse

C148

8

20.95 2.63 0.53 9.48 1.02 0.061 0.056 22.1 14.0 53.1 – 2.80 – –

Electric Power Research Institute

EPRI

8

22.04 0.23 0.84 7.93 0.74 0.030 0.045 36.0 32.0 30.0 0.564 2.10 0.60 225.0

Correlations were also developed for estimating changes in tensile stress and Ramberg/Osgood 

parameters for tensile strain hardening.52 However, only fracture toughness J-R curve data 

have been reevaluated and updated in this report.
24

2.4.1 Charpy-Impact Energy

In the Argonne studies,48 different correlations were developed to estimate the saturation RT 

impact energy of the various grades of cast SS. To ensure that the estimates are either 

accurate or conservative for all heats, the saturation RT impact energy for a specific cast SS

was determined by two different expressions between RT Charpy-impact energy and a material 

parameter that depends on material ferrite content and chemical composition. The lower value 

was used for estimating mechanical properties of thermally aged CASS materials. For CF-3 

and CF-8 steels, the saturation value of RT impact energy, CVsat, is the lower value determined

from

log10CVsat = 1.15 + 1.36exp(–0.035Φ), (11)

where the material parameter Φ is expressed as

Φ = δc(Cr + Si)(C + 0.4N), (12)

and from

log10CVsat = 5.64 – 0.006δc – 0.185Cr + 0.273Mo – 0.204Si 

+ 0.044Ni – 2.12(C + 0.4N). (13)

For CF-8M steel with <10% Ni, the saturation value of RT impact energy CVsat is the lower value 

determined from

log10CVsat = 1.10 + 2.12exp(–0.041Φ), (14)

where the material parameter Φ is expressed as

Φ = δc(Ni + Si + Mn)2(C + 0.4N)/5; (15)

and from 

log10CVsat = 7.28 – 0.011 δc – 0.185Cr – 0.369Mo – 0.451Si 

– 0.007Ni – 4.71(C + 0.4N). (16)

For CF-8M steel with ≥10% Ni, the saturation value of RT impact energy CVsat is the lower value 

determined from

log10CVsat = 1.10 + 2.64exp(–0.064Φ), (17)

where the material parameter Φ is expressed as

Φ = δc(Ni + Si + Mn)2(C + 0.4N)/5; (18)

and from

log10CVsat = 7.28 – 0.011 δc – 0.185Cr – 0.369Mo – 0.451Si 

– 0.007Ni – 4.71(C + 0.4N). (19)
25

If not known, the N content was assumed to be 0.04 wt.%. The correlations were optimized by 

using a larger database (≈80 compositions of CASS materials) and mechanical property results 

on materials that were aged up to ≈58,000 h at 290–350°C (554–662°F). The saturation values 

of RT impact energy for CF-3, CF-8, and CF-8M steels observed experimentally at 

Argonne,

48,77,78 GF,30 Westinghouse (WH),31 EdF,43 NP,44 FRA,34 and EPRI35 are plotted as a 

function of the material parameter φ in Fig. 11. The solid lines represent best-fit curves for the 

data (i.e., Eq. 11 for CF-3 or CF-8 steels and Eqs. 14 and 17 for CF-8M steel). The chemical 

composition, ferrite content, and saturation RT Charpy-impact energy of most of the materials 

are given in Table 1. The measured values of CVsat and those estimated from Eqs. 11–19 are 

plotted in Fig. 12. The difference between the predicted and observed values is less than ±15% 

for most of the CF-3 and CF-8 steels and less than ±25% for the CF-8M steels. 

The expressions given in Eqs. 11–19 for estimating the saturation RT Charpy-impact energy for 

a specific CASS material, were correlated to the material ferrite and the chemical composition 

because a review of the thermal embrittlement data for aged CASS materials indicated that a 

correlation between Charpy-impact energy and ferrite content alone did not yield good results. 

The saturation RT impact energies for CF-3, CF-8, and CF-8M steels shown in Fig. 11 are

plotted as a function of the material ferrite content in Fig. 13; the results show poor correlation. 

2.4.2 Fracture Toughness J-R Curve 

The saturation fracture toughness J-R curve for a specific CASS material can be estimated from 

its RT Charpy-impact energy at saturation. The J-R curve is expressed by the power-law 

relation Jd = CΔan, where Jd is deformation J per ASTM Specifications E 813-85 and E 1152-87, 

Δa is crack extension, and C and n are constants. The coefficient C at room temperature or are 

plotted in Fig. 14 using an updated fracture toughness database. Fracture toughness data from 

studies at Argonne,

48,77–79 FRA,34,36 EdF,38 EPRI,35 The Welding Institute (TWI),80 and Materials 

Engineering Associates, Inc. (MEA),

81 are included in the figure. 

At both RT and reactor temperatures, the coefficient C decreases with a decrease in the RT 

Charpy-impact energy. Separate correlations are obtained for CF-3 and CF-8 steels and for 

CF-8M steels; the latter show larger decrease in fracture toughness for a given impact energy. 

Furthermore, for CF-8M steels, the decrease in the values of coefficient C is much greater for 

RT Charpy-impact energy values less than about 40 J/cm2. As discussed later in this section, 

data on chemical compositions of CASS piping materials from a select sample of nuclear power 

plants indicate that at least 9% of the CF-8M materials that are currently in use in the U.S.

nuclear power plants contain more than 25% ferrite. However, the methodology developed 

earlier in NUREG/CR-4513, Rev. 1 was not applicable for CASS materials with greater than 

25% ferrite. The methodology for estimating thermal embrittlement has been extended to cover 

CASS materials containing greater than 25% ferrite. For CF-8M materials, a bilinear expression 

is developed between RT Charpy-impact energy and coefficient C of the power-law J-R curve. 

To help ensure that the estimated J-R curve was conservative for all material and aging 

conditions, the correlations for estimating the J-R curves for static-cast materials are obtained 

by subtracting the value of standard deviation for the fit to the data from the best-fit curve in 

Fig. 14; these curves are shown as chain dot curves. For centrifugally-cast materials, the bestfit correlations are used. For CF-8M materials, the value of the RT Charpy-impact energy for 
26

Figure 11. Correlation between RT Charpy-impact energy at saturation and the material 

parameter φ for CF-3, CF-8, and CF-8M steels. 

10

100

0 10 20 30 40 50

ANL

EdF

GF

FRA

CEGB

EPRI

RT Charpy-Impact Energy, CVsat (J/cm2)

Φ = δc(Cr+Si)(C+0.4N)

CF-3 & CF-8

log10(CVsat) = 1.05 + 1.36exp(- 0.035φ)

log10(CVsat) =1.15 + 1.36exp(- 0.035φ)

10

100

0 10 20 30 40 50

ANL

EdF

GF

WH

RT Charpy-Impact Energy, CVsat (J/cm2)

Φ = δc(Ni+Si+Mn)2(C+0.4N)/5

CF-8M

Ni < 10 wt.%

log10(CVsat) = 0.99 + 2.12exp(- 0.041φ)

log10(CVsat) =1.10 + 2.12exp(- 0.041φ)

10

100

0 10 20 30 40 50

ANL

EdF

FRA

RT Charpy-Impact Energy, CVsat (J/cm2)

Φ = δc(Ni+Si+Mn)2(C+0.4N)/5

CF-8M

Ni ≥ 10 wt.%

log10(CVsat) = 0.96 + 2.64exp(- 0.064φ)

log10(CVsat) =1.10 + 2.64exp(- 0.064φ)
27

Figure 12. Measured and estimated values of saturation RT Charpy-impact energy for CF-3, 

CF-8, and CF-8M CASS materials. 

0

50

100

150

200

0 50 100 150 200

CF-3

CF-8

Estimated CVsat (J/cm2)

Measured CVsat (J/cm2)

Thermally Aged CASS Materials

0

50

100

150

200

0 50 100 150 200

CF-8M Ni <10%

CF-8M Ni ≥10%

Estimated CVsat (J/cm2)

Measured CVsat (J/cm2)

Thermally Aged CASS Materials
28

Figure 13. Correlation between RT Charpy-impact energy at saturation and the ferrite content 

for CF-3, CF-8, and CF-8M steels. 

10

100

0 10 20 30 40 50

ANL

EdF

GF

FRA

CEGB

EPRI

Impact Energy, CVsat (J/cm2)

Ferrite Content (%)

CF-3 & CF-8

10

100

10 15 20 25 30 35 40 45 50

ANL

GF

EdF

WH

Impact Energy, CVsat (J/cm2)

Ferrite Content (%)

CF-8M

Ni < 10 wt.%

10

100

0 10 20 30 40 50

ANL

EdF

FRA

Impact Energy, CVsat (J/cm2)

Ferrite Content (%)

CF-8M

Ni ≥ 10 wt.%
29

Figure 14. Correlation between RT Charpy-impact energy and coefficient C at 290–320°C and 

RT for CF-3, CF-8, and CF-8M CASS materials.

the transition from one expression to the other varies between 35 and 45 J/cm2 because of the 

differences in the standard deviation for the fit to the individual set of data. For static-cast CASS 

materials, the coefficient C of the J-R curve at RT for CF-3 or CF-8 steels is expressed as

C = 49[CV]

0.52, (20)

and for CF-8M steels, for RT Charpy-impact energy values greater than 35 J/cm2 as 

10

100

1000

10 100

ANL

EPRI

KRB

ANL

FRA/EdF

Coefficient C

RT Charpy-Impact Energy, CV (J/m2)

290 - 320°C

Open Symbols: CF-3, CF-8

Closed Symbols: CF-8M

C = 134CV

0.28

C = 102CV

0.28

C = 57CV

0.41

C = 49CV

0.41

C = 5.5CV

0.98

C = 6.9CV

0.98

10

100

1000

10 100

ANL

EPRI

KRB

MEA

ANL

FRA/EdF

Coefficient C

RT Charpy-Impact Energy, CV (J/m2)

Room Temperature

Open Symbols: CF-3, CF-8

Closed Symbols: CF-8M

C = 57CV

0.52

C = 49CV

0.52

C = 20CV

0.67

C = 16CV

0.67

C = 1.44CV

1.35

C = 1.78CV

1.35
30

C = 16[CV]

0.67, (21)

and for RT Charpy-impact energy values equal to or less than 35 J/cm2 as

C = 1.44[CV]

1.35. (22)

For static-cast CASS materials, the coefficient C of the J-R curve at 290–320°C for CF-3 or 

CF-8 steels is expressed as

C = 102[CV]

0.28, (23)

and for CF-8M steels, for RT Charpy-impact energy values greater than 45 J/cm2 as

C = 49[CV]

0.41, (24)

and for RT Charpy-impact energy values equal to or less than 45 J/cm2 as

C = 5.5[CV]

0.98. (25)

For centrifugally-cast CASS materials, the coefficient C of the J-R curve at RT for CF-3 or CF-8 

steels is expressed as

C = 57[CV]

0.52, (26)

and for CF-8M steels, for RT Charpy-impact energy values greater than 35 J/cm2 as

C = 20[CV]

0.67, (27)

and for RT Charpy-impact energy values equal to or less than 35 J/cm2 as

C = 1.78[CV]

1.35. (28)

For centrifugally-cast CASS materials, the coefficient C of the J-R curve at 290–320°C for CF-3 

or CF-8 steels is expressed as

C = 134[CV]

0.28, (29)

and for CF-8M steels, for RT Charpy-impact energy values greater than 40 J/cm2 as

C = 57[CV]

0.41, (30)

and for RT Charpy-impact energy values equal to or less than 40 J/cm2 as

C = 6.9[CV]

0.98. (31)

The exponent n of Δa has also been correlated with the RT Charpy-impact energy CV. To 

ensure that the estimated J-R curves were conservative, the correlations represented the lower 

bound values of the exponent n versus Charpy-impact energy experimental data. A review of 

the updated fracture toughness data for thermally aged CF-3, CF-8, CF-8M steels indicated that 
31

Figure 15. Correlation between RT Charpy-impact energy and exponent n of the power-law 

J-R curve at RT and 290°C for CF-8M materials. The solid line bounds the 

existing data for exponent n as a function of RT Charpy-impact energy. 

the existing correlation between RT Charpy-impact energy and the exponent n in the power-law 

J-R curve50 were adequate and did not need to be updated. For example, the existing

correlation for CF-8M bound the updated data for RT Charpy-impact energy and exponent n 

shown in Fig. 15. 

For static-cast or centrifugally-cast CASS materials, the exponent n at RT is given by the value 

larger of 0.34 and that determined for CF-3 steels from the expression 

n = 0.15 + 0.16log10[CV], (32)

0.2

0.3

0.4

0.5

0.6

0.7

0.8

101 102

Exponent n

Charpy-Impact Energy CV (J/cm1)

CF-8M Steel

Room Temp.

n = 0.23 +0.08log(CV)

0.2

0.3

0.4

0.5

0.6

0.7

0.8

101 102

Exponent n

Charpy-Impact Energy CV (J/cm1)

CF-8M Steel

290°C

n = 0.23 +0.06log(CV)
32

for CF-8 steels from 

n = 0.20 + 0.12log10[CV], (33)

and for CF-8M steels from 

n = 0.23 + 0.08log10[CV]. (34)

For static-cast or centrifugally-cast CASS materials, the exponent n at 290°C is given by the 

value larger of 0.34 and that determined for CF-3 steels from the expression 

n = 0.17 + 0.12log10[CV], (35)

for CF-8 steels from

n = 0.21 + 0.09 log10[CV], (36)

and for CF-8M steels from

n = 0.23 + 0.06log10[CV]. (37)

Equations 20–37 may be used to determine the fracture toughness J-R curve of static- or 

centrifugally-cast CF-3, CF-8, and CF-8M CASS materials from the RT Charpy-impact energy of 

the material. If the RT Charpy-impact energy is not known, then the saturation fracture 

toughness J-R curves for these CASS materials can be determined from their chemical 

composition available in the CMTRs using Eqs. 11–19 and Eqs. 20–37. In general, the 

following observations may be drawn from the study of thermal embrittlement of CASS 

materials:82

(a) Among the grades considered, CF-8M steels have the smallest Jd value for a given 

crack extension. The value of Jd at a given crack extension for CF-8M steels is 

generally about half of that of the other grades. 

(b) CF-3, CF-3A, CF-8, and CF-8A steels show a similar extent of thermal aging. The value 

of Jd at a given crack extension for CF-8 and CF-8A steels is generally less than 10% 

lower than that for CF-3 and CF-3A steels.

(c) Static-cast CASS materials are more susceptible to thermal aging than the centrifugallycast materials. The value of Jd at a given crack extension for static-cast SS is generally 

about 20% lower than that for the centrifugally-cast material with a similar ferrite content. 

(d) Although CF-8M steels show a similar extent of thermal aging when the J-R curve is 

measured at both RT and at reactor operating temperature, the other grades considered 

show a decreased value of Jd at a given crack extension when the J-R curve is 

measured at reactor operating temperature. The value of Jd at a given crack extension 

for CF-3, CF-3A, CF-8, and CF-8A steels is generally about 20% lower when measured 

at reactor operating temperature than when measured at RT.

These correlations are valid for static- and centrifugally-cast CF-3, CF-3A, CF-8, CF-8A, and 

CF-8M steels defined by ASTM Specification A351. The criteria used in developing these 

correlations ensure that the estimated mechanical properties are adequately conservative for 

compositions of CASS materials within ASTM A351. The updated correlations are applicable to 

all compositions of CF-8M materials, including materials with ferrite contents above 25%. 

However, in the updated database, there were few or no fracture toughness J-R curve data for 
33

CF-3 and CF-8 materials with CVsat less than 30 J/cm2 (17.7 ft·lb). Therefore, the correlations 

presented in this report may not be applicable for those compositions of CF-3 and CF-8 

materials for which the estimated value of CVsat is less than 30 J/cm2 (<17.7 ft·lb). Typically,

such compositions would contain more than 30% ferrite. Furthermore, the correlations may not 

encompass all metallurgical factors that can arise from differences in production heat treatment 

or casting processes and may be overly conservative for some steels.

It should be noted that these correlations account for the degradation of mechanical properties 

of typical heats of CASS materials. They do not consider the initial fracture properties of the 

unaged material. Some CASS materials may have low initial fracture toughness and the 

estimated J-R curves may be higher than the initial value. Therefore, some knowledge 

regarding the initial fracture toughness of the material is needed to justify the use of the 

estimated fracture toughness. The initial fracture toughness J-R curves may be estimated from 

the RT Charpy-impact energy of the unaged material. 

Flaw tolerance methods are often used to develop ASME Code Section XI flaw acceptance 

standards or to justify alternatives to the ASME Code Section XI in-service inspection (ISI) 

requirements. Recently, an EPRI Report 1019128 (December 2009),83 presented a flaw 

tolerance approach based on elastic-plastic fracture mechanics considerations that could be 

used in combination with a demonstrated inspection method for managing the effects of aging of 

CASS piping, particularly with greater than 20% ferrite content. The report included a database 

of CASS CF-3, CF-8, CF-8A, and CF-8M piping materials. The sources of the CASS data 

Table 2. The primary circuit piping CASS material in Westinghouse designed plants.

Plant Name Size (Mwe) System 

Loops

Material 

Type

Beaver Valley 2 852 3 CF-8A

Callaway 1 1157 4 CF-8A

Catawba 1 1153 4 CF-8A

Catawba 2 1153 4 CF-8A

Comanche Peak 1 1150 4 CF-8A

Farley 1 829 3 CF-8A

Farley 2 829 3 CF-8A

McGuire 1 1180 4 CF-8A

McGuire 2 1180 4 CF-8A

Millstone 3 1150 4 CF-8A

North Anna 1 934 3 CF-8A

North Anna 2 788 3 CF-8A

South Texas 1 1250 4 CF-8A

South Texas 2 1250 4 CF-8A

Vogtle 1 1113 4 CF-8A

Vogtle 2 1113 4 CF-8A

Watts Bar 1 1177 4 CF-8A

Watts Bar 2 1177 4 CF-8A

Wolf Creek 1158 4 CF-8A

Beaver Valley 1 852 3 CF-8M

Cook 1 1090 4 CF-8M

Cook 2 1054 4 CF-8M

Kewaunee 560 2 CF-8M

Prairie Island 2 530 2 CF-8M

Sequoyah 1 1140 4 CF-8M

Sequoyah 2 1140 4 CF-8M
34

included: (a) information obtained from a Westinghouse data search based on random sampling 

of heats of CASS material from 15 plants (see Table 2) and (b) data packages from the 

Sandusky Foundry and Machine Co., NUREG/CR-5024, and Structural Integrity Associates. 

The CASS data was used to estimate the ferrite contents of a representative sample of CASS 

CF-3, CF-8, CF-8A, and CF-8M materials, and perform statistical analyses to study the 

distribution of the ferrite contents of these materials. When information for the Mo and N 

contents was not available, values of 0.5 and 0.04 (wt.%), respectively, were assumed. 

The results yielded mean ferrite content and standard deviation values of 15.9% and 4.9% for 

CF-8 material (total 333 data points), and 17.6% and 5.4% for CF-8M material (total 147 data 

points). The total data for CF-3 and CF-8 were too small (less than 15 data points) to provide 

meaningful estimates. The distribution of the ferrite content for CF-8M materials (Fig. 16) 

indicated that nearly 9% of the heats of CF-8M material contained greater than 25% ferrite. 

Therefore, the Argonne correlations presented earlier in NUREG/CR-4513, Rev. 1 would not be 

applicable to at least some of the heats of CF-8M material currently in use in the U.S. PWRs. 

The updated expressions given in Eqs. 20–31 should be used for these materials. However, 

since the updated database used in this report contained little or no fracture toughness J-R 

curve data for thermally aged CF-3 and CF-8 materials with more than 25% ferrite, the 

expressions presented in Eqs. 20–31 may not be applicable to such materials. The use of 

these expressions for CF-3 and CF-8 materials with more than 25% ferrite needs to be justified. 

Figure 16.

Distribution of ferrite content in CASS 

CF-8M piping material in Westinghouse 

PWRs (Ref. EPRI 1019128).

2.5 Assessment of Thermal Embrittlement of CASS Materials

2.5.1 Estimation of Thermal Embrittlement of CASS Materials of Known Composition 

and Service Condition – Service Time Values

Room-temperature impact energy as a function of time and temperature of aging is estimated 

from the RT saturation impact energy CVsat and the kinetics of embrittlement. The decrease in 

0

5

10

15

20

25

30

1 3 5 7 9 11 13 15 17 19 21 23 25 27 29 31

Count

Ferrite Content (%)

CF-8M Piping Material
35

RT Charpy-impact energy CV with time is expressed as 

log10CV = log10CVsat + β{1 – tanh [(P – θ)/α]}, (38)

where the aging parameter P is determined from Eq. 9 The constants α and β can be 

determined from CVint and CVsat as follows:

α = –0.585 + 0.795log10CVsat (39)

and

β = (log10 CVint – log10CVsat)/2. (40)

The CMTR for a specific CASS component provides information on chemical composition, 

tensile strength, and possibly Charpy-impact energy of the material. If CVint is not known, a 

typical value of 200 J/cm2 [or 160 J (118 ft·lb) for a standard Charpy V-notch specimen] may be 

used. The value of θ is not available for CASS components in the field, and can only be 

obtained from aging archival material for 5,000–10,000 h at 400°C (752°F). However, 

parametric studies show that the aging response at reactor temperatures is relatively insensitive 

to the values of θ.

84 A value of 2.9 for θ (i.e., mean of the experimental data) is used to estimate 

thermal embrittlement at 280–330°C (536–626°F). A θ value of 3.3 is used for estimates at 

temperatures <280°C (<536°F) and 2.5 for estimates at 330–360°C (626–680°F). In the 

Argonne study, the activation energy for thermal embrittlement is expressed in terms of both 

chemical composition and the constant θ.

50 The activation energy Q in kJ/mole for CF-3 and

CF-8 steels is given by

Q = 10 [74.52 – 7.20 θ – 3.46 Si – 1.78 Cr + 148 N – 61 C], (41)

and for CF–8M steels by 

Q = 10 [74.52 – 7.20 θ – 3.46 Si – 1.78 Cr – 4.35 Mn+ 23 N]. (42)

Equations 41 and 42 are slightly different from the expression proposed by Slama et al. 1983 

(i.e. Eq. 10 presented earlier).34 These equations are applicable to compositions within ASTM 

Specification A351, with an upper limit of 1.2 wt.% for Mn content. Actual Mn content is used 

when CASS materials that contain up to 1.2 wt.% Mn; for steels containing greater than 

1.2 wt.% Mn, 1.2 wt.% is assumed. Furthermore, the values of Q predicted from Eqs. 41 and 42

should be between 65 kJ/mole (15.5 kcal/mole) minimum and 250 kJ/mole (59.8 kcal/mole) 

maximum; Q is assumed to be 65 kJ/mole if the predicted values are lower, and 250 kJ/mole if 

the predicted values are higher. The above expressions for estimating activation energy Q for 

thermal aging embrittlement of CASS materials agree qualitatively with the microstructural and 

mechanical-property data. For example, an increase in the value of θ decreases the activation 

energy, as expected. The contribution of Si for all grades of CASS steels and of Mn for CF-8M 

steels are consistent with their effect on the formation of G-phase. These elements should 

promote precipitation of G phase: hence, the coefficients for these elements should have a 

negative sign, because activation energy for thermal embrittlement is low for steels that show 

G-phase precipitation. An increase in C or N in the steel will promote carbide or nitride 

precipitation at high temperatures and thus increase the activation energy. The positive sign of 

the constant for the N content agrees with this behavior. The constant for the C content in steel, 

however, has a negative sign. It is likely that C also promotes precipitation of G phase, a 

multicomponent phase consisting of Ni, Si, Mo, Cr, Fe, and some Mn and C.61,66
36

The RT Charpy-impact energy of a specific CASS material as a function of service time and 

temperature can be obtained from estimated CVsat (Eqs. 3–5 and 11–19) and the kinetics of 

embrittlement (Eqs. 38–42). The initial Charpy-impact energy of the unaged steel is needed for 

estimating the decrease in impact energy as a function of time and temperature of service. The 

RT Charpy-impact energy observed experimentally and that estimated from the chemical 

composition and initial impact energy of the CASS materials used to develop these correlations 

are presented in Figs. 19 and 20 of NUREG/CR-4513, Rev. 1. A θ value of 2.9 was used for all 

temperatures of thermal aging. The estimated change in impact energy at temperatures less 

than or equal to 330°C (626°F) is either accurate or slightly conservative for most of the heats. 

A few heats show poor agreement because either the estimated CVsat is higher than the 

experimental value (Framatome Heat D and Argonne Heat 47) or the estimated activation 

energy is high (Framatome Heat C and Georg Fischer Heat 278). Even at 350°C, the estimated 

impact energies show good agreement with the experimental results because the θ values for 

most of the heats shown in the figures are either greater or only slightly lower than 2.9. The 

EPRI heat and an EdF heat (experimental θ is 2.1 for both heats) alone show nonconservative 

estimates at 350°C. A θ value of 2.5 rather than 2.9 should be used to ensure that the 

estimates at 330–360°C (626–680°F) are conservative.

Once the RT Charpy-impact energy CV is known, the service-time coefficient C and exponent n 

of the fracture toughness J-R curve are determined from Eqs. 20–31 and 32–37, respectively. 

The methodology and expressions for estimating the change in tensile flow and yield stresses 

and the engineering stress-strain curve of CASS materials as a function of time and 

temperature of service, have been described earlier in NUREG/CR-6142.52 The tensile 

properties of aged CASS materials are determined from known material information, (i.e., 

chemical composition and initial tensile strength of the steel). The fracture toughness JIc values 

for the service-aged CASS material can be determined from the estimated values of the fracture 

toughness J-R curve and flow stress. In Figs. 17–20, the change in the measured values of 

coefficient C as a function of aging time and temperature, are compared with those estimated 

from material composition and the initial RT Charpy-impact energy, for some of the CASS 

materials. 

The results indicate that in general, the estimates at aging temperatures of 320 and 350°C, are 

either accurate or conservative for all of the materials. The proposed methodology can be used 

to estimate the change in fracture toughness of CASS materials for known material composition 

and service conditions, particularly for service temperature between 290 and 320°C. 

2.5.2 Estimation of Thermal Embrittlement of CASS Materials of Unknown Composition 

– Lower-bound values

The Argonne methodology discussed above provides the expressions for estimating the fracture 

toughness J-R curves of unaged and aged CASS materials as a function of the RT Charpyimpact energy, which in turn is determined from the a material parameter that depends on the 

chemical composition and ferrite content of the material. However, for convenience, lowerbound J-R curves are defined as a function of ferrite content of the CASS material. The fracture 

toughness is defined by the J-R curve expressed by the power-law relation Jd = CΔan, where Jd

is deformation J per ASTM Specifications E 813-85 and E 1152-87, Δa is crack extension, and 

C and n are constants. The coefficient C and exponent n for the lower-bound fracture 

toughness J-R curve are determined from the bounding value of CVsat for a range of ferrite 

content in the CASS material. 
37

(a)

(b)

Figure 17. Comparison of the experimental values of coefficient C of the power-law J-R curve 

at (a) RT and (b) 290 or 320°C as a function of aging time and temperature for

static-cast CF-3 plate with the values estimated from material composition and the 

initial RT Charpy-impact energy. 

200

400

600

800

1000

1200

101 102 103 104 105 106

400°C

350°C

320°C

Coefficient C

Aging Time (h)

Static Cast CF-3 Plate

Heat 69, δc=21.0% δm=23.6%

Aging Temp.

All specimens precraked & tested in air at RT

400°C

350°C

320°C

200

300

400

500

600

700

800

101 102 103 104 105 106

400°C

350°C

320°C

400°C

Coefficient C

Aging Time (h)

Static Cast CF-3 Plate

Heat 69, δc=21.0% δm=23.6%

Aging Temp.

Irradiated to 0.08 dpa

All specimens precraked & tested in

Open Symbols: Air at 290°C

Closed Symbols: PWR water at 320°C

400°C

350°C

320°C
38

(a)

(b)

Figure 18. Comparison of the experimental values of coefficient C of the power-law J-R curve 

at (a) RT or 54°C and (b) 290 or 315–320°C as a function of aging time and 

temperature for static-cast CF-8 plate with the values estimated from material 

composition and the initial RT Charpy-impact energy. 

0

200

400

600

800

1000

1200

101 102 103 104 105 106

400°C

350°C

320°C

350°C

Coefficient C

Aging Time (h)

Static Cast CF-8 Plate

Heat 68, δc=14.9% δm=23.4%

Aging Temp.

Tested at RT

Tested at 54°C

All specimens precraked in air

Fracture toughness J-R test in

Open Symbols: Air

Closed Symbols: PWR water

400°C

350°C 320°C

Precracked in

315°C PWR water

0

100

200

300

400

500

600

700

800

101 102 103 104 105 106

400°C

350°C

320°C

350°C

400°C

Coefficient C

Aging Time (h)

Static Cast CF-8 Plate

Heat 68, δc=14.9% δm=23.4%

Aging Temp.

All specimens precraked & testedt in

Open Symbols: Air at 290°C

Closed Symbols: PWR water at 315-320°C

400°C

350°C

320°C

Irradiated to 0.08 dpa
39

(a)

(b)

Figure 19. Comparison of the experimental values of coefficient C of the power-law J-R curve 

at (a) RT and (b) 290 or 320°C as a function of aging time and temperature for 

static-cast CF-8M plate with the values estimated from material composition and 

the initial RT Charpy-impact energy. 

0

200

400

600

800

1000

1200

101 102 103 104 105 106

400°C

350°C

320°C

Coefficient C

Aging Time (h)

Static Cast CF-8M Plate

Heat 75, δc=24.8% δm=27.8%

Aging Temp.

All specimens precraked & tested in air at RT

400°C

350°C

320°C

0

100

200

300

400

500

600

700

800

101 102 103 104 105 106

400°C

350°C

320°C

400°C

Coefficient C

Aging Time (h)

Static Cast CF-8M Plate

Heat 75, δc=24.8% δm=27.8%

Aging Temp.

All specimens precraked & testedt in

Open Symbols: Air at 290°C

Closed Symbols: PWR water at 320°C

400°C

350°C

320°C

Irradiated to 0.08 dpa
40

(a)

(b)

Figure 20. Comparison of the experimental values of coefficient C of the power-law J-R curve 

at (a) RT and (b) 290°C as a function of aging time and temperature for 

centrifugally-cast CF-8 pipe with the values estimated from material composition 

and the initial RT Charpy-impact energy. 

Therefore, first, the bounding values of saturation RT Charpy-impact energy, CVsat, for ferrite 

contents of <10%, 10–15%, 15–25%, and 25–40% are established for each grade of CASS 

material, from the plots of estimated CVsat and ferrite content (Fig. 21). As discussed earlier, for 

the different grades of CASS materials, CVsat is the lower value determined from the sets of 

expressions given in Eqs. 11–19. Separate expressions are proposed for CF-8M steels 

0

200

400

600

800

1000

1200

101 102 103 104 105 106

400°C

350°C

320°C

Coefficient C

Aging Time (h)

Centrifugally Cast CF-8 Pipe

Heat P1, δc=17.6% δm=24.1%

Aging Temp.

All specimens precraked & tested in air at RT

400°C

350°C

320°C

0

100

200

300

400

500

600

700

800

101 102 103 104 105 106

400°C

350°C

320°C

Coefficient C

Aging Time (h)

Centrifugally Cast CF-8 Pipe

Heat P1, δc=17.6% δm=24.1%

Aging Temp.

All specimens precraked & tested in air at 290°C

400°C

350°C

320°C
41

containing less than 10% Ni and those containing equal to or greater than 10% Ni. The latter

are most susceptible to thermal embrittlement than the other grades. Note that in Fig. 21, the 

CVsat versus Fc curves bound approximately 95% of the data, and represent a conservative 

estimate of CVsat. 

A value of 10% Ni is used in this study to differentiate whether primary γ (austenite) form during 

solidification of the castings from the liquid. For CASS CF-8M compositions containing less 

than 10% Ni, the solidification sequence is as follows: Liq. → Liq. + δ → δ; with δ → γ in the 

solid.85 Austenite nucleates predominantly at the δ ferrite grain boundaries and to a lesser 

extent at interdendritic locations within the ferrite. For CF-8M compositions containing equal to 

or more than 10% Ni, the solidification sequence is likely to be as follows: Liq. → Liq. + δ → Liq. 

+ δ + γ → δ + γ; with δ → γ continuing in the solid.85 The austenite forms first in the liquid as a 

secondary phase enveloping the primary δ ferrite. On further cooling, it grows in the remaining 

liquid as well as into ferrite. The latter reaction continues below the solidus line. The actual 

microstructures of the casting depend on the Cr/Ni ratio for the specific composition. 

The lower-bound J-R curves at RT and 290–320°C determined from the estimated CVsat for 

static-cast and centrifugally-cast CASS materials with 25–40%, 15–25%, 10–15%, and <10% 

ferrite, are plotted in Figs. 22a, 22b, 23a and 23b. These plots are updated versions of Figs. 3 

and 4 of NUREG/CR-4513, Rev. 1.50 These lower-bound fracture toughness J-R curves can be 

used for completely embrittled CASS materials of unknown composition. The estimated CVsat

and the corresponding values of coefficient C and exponent n of the J-R curve at RT and 290–

320°C (550–608°F) for CF-3, CF-8, and CF-8M with <10 and ≥10% Ni are listed in Table 3. 

Table 3. Coefficient C and exponent n for the lower-bound J-R curve at RT and at reactor 

operating temperatures for thermally aged CASS materials. 

Estimated Static Cast Material Centrifugally-cast Material

Material CVsat Room Temp. 290–320°C Room Temp. 290–320°C 

Grade (J/cm2) C n C n C n C n

Ferrite content 25–40%

CF-3/CF-3A 27 270 0.38 256 0.34 314 0.38 336 0.34

CF-8/CF-8A 22 242 0.36 241 0.34 282 0.36 317 0.34

CF-8M (Ni <10%) 16 58 0.34 81 0.34 72 0.34 101 0.34

CF-8M (Ni ≥10%) 9 27 0.34 46 0.34 33 0.34 57 0.34

Ferrite content 15–25%

CF-3/CF-3A 37 321 0.40 281 0.36 373 0.40 369 0.36

CF-8/CF-8A 28 279 0.37 260 0.34 325 0.37 342 0.34

CF-8M (Ni <10%) 23 100 0.34 119 0.34 123 0.34 150 0.34

CF-8M (Ni ≥10%) 13 47 0.34 70 0.34 59 0.34 87 0.34

Ferrite content 10–15%

CF-3/CF-3A 61 416 0.44 323 0.38 484 0.44 424 0.38

CF-8/CF-8A 43 346 0.40 292 0.36 403 0.40 384 0.36

CF-8M (Ni <10%) 42 196 0.36 215 0.34 245 0.36 264 0.34

CF-8M (Ni ≥10%) 26 115 0.34 132 0.34 142 0.34 166 0.34

Ferrite content <10%

CF-3/CF-3A 91 512 0.46 361 0.41 595 0.46 474 0.41

CF-8/CF-8A 60 411 0.41 320 0.37 478 0.41 421 0.37

CF-8M (Ni <10%) 68 270 0.38 276 0.34 338 0.38 322 0.34

CF-8M (Ni ≥10%) 43 199 0.36 220 0.34 249 0.36 267 0.34
42

Figure 21. Correlation between saturation RT Charpy-impact energy and ferrite content for 

CF-3, CF-8, and CF-8M CASS materials.

10

100

0 5 10 15 20 25 30 35 40

ANL

EPRI

GF

NP

TWI

FRA

Estimated RT Charpy-Impact Energy, CVsat (J/cm2)

Ferrite Content, Fc (%)

CF-3 Steel

CV = 870Fc-0.98

10

100

0 5 10 15 20 25 30 35 40

ANL

MHI

GF

EdF

Shippingport

FRA

Estimated RT Charpy-Impact Energy, CVsat (J/cm2)

Ferrite Content, Fc (%)

CF-8 Steel

CV = 385Fc-0.81
43

Figure 21 Continued.

10

100

0 5 10 15 20 25 30 35 40

ANL

MHI

GF

Westinghouse

EdF

FRA

Estimated RT Charpy-Impact Energy, CVsat (J/cm2)

Ferrite Content, Fc (%)

CF-8M Steel (Ni <10%)

CV = 1030Fc-1.18

10

100

0 5 10 15 20 25 30 35 40

ANL

MHI

Ringhals

Westinghouse

EdF

FRA

Estimated RT Charpy-Impact Energy, CVsat (J/cm2)

Ferrite Content, Fc (%)

CF-8M Steel (Ni ≥10%)

CV = 820Fc-1.28
44

Figure 22a. Estimated lower-bound J-R curves at RT for static-cast CASS materials 

with ferrite contents of 25–40%, 15–25%, 10–15%, and <10%. 

0

200

400

600

800

1000

0 1 2 3 4 5

J (kJ/m2)

Crack Extension (mm)

Lower-Bound J-R Curve

Static-Cast Steels 25-40% Ferrite

Room Temp.

Jd=27Δa0.34

CF-8M ≥10% Ni

Jd=58Δa0.34

CF-8M <10% Ni

Jd=242Δa0.36

CF-8 Jd=270Δa0.38

CF-3 

0

200

400

600

800

1000

0 1 2 3 4 5

J (kJ/m2)

Crack Extension (mm)

Lower-Bound J-R Curve

Static-Cast Steels 15-25% Ferrite

Room Temp.

Jd=47Δa0.34

CF-8M ≥10% Ni

Jd=100Δa0.34

CF-8M <10% Ni

Jd=279Δa0.37

CF-8 

Jd=321Δa0.40

CF-3 
45

Figure 22a Continued. 

0

200

400

600

800

1000

0 1 2 3 4 5

J (kJ/m2)

Crack Extension (mm)

Lower-Bound J-R Curve

Static-Cast Steels 10-15% Ferrite

Room Temp.

Jd=115Δa0.34

CF-8M ≥10% Ni

Jd=196Δa0.36

CF-8M <10% Ni

Jd=346Δa0.40

CF-8 

Jd=416Δa0.44

CF-3 

0

200

400

600

800

1000

0 1 2 3 4 5

J (kJ/m2)

Crack Extension (mm)

Lower-Bound J-R Curve

Static-Cast Steels <10% Ferrite

Room Temp.

Jd=199Δa0.36

CF-8M ≥10% Ni

Jd=270Δa0.38

CF-8M <10% Ni

Jd=411Δa0.41

CF-8 

Jd=512Δa0.46

CF-3 
46

Figure 22b Estimated lower-bound J-R curves at 290–320°C for static-cast CASS materials 

with ferrite contents of 25–40%, 15–25%, 10–15%, <10%. 

0

200

400

600

800

1000

0 1 2 3 4 5

J (kJ/m2)

Crack Extension (mm)

Lower-Bound J-R Curve

Static-Cast Steels 25-40% Ferrite

290-320°C

Jd=46Δa0.34

CF-8M ≥10% Ni

Jd=81Δa0.34

CF-8M <10% Ni

Jd=241Δa0.34

CF-8 

Jd=256Δa0.34

CF-3 

0

200

400

600

800

1000

0 1 2 3 4 5

J (kJ/m2)

Crack Extension (mm)

Lower-Bound J-R Curve

Static-Cast Steels 15-25% Ferrite

290-320°C

Jd=70Δa0.34

CF-8M ≥10% Ni

Jd=119Δa0.34

CF-8M <10% Ni

Jd=260Δa0.34

CF-8 

Jd=281Δa0.36

CF-3 
47

 

 

Figure 22b Continued.

0

200

400

600

800

1000

0 1 2 3 4 5

J (kJ/m2)

Crack Extension (mm)

Lower-Bound J-R Curve

Static-Cast Steels 10-15% Ferrite

290-320°C

Jd=132Δa0.34

CF-8M ≥10% Ni

Jd=215Δa0.34

CF-8M <10% Ni

Jd=292Δa0.36

CF-8 Jd=323Δa0.38

CF-3 

0

200

400

600

800

1000

0 1 2 3 4 5

J (kJ/m2)

Crack Extension (mm)

Lower-Bound J-R Curve

Static-Cast Steels <10% Ferrite

290-320°C

Jd=220Δa0.34

CF-8M ≥10% Ni

Jd=276Δa0.34

CF-8M <10% Ni

Jd=320Δa0.37

CF-8 Jd=361Δa0.41

CF-3 
48

Figure 23a. Estimated lower-bound J-R curves at RT for centrifugally-cast CASS materials with 

ferrite contents of 25–40%, 15–25%, 10–15%, and <10%.

0

200

400

600

800

1000

0 1 2 3 4 5

J (kJ/m2)

Crack Extension (mm)

Lower-Bound J-R Curve

Centrifugally-Cast Steels 25-40% Ferrite

Room Temp.

Jd=33Δa0.34

CF-8M ≥10% Ni

Jd=72Δa0.34

CF-8M <10% Ni

Jd=282Δa0.36

CF-8 

Jd=314Δa0.38

CF-3 

0

200

400

600

800

1000

0 1 2 3 4 5

J (kJ/m2)

Crack Extension (mm)

Lower-Bound J-R Curve

Centrifugally-Cast Steels 15-25% Ferrite

Room Temp.

Jd=59Δa0.34

CF-8M ≥10% Ni

123Δa0.34

CF-8M <10% Ni

Jd=325Δa0.37

J CF-8 d=373Δa0.40

CF-3 
49

Figure 23a. Continued.

0

200

400

600

800

1000

0 1 2 3 4 5

J (kJ/m2)

Crack Extension (mm)

Lower-Bound J-R Curve

Centrifugally-Cast Steels 10-15% Ferrite

Room Temp.

Jd=142Δa0.34

CF-8M ≥10% Ni

Jd=245Δa0.36

CF-8M <10% Ni

Jd=403Δa0.40

CF-8 

Jd=484Δa0.44

CF-3 

0

200

400

600

800

1000

0 1 2 3 4 5

J (kJ/m2)

Crack Extension (mm)

Lower-Bound J-R Curve

Centrifugally-Cast Steels <10% Ferrite

Room Temp.

Jd=249Δa0.36

CF-8M ≥10% Ni

Jd=338Δa0.38

CF-8M <10% Ni

Jd=478Δa0.41

CF-8 

Jd=595Δa0.46

CF-3 
50

Figure 23b Estimated lower-bound J-R curves at 290–320°C for centrifugally-cast CASS 

materials with ferrite contents of 25–40%, 15–25%, 10–15%, and <10%. 

0

200

400

600

800

1000

0 1 2 3 4 5

J (kJ/m2)

Crack Extension (mm)

Lower-Bound J-R Curve

Centrifugally-Cast Steels 25-40% Ferrite

290-320°C

Jd=57Δa0.34

CF-8M ≥10% Ni

Jd=101Δa0.34

CF-8M <10% Ni

Jd=317Δa0.34

CF-8 Jd=336Δa0.34

CF-3 

0

200

400

600

800

1000

0 1 2 3 4 5

J (kJ/m2)

Crack Extension (mm)

Lower-Bound J-R Curve

Centrifugally-Cast Steels 15-25% Ferrite

290-320°C

Jd=87Δa0.34

CF-8M ≥10% Ni

Jd=150Δa0.34

CF-8M <10% Ni

Jd=342Δa0.34

CF-8 

Jd=369Δa0.36

CF-3 
51

/

Figure 23b Continued.

0

200

400

600

800

1000

0 1 2 3 4 5

J (kJ/m2)

Crack Extension (mm)

Lower-Bound J-R Curve

Centrifugally-Cast Steels 10-15% Ferrite

290-320°C

Jd=166Δa0.34

CF-8M ≥10% Ni

Jd=264Δa0.34

CF-8M <10% Ni

Jd=384Δa0.36

CF-8 

Jd=424Δa0.38

CF-3 

0

200

400

600

800

1000

0 1 2 3 4 5

J (kJ/m2)

Crack Extension (mm)

Lower-Bound J-R Curve

Centrifugally-Cast Steels <10% Ferrite

290-320°C

Jd=267Δa0.34

CF-8M ≥10% Ni

Jd=322Δa0.34

CF-8M <10% Ni

Jd=421Δa0.37

CF-8 

Jd=474Δa0.41

CF-3 
52

Figure 24. Comparison of the lower-bound fracture toughness J-R curve for thermally aged 

(a) static-cast and (b) centrifugally-cast CF-8M steels at 290°C estimated using the 

NUREG/CR-4513, Rev. 1 expressions, with the fracture toughness of SA welds 

used to develop ASME Section XI Subsection IWB-3640 evaluations.

The NRC staff compared the estimated lower-bound fracture toughness J-R curves developed

by Argonne for thermally aged CASS materials with the fracture toughness used in American 

Society of Mechanical Engineers (ASME) Section XI Subsection IWB-3640 for evaluating flux 

welds.82 The comparison is shown in Fig. 24 for static-cast and centrifugally cast CASS

materials. The results show that the lower-bound fracture toughness of thermally aged CASS 

material with up to 25% ferrite is similar to the fracture toughness J-R curve used in Section XI

0

100

200

300

400

500

0 1 2 3 4 5

Type 316 SAW

Type 304 SAW

J (kJ/m2)

Crack Extension (mm)

Jd=178.2Δa0.513

Reconstructed J–R Curve for

ASME IWB–3640 SAW 

Lower Bound Curve 

10-15% Ferrite

Lower-bound Curve

15-25% Ferrite

Lower Bound Curve 

<10% Ferrite

Static-Cast CF-8M Steels (a)

at 290°C

0

100

200

300

400

500

0 1 2 3 4 5

Type 316 SAW

Type 304 SAW

J (kJ/m2)

Crack Extension (mm)

Centrifugally-Cast CF-8M Steels

at 290°C

Jd=178.2Δa0.513

Reconstructed J–R Curve for

ASME IWB–3640 SAW 

Lower Bound Curve 

10-15% Ferrite

Lower-bound Curve

15-25% Ferrite

Lower-bound Curve

<10% Ferrite

(b)
53

IWB-3640 to evaluate SA welds. The procedures in ASME Section XI Subsection IWB-3640 

reduce the load-bearing capacity of the SS component to account for the reduced fracture 

toughness of the SA welds, based on elastic-plastic fracture mechanics analyses.86,87 Because 

the lower-bound fracture toughness of thermally aged CASS steel is similar to the fracture 

toughness of SA welds used in the elastic-plastic fracture mechanics analyses of IWB-3640, the

procedures in IWB-3640 for SA welds are directly applicable to CASS materials. 

Furthermore, since the flaw evaluation is based on the unaged material ultimate stress, this 

approach would be conservative. Experimental data indicate that the flow stress, that is, half of 

the sum of the ultimate and yield stresses, is increased by about 10%, 14%, and 24%, for CF-3 

and CF-3A, CF-8 and CF-8A, and CF-8M materials, respectively.50 A higher ultimate stress 

would increase the load bearing capability of a component. Therefore, the NRC staff has 

accepted the procedures developed in IWB-3640 of Section XI for SA welds for evaluating flaws 

in thermally aged CASS materials for addressing aging degradation concerns in license renewal

applications. 

The use of the IWB-3640 SA weld procedures for evaluating flaws in thermally aged CASS

components are considered conservative because:

1. The lower-bound fracture toughness of thermally aged CASS is similar to the fracture 

toughness used in IWB-3640 to evaluate SA welds.

2. The actual fracture toughness of a thermally aged CASS component in a U.S. utility plant 

would likely be higher than the Argonne lower-bound fracture toughness.

3. The thermally aged materials would be able to bear more load due to the increased ultimate 

stress resulting from thermal aging. This is not credited in this flaw evaluation procedure.

Thus, the use of the IWB-3640 SA weld procedures for evaluating flaws in thermally aged CASS

components is considered as a “screening” step to determine whether further detailed flaw 

evaluation accounting for actual plant-specific material properties should be performed. The 

procedures in IWB-3640 for SA welds have been available since the winter 1985 Addenda of 

Section XI, and have been applied successfully by utilities without resulting in unnecessary 

component repairs or replacements. However, the Argonne methodology of NUREG/CR-4513,

Rev. 1 was not applicable to CASS CF-8M materials with greater than 25% ferrite or for CF-8M 

materials containing more than trace amounts of Nb. Even then, the proposed flaw evaluation 

procedures based on the lower-bound fracture toughness of thermally aged CASS materials 

was expected to be sufficient for the vast majority of cases. 

In this report, the plots in Fig. 24 have been updated to include a comparison of the new lowerbound fracture toughness J-R curve presented in Figs. 22 and 23 with the fracture toughness 

J-R curve of SA welds used to develop ASME Section XI Subsection IWB-3640 evaluations. 

The results are shown in Figs. 25 and 26 for static-cast and centrifugally cast CASS materials, 

respectively. The results indicate that the fracture toughness J-R curve used in Section XI,

Subsection IWB-3640 to evaluate SA welds, bounds the lower-bound fracture toughness J-R 

curve for thermally aged CF-8M material with <10% Ni and containing 10-15% ferrite. The 

lower-bound J-R curve for materials with 15-25% ferrite is lower, particularly for static-cast 

materials. However, the updated lower-bound J-R curves represent a conservative estimate of

fracture toughness. Therefore, since the thermally aged components would be able to bear 

more load due to the increased ultimate stress resulting from thermal aging, the J-R curve in 
54

Figure 25. Comparison of the lower-bound fracture toughness J-R curve for thermally aged 

static-cast CF-8M steels containing (a) <10% Ni and (b) ≥10% Ni at 290°C 

estimated using the updated expressions presented in this report, with the fracture 

toughness of SA welds used to develop ASME Section XI Subsection IWB-3640 

evaluations.

Section XI, Subsection IWB-3640 to evaluate SA welds may be adequate for thermally aged 

CF-8M material with <10% Ni and containing up to 25% ferrite. The SA weld J-R curve is 

inadequate for CF-8M materials with <10% Ni and containing greater than 25% ferrite.

0

100

200

300

400

500

0 1 2 3 4 5

Type 316 SAW

Type 304 SAW

J (kJ/m2)

Crack Extension (mm)

Static-Cast CF-8M (Ni <10%)

at 290°C

Jd=178.2Δa0.513

Reconstructed J–R Curve for

ASME IWB–3640 SAW 

Lower Bound Curve 

15-25% Ferrite

Lower-bound Curve

10-15% Ferrite

(a)

Lower-bound Curve

25-40% Ferrite

0

100

200

300

400

500

0 1 2 3 4 5

Type 316 SAW

Type 304 SAW

J (kJ/m2)

Crack Extension (mm)

Static-Cast CF-8M (Ni ≥10%)

at 290°C

Jd=178.2Δa0.513

Reconstructed J–R Curve for

ASME IWB–3640 SAW 

Lower Bound Curve 

15-25% Ferrite

Lower-bound Curve

25-40% Ferrite

Lower-bound Curve

10-15% Ferrite

(b)
55

Figure 26. Comparison of the lower-bound fracture toughness J-R curve for thermally aged 

centrifugally-cast CF-8M steels containing (a) <10% Ni and (b) ≥10% Ni at 290°C 

estimated using the updated expressions presented in this report, with the fracture 

toughness of SA welds used to develop ASME Section XI Subsection IWB-3640 

evaluations.

Figures 25 and 26 indicate that for CF-8M materials containing ≥10% Ni, the J-R curve of SA 

weld may be adequate only for materials containing up to 15% ferrite. The lower-bound J-R 

curves presented in Figs. 22 and 23 should be used for materials containing greater than 15% 

ferrite.

0

100

200

300

400

500

0 1 2 3 4 5

Type 316 SAW

Type 304 SAW

J (kJ/m2)

Crack Extension (mm)

Centrifugally-Cast CF-8M (Ni <10%)

at 290°C

Jd=178.2Δa0.513

Reconstructed J–R Curve for

Lower-bound Curve ASME IWB–3640 SAW 

>25% Ferrite

Lower-bound Curve

10-15% Ferrite

(a)

Lower Bound Curve 

15-25% Ferrite

0

100

200

300

400

500

0 1 2 3 4 5

Type 316 SAW

Type 304 SAW

J (kJ/m2)

Crack Extension (mm)

Centrifugally-Cast CF-8M (Ni ≥10%)

at 290°C

Lower Bound Curve 

15-25% Ferrite

Lower-bound Curve

>25% Ferrite

Lower-bound Curve

10-15% Ferrite

(b)

Jd=178.2Δa0.513

Reconstructed J–R Curve for

ASME IWB–3640 SAW 
56

2.5.3 Updated Fracture Toughness Data for CASS Materials 

Fracture toughness J-R curve data obtained on four heats of CF-8M material at EdF,38 after the 

publication of NUREG/CR-4513, Rev. 1, are shown in Figs. 27 and 28. All heats except Heat 

EK contained ≥10% Ni, and all heats except DI contained >25% ferrite. The corresponding 

updated lower-bound J-R curve for the grade of material and range of ferrite content is also 

;

Figure 27. Fracture toughness J-R curve data at 320°C for static-cast and thermally aged 

plates (Heats CC and DI) and an elbow (Heat EL) of CF-8M material with ≥10% Ni 

and (a) >25% and (b) 15–25% ferrite (Ref. 38). 

0

50

100

150

200

0 0.5 1 1.5 2 2.5 3 3.5 4

EL 3k @ 400°C

CC 1k @ 400°C

J (kJ/m2)

Crack Extension (mm)

Aged CF-8M CASS ≥10% Ni

Tested at 320°C 

Aging Condition 

Ferrite Content 30%

(a)

Lower Bound Curve

J = 46Δa0.34 

0

100

200

300

400

0 0.5 1 1.5 2 2.5 3 3.5 4

DI 10k @ 400°C

DI 0.7k @ 400°C

J (kJ/m2)

Crack Extension (mm)

Aged CF-8M CASS ≥10% Ni

Tested at 320°C 

Aging Condition

Ferrite Content 19%

(b)

Lower Bound Curve

J = 70Δa0.34
57

Figure 28. Fracture toughness J-R curve data at 320°C for static-cast and thermally aged 

elbow (Heat EK) of CF-8M material with <10% Ni and >25% ferrite (Ref. 38). 

plotted in the figure (i.e., curves shown in Figs. 22 and 23 and associated power-law J-R curve 

parameters listed in Table 3). The results indicate that for all CF-8M materials, the applicable 

lower bound J-R curve bounds all of the J-R curve data. 

A recent scoping study on low-temperature crack propagation for CASS CF-8 material in PWR 

environments investigated the potential synergy between hydrogen embrittlement and thermal 

embrittlement associated with decomposition of the ferrite at reactor temperatures.88 Fracture 

toughness J-R curve tests were conducted in air and PWR shutdown water chemistry (SWC) at 

54°C on 1-T CT specimens of CF-8 material (Argonne Heat 68) that was thermally aged for 

about 138,000 h (about 15.8 yr) at 350°C. The ferrite content calculated from Hull’s equivalent 

factors was 15%, and measured using a ferritescope was 23%. The specimens were either 

fatigue precracked in air at 54°C or fatigue plus SCC precracked in PWR water at 315°C. The 

fracture toughness J-R curve data for these tests are shown in Fig. 29. Note that one of the 

specimens that was precracked in PWR water at 315°C was inadvertently tested at elevated 

temperature of 315°C in PWR primary water chemistry. The updated lower-bound J-R curve at 

room temperature for CF-8 material with 15–25% ferrite is also shown in the figure for 

comparison. 

The results indicate that at 54°C, the J-R curve data in air are bounded, with additional margin, 

by the updated lower-bound J-R curve. However, the J-R curve data for specimens precracked 

in air at 54°C and then tested in PWR (SWC) at 54°C are significantly below the lower-bound 

curve. These results show an apparent large effect of the coolant environment on fracture 

toughness; every material and test condition is identical for the two sets of duplicate tests 

except the one was tested in air and the other in water. The specimens that were precracked in 

PWR water at 315°C and then tested in PWR water at 54°C or 315°C also show reduced 

fracture toughness relative to that in air, but the difference is less. The J-R curve data for the 

specimen that was cooled down to from 315°C and then tested at 54°C SWC are only 

0

100

200

300

400

0 0.5 1 1.5 2 2.5 3 3.5 4

EK 30k @ 350°C

EK 30k @ 325°C

J (kJ/m2)

Crack Extension (mm)

Aged CF-8M CASS <10% Ni

Tested at 320°C 

Aging Condition 

Ferrite Content 30%

Lower Bound Curve

J = 81Δa0.34 
58

marginally below the lower-bound J-R curve, and the data for the specimen that was precracked 

and tested at 315°C PWR primary water chemistry are slightly above the lower-bound curve. 

The reason for the drastic reduction in fracture toughness in PWR SWC at 54°C relative to that 

in air at 54°C is not clear. Additional fracture toughness tests on thermally aged CASS CF-8 

materials in air and LWR environments should be conducted to better understand the combined

effects of hydrogen embrittlement and thermal embrittlement in LWR environments.

2.5.4 Screening Criteria for Susceptibility of CASS Materials to Thermal Embrittlement 

An EPRI report on thermal aging embrittlement of CASS 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 thermally aged CASS materials.46 

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 found that 

using J2.5 = 255 kJ/m2 is an acceptable screening approach for fracture toughness of CASS 

materials.45 The screening criteria to determine the susceptibility of CASS components to 

thermal aging embrittlement are outlined in Table 4.

45 

The expressions presented in Fig. 21, and the methodology described in Section 2.4.2 for 

estimating the corresponding fracture toughness J-R curve for the RT Charpy-impact energy 

trend curves in Fig. 21, can be used to check the validity of the ferrite content criteria given in 

Table 4. The saturation values of J at 2.5 mm crack extension corresponding to the RT Charpyimpact energy trend curves in Fig. 21 are shown in Figs. 30 and 31 for static-cast and 

Figure 29. Fracture toughness J-R curve data for thermally aged Heat 68 of CF-8M plate at 

54°C. The curve represent the lower bound curve at room temperature for staticcast CF-8 material (Ref. 88).

0

200

400

600

800

1000

0 1 2 3 4 5

54°C Air

54°C Air

54°C Air

54°C Air

315°C PWR Env.*

315°C PWR Env.

J (kJ/m2)

Crack Extension (mm)

CF-8 CASS Ht. 68 (15% ferrite)

Aged 138,000 h at 350°C

Tested at 54°C in 

Open Symbols: Air

Closed Symbols: PWR SWC

Precracked at 

Lower-bound CF-8 

at Room Temp.

J = 346Δa0.40

* Tested at 315°C in PWR Env.
59

Figure 30. The change in saturation J at 2.5 mm crack extension determined at (a) RT and 

(b) 290-320°C as a function of ferrite content for static-cast CF-3, CF-8, and CF8M CASS materials.

10

100

1000

0 5 10 15 20 25 30 35 40

Saturation J at 2.5 mm Crack Extension (kJ/m2)

Ferrite Content (%)

Static-Cast CASS Materials

Thermally Aged to Saturation

CF-3

CF-8

CF-8M (Ni <10%)

CF-8M (Ni ≥10%)

J = 255 kJ/m2

J values at Room Temp.

(a)

10

100

1000

0 5 10 15 20 25 30 35 40

Saturation J at 2.5 mm Crack Extension (kJ/m2)

Ferrite Content (%)

Static-Cast CASS Materials

Thermally Aged to Saturation

CF-3

CF-8

CF-8M (Ni <10%)

CF-8M (Ni ≥10%)

J values at 290-320°C

J = 255 kJ/m2

(b)
60

Figure 31. The change in saturation J at 2.5 mm crack extension determined at (a) RT and 

(b) 290-320°C as a function of ferrite content for centrifugally-cast CF-3, CF-8, 

and CF-8M CASS materials.

10

100

1000

0 5 10 15 20 25 30 35 40

Saturation J at 2.5 mm Crack Extension (kJ/m2)

Ferrite Content (%)

Centrifugally-Cast CASS Materials

Thermally Aged to Saturation

CF-3

CF-8

CF-8M (Ni <10%)

CF-8M (Ni ≥10%)

J = 255 kJ/m2

J values at Room Temp.

(a)

10

100

1000

0 5 10 15 20 25 30 35 40

Saturation J at 2.5 mm Crack Extension (kJ/m2)

Ferrite Content (%)

CF-3

CF-8

CF-8M (Ni <10%)

CF-8M (Ni ≥10%)

J values at 290-320°C

J = 255 kJ/m2

Centrifugally-Cast CASS Materials (b)

Thermally Aged to Saturation
61

centrifugally-cast CASS materials, respectively. The results indicate that for CF-3 and CF-8 

steels, values of J2.5 for both static-cast and centrifugally-cast materials are significantly above 

the threshold value of 255 kJ/m2, particularly for centrifugally-cast materials. Therefore, the 

screening criterion that the reduction in fracture toughness due to thermal embrittlement of 

centrifugally-cast CF-3 and CF-8 steels is "nonsignificant" is valid, and that the reduction in 

fracture toughness of static-cast CF-3 and CF-8 steels is "significant" only for steels with greater 

than 20% is a conservative criterion. 

Table 4. Screening criteria for thermal-aging susceptibility of CASS CF-3, CF-8, and CF-8M

materials.

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

The results for CF-8M steels with <10% Ni indicate that the J2.5 values are above the 255 kJ/m2

threshold value for ferrite contents less than 16% for static-cast materials, and less than about 

19% for centrifugally-cast materials. Therefore, for static cast CF-8M materials with <10% Ni, 

the existing screening criterion that materials with >14% ferrite are considered susceptible to 

thermal embrittlement is a conservative criterion. However, for centrifugally cast CF-8M 

materials with <10% Ni, the existing screening criterion that materials with >20% ferrite are 

considered susceptible to thermal embrittlement should be decreased to >19% ferrite.

Similarly for CF-8M steels with ≥10% Ni, the J2.5 values are above the 255 kJ/m2 threshold 

value for ferrite contents less than 11% for static-cast materials, and less than about 13% for 

centrifugally-cast materials. Therefore, the existing screening criteria for the susceptibility of 

CF-8M materials with ≥10% Ni, to thermal embrittlement needs to be revised. The threshold 

ferrite content above which CF-8M materials with ≥10% Ni are considered “potentially 

susceptible” to thermal embrittlement, should be decreased from 14% to 11% ferrite for staticcast materials, and from 20% to 13% for centrifugally-cast materials. The updated screening 

criteria for thermal-aging susceptibility of CASS CF-3, CF-8, and CF-8M materials are given in 

Table 5.

2.5.5 Methodology for Estimating Thermal Embrittlement of CASS Materials

A flow diagram for estimating mechanical properties of CASS materials during reactor service is 

shown in Fig. 32. The estimation scheme is divided into three sections based on available 

material information. In Section A of the flow diagram, “predicted lower-bound” fracture 

toughness is defined for CF-3, CF-8, and CF-8M steels of unknown composition (i.e., CMTR is 

not available). For materials with unknown composition, when the ferrite content of the steel is 

known, a different lower-bound fracture toughness and impact energy are defined for steels 

containing <10%, 10–15%, 15–25%, or 25–40% ferrite. However, as discussed in Section 2.2, 

the ferrite content used in this methodology should be representative of the material.
62

Figure 32. Flow diagram for estimating mechanical properties of thermally aged CASS 

materials in LWR systems.

Estimate for known composition and service history

Yes Yes

No Sat. J–R

> Min. J–R

Unaged

No

Minimum J–R

Curve Unaged

Estimate for known composition unknown service history

Sat. J–R

> Initial J–R

Unaged

No

No Saturation

J–R Curve &

Impact Energy

Input Initial Yes

Impact Energy

Initial

Impact

Energy

Estimate Initial

J–R Curve

Minimum J–R

Curve Unaged

Initial

J–R Curve

Unaged

Minimum

J–R Curve

Unaged

Yes Yes

Estd. J–R

> Min. J–R

Unaged

No

Initial Impact

Energy 160 J 

No

Estimate

 ServiceTime

Impact Energy &

J–R Curve

Yes

Section B: 

SATURATION

Yes

CMTR

ESTIMATION SCHEME

Input 

Composition &

Casting Process

Estimate Saturation

Impact Energy & 

J–R Curve

Service

Time &

Temp.

Input Service

Temp. & Time

Section C: 

SERVICE TIME

Service Time

J–R Curve &

Impact Energy

Estimate Initial

J–R Curve

Estimate

 ServiceTime

Impact Energy &

J–R Curve

Estd. J–R

> Initial J–R

Unaged

Initial

Impact

Energy

Yes Input Initial

Impact Energy

No

Estimate for unknown composition

Lower–Bound

J–R Curve for

CF–3, 8, & 8M

Lower–Bound

J–R Curve for

Ferrite Range

No

Yes

Section A: 

LOWER 

BOUND

Ferrite No

content 

Minimum

J–R Curve

Unaged

Initial

J–R Curve

Unaged

Initial

Flow

Stress

Yes

Estimate Service

 Time Flow Stress,

Yield Stress, & 

Ramberg/Osgood

Parameters

No

Table 5. Updated screening criteria for thermal-aging susceptibility of CASS materials.

Mo Content (wt.%) Casting Method Ferrite Content Susceptibility Determination

High (2.0–3.0) Static ≤14% (≤11%) Not susceptible

With <10% Ni (≥10% Ni) >14% (>11%) Potentially susceptible

Centrifugal ≤19% (≤13%) Not susceptible

>19% (>13%) Potentially susceptible

Low (0.5 max.) Static ≤20% Not susceptible

>20% Potentially susceptible

Centrifugal All Not susceptible
63

The methodology for estimating mechanical properties when a CMTR is available (i.e., the 

chemical composition of the CASS material is known) is presented in Sections B and C of the 

flow diagram. Section B describes the estimation of “saturation” impact energy and fracture 

toughness J-R curve (i.e., the lowest values that would be achieved for a completely thermallyembrittled material). The only information needed for these estimations is the chemical 

composition of the material, which is used to estimate the saturation J-R curve for the thermally 

aged material. The correlations presented in this report account for the degradation of 

mechanical properties due to thermal aging. They do not explicitly consider the initial fracture 

properties of the unaged material. Some heats of inherently weak CASS materials may have 

poor fracture properties in the unaged condition, and estimations of saturation fracture 

toughness based on the methodology presented in this report may be higher than the fracture 

toughness of the unaged material. The available fracture toughness J-R curve data at 290–

320°C (554–608°F) for unaged CASS materials indicate that the J-R curves for a few “weak” 

heats of static-cast steels may be lower than those for wrought SSs. Therefore, the saturation 

fracture toughness properties that should be used for design analyses for thermally aged CASS 

materials depend on whether or not the estimated saturation fracture toughness is lower than 

the initial fracture toughness of the unaged materials. The initial fracture toughness of the 

material can be estimated from its RT Charpy-impact energy. 

The fracture toughness of the CASS material is not available in CMTRs. Two different options 

are used to establish the saturation fracture toughness of the thermally aged material. In the 

first option, the initial RT Charpy-impact energy of the unaged material is known, and this value 

is used to estimate the initial fracture toughness of the material. CASS materials with poor 

fracture properties are relatively insensitive to thermal aging and, typically, the fracture 

toughness of the material does not change significantly due to thermal aging during reactor 

service. Therefore, if the estimated saturation fracture toughness of aged material is higher 

than the fracture toughness of the unaged material, the latter is used as the worst-case 

toughness for that material. 

In the second option, the initial RT Charpy-impact energy is not known and, therefore, the 

minimum fracture toughness of all unaged heats and heat treatment conditions of CASS 

materials is used as an upper bound for the estimated fracture toughness. This minimum 

upper-bound fracture toughness of unaged CASS materials is used as the worst-case fracture 

toughness for the material if the estimated saturation fracture toughness of thermally aged 

material is higher. The fracture toughness of unaged CASS materials is slightly higher at RT 

than at 290–320°C. However, for convenience, the minimum fracture toughness of unaged 

static-cast CASS materials at temperatures between RT and 320°C, can be expressed as 

Jd = 400[Δa]0.40, (43)

and of centrifugally-cast CASS materials, as

Jd = 650[Δa]0.43. (44)

Estimation of mechanical properties at any given time and temperature of service (i.e., service 

time properties) is described in Section C. The initial impact energy of the unaged material is 

required for these estimations. If not known, the initial impact energy of 200 J/cm2 (118 ft·lb) is 

assumed. However, similar to Section B, initial fracture toughness of the unaged material or the 

minimum fracture toughness of unaged cast stainless steels is used as an upper bound for the 

estimations. The initial tensile properties of the unaged material are needed for estimations of 

the tensile strength and Ramberg/Osgood strain hardening parameters. If the initial flow stress 
64

of the CASS material is known, the JIc value and tearing modulus of the thermally aged material 

can then be determined from the estimated values of J-R curve and flow stress of the aged 

material. 

However, as mentioned earlier, the above methodology for estimating fracture toughness of 

CASS materials in reactor service is not applicable for CF-8M materials with more than a trace 

amount of Nb (e.g., greater than 0.05%). The chemical requirements for ASTM A351 or A451 

do not specify any upper limit for Nb. Typically, CF-8M steels contain only trace amounts of Nb. 

However, the Nb content could be high in castings produced using Type 347 scrap metal. For 

example, In the Argonne study, the measured RT Charpy-impact energy of a thermally aged 

heat of CF-8M steel with about 23% ferrite and containing 0.2% Nb, was significantly lower than 

that predicted from the above methodology. The fracture surface of the Charpy specimen 

showed that the phase boundaries were decorated with large Nb carbide particles that cracked 

easily. The available thermal embrittlement data on CF-8M materials is inadequate to 

accurately establish the threshold Nb content above which it contributes to increased 

embrittlement of the material. Therefore, for CF-8M materials with greater than 15% ferrite, it 

would be advisable to determine the Nb content of the material, and take appropriate action if 

the it is greater than 0.05%.

2.6 Extent of Thermal Embrittlement of Austenitic Stainless Steel Welds

In an earlier study at Argonne48–51 on thermal embrittlement of CASS materials the 

methodology and correlations were developed for estimating fracture toughness, tensile, and 

Charpy-impact properties of CASS components during service from known material information. 

Although austenitic SS welds have a duplex structure and their chemical compositions are 

similar to those of CASS materials, the estimation scheme was not applicable to austenitic SS 

welds. The degradation of fracture toughness, tensile, and Charpy-impact properties of welds 

due to thermal aging was characterized in another report.53 Five Type 304/308 or 308L welds 

were aged for 7,000–1,0000 h at 400°C to simulate saturation conditions (i.e., lowest impact 

energy that would be achieved by the material after long-term aging) and the results were 

compared with fracture toughness data from other studies. The ferrite morphologies of four of 

the five welds that were investigated are shown in Fig. 4. The J-R curves are expressed by the 

power-law relation Jd = C(Δa)n per ASTM Specifications E 813-85 and E 1152-87. 

The results indicated that for austenitic SS welds, thermal aging resulted in moderate decreases 

in Charpy-impact strength and fracture toughness at both room temperature and 290°C. For the 

various welds, upper-shelf energy (USE) decreased by 50–80 J/cm2 (30–47 ft·lb.). The 

decrease in the fracture toughness J-R curve or JIc was relatively small. Metallographic 

examination of the specimens indicated that failure occurred by the formation and growth of 

microvoids near hard inclusions. Such processes are relatively insensitive to thermal aging. 

Fracture resistance of the PWWO weld was inferior to that of the PWCE weld because of a 

higher density and a larger size of inclusions. Furthermore, the ferrite phase had little or no 

effect on the fracture properties of the welds because of its thin vermicular morphology and 

relatively low content (e.g., 6–8% ferrite). 

The fracture toughness J-R curve data from the work conducted for the NRC and compiled in 

the Pipe Fracture (PIFRAC) Database* and from a few other sources,40,89–93 are shown in 

Fig. 33. The PIFRAC database, consisting of the data from Refs. 94–98 was originally compiled 

at MEA,

99 and updated later by Battelle Memorial Institute.100 The results indicate that fracture 

 * G. Wilkowski and N. Ghadiali, “Short Crack in Piping and Piping Welds,” in Technical Data CD-ROM, Battelle 

Columbus Division, Columbus, OH (May 1995).
65

properties of SS welds are relatively insensitive to filler metal.40 However, the welding process 

significantly affects fracture toughness. In general, GTA welds exhibit higher fracture resistance 

than SMAs or SA welds. The statistical differences in SA and SMA weld fracture toughness J-R 

curves have also been evaluated101 and results indicate no difference between SA and SMA 

welds J-R curves. At 288°C, the lower-bound fracture toughness J-R curve for SA and SMA 

welds101 is represented by 

J(kJ/m2) = 73.4 + 83.5 Δa(mm)0.643 (45)

Figure 33. Fracture toughness J-R curves for stainless steel welds at (a) RT and 

(b) 288–427°C. Solid line represents lower-bound curve. 

0

500

1000

1500

2000

2500

3000

0 2 4 6 8 10

SAW

SMAW

GTAW 

J (kJ/m2)

Crack Extension (mm)

Austenitic Stainless Steel Welds

Open Symbols: Type 308

Closed Symbols: Type 316

Tested at Room Temperature (a)

0

500

1000

1500

2000

2500

0 2 4 6 8 10

308 SAW

308 SMAW

308 GTAW

J (kJ/m2)

Crack Extension (mm)

Austenitic Stainless Steel Welds

Open Symbols: Type 308

Closed Symbols: Type 316

Tested at 288-427°C

J = 73.4 + 83.5Δa0.643

(b)
66

where 73.4 kJ/m2 is the fracture toughness JIc. The lower-bound curve for SA and SMA welds 

shows very good agreement with the data in Fig. 33. The fracture toughness data used in the 

analysis performed for the technical basis document for ASME Section XI Article IWB-3640 are 

somewhat higher than the curve given by Eq. 45. The data scatter in the fracture toughness 

J-R curve obtained at 200 and 288°C (392 and 550°F) for Linde 80 weld metal102 is shown in 

Fig. 34. The J-R curve reconstructed form the data for the ASME IWB-3640 SA weld and the 

lower-bound J-R curve proposed in NUREG/CR-6004 for austenitic SS welds (i.e., Eq. 45), are 

also shown in the figure. The results indicate that most of the Linde 80 data are bounded by the 

lower-bound curve. 

Figure 34. Fracture toughness J-R curves for Linde 80 welds at 200 and 288°C and the lowerbound J-R curve for unaged austenitic SS welds. 

The available fracture toughness J-R curves for aged SMA, SA, and GTA welds are shown in 

Fig. 35.

41,90,103 In these studies, the time and temperature of aging was sufficient to achieve

saturation toughness (i.e., the minimum value that could be achieved after long-term aging). 

The JIc values for unaged and aged welds are plotted in Fig. 36. The results indicate that the JIc

values for SA or SMA welds are generally lower than for the GTA welds. The one exception is 

the values for a Type 316L GTA weld, which are significantly lower than typically observed for 

GTA weld, both at RT and at 288°C. The reason for the low fracture toughness values is not 

known. The data also indicate that at reactor temperatures, the fracture toughness JIc of SS 

welds can be as low as 40 kJ/m2. Hence, the fracture toughness J-R curves for fully embrittled 

SMA and SA welds can be lower than that predicted by Eq. 45; a conservative estimate for 

aged welds may be expressed as 

J(kJ/m2) = 40 + 83.5 Δa(mm)0.643. (46)

This curve is plotted in Fig. 35. The fracture toughness J-R curves for unaged and aged SS 

welds (i.e., Eqs. 45 and 46, respectively) are compared in Fig. 37 with the data for aged 316L 

0

100

200

300

400

500

600

0 1 2 3 4 5 6 7

62W-121

62W-23

62W-24

62W-33

62W-45

62W-137

62W-157

62W-50

J (kJ/m2)

Crack Extension (mm)

Linde 80 Weld

Open Symbols: 200°C

Closed Symbols: 288°C 

J = 73.4 + 83.5Δa0.643

Lower-bound curve 

in NUREG/CR-6004

Jd=178.2Δa0.513

Reconstructed J-R Curve for

ASME IWB-3640 SAW
67

and CF-3 welds90,104 and the data in the technical basis document for ASME Section XI Article 

IWB-3640.95 Note that the data from Ref. 95 are modified J, and the deformation J values 

would be slightly lower. 

Figure 35. Fracture toughness J-R curves for aged stainless steel welds at (a) RT and 

(b) 288°C. Solid line represents lower-bound curve. 

0

500

1000

1500

2000

2500

3000

0 2 4 6 8 10

SMAW

GTAW

J (kJ/m2)

Crack Extension (mm)

Austenitic Stainless Steel Welds

Open Symbols: Type 308

Closed Symbols: Type 316

Tested at Room Temperature

J = 40 + 83.5Δa0.643

(a)

0

500

1000

1500

2000

2500

3000

0 2 4 6 8 10

SMAW 

J (kJ/m2)

Crack Extension (mm)

Austenitic Stainless Steel Welds

Open Symbols: Type 308

Closed Symbols: Type 316

Tested at 288°C

J = 40 + 83.5Δa0.643

(b)
68

Figure 36.

Fracture toughness JIc for 

unaged and aged stainless 

steel welds.

Figure 37. Fracture toughness lower-bound J-R curves and the data on unaged and aged 304, 

316L, and CF-3 welds used to develop the ASME Code IWB-3640 analysis.

0

100

200

300

400

500

600

700

800

0 100 200 300 400 500 600

308L GTAW

308 SMAW

316L MMAW

308 SMAW

308 SAW

308 GTAW

316 SMAW

316 SAW

316 GTAW

316L GTAW

Fracture Toughness JIc (kJ/m2)

Test Temperature (°C)

Austenitic Stainless Steels

Closed Symbols: Aged

0

100

200

300

400

500

600

700

0 2 4 6 8 10

CF–3 MMAW Hale & Garwood

(Aged 10,000 h at 400°C)

316L SAW Chipperfield

(Aged 10,000 h at 400°C)

304 SAW Landes & McCabe

316 SAW Landes & McCabe

J (kJ/m2)

Crack Extension (mm)

Austenitic Stainless Steel Welds At 288°C

J = 40 + 83.5Δa0.643

Aged to Saturation

J = 73.4 + 83.5Δa0.643

Unaged Material

J-R Curve 316 SAW

Landes & McCabe
69

2.6.1 Update of Thermal Embrittlement Data for Austenitic Stainless Steel Welds

The J-R curve data for GTA, SMA, and SA welds of Type 308 and 16-8-1 SSs in air at 24 and 

427°C is shown in Fig. 38.

40 The ferrite contents measured in the Type 16-8-1 GTA and SA 

welds were 5.7% and 9.0% respectively, and in the Type 308 GTA, SA, and SMA welds were 

9.9%, 10.7%, and 7.0%, respectively. The results indicate that the fracture toughness of GTA 

weld is superior to that of SMA weld, which is superior to that of the SA weld. The fracture 

toughness at room temperature (i.e., closed symbols in the figure) is slightly higher than that at 

427°C (i.e., open symbols). All of the data are bounded by the lower-bound J-R curve given by 

Eq. 45. 

The fracture toughness J-R curve data for unaged and aged Type 308 SS GTA and SMA welds 

in air at 427°C is shown in Fig. 39.

41 The results indicate little or no effects of long-term aging 

(i.e., 10,000 h at 400°C) on the fracture toughness of GTA and SMA welds in air at 427°C. Little 

or no effect of thermal aging at 400°C on the fracture toughness of Type 304/308 SS pipe welds 

was also observed in an earlier study at Argonne.

53 The fracture toughness J-R curve data at 

25 and 290°C for unaged and aged Type 304/308 pipe weld and a Type 304/308L pipe weld are 

shown in Figs. 40a and 40b, respectively.53 The Type 308 weld was aged 10,000 h at 400°C 

and the Type 308L weld aged 7,700 h at 400°C. The measured ferrite content was 6.1% in the 

Type 304/308 and 6.8% in the Type 304/308L weld. The difference between the J-R curve data 

at both room temperature and 290°C, for unaged and aged welds is within the typical data 

scatter for fracture toughness tests. The results suggest that for Type 308 weld metal, thermal 

aging may not have a significant effect on its fracture toughness. However, the data for CF-3 

MMA and 316L SA welds aged for 10,000 h at 400°C, the measured fracture toughness J-R 

curve data are both below the lower-bound J-R curve given by Eq. 45. A lower-bound J-R curve 

expressed by Eq. 46 was proposed in NUREG/CR-6428 for thermally aged austenitic SS welds. 

Figure 38. Fracture toughness J-R curve data for GTA, SMA, and SA welds in air at 

RT and 427°C (Ref. 40).

0

200

400

600

800

1000

0 0.5 1 1.5 2 2.5

16-8-1 SS GTA

308 SS GTA

308 SS GTA

16-8-1 SS SA

308 SS SA

308 SS SMA

J (kJ/m2)

Crack Extension (mm)

Austenitic Stainless Steel Welds

Open Symbols: 427°C

Closed Symbols: 24°C J = 73.4 + 83.5Δa0.643

Tested in Air
70

Figure 39. Fracture toughness J-R curve data for unaged and aged Type 308 SS GTA 

and SMA welds in air at 427°C (Ref. 40).

0

200

400

600

800

1000

0 0.5 1 1.5 2 2.5

GTA

SMA

J (kJ/m2)

Crack Extension (mm)

Type 308 SS Weld

Open Symbols: As welded

Closed Symbols: Aged 

J = 73.4 + 83.5Δa0.643

Tested at 427°C in Air

Welds aged 10,000 h at 427°C
71

Figure 40. Fracture toughness J-R curve data for unaged and aged (a) Type 304/308 

and (b) Type 304/308L pipe welds in air at 25 and 290°C (Ref. 53). 

The fracture toughness J-R data obtained on two heats of Type 316L GTA weld obtained in air 

and in high-purity water with 300 ppb DO are plotted in Fig. 41.

105,106 The measured ferrite 

contents in the two welds were 10% and 14%. The results indicate that in-situ fracture 

toughness in the reactor coolant environment can result in a reduction of toughness by up to 

40% relative to that in air. The in-situ properties degradation was attributed to absorption of 

hydrogen in the material during exposure to the high temperature aqueous environment.105 The 

results also show that the fracture toughness of unaged austenitic SS welds in LWR 

0

400

800

1200

1600

0 0.5 1 1.5 2 2.5 3

25°C

290°C

J (kJ/m2)

Crack Extension (mm)

Type 304/308 Pipe Weld 

0.71 m (28 in.) diameter

Open Symbols: As welded

Closed Symbols: Aged 10k h at 400°C 

J = 73.4 + 83.5Δa0.643

Tested in Air at

(a)

0

200

400

600

800

1000

0 0.5 1 1.5 2 2.5 3

25°C

290°C

J (kJ/m2)

Crack Extension (mm)

Type 304/308L Pipe Weld 

0.31 m (12 in.) diameter

Open Symbols: As welded

Closed Symbols: Aged 7.7k h at 400°C 

J = 73.4 + 83.5Δa0.643

Tested in Air at

(b)
72

environments may be below the lower-bound J-R curve for aged welds given by Eq. 46. 

Additional data are needed to establish the in-situ fracture toughness degradation of SS welds 

in LWR environments.

Figure 41. Fracture toughness J-R curve data for as-welded Type 316L GTA weld at 288°C in 

air and BWR environment with 300 ppb DO (Ref. 105).

0

50

100

150

200

250

0 0.5 1 1.5 2 2.5

10% Ferrite

14% Ferrite

J (kJ/m2)

Crack Extension (mm)

Type 316 GTA Welds

Open Symbols: Air

Closed Symbols: 300 ppb DO

Tested at 288°C

J = 40 + 83.5Δa0.643
73

3. IASCC CRACK GROWTH RATES 

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

components is IASCC of structural materials (i.e., primarily austenitic SSs). Several incidents of 

IASCC have occurred since the mid-1970s in control blade handles and instrumentation tubes 

of BWRs, and since the 1990s, in BWR core shroud and PWR baffle bolts. As the name 

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

materials. The susceptibility of austenitic SSs to IASCC has been investigated by conducting 

slow-strain-rate-tensile (SSRT), CGR, and crack initiation tests on irradiated material in 

simulated LWR environments107–117. The factors that influence IASCC of austenitic SSs include 

neutron irradiation conditions such as neutron fluence, flux, and energy spectrum, cold work, 

material composition, corrosion potential, water purity, temperature, and loading conditions. 

The susceptibility of austenitic SSs to IASCC increases with neutron fluence, corrosion 

potential, and water conductivity. 

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. 

Experimental data on Type 304 and 316 SSs irradiated up to 4.0 × 1021 n/cm2 (6.0 dpa) show a 

beneficial effect of reducing the corrosion potential of the environment on IASCC 

susceptibility.27,28 For example, the addition of H2 to the reactor water greatly reduces the effect 

of radiolysis by scavenging the radiolysis products.

117 Since PWR coolants typically contain 

2 ppm H2 by weight or 30 cc/kg, radiolysis has no effect on the corrosion potential in PWRs.

However, a low corrosion potential does not provide immunity to IASCC if the fluence is high 

enough; intergranular stress corrosion cracking (IGSCC) has been observed in baffle bolts in 

PWRs. Threshold fluence for IASCC is higher under low-potential BWR hydrogen water 

chemistry (HWC) or PWR primary water chemistry. It also varies with material composition and 

thermo-mechanical treatment.

27

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.4,27,118,119 A threshold fluence level of 5 × 1020 n/cm2 (E >1 MeV)∗

(0.75 dpa)† has been proposed for austenitic SSs in NWC BWR environments. However, the 

results in Fig. 42 indicate that the intergranular cracking susceptibility in some commercial-purity 

SSs increases rapidly at fluence levels above about 2 × 1020 n/cm2 (0.3 dpa) and in high-purity 

heats of SSs at even lower fluence levels.

3.1 Effects of Material and Environmental Parameters on IASCC Susceptibility

In an earlier Argonne study,21 the existing data have been reviewed to identify the key 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) that influence IASCC, and to 

establish their effects on IASCC susceptibility of austenitic SSs. Discussions of some of those 

parameters follow.

 ∗ Unless otherwise mentioned otherwise all references to fluence levels are calculated for E ≥1 MeV.

† Conversion to dpa for LWRs, is E > 1 MeV and 1022 n/cm2 ~ 15 dpa.
74

Figure 42.

Susceptibility of irradiated 

austenitic SSs to IGSCC as a 

function of fluence in high-DO 

water.

3.1.1 Microstructure

The microstructural changes in austenitic SSs due to neutron irradiation vary with the material 

composition, irradiation temperature, and neutron fluence, flux, and energy spectrum. 

Irradiation damage is characterized by either the neutron fluence (n/cm2) or the average number 

of displacements per atom (dpa). For BWR-irradiated SSs (i.e., irradiated at 275–300°C), the 

loop density saturates at about 1 dpa, and the average loop diameter saturates at 5 dpa. In 

general, the loop size increases and loop density decreases with irradiation temperature. 

Irradiation temperatures above 350°C lead to the formation of second phase particles.5,23,120–123 

Cavities and voids form at high doses and high temperatures. Under LWR conditions, metal 

carbides are the primary stable precipitates in 300-series SSs, 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; Si, P, and Ni are enriched at regions 

such as grain boundaries, while Cr, Mo, and Fe are depleted. The extent of RIS depends on 

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. 

3.1.2 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.2,4,5 The extent of segregation or depletion depends on the 

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/m2, 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
75

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. For 

a specific neutron dose, RIS is greater at lower dose rate. 

Most austenitic SSs show a rapid decrease in Cr to about 13 wt.% and an increase in Si to 

about 4 wt.% at 5 dpa.

23 The data at dose levels above 10 dpa are limited, but they indicate 

that Cr content can decrease to 8–10 wt.% and Si content increase to 6 wt.% in SSs irradiated 

in LWRs at 300–320°C up to 65 dpa.

23 Under similar irradiation conditions, grain boundary Ni 

concentrations increase to about 22 wt.% at 5 dpa and can be as high as 30 wt.% at 65 dpa. 

The RIS behavior of minor elements such as P, C, N, and B, all of which segregate at the 

boundary, is not well established. 

3.1.3 Radiation Hardening

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

obstacles to 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 to dislocation motion.1 The yield strength of irradiated SSs can increase up to five 

times that of the non-irradiated material after a neutron dose of about 5 dpa.3 The yield and 

ultimate stresses increase and ductility decreases with irradiation. 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 greatest increase 

in yield strength for a given dose occurs at irradiation temperatures near 300°C (572°F). For a 

given neutron dose and irradiation temperature, the yield strength decreases with increasing 

test temperature.

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

material, there is a change in the deformation mode. Deformation by a planar slip mechanism is 

promoted, and the material exhibits strain softening.

114 This process can be explained by

“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 irradiation dose above 3–5 dpa.115,116 The enhanced planar 

slip also leads to a pronounced degradation in the fracture toughness of austenitic SSs.

3.1.4 Yield Strength

Three elements in combination influence IGSCC of austenitic SSs in BWR environments: a 

susceptible material, a significant tensile stress, and an aggressive environment. Low-C 

wrought austenitic SSs and weld metals are considered adequately resistant to sensitization by 

welding126 and therefore are recommended to reduce the susceptibility to IGSCC. However, 

Andresen has shown that nonsensitized SSs are not immune to SCC; SSs with high yield 

strength are susceptible to IGSCC in ultra-high purity water.117 The CGRs of austenitic SSs

increase with increasing yield strength. The effect of yield strength on the SCC growth rates of 

LWR irradiated austenitic SSs in NWC and HWC BWR environments at 289°C is shown in 

Fig. 43.

20,124 Increases in yield strength can originate from surface or bulk cold work, weld 

shrinkage strain, precipitation hardening, oxide dispersion hardening, or even irradiation 

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; however, the 
76

CGRs in low potential water are an order of magnitude lower.117 There is no significant 

difference among the different grades of SSs. The yield strength of irradiated austenitic SSs 

can increase up to five times that of the nonirradiated material after a neutron dose of about 

5 dpa.29 For austenitic SSs, the greatest increase in yield strength for a given irradiation level 

occurs at irradiation temperatures near 300°C (572°F). 

Figure 43. Change in SCC growth rates at 289°C of LWR irradiated austenitic SSs as a 

function of yield strength in NWC and HWC BWR environments (Refs. 20, 124).

3.1.5 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 detrimental because at all relevant pH/potential conditions, Si oxidizes to SiO2, 

which is highly soluble in hot water. Unusually high SCC growth rates have been observed in 

cold worked (CW) SSs, containing 1.5 to 5 wt.% Si, in high-purity water with 2 ppm (by weight) 

DO.

125 The CGR in 15% CW Type 304L SS with 5% Si (no Mo, Nb, or Ti additions) at 220°C 

(yield strength about 500 MPa) was 2.2 × 10-9 m/s at a stress intensity factor (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 resulted in an initial decrease in CGR, from 2.2 to 1 × 10-9 m/s, but thereafter the 

CGR remained constant as the K decreased from 30 to 15 MPa m1/2 (Figs. 8 and 9 in Ref. 125). 

In contrast to the increased susceptibility of high-Si materials indicated by crack growth tests, 

SSRT test data on irradiated Type 304 and 316 SSs with 0.5–1.8 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.116 Because most 

commercial SSs contain about 0.5% Si, its effect on irradiation hardening is not likely to be 

significant. Available data suggest that high concentrations of Si at grain boundaries due to RIS 

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

1/2

289°C High Purity Water 

Open Symbols: NWC BWR 

Closed Symbols: HWC BWR 

Ooki et al., 2005

NUREG/CR-6960
77

could increase susceptibility to IASCC. The absence of γ’ silicide (Ni3Si) at the grain boundaries 

suggests that RIS of Ni and Si probably saturates at higher doses; however, there is no 

evidence that Si segregation saturates at high dose. The significance of Si segregation at grain 

boundaries on the SCC behavior of irradiated SSs is not clear, but Si segregation could be 

important. 

3.1.6 Stacking Fault Energy

The stacking fault energy (SFE) is an important parameter that determines the deformation 

mode. Alloys with Ni concentration >18 wt.% are highly 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 (with 

>20 wt.% Ni). The increased susceptibility to IASCC is attributed to absorption of dislocations 

at the grain boundaries, which 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. 

3.2 IASCC Crack Growth Rates

3.2.1 In BWR and PWR Environments

In an earlier study at Argonne,

21,29 the SCC CGR data for irradiated wrought and cast austenitic 

SSs as well as austenitic SS weld metal and weld HAZ materials in LWR environments were 

compiled and evaluated to define 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 were discussed. The SCC CGRs for various grades and heats of 

austenitic SSs irradiated from 5 × 1020 to 2.5 × 1022 n/cm2 (0.75 to 37.5 dpa) and tested in NWC 

and HWC BWR environments are shown in Figs. 44 and 45, respectively. The results are 

compared with the K versus CGR disposition curves proposed in the NRC report NUREG-0313 

for nonirradiated sensitized austenitic SSs in high-purity water126 and the curve proposed by 

EPRI for austenitic SS BWR core internal components.127 The NUREG-0313 disposition curve 

is expressed as

da/dt (m/s) = A1 (K)2.161, (47)

where the stress intensity factor, K, is in MPa m1/2, and the magnitude of A1 depends on the 

water chemistry. The value of A1 is 2.1 × 10-13 in water with 8 ppm DO and 7.0 × 10-14 in water 

with 0.2 ppm DO. It is smaller in low-potential HWC BWR and PWR primary water 

environments. The EPRI disposition curve for use in BWR core environments is expressed as 

da/dt (m/s) = A2 (K)2.5, (48)

where the constant A2 is 4.564 × 10-13 in NWC BWR and 1.512 × 10-13 in HWC BWR 

environments. The EPRI correlations are based on two datasets; General Electric (GE) and 

Japan Power Engineering and Inspection Corp. (JAPEIC) data for SSs irradiated to 4.0–

4.5 dpa,127 and the Halden reactor data for Type 304, 347, and 316NG SSs that were irradiated 

to 13.5, 2.25, and 1.35 dpa.128–132
78

 

Figure 44. SCC growth rates in NWC BWR environment on austenitic SSs 

irradiated to 0.75–37.5 dpa (Refs. 20, 26, 128–133). 

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 1.35 dpa

316L 3.0 dpa

304L 3.0 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

304L 12.9 dpa 288°C

Experimental CGR (m/s)

Stress Intensity K (MPa m1/2)

Material & Dose

(ANL data)

NUREG-0313

Curve (8 ppm DO)

Irradiated Stainless Steels 

BWR Environment

EPRI

NWC BWR

(a)

10 x NUREG-0313

Curve

(Studsvik data)

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 #2

347 SS 2.3-3.3 dpa #3

347 SS 0.7 dpa

347 SS 0.95 dpa

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

304 SS 4.25 dpa

Experimental CGR (m/s)

Stress Intensity K (MPa m1/2)

Material & Dose

(Halden data)

NUREG-0313

Curve (8 ppm DO)

Irradiated Stainless Steels 

BWR Environment

EPRI

NWC BWR

(b)

20 x NUREG-0313

Curve

(GE/JAPEIC data)
79

Figure 45. SCC growth rates in HWC BWR environment on austenitic SSs 

irradiated to 0.75–37.5 dpa (Refs. 20, 26, 128–133). 

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

304L 12.9 dpa 288°C

Experimental CGR (m/s)

Stress Intensity K (MPa m1/2)

Material & Dose

(ANL data)

NUREG-0313

Curve (0.2 ppm DO)

Irradiated Stainless Steels 

BWR Environment

EPRI

HWC BWR

(a)

20 x NUREG-0313

Curve

(Studsvik data)

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 SS 13.6-14.4 dpa 280°C

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

Experimental CGR (m/s)

Stress Intensity K (MPa m1/2)

Material & Dose

(Halden data)

NUREG-0313

Curve (0.2 ppm DO)

Irradiated Stainless Steels 

BWR Environment

EPRI

HWC BWR

(b)

(GE/JAPEIC data)

(Studsvik data)

20 x NUREG-0313

Curve
80

In the NWC BWR environment, the SCC CGRs of austenitic SSs irradiated to 5 dpa could be a 

factor of 10 higher and those irradiated to 10–12 dpa could be a factor of 25 higher than the 

NUREG-0313 disposition curve for nonirradiated SSs in 8 ppm water. The CGRs for SSs 

irradiated to 5–8 dpa (corresponding to a 60-year end-of-life neutron dose for a BWR) are a 

factor of 20 higher than the NUREG-0313 curve. However, the growth rates for highly irradiated 

materials show anomalous behavior. The CGRs for SSs irradiated in LWRs 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 materials in high-DO environment are not clear. 

In the HWC BWR environment, the SCC CGRs 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 

of the materials irradiated to 3–4 dpa are above the CGR disposition curve proposed by EPRI 

for the HWC BWR environment. The tests that did not show the benefit of reduced corrosion 

potential on growth rates are often screened out because their loading conditions only 

marginally exceeded the proposed K/size criterion based on an effective yield stress. However, 

as discussed earlier in NUREG/CR-7027, none of these tests violated the K/size criteria and 

should be considered as valid tests. 

Data on the SCC CGR for irradiated austenitic SSs in the PWR environment have been 

obtained at Halden129–131 on BWR irradiated materials, at Studsvik26 on BOR-60 irradiated 

materials, and from EPRI-sponsored studies on BWR and PWR irradiated materials. The 

materials were irradiated to 3.0–37.5 dpa at 288–340°C. The CGRs in the PWR environment 

show significant variability. For the same material and irradiation condition, the CGR increase 

with increasing temperature. The temperature dependence of SCC CGRs in irradiated and 

nonirradiated SSs yields activation energies between 60 and 150 kJ/mol. The data for CGRs of 

CW Type 316+Ti SS irradiated in a PWR yield an activation energy of 105 kJ/mol.26 The 

experimental SCC CGRs for various grades and heats of austenitic SSs in a PWR environment 

are shown in Fig. 46; the data on materials irradiated in fast reactors are not included. The 

CGRs obtained at different temperatures have been normalized to a temperature of 320°C 

using an activation energy of 100 kJ/mol. Normalization did not significantly change the data 

scatter; it changed only the magnitude of the CGRs. 

In PWR water at 320°C, the CGR data for SSs irradiated to 3 dpa are up to a factor of 12 above 

the NUREG-0313 curve for nonirradiated materials in high-purity water with 0.2 ppm DO. In 

addition, CGRs for SSs irradiated to 11–25 dpa are nearly two orders of magnitude above the 

same curve. As observed earlier for HWC BWR water, the CGRs in the PWR environment of 

SSs irradiated to 37.5 dpa are marginally above the NUREG-0313 curve for nonirradiated SSs 

in high-purity water with 0.2 ppm DO. The CGRs for SSs irradiated to 18 dpa show significant 

data scatter; CGRs obtained at K values between 7 and 9 MPa m1/2 vary more than an order of 

magnitude; most values are above the NUREG-0313 curve. The reason is not known for the 

high CGRs for some SSs irradiated to 11.4 dpa and relatively low CGRs in SSs irradiated to 18

or 37.5 dpa. Additional data on PWR materials irradiated at temperatures of 300–350°C are 

needed to better understand the IASCC susceptibility of austenitic SSs in PWR environments.

All wrought and cast austenitic SSs and their welds, irradiated below the 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 × 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 
81

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. 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. There are no data on CASS materials and austenitic SS weld and weld HAZ materials. 

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.

In addition, 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. Furthermore, 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. 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. 

3.2.2 Update of the CGR Data

The SCC CGRs for irradiated and nonirradiated Types 304 or 304L austenitic SS, SA or SMA 

weld HAZ materials, tested at 289°C in high-purity water with 600 ppb or <50 ppb DO are 

shown in Fig. 47.

20,124,134,135 The materials were irradiated between 5 × 1020 and 

Figure 46. Normalized SCC growth rates in PWR environment for austenitic 

SSs irradiated to 0.75–37.5 dpa (Refs. 20, 26, 131). 

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 305-320°C

304L 3.0 dpa 305-320°C

304L 11.4 dpa 305°C

304L 11.4 dpa 305°C

304 37.7 dpa 335°C

304 18.7 dpa 335°C

304L 14.0 dpa 335°C

Experimental CGR (m/s)

304 6.3 dpa 340°C

Stress Intensity K (MPa m1/2)

Material & Dose

(Halden data)

NUREG-0313

Curve (0.2ppm DO)

Irradiated Stainless Steels 

PWR Environment 

EPRI

HWC BWR

Normalized to 320°C

100 x 

NUREG-0313 Curve

(GE data)
82

Figure 47. SCC growth rates for nonirradiated and irradiated austenitic SS (a) SA and 

(b) SMA weld HAZ materials at 289°C in high-purity water with 600 or 

<50 ppb DO (Refs. 20, 124, 134, 135).

10-12

10-11

10-10

10-9

10-8

10-7

5 10 15 20 25 30 35 40

AW+TT

AW+TT + 2.43 dpa

As Welded

AW+0.75 dpa

AW+2.16 dpa

Experimental CGR (m/s)

Stress Intensity K (MPa m1/2)

Condition & Dose

NUREG-0313

Curve (8ppm DO)

20 x NUREG-0313

Curve (8ppm DO)

Type 304L SA Weld HAZ at 289°C

Open Symbols: 600 ppb water

Closed Symbols: < 50 ppb water

All tests performed approx. at constant K

Unirr. materials with partial unloading every 1 h

NUR0.2G-0313

Curve (0.2ppm DO)

(a)

10-12

10-11

10-10

10-9

10-8

10-7

5 10 15 20 25 30 35 40

AW+TT

AW+TT 

 + 2.15 dpa

As Welded

AW+0.75 dpa

AW+2.16 dpa

Experimental CGR (m/s)

Stress Intensity K (MPa m1/2)

Condition & Dose

NUREG-0313

Curve (8ppm DO)

20 x NUREG-0313

Curve (8ppm DO)

Type 304 SMA Weld HAZ at 289°C

Open Symbols: 600 ppb water

Closed Symbols: < 50 ppb water

NUREG-0313

Curve (0.2ppm DO)

(b)
83

2.5 × 1022 n/cm2 (0.75 to 37.5 dpa). As before, the results are compared with the K versus

CGR disposition curves proposed in the NRC report NUREG-0313 for nonirradiated sensitized 

austenitic SSs in high-purity water with 8 ppm or 0.2 ppm DO.126 

The results indicate that in NWC BWR environment, the CGRs of Types 304 or 304L austenitic 

SS, SA, or SMA weld HAZ materials irradiated up to about 2.5 dpa are a factor of up to 20 

higher than those predicted by the NUREG-0313 curve for nonirradiated sensitized austenitic 

SSs in high-purity water with 8 ppm DO. The CGRs in HWC BWR environment are generally 

below the NUREG-0313 curve for nonirradiated sensitized austenitic SSs in high-purity water 

with 0.2 ppm DO. These results suggest that the SCC CGRs for SA or SMA weld HAZ 

materials in NWC BWR environments are comparable to those for wrought austenitic SSs. For 

these materials, the CGRs in HWC BWR environments are mostly below the 0.2 ppm DO 

NUREG-0313 curve. However, the data are too limited to draw a conclusion.

The SCC CGRs for thermally aged and unaged Type 316L weld metal and irradiated and 

nonirradiated CASS materials, tested at 289 or 320°C in high-purity water with 300–2,000 or 

<50 ppb DO or PWR environments are shown in Fig. 48.20,88,105,135,136 The materials were 

irradiated between 5 × 1020 and 2.5 × 1022 n/cm2 (0.75 to 37.5 dpa). As before, the results are 

compared with the K versus CGR disposition curves proposed in the NRC report NUREG-0313 

for nonirradiated sensitized austenitic SSs in high-purity water with 8 ppm or 0.2 ppm DO.126 

The results indicate that the SCC CGRs for thermally aged and unaged Type 316L weld metal 

at 289°C in high-purity water with 300–2,000 ppb DO are bounded by the NUREG-0313 curve 

for nonirradiated sensitized austenitic SSs in high-purity water with 8 ppm DO. All of the data 

are below this curve. However, similar data for irradiated Type 316L weld metal are not 

available. The results also indicate that the CGRs for CF-8M CASS material irradiated to about 

2.5 dpa and tested in high-DO water at 289°C are up to a factor of 6 higher than the NUREG0313 curve for 8 ppm DO environment. The data in high-purity water with <50 ppb DO at 320°C 

indicate that the SCC CGRs for thermally aged and unaged CF-8M CASS materials irradiated to 

0.08 dpa are all below the NUREG-0313 curve for nonirradiated sensitized austenitic SSs in 

high-purity water with 0.2 ppm DO. The limited data for CASS materials suggests that the 

0.2 ppm DO curve for NUREG-0313 can be used for the CGRs of CASS materials irradiated up 

to 0.08 dpa in HWC BWR environments. 

3.2.3 Effect of Neutron Fluence

The SCC growth rates for various grades and heats of austenitic SSs, irradiated from 

5 × 1020 n/cm2 to 2.5 × 1022 n/cm2 (0.75 to 37.5 dpa) in NWC BWR and HWC BWR 

environments at 288°C and K = 20 MPa m1/2 as a function of neutron dose, are shown in 

Figs. 49a and 49b, respectively.21 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. 

In the NWC BWR environment, the CGR increases above the 8 ppm DO NUREG-0313 curve

for neutron doses above 0.42 dpa. The CGR also increases with dose out to about 10 dpa, 

although there is variability in the data. In the HWC BWR environment, the CGRs do not 

exceed the 0.2 ppm DO NUREG-0313 curve for neutron doses above 2.2 dpa. The CGRs in 

HWC environment for materials irradiated 12–15 dpa are comparable to those observed in the 

NWC BWR environment. However, as noted previously, there are some anomalously data with 

very low growth rates; the CGRs of SSs irradiated to about 19 and 38 dpa are comparable to 

those for nonirradiated sensitized SSs in high-DO water. The curves in Fig. 49 bound 

approximately 75th percentile of the data in the NWC environment and even more for the HWC 
84

Figure 48. SCC growth rates for aged and unaged (a) Type 316L welds and (b) CASS 

materials in irradiated or nonirradiated conditions and tested at 289 or 320°C in 

high-purity water with 300–2,000 or <50 ppb DO (Refs. 20, 88, 105, 135, 136).

10-12

10-11

10-10

10-9

10-8

10-7

5 10 15 20 25 30 35 40

316L Weld Lu et al.

316L Weld 10% ferrite

316L Weld 14% ferrite

CF-8M Aged+2.46 dpa

CF-8M Aged+0.08 dpa

CF-8 Aged

Experimental CGR (m/s)

Stress Intensity K (MPa m1/2)

NUREG-0313

Curve (8ppm DO)

6 x NUREG-0313

Curve (8ppm DO)

316L Weld & CASS Materials at 289°C

Open Symbols: Aged 

Closed Symbols: Unaged

300 - 2000 ppb DO

PWR or <50 ppb DO

NUREG-0313

Curve (0.2ppm DO)

(a)

10-12

10-11

10-10

10-9

10-8

10-7

5 10 15 20 25 30 35 40

CF-3 10k@400°C

CF-8 10k@400°C

CF-8M 10k@400°C

Experimental CGR (m/s)

Stress Intensity K (MPa m1/2)

CASS Materials Irradiated to 0.08 dpa

Tested at 320°C in <50 ppb high-purity water

 using partial unloading every 1 or 2 h

Open Symbols: Aged 

Closed Symbols: Unaged

NUREG-0313

Curve (0.2ppm DO)

(b)
85

Figure 49. SCC growth rates of irradiated 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.

environment. The curves assume a threshold dose of 0.42 dpa and 2.2 dpa (or 2.8 × 1020 and 

approximately 1.47 × 1021 n/cm2) in the NWC and HWC environment, respectively. The curves 

assume that constant A1 in the NUREG-0313 disposition curve (Eq. 47) varies with neutron 

dose. In the NWC environment A1 is taken as 

A1 = 1.21 × 10–12 + 2.65 × 10–12 log(dpa), (49)

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 

Curve (8ppm DO)

(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 m 1/2

HWC BWR Water at 288°C

NUREG-0313 

Curve (0.2ppm DO)

(b)

CGR = -1.34x10-9

 + 4.0x10-09 log(dpa)
86

and in the HWC environment as 

A1 = -2.07 × 10–12 + 6.17 × 10–12 log(dpa). (50)

Equation 50 is an updated version of the expression published earlier in NUREG/CR-2027 and 

Ref. 29. In addition, note that because Eqs. 49 and 50 define the constant A1 in Eq. 47, these 

values must be multiplied by (20)2.161 to obtain the CGRs shown in Fig. 10.

The available SCC CGR data for austenitic SS welds and weld HAZ as well as CASS materials 

in NWC and HWC BWR environments are shown in Fig. 50.

20,105,124,134,135 The data plotted at 

0.05 dpa are for nonirradiated material. These limited data indicate that in NWC BWR 

environment, the CGRs for nonirradiated materials or materials irradiated up to 0.42 dpa are 

below the NUREG-0313 curve for nonirradiated sensitized austenitic SSs in high-purity water 

with 8 ppm DO. For materials irradiated above 0.42 dpa, most of the data are bounded by the 

curve represented by Eqs. 47 and 49. In HWC BWR environment, the CGRs for nonirradiated 

materials or materials irradiated up to 2.2 dpa are below the NUREG-0313 curve for 

nonirradiated sensitized austenitic SSs in high-purity water with 0.2 ppm DO. However, there is 

little or no IASCC CGR data in HWC BWR environment on CASS materials or austenitic SS 

welds and HAZ materials irradiated above 2.2 dpa. 

Figure 51a 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 HWC BWR water at 288°C 

or the PWR environment at 320°C, with K = 20 MPa m1/2(from Ref. 21). 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. 47) with neutron dose that is expressed by 

Eq. 50. 

A few SCC CGR results for aged CF-8M CASS material in the nonirradiated condition and 

unaged and aged CF-3 CASS material irradiated to 0.08 dpa in PWR environment are shown in 

Fig. 51b. 88,136 The results are below the NUREG-0313 curve for nonirradiated sensitized 

austenitic SSs in high-purity water with 0.2 ppm DO. 
87

Figure 50. SCC growth rates of nonirradiated and irradiated austenitic SS weld HAZs and 

welds and CASS materials 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. The data points plotted 

at 0.05 dpa are for nonirradiated materials. 

10-11

10-10

10-9

10-8

10-1 100 101

304 SMAW HAZ 

304L SAW HAZ 

304 Weld HAZ

316L Weld HAZ

316L Weld

CF8M Aged

316L GTAW 

Crack Growth Rate (m/s)

Neutron Dose (dpa)

Irradiated Austenitic SSs

Weld HAZ & Welds, & CASS

Normalized to 20 MPa m 1/2

NWC BWR Water at 288°C

CGR = 7.84x10-10

 + 1.72x10-09log(dpa)

(a)

Open Symbols: as welded or as cast

Closed Symbols: thermally treated or aged

NUREG-0313 

Curve (8ppm DO)

10-11

10-10

10-9

10-8

10-1 100 101

304 SMAW HAZ 

304L SAW HAZ 

CF-3

CF-8

CF8M

Crack Growth Rate (m/s)

Neutron Dose (dpa)

Irradiated Austenitic SSs

Weld HAZ & CASS Materials

HWC BWR at 288°C

CGR = -1.34x10-9

 + 4.0x10-09 log(dpa)

Normalized to 20 MPa m1/2

Open Symbols: Unaged

Closed Symbols: Aged

NUREG-0313 

Curve (0.2ppm DO)

(b)
88

(a)

(b)

Figure 51. Crack growth rates for irradiated (a) austenitic SSs and (b) CASS materials in HWC 

BWR water at 288°C or PWR environment at 320°C plotted as a function of 

neutron dose.

10-12

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 CGR = -1.34x10-9

 + 4.0x10-09 log(dpa)

Normalized to 20 MPa m1/2

Open Symbols: HWC BWR at 288°C

Closed Symbols: PWR Env. 320°C

NUREG-0313 

Curve (0.2ppm DO)

10-12

10-11

10-10

10-9

10-8

10-1 100 101

CF-3

CF-3 Aged

CF-8 Aged

Crack Growth Rate (m/s)

Neutron Dose (dpa)

Irradiated CASS Materials

PWR Environment at 320°C

CGR = -1.34x10-9

 + 4.0x10-09 log(dpa)

NUREG-0313 

Curve (0.2ppm DO)

Normalized to 20 MPa m1/2

Open Symbols: Unaged

Closed Symbols: Aged
89

4. NEUTRON EMBRITTLEMENT 

The fracture toughness of unaged and nonirradiated wrought and cast austenitic SSs and their 

welds fall in Category III, with JIc above 150 kJ/m2 (857 in.-lb/in.2). In these materials, fracture 

occurs after stable crack extension at stresses well above the yield stress. However, neutron 

irradiation can degrade fracture toughness of these materials to the level of Category II (JIc in 

the range of 30–150 kJ/m2 [171–857 in.-lb/in.2]), or Category I (JIc < 30 kJ/m2 [< 171 in.-lb/in.2]).

Therefore, failure in neutron-irradiated materials may occur without general yielding. In such 

instances, either EPFM or LEFM is used for analysis of structural integrity and development of 

inspection guidelines. The EPFM analysis is used where failure is caused by plastic 

deformation, and involves the J-R curve approach. 

As discussed earlier in Section 2, austenitic SS welds and CASS materials have a duplex 

microstructure consisting of austenite and ferrite phases. Embrittlement of the ferrite phase due 

to neutron irradiation occurs much faster than for austenitic SSs. The irradiation temperature is 

an important factor in establishing the extent of embrittlement of the ferrite. However, the ferrite 

phase is susceptible to thermal embrittlement after extended service at reactor operating 

temperatures. The formation of the Cr-rich α′ phase in the ferrite is the primary mechanism for 

thermal embrittlement; thermal aging has little or no effect on the austenite phase. Although 

both the thermal aging embrittlement of ferrite and the neutron irradiation embrittlement of ferrite 

are well characterized, experimental data on the combined effect of thermal aging and neutron 

irradiation on the embrittlement of ferrite in SS welds and CASS materials are rather limited. 

The possible combined effects of thermal aging and neutron irradiation embrittlement of these

materials are discussed in Section 5. 

4.1 Fracture Toughness of Irradiated Austenitic Stainless Steels

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. In these studies, the 

embrittlement of the materials has been characterized in terms of tensile properties, Charpyimpact 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 

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-mmthick 1⁄4-T CT) have been used in several studies. For these specimens, J values above 

150 kJ/m2 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 versus 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/m2, 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 
90

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. 51 is satisfied. The blunting line is defined as 

J = MσfΔa, (51)

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 over predicts the crack extension due to crack tip blunting, and, therefore, it 

yields a non-conservative value of JIc.

8,137 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). (52)

The use of a higher value for M in Eq. 51 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.8 Thus, 

for such materials, the crack tip blunting line is given by

J = Mσy(CTOD) ≈ 2(2σf/3)(3Δa) = 4σfΔa (53)

(i.e., Eq. 20 with M = 4). This relationship has been used to determine JIc in most investigations 

on neutron embrittlement.

12,20 A value of 2 for M has also been used by some investigators.10 

The latter typically yields a higher value of JIc for Category III materials (i.e., with JIc above 

150 kJ/m2). 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/m2). 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 the use of an effective yield stress 

instead of the measured yield stress. 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.

20 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. 
91

4.1.1 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. 52.

8,14,15,138–151 The irradiation and test temperatures were 100–155°C and 125°C for the 

data obtained by Sindelar et al.,146 90–250°C and 25–250°C for the data obtained by Alexander 

et al.,145 and 325°C and 25°C for the data obtained by Kim et al.15 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 

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]). In addition, the failure mode changes from dimple fracture to channel fracture. 

Figure 52. 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. 8, 14, 15, 138–151).

The fracture toughness trend for the LWR data6,9–12,16–20 is similar to that observed for fast 

reactor data (Fig. 53). 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 

CASS material and austenitic SS 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 

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 (Kim 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
92

reactor data in Fig. 52. The JIc values of austenitic SS welds and HAZ materials are 

consistently lower than those for the solution-annealed and CW materials. The data for CF-8 

steel are relatively high because they were obtained at RT; 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 (indicated by crosses in Fig. 53), nine out of 10 CT 

specimens showed no ductile crack extension and the KIc values were 52.5–67.5 MPa m1/2

(47.7–61.4 ksi in.1/2).9 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 in.1/2), was for a Type 347 SS irradiated to 16.5 dpa in a 

PWR9 and for a Type 304 SS irradiated to 7.4–8.4 dpa in a BWR.18 

Figure 53. Change in fracture toughness JIc as a function of neutron exposure for LWR 

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. 6,

9–12, 16–20). JAPEIC = Japan Power Electric Engineering and Inspection 

Corp.; BB = bend bar specimen; and SR = short rod specimen.

Another significant result is a strong orientation effect on fracture toughness. Fracture 

toughness J-R tests and microstructural and microchemistry characterization were performed 

on a Type 304 control-rod and Type 304L top guide materials irradiated to 4.7–12.0 dpa in a 

BWR, and on Type 304 control-rod material irradiated to 7.4 and 8.4 dpa in another BWR. All 

0

100

200

300

400

500

600

0 5 10 15 20

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/m2

Closed Symbols: BWR Water

Open Symbols: Air
93

materials consistently showed lower fracture toughness in the T-L orientation than in the L-T 

orientation,10 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 × 1021 n/cm2).10,13 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 quasicleavage structure parallel to the crack advance, thereby accelerating the crack advance.10 

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.10 The density was 1–3 × 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.

152 The contribution of additional 

precipitate phases, voids, and cavities on fracture toughness needs to be investigated. 

4.1.2 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,9–13,16–20 and in 

fast reactors to much higher dose levels. The change in the J-R curve with neutron dose is 

shown in Fig. 54. The decrease in fracture toughness is quite rapid up to about 6 dpa, and the 

toughness continues to decrease 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. 52). 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 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. In 

addition, most of the data for SS welds are on Type 308L material irradiated to <1 dpa or 

12 dpa, and a few tests on 316 weld and one test on 316LN weld. Similarly, there are only a 
94

few J-R curve tests on weld HAZ materials and CF-8M steel 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

Figure 54. Change in fracture toughness J-R curve and JIc as a function of neutron 

exposure for Type 304 SS at (a) 289 and (b) 320°C (Refs. 18,20). 

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 = 575Δa0.17

JIC = 503 kJ/m2

J = 438Δa0.33

JIC = 308 kJ/m2

J = 265Δa0.29

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 = 85.7Δa0.19

JIC = 65 kJ/m2

J = 651Δa0.45

JIC = 744 kJ/m2

J = 143Δa0.21

JIC = 109 kJ/m2

(b)
95

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. 53, respectively). In addition, 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. 53, respectively), and the toughness of austenitic SS 

welds and thermally aged CASS materials is lower than that of the HAZ material (Fig. 55). 

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 JIc value at saturation is higher than those of irradiated solution-annealed steels 

(Fig. 52). 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 

between the two are likely to be small. 

Nonirradiated weld metals and thermally aged CASS materials 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 wrought materials as well as austenitic welds 

and CASS materials. Although LWR core internals are typically constructed of CF-8 or CF-3 

steels, the only data for LWR irradiation of CASS materials are for CF-8M steel. The data for 

thermally aged CF-8 steel shown in Fig. 52 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 CASS materials,

Figure 55. Fracture toughness J-R curves for sensitized Type 304 SS, weld HAZ materials of 

Type 304 and 304L SS, and CF-8M steel in high-purity water at 289°C (Ref. 21).

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
96

the fracture toughness of CF-8M steel represents the worst-case scenario.50,53 It thus might 

also represent a bounding case also for the combined 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 RT air. However, in reactor core components, cracks are initiated primarily by SCC 

and have intergranular (IG) morphology, whereas the fatigue precracks in fracture toughness 

tests are always transgranular (TG). In addition, 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. As discussed earlier in Section 2.5.3 for CF-8M material aged for more than 

15 years at 350°C, the J-R curve data for specimens precracked in air at 54°C and then tested 

in PWR (SWC) at 54°C were lower relative to those tested in air. The specimens that were 

precracked in PWR water at 315°C and then tested in PWR water also show reduced fracture 

toughness but the effect was less.88 Similarly, J-R curve data for Type 316L GTA weld 

indicated that the reactor coolant environment could decrease fracture toughness by up to 40% 

relative to that in air (Section 2.6.1).105

To investigate potential effects of reactor coolant environment on the fracture toughness of 

irradiated austenitic SSs, J-R curve tests were conducted in a BWR NWC environment.20 The 

J-R data on irradiated SS weld HAZ materials (Figs. 56a and 56b) 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 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. 

Figure 57 shows the J-R curves for a sensitized Type 304 SS in air and water environments. 

The results indicate slightly lower fracture toughness in water. In addition, 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 might be greater than that indicated by Fig. 57. Thus, the lower J-R curve in 

water may be just data scatter or due to the effect of reactor coolant environment. Although 

these results indicate little or no effect of environment, the data presented earlier in Fig. 29, 

indicate a significant decrease in fracture toughness J-R curve of aged CF-8 material in PWR 

SWC environment at 54°C. Duplicate tests under identical material and test conditions except 

the environment showed significantly lower fracture toughness in the coolant environment. 

The J-R curves and the load versus load-line displacement curves for the two tests on thermally 

aged and irradiated CF-8M steel in NWC BWR water are shown in Fig. 58. 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.50 Fracture toughness J-R curve tests have also been conducted recently at 

Argonne on unaged and aged CF-3 (Heat 69) and CF-8 (Heat 68) materials in air and low 

corrosion potential high–purity water.136 The tests in air were performed on 1-T CT specimens 

and in water on 1⁄4-T CT specimens. The results showed lower fracture toughness of the 

unaged material in water environment compared to that in air. The fracture toughness J-R 

curves for these materials aged 10,000 h at 400°C showed little or no environmental effects; the 

J-R curves in in air and water environments were comparable. However, since 1⁄4-T CT 

specimens were used for the tests in water, considering size effects, the J-R curves for tests in 

water on a 1-T CT specimen are likely to be lower than those observed for a 1⁄4-T CT specimen. 
97

Figure 56a. Fracture toughness J-R curves for irradiated specimens of Type 304 SS SMA weld 

HAZ in air and NWC BWR environment (Ref. 21).

Figure 56b. Fracture toughness J-R curves for irradiated specimens of Type 304L SA weld HAZ 

in air and NWC BWR environment (Ref. 21).

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 1025 n/m2

289°C

Estimated Effective 

Flow Stress: 528 MPa

J = 219Δa0.43

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 1025 n/m2

289°C Air

Estimated Effective 

Flow Stress: 502 MPa

(b)
98

Figure 57. Fracture toughness J-R curves for sensitized Type 304 SS in air and NWC BWR 

water at 289°C (Ref. 21).

Thus, although available fracture toughness data in reactor coolant environments are 

inconclusive, the limited data suggest possible environmental effects on fracture toughness of 

CASS materials and austenitic SS welds, at least for materials with moderate or superior 

fracture toughness. Additional tests on irradiated CASS materials 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 temperature of ≈300°C (≈572°F).114 Thus, the JIc values for all materials 

irradiated above 350°C (662°F) (e.g., fast reactor irradiations), particularly for neutron 

exposures greater than 20 dpa, would be greater than for materials irradiated at temperatures of 

290–320°C (554–608°F). 

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. 59 for several grades of SSs and welds irradiated in LWRs 

and fast reactors. The fracture toughness of steels irradiated to relatively low doses (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 

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 1025 n/m2

290°C 

J = 316Δa0.45

JIC = 176 kJ/m2

Open Symbols: Water

Closed Symbols: Air

Heat Treatment

J = 376Δa0.38

JIC = 238 kJ/m2
99

 

Figure 58. Fracture toughness J-R curves and load vs. loadline displacement curves for 

thermally aged and irradiated CF-8M steel (Ref. 21). 

this fluence level the toughness value is already low, which makes it difficult to discern definitive 

trends. 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 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 

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

Irradiated to 2.46 dpa

289°C BWR Water

Specimen No.
100

along the narrow band clears away the irradiation-induced defect structure, creating a defectfree channel that offers less resistance to subsequent dislocation motion. The localization of the 

deformation ultimately leads to channel failure. 

4.2 Fracture Toughness Trend Curve

The change in initiation toughness JIc of wrought austenitic SSs (including weld HAZ materials 

and sensitized SSs), CASS materials, and weld metals is shown in Fig. 60 as a function of 

neutron dose. 20 The fracture toughness data from both fast reactor and LWR irradiations are 

included. The irradiation temperatures range from 90 to 427°C (194 to 800°F) and test 

temperatures generally from 250 to 427°C (212 to 800°F). Only the data for CF-8 steel

irradiated in a fast reactor to 10–11 dpa (inverted triangles in the Fig. 60b) were obtained at 

room temperature. The data in Fig. 59 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. In addition, as discussed previously in Section 3.1, the procedures for 

determining JIc vary 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. 60a. The 

threshold dose and 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 2.2 dpa. The two data 

points at 2.25 dpa with very low JIc values are for instrumentation dry tube plunger material from 

an operating BWR tested using CNSR specimens.6,9 Several data points on Type 304 SS 

Figure 59.

Fracture toughness JIc of 

irradiated austenitic stainless 

steels and welds as a function 

of test temperature (Ref. 20). 

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
101

irradiated to 2.25 and 3 dpa are below the lower bound curve for JIc values. In addition, for 

Type 316L SS the decrease in fracture toughness appears to occur at a somewhat higher 

fluence range (3 dpa to 10 dpa), and for Type 304L SS it occurs at a somewhat lower fluence. 

Therefore, the decrease in toughness with increasing fluence occurs earliest in Type 304L SS, 

followed by Type 304 SS and then Type 316 SS. The results indicate that for Type 304 SS, 

either the threshold for the decrease in fracture toughness is earlier or the decrease occurs 

faster.

Figure 60. The change in initiation toughness JIc of (a) wrought austenitic SSs and (b) CASS 

materials and weld metals as a function of neutron exposure. The data points 

plotted at 0.005 dpa are for nonirradiated materials (Ref. 21).

0

50

100

150

200

250

300

350

400

0.01 0.1 1 10 100

304

304

304L

304CW

316

316L

316CW

347

348

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: LWRs

Closed Symbols: Fast reactors

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)
102

The fracture toughness data in Fig. 60b for CASS materials and SS 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 CASS materials are extremely 

limited. There are no data on any of these materials for fluences above 20 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. The existing data for CASS 

materials are inadequate to define a threshold dose for irradiation effects on fracture toughness;

there are no fracture toughness data for dose levels of 0.1–2.0 dpa. Furthermore, the potential 

combined interactions of thermal aging and neutron irradiation embrittlement of CASS materials 

and SS welds50–49 have not been addressed. 

The two curves shown in Fig. 60 represent a disposition curve proposed by EPRI10 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. 10, 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]. (54)

This lower bound curve represents a threshold dose of about 0.3 dpa for neutron embrittlement, 

a minimum fracture toughness JIc of ≈118 kJ/m2 for neutron doses below the threshold value, 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 for 

neutron doses below the threshold dose is appropriate for thermally aged and unaged CASS 

materials and SS flux welds. A value higher than 118 kJ/m2 may be 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. 54 is consistent with the 

Materials Reliability Program (MRP) lower-bound model proposed for PWRs.16 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)]. (55) 

Equations. 54 and 55 both predict a saturation fracture toughness KIc of 38 MPa m1/2. For 

materials irradiated below the threshold dose for irradiation embrittlement, Eq. 54 predicts a 

minimum KJc of about 151 MPa m1/2. However, for materials such as SS welds or weld HAZ, 

the fracture toughness values predicted by the MRP expression may be higher than the 

minimum toughness of the nonirradiated materials. Therefore, the disposition curve proposed 

by EPRI 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 CASS materials with >15% ferrite.50 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.16 The saturation KIc for the EPRI curve was based on data for which the 
103

specimen orientation was unknown. Recent data indicate that fracture toughness in the 

transverse orientation is nearly half of that in the longitudinal orientation.10 Therefore, the 

bounding KIc values above 4.5 dpa are likely to be lower than 55 MPa m1/2. Furthermore, as 

seen in Fig. 60b, some of the data for SS welds irradiated to 2–4 dpa are also below the EPRI 

curve. 

The existing fracture toughness JIc data for irradiated austenitic SS welds, CASS materials, and 

austenitic SS HAZ, including some recent data published after the plot in Fig. 60 was prepared,

are plotted as a function of neutron dose in Fig. 61.

9,15,16,20,53,88,105,134,138 142,148 Although most 

of the data were obtained in air, some of the fracture toughness J-R curve tests were performed 

in LWR coolant environments. As discussed before, the fracture toughness obtained from tests 

in LWR environments may be lower than that from tests in air; particularly for materials with high 

fracture toughness. The plots in Fig. 61 indicate that the lower-bound curve for JIc values may 

not be adequate for austenitic SS welds; some of the data for both irradiated and nonirradiated

welds falls below the lower-bound curve. As discussed earlier in Section 2.6, the lower-bound 

fracture toughness of unaged and aged austenitic SS welds are given by Eqs. 45 and 46, 

respectively. The fracture toughness JIc for unaged and aged SS welds are better represented 

by 73.4 and 40 kJ/m2, respectively (the lower-bound JIc for unaged welds is shown by dashed 

line in Fig. 61). Furthermore, some of the data obtained in LWR environments for austenitic SS 

HAZ, welds, and CASS materials also falls below the lower-bound curve. The data for aged 

CF-8 material in PWR SWC environment at 54°C are significantly lower. Therefore, based on 

these results the minimum acceptable fracture toughness JIc values for nonirradiated wrought 

and CASS materials and their welds need to be established. 

Figure 61. Plots of fracture toughness JIc values as a function of neutron dose for 

(a) austenitic SS welds, (b) CASS materials, and (c) austenitic SS HAZ. Solid line 

represents the lower-bound JIc values proposed in NUREG/CR-7027. The data 

points plotted at 0.007 dpa are for nonirradiated materials. 

0

100

200

300

400

500

600

0.01 0.1 1 10

Haggag 1990 288/120

Sindelar 1993 100-155/125

MRP-79 280/150-259

Haggag 1990 288/288

NUREG/CR-6428 -/290

Chen ANL Unpublished 315/320

Michel&Gray 1983 427/427

Picker et al 1983 370/370

Lucas 10% ferrite

Lucas 14% ferrite JIc (kJ/m2)

Neutron Exposure (dpa)

Austenitic SS Weld Metals

Type 316/316L Weld

Ref. & Irr./Test temperature °C

 Type 308/309 Welds

EPRI Curve

Lower bound (Demma et al 2007)

Unaged Welds

Open Symbols: Air

Closed Symbols: LWR Water
104

Figure 61. Continued.

0

100

200

300

400

500

600

0.01 0.1 1 10

CF-3 PWR Fuel Nozzle

CF-3 Ht 69 Unaged

CF-3 Ht 69 10k @ 400°C

CF-8 Ht 68 Unaged

CF-8 Ht 68 10k @ 400°C

CF-8M Ht 75 Unaged

CF-8M 10k @ 400°C

CF-8 Unaged

54°C Air / 54°C PWR

315°C PWR / 315°C PWR

315°C PWR / 54°C PWR

Neutron Exposure (dpa)

CASS Materials

Irradiated 315-320°C

Specimens precracked & tested in

Closed Symbols: PWR Water at 320°C

Open Symbols: Air at 290°C

Irr./tested 

at 427°C }

Lower bound Irr./tested at 290°C 

in BWR water

EPRI Curve

(Demma et al 2007)

CF-8 Aged 134.8k @ 350°C

Precracked/Tested

JIc (kJ/m2)

0

100

200

300

400

500

600

0.01 0.1 1 10

304 HAZ, Sindelar et al. 1993, C-L

304 HAZ, Sindelar et al. 1993, L-C

304 HAZ, Fyfitch

304 HAZ, ANL 

304L HAZ, ANL

JIc (kJ/m2)

Neutron Exposure (dpa)

Austenitic SSs HAZ Materials

Closed Symbols: BWR Water

Open Symbols: Air

Irradiated & tested at 288°C

Irradiated 100-150°C, tested 125°C

JIc = 7.5 + 110exp[-0.35(dpa)1.4]

EPRI Curve 

(Demma et al 2007)
105

Figure 62. Change in coefficient C of the power-law J-R curve for (a) wrought austenitic SSs 

and (b) CASS materials and weld metals as a function of neutron exposure. The 

data points plotted at 0.005 dpa are for nonirradiated materials (Ref. 21).

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 variation of C for wrought and CASS materials and welds as a 

function of neutron dose is shown in Fig. 62.

21 Except for the results for CF-3 (green right-angle 

triangle in Fig. 62b) and CF-8 irradiated in BOR-60 reactor (filled inverted triangle in Fig. 62b), 

0

100

200

300

400

500

600

700

0.01 0.1 1 10 100

304

304L

304CW

316

316L

316CW

348

Power-law Constant C

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: LWRs

Closed Symbols: Fast reactors

EPRI Curve 

(Demma et al. 2007)

(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

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)
106

the remaining data were obtained at temperatures above 250°C. Based on the data trends in 

Fig. 59, test temperature should have little or no effect on the data for CF-8 steel. However, the 

constant C for CF-3 steel may be more than a factor of two lower at LWR operating 

temperatures. The two curves in Fig. 62 represent the disposition curve proposed by EPRI for 

BWRs,13 and a trend curve for coefficient C that bounds the existing data.

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

CASS materials 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. 62, the existing 

fracture toughness data can be bounded by the following expression for C:

C = 25 + 175 exp[–0.35(dpa)1.4], (56)

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,50 no obvious trend of n 

with fluence is evident. For irradiated materials, the median value of n is 0.37. 

Equation 56 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. 60. 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 

diamonds in Fig. 60a).

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). 

The existing data for fracture toughness power-law constant C, including some recent data 

published after the plot in Fig. 62 was prepared, for irradiated austenitic SS welds, CASS 

materials, and austenitic SS HAZ, are plotted as a function of neutron dose in Fig. 63. The data

trends for power-law constant C are similar to those observed for fracture toughness JIc values. 

As discussed earlier, for some materials, the fracture toughness in LWR environments may be

lower than in air. The plots in Fig. 63 indicate that the lower-bound curve for constant C based 

on Eq. 56 may not be adequate for austenitic SS welds and CASS materials. Therefore, the 

minimum acceptable fracture toughness J-R curves for irradiated and even nonirradiated

wrought and CASS materials and their welds need to be established.
107

Figure 63. Plots of constant C of the power-law J-R curve as a function of neutron dose for 

(a) austenitic SS welds, (b) CASS materials, and (c) austenitic SS HAZ. The solid 

line represents the lower-bound C values proposed in NUREG/CR-7027. The data 

points plotted at 0.007 dpa are for nonirradiated materials.

0

200

400

600

800

0.01 0.1 1 10

O'Donnell 1991 370/370

Sindelar 1993 100-155/125

MRP-79 280/150-259

Haggag 1990 288/288

NUREG/CR-6428 -/290

Chen ANL Unpublished 315/320

Mills 1988 427/427

Picker et al 1983 370/370

Power-law Constant C

Neutron Exposure (dpa)

SS Weld Metals

Type 316 Weld

Ref. & Irr./Test temperature °C

 Type 308/309 Welds

EPRI Curve

(Carter & Gamble 2002)

C = 25 + 175exp[-0.35(dpa)1.4]

0

200

400

600

800

0.01 0.1 1 10

CF-3 PWR Fuel Nozzle

CF-3 Ht 69 Unaged

CF-3 Ht 69 10k @ 400°C

CF-8 Ht 68 Unaged

CF-8 Ht 68 10k @ 400°C

CF-8M Ht 75 Unaged

CF-8M 10k @ 400°C

CF-8 Unaged

54°C Air / 54°C PWR

315°C PWR / 315°C PWR

315°C PWR / 54°C PWR

Power-law Constant C

Neutron Exposure (dpa)

CASS Materials 

Irradiated 315-320°C

Specimens precracked & tested in

Closed Symbols: PWR Water at 320°C

Open Symbols: Air at 290°C

Irr./tested 

at 427°C

Lower bound

Irr./tested at 290°C 

in BWR water

EPRI Curve

(Carter & Gamble 2002)

CF-8 Aged 134.8k @ 350°C

Precracked/Tested
108

Figure 63. Continued.

0

200

400

600

800

0.01 0.1 1 10

304 HAZ, Sindelar et al. 1993, C-L

304 HAZ, Sindelar et al. 1993, L-C

304 HAZ, Fyfitch

304 HAZ, ANL 

304L HAZ, ANL 

Power-law Constant C

Neutron Exposure (dpa)

Austenitic SSs HAZ Materials

Closed Symbols: BWR Water

Open Symbols: Air

Irradiated & tested at 288°C

Irradiated 100-150°C, tested 125°C

C = 25 + 175exp[-0.35(dpa)1.4]

EPRI Curve 

(Carter & Gamble 2002)
109

5. COMBINED EFFECTS OF THERMAL AND NEUTRON 

EMBRITTLEMENT 

It is well known that thermal aging of CASS materials and austenitic SS welds at reactor 

operating temperatures leads to degradation of their fracture properties.

43,50–53 Typically 

wrought SSs are completely austenitic. However, austenitic SS welds and CASS materials 

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 

CASS materials at 250–400°C (482–752°F) leads to precipitation of additional phases in the 

ferrite.60,64,66 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 CASS materials and SS welds leads to the 

development of a material with a brittle phase dispersed in a ductile matrix. 

The minimum fracture toughness that can occur due to thermal aging embrittlement of CASS 

materials depends strongly on the chemical composition of the steel, the casting process used 

to construct the component, and the ferrite content and morphology.50 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 10% ferrite, where a 

continuous fracture path through the embrittled ferrite is possible. The minimum toughness 

occurs when the ferrite is fully embrittled, and the remaining toughness depends on the 

toughness provided by the ductile matrix surrounding the embrittled phase. Furthermore, CASS 

materials 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. Based on this update 

of the previous study,50 the predicted coefficient C and exponent n for the saturation fracture 

toughness J-R curves (i.e., the lowest fracture toughness that could be achieved for the steel 

after thermal aging) for thermally-aged CASS materials with <10%, 10–15%, 15–25%, and 25–

40% ferrite were given previously in Table 3 in Section 2.5.2. 

The results for austenitic SS welds indicate that the decrease in fracture toughness due to 

thermal aging also depends on the ferrite content and initial toughness of the weld.

53 

Differences in the fracture toughness of austenitic SS welds arise from differences in the density 

and size of inclusions in the material. In austenitic SS welds, failure typically occurs by the 

formation and growth of microvoids near hard inclusions. Welds with relatively high fracture 

toughness (e.g., GTA or tungsten inert gas [TIG] weld) show a significant decrease due to 

thermal aging, whereas welds with poor fracture toughness (e.g., SA, SMA, or manual metal arc 

[MMA] welds) show minimal change. In the latter case, 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. As shown in Fig. 35, a conservative estimate of the J-R curve for aged 

austenitic SS welds53 is given by J = 40 + 83.5 Δa0.643.

However, reactor core internal components are subjected to both elevated temperatures and

neutron irradiation. The mechanisms of thermal embrittlement and neutron embrittlement have 

been discussed in detail in Sections 2 and 4, respectively. Although both the thermal aging 

embrittlement and the neutron irradiation embrittlement of ferrite have been well characterized, 

an issue that has been a concern for reactor core internal components is the possibility of a 

combined interaction between thermal embrittlement and irradiation embrittlement of materials 
110

with duplex microstructure such as CASS materials and austenitic SS welds. The concurrent 

exposure to elevated temperatures and high neutron fluence levels could result in a combined

effect wherein the service-degraded fracture toughness is reduced from the levels predicted 

independently for either of the two mechanisms.

The screening criteria for susceptibility of CASS materials to thermal embrittlement and the 

methodology for estimating thermal embrittlement of CASS materials are presented in 

Sections 2.5.4 and 2.5.5, respectively. An assessment of thermal embrittlement of austenitic 

SS welds is described in Section 2.6. The effect of neutron irradiation on the fracture toughness 

of wrought and CASS materials and austenitic SS welds is discussed in Section 4, and the 

change in fracture toughness JIc and coefficient C of the power-law J-R curve with neutron dose 

is shown in Figs. 60 and 62, respectively. 

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 an aging management program (AMP). The program is implemented for CASS 

components that either have (a) a fluence of greater than 1 × 1017 n/cm2 (0.00015 dpa) or 

(b) are determined to be susceptible to thermal embrittlement. The AMP should consist of 

either a supplemental examination of the affected components as part of the applicant’s 10-year 

ISI program during the license renewal term, or a component-specific evaluation to determine 

the susceptibility to loss of fracture toughness.45 Furthermore, the program should provide for 

the consideration of the combined loss of fracture toughness due to neutron irradiation 

embrittlement and thermal aging embrittlement. The screening criteria to determine the 

susceptibility of CASS components to thermal aging embrittlement are outlined in Table 4.45

In NUREG/CR-7027, for the cast and wrought austenitic SSs and welds listed in Fig. 62, the 

experimental J-integral values at a crack extension of 2.5 mm are plotted as a function of 

neutron dose. The plots are reproduced in Fig. 64. The solid curve in Fig. 64 represents the 

predicted values of J2.5 at 2.5-mm crack extension that are expected to bound the existing 

experimental data shown in Fig. 62. The curve was obtained using the power-law J-R curve 

relationship, coefficient C determined from Eq. 56 and exponent n of 0.37. The lower bound 

curve indicates that for CASS materials and austenitic SS welds irradiated up to 1.0 dpa, the 

predicted J2.5 values are above the screening value of 255 kJ/m2 (1456 in.-lb/in.2) for all 

materials except some of the data for Type 316 welds are below the lower-bound curve. 

In addition to possibly altering the threshold dose for neutron embrittlement, the combined 

effects of neutron and thermal aging embrittlement could decrease the saturation fracture 

toughness of irradiated austenitic SS welded and CASS materials and accelerate the change in 

fracture toughness between the threshold and saturation neutron exposures. 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 (right-angle triangles 

in Fig. 64b). The resulting J2.5 values for these two tests are bounded by the curve for other 

SSs irradiated to a similar level, i.e., neutron irradiation of thermal aged material does not seem 

to lower the toughness below that expected for irradiation alone at these neutron dose levels. 

However, at the time NUREG/CR-7027 was published, fracture toughness data on irradiated SS 

welds and CASS materials, particularly at 0.1–2.0 dpa, were inadequate to accurately define the 

threshold dose beyond which the fracture toughness of austenitic SSs begins to significantly 

decrease. Furthermore, the evaluation in NUREG/CR-7027 did not consider the combined

interaction of neutron irradiation and thermal aging embrittlement. Embrittlement of the ferrite 

phase from neutron irradiation occurs at lower dose levels than does embrittlement of the 
111

Figure 64. Experimental values of J-integral at a crack extension of 2.5 mm for (a) wrought 

austenitic SSs and (b) cast materials and SS weld metals plotted as a function of 

neutron dose. The data points plotted at 0.005 dpa are for nonirradiated materials.

austenite phase. The NRC staff has proposed that CASS materials neutron dose levels greater 

than 1017 n/cm2 (0.00015 dpa) be considered susceptible to neutron embrittlement.45 A shift in 

the nil-ductility transition temperature of up to 150°C (302°F) has been observed in pressure 

vessel steels irradiated to 0.07–0.15 dpa.153 Thus, embrittlement of ferrite is expected to occur 

at 0.05–1.0 dpa, whereas any significant neutron embrittlement of the austenite phase occurs 

only above ≈0.5 dpa (Fig. 60). 

The available data on fracture toughness J2.5 for austenitic SS welds, CASS materials, and SS 

HAZ materials are plotted as a function of neutron dose in Figs. 65a, 65b, and 65c, respectively. 

0

200

400

600

800

1000

0.01 0.1 1 10 100

304

348

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: 250–427°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: 250–427°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)
112

The solid curve in these figures represent the J-integral values based on the lower bound trend

curve for coefficient C given in Fig. 62 and median value of 0.37 for exponent n of the J-R curve. 

The open symbols in the figure represent the J2.5 values in air and solid symbols represent the 

values in LWR coolant environment. The results indicate that the J2.5 values of some of the 

austenitic SS welds, particularly the values for Type 316 weld, may not be bounded by the lower 

bound trend curve. 

The limited fracture toughness data obtained from tests in LWR coolant environments indicate 

that the LWR coolant environment can further decrease fracture toughness of these materials 

relative to that obtained in air. For some heats of wrought and cast austenitic SSs and welds, 

some values of fracture toughness in coolant environment are likely to be below the lowerbound trend curve. The two tests on thermally aged CF-8 material tested in PWR environment 

at 54°C are significantly below the curve. Additional fracture toughness data in LWR coolant 

environments are needed to define the lower bound fracture toughness of these materials in 

LWR environments.

0

200

400

600

800

1000

0.01 0.1 1 10

Westinghouse 2006 330/25

Sindelar 1993 100-155/125

MRP-79 280/150-259

O'Donnell 1991 370/370

NUREG/CR-6428 -/290

Chen ANL Unpublished 315/320

Michel&Gray 1983 427/427

Picker et al 1983 370/370

J at 2.5 mm Crack Extension (kJ/m2)

Neutron Exposure (dpa)

Austenitic SS Weld Metals

Type 316 Weld

Ref. & Irr./Test temperature °C

 Type 308/309 Welds

Lower bound

J = 255 kJ/m2

Open Symbols: Air

Closed Symbols: LWR Water
113

Figure 65. Plots of fracture toughness J2.5 values as a function of neutron dose for 

(a) austenitic SS welds, (b) CASS materials, and (c) austenitic SS HAZ. Solid line 

represents the lower-bound J2.5 values proposed in NUREG/CR-7027. The data 

points plotted at 0.007 dpa are for nonirradiated materials.

0

200

400

600

800

1000

0.01 0.1 1 10

CF-3 PWR Fuel Nozzle

CF-3 Ht 69 Unaged

CF-3 Ht 69 10k @ 400°C

CF-8 Ht 68 Unaged

CF-8 Ht 68 10k @ 400°C

CF-8M Ht 75 Unaged

CF-8M 10k @ 400°C

CF-8 Unaged

54°C Air / 54°C PWR

315°C PWR / 315°C PWR

315°C PWR / 54°C PWR

Neutron Exposure (dpa)

CASS Materials

Irradiated 315-320°C

Specimens precracked & tested in

Closed Symbols: PWR Water at 320°C

Open Symbols: Air at 290°C

Irr./tested 

Lower bound at 427°C

Irr./tested at 290°C 

in BWR water

J = 255 kJ/m2

CF-8 Aged 134.8k @ 350°C

Precracked/Tested

J at 2.5 mm Crack Extension (kJ/m2)

0

200

400

600

800

1000

0.01 0.1 1 10

304 HAZ, Sindelar et al. 1993, C-L

304 HAZ, Sindelar et al. 1993, L-C

304 HAZ, Fyfitch

304 HAZ, ANL

304L HAZ, ANL

J at 2.5 mm CRack Extension (kJ/m2)

Neutron Exposure (dpa)

Austenitic SSs HAZ Materials

Closed Symbols: BWR Water

Open Symbols: Air

Irradiated & tested at 288°C

Irradiated 100-150°C, tested 125°C

J = 255 kJ/m2

Lower bound

115

6. SUMMARY 

The significant results for thermal embrittlement; IASCC susceptibility and CGRs; and neutron 

embrittlement, including combined effects of thermal and neutron embrittlement, for austenitic 

SS welds, HAZ materials, and CASS materials are summarized below. In addition, information 

gaps in the data needed to address the effects of thermal and neutron embrittlement on the 

structural and functional integrity of CASS components and austenitic SS welds in LWR 

systems are identified.

6.1 Significant Results

6.1.1 Thermal Embrittlement of Duplex Stainless Steels

Cast Austenitic Stainless Steels

Stainless steels are completely austenitic in the wrought condition, but welded SSs and CASS 

materials have a duplex structure consisting of austenite and ferrite phases. The ferrite phase 

increases tensile strength and improves resistance to SCC. The ferrite content generally varies 

significantly in CASS components, particularly static cast components. Variations have been 

observed in the range of 3.0–22.5% in a CF-8M hot-leg elbow and 1.5–15.0% in a crossover-leg 

elbow. For centrifugally cast pipes, the ferrite content varies across the wall thickness. These 

observations indicate that when measured ferrite content is used for design analyses, it is 

important to ensure that it is representative of the entire component. 

The quantitative metallographic method provides the most accurate estimate of ferrite content in 

CASS components. However, it may not be practical because of large variations in ferrite within 

a casting. Therefore, determination of ferrite from chemical composition of the material offers 

the most useful and common method for estimating ferrite in castings or welds. For CASS 

materials, the average measured ferrite is comparable with the values estimated from the 

material chemical composition using the Hull’s equivalent factors or from the ASTM 

A800/A800M methodology, although the latter seems to underestimate ferrite contents above 

20%. The difference between ferrite contents calculated from Hull's equivalent factor and 

measured values is ±6% ferrite. Both methods underestimate the ferrite contents of austenitic 

SS welds. The most commonly used instruments for measuring ferrite contents in CASS 

materials and austenitic SS welds are the Magne-Gage and Ferritescope. However, the surface 

roughness or curvature of the sample can introduce error because of variations in the magnetic 

linkage between the measuring probe and the test specimen. 

It is well established that CASS materials are susceptible to thermal embrittlement when aged at 

temperatures between 290 and 450°C (554 and 842°F). The formation of Cr-rich α' regions in 

the ferrite is the primary mechanism for thermal embrittlement. The α' phase strengthens the 

ferrite matrix and increases strain hardening, which raises the local tensile stresses above the 

critical value for brittle fracture. In addition, the formation of G phase and precipitation and/or 

growth of carbides or nitrides at the phase boundaries lead to a brittle failure by phase boundary 

separation. These particles may also facilitate cleavage of the ferrite by particle cracking. The 

microstructural changes result in brittle transgranular fracture of the ferrite or separation of the 

ferrite/austenite phase boundary. The degree of embrittlement and, hence, the toughness of the 

material, is controlled by the amount of brittle fracture. A predominantly brittle failure occurs when 

either the ferrite phase is continuous (i.e., high ferrite content) or the ferrite/austenite phase 
116

boundary provides an easy path for crack propagation (i.e., high-C grades of steel). 

Furthermore, the microstructural changes in the ferrite phase by spinodal decomposition are

very sensitive to the initial state of the ferrite. Thus, any manufacturing process parameter that 

improves the homogeneity of the ferrite solid solution, such as long homogenizing heat 

treatments and rapid quench, delay the beginning of the ferrite decomposition, which results in a

more aging-resistant CASS material.

Studies at Argonne have characterized the extent of thermal embrittlement of CF-3, CF-8, and 

CF-8M steels in terms of the RT Charpy-impact energy. All CASS materials attain “saturation” 

RT Charpy-impact energy (i.e., a minimum value that would be achieved by the material) after 

long-term thermal aging. The actual saturation value is independent of aging temperature but 

depends strongly on the chemical composition of the steel. It is lower for the CF-8M steel than 

for the CF-3 or CF-8 steels, and decreases with an increase in ferrite content or the 

concentration of C or N in the steel. The kinetics of embrittlement are controlled by three 

processes: spinodal decomposition, precipitation and growth of phase boundary carbides, and 

precipitation of G phase in ferrite. Small changes in the composition cause the kinetics to vary 

significantly. The activation energies range from 65 to 230 kJ/mole (15 to 55 kcal/mole). The 

low values are most likely due to the formation of carbides/nitrides at the phase boundaries or Gphase and/or γ2 precipitation in ferrite. The presence of Ni-Si-Mo clusters in the ferrite matrix of 

an unaged material is considered a signature of steels that show low activation energy (i.e., fast 

embrittlement). Such materials contain G-phase particles after aging.

The laboratory data obtained for materials aged at 400°C (752°F) and the activation energy for 

the process of embrittlement are typically used to predict the end-of-life impact strength at 

reactor temperatures. This methodology, however, assumes a unique aging behavior at 400°C, 

which is not observed for CASS materials. For a given decrease in the RT Charpy-impact 

energy, the time for aging at 400°C varies by more than two orders of magnitude for the various 

heats of CF-3, CF-8 and CF-8M steels. In general, CASS materials that take longer for thermal 

embrittlement to start have low activation energy (i.e., embrittlement is fast) and materials that 

take short time for embrittlement to start have high activation energy (i.e., take longer to 

embrittle). Therefore, estimates of embrittlement at reactor temperatures (290–320°C [554–

608°F]) are not likely to be significantly affected by which set of activation energy and kinetics 

data are used for estimations. 

An earlier report, NUREG/CR-4513, Rev. 1, presented a methodology for estimating the fracture 

toughness J-R curves of CASS materials under LWR operating conditions. The saturation RT 

Charpy-impact energy was correlated to the ferrite content and chemical composition of the 

CASS material. The existing thermal aging data indicated that a correlation between RT 

Charpy-impact energy and ferrite content alone did not yield good results. The extent of 

embrittlement as a function of time and temperature of reactor service was estimated from the 

extent of embrittlement at saturation and the correlations describing the kinetics of 

embrittlement, which were also given in terms of chemical composition. 

However, the methodology of NUREG/CR-4513, Rev. 1, is not applicable to CASS materials 

containing greater than 25% ferrite. A recent EPRI Report 1019128 presented a Westinghouse 

database of random sampling of heats of CASS material from 15 nuclear power plants. The 

data indicated that about 9% of the heats of CF-8M material currently in use in operating 

nuclear power plants contain greater than 25% ferrite. There is no acceptable methodology for 

estimating the reduction in fracture toughness of these materials with reactor service. 
117

In this report, the correlations presented in NUREG/CR-4513, Rev. 1, for estimating the 

reduction in fracture toughness of CASS materials as a function of reactor service have been 

updated to cover CASS materials containing greater than 25% ferrite. The NUREG/CR-4513 

expressions for estimating fracture toughness coefficient C and exponent n of the J-R curve at 

RT and 290–320°C (554–608°F) from RT Charpy-impact energy have been reevaluated and, if 

needed, updated to cover ferrite contents up to 40%. Common lower-bound J-R curves at RT 

and 290–320°C for thermally aged static and centrifugally cast CF-3, CF-8, and CF-8M steels 

with four ranges of ferrite content, <10%, 10–15%, 15–25%, and 25–40% ferrite, are also 

defined. Separate expressions are proposed for CF-8M steels containing <10% Ni and those 

containing ≥10% Ni. The latter are more susceptible to thermal embrittlement than the other 

grades.

A recent scoping study on low-temperature crack propagation for CASS CF-8 material in PWR 

environments investigated the potential synergy between hydrogen embrittlement and thermal 

embrittlement associated with decomposition of the ferrite at reactor temperatures. Fracture 

toughness J-R curve tests were conducted in air and PWR SWC at 54°C on 1-T CT specimens 

of CF-8 material (Argonne Heat 68) that was thermally aged for about 15.8 yr at 350°C (662°F). 

The results show a large effect of PWR SWC environment on fracture toughness; every material 

and test condition was identical for the two sets of duplicate tests except that one set of 

sepcimens was tested in air and the other set in water. The reason for the drastic reduction in 

fracture toughness is not clear. 

Correlations for estimating changes in tensile stress and Ramberg/Osgood parameters for 

tensile strain hardening are not presented in this report; they are discussed in detail in 

NUREG/CR-6142, ANL-93/35 (1994). The information may be used to determine fracture 

toughness JIc for the aged materials. A flow diagram of the methodology for estimating 

mechanical properties of thermally aged CASS materials in LWR systems is discussed. Three 

schemes are presented for estimation of mechanical properties of CASS material of 

(a) unknown composition, (b) known composition and unknown service history, and (c) known 

composition and service history. 

The correlations for estimating the degradation of mechanical properties of typical heats of 

CASS materials, however, do not consider the initial fracture properties of the unaged material. 

Some CASS materials may have low initial fracture toughness and the estimated J-R curves 

may be higher than the initial value. Therefore, some knowledge regarding the initial fracture 

toughness of the material is needed to justify the use of the estimated fracture toughness. The 

initial fracture toughness J-R curves may be estimated from the RT Charpy-impact energy of the 

unaged material, if known. Otherwise, minimum fracture toughness of the unaged material is 

used (i.e., Eqs. 43 and 44 of this report).

An evaluation of the NRC screening criteria established to determine the susceptibility of CASS 

materials to thermal aging embrittlement indicate that the existing criteria are valid, except those 

for CF-8M materials. The present reevaluation of the thermal embrittlement data indicates that 

to maintain the same level of safety margin, the existing criterion for potentially susceptible 

material of greater than 14% ferrite for static-cast CF-8M material and greater than 20% ferrite 

for centrifugally-cast CF-8M material, needs to be revised. The updated criterion for static-cast 

and centrifugally-cast CF-8M materials has been decreased to 13% ferrite and 17% ferrite, 

respectively, for materials with less than 10% Ni, and to 10% ferrite and 13% ferrite, 

respectively, for materials with equal to or greater than 10% Ni.
118

Austenitic Stainless Steel Welds

Austenitic SS welds also have a duplex structure and their chemical compositions are similar to 

those of CASS materials. Four distinct morphologies are observed in SS welds: vermicular, 

lacy, acicular, and globular. A vermicular morphology is typically observed in austenitic SS 

welds with FN 5–15, and lacy or acicular morphologies are observed in welds with FN 13–15.

Austenitic SS welds generally contain 5–15% ferrite, but their mechanical properties differ from 

those of CASS materials. For a given ferrite content, their tensile strength is higher and fracture 

toughness is lower than that of CASS materials. Generally, ferrite phase has little effect on the 

fracture toughness of austenitic SS welds because of the thin vermicular morphology and 

relatively low ferrite content. The welding process, however, significantly affects fracture 

toughness. In general, GTA welds exhibit higher fracture resistance than SMAs or SA welds. 

The J-R curve proposed earlier in NUREG/CR-6004 (1995) (i.e., Eq. 45 of this report) for SA 

and SMA welds, is defined as the lower-bound J-R curve at 288°C (550°F) for unaged austenitic 

SS welds.

The mechanism for thermal embrittlement of austenitic SS welds is somewhat different that of 

CASS materials. Thermal aging of austenitic SS welds at reactor operating temperatures 

results in moderate decreases in Charpy-impact strength and fracture toughness at both RT and 

290°C. In SS welds, failure occurred by the formation and growth of microvoids near hard 

inclusions. However, since the ferrite content in austenitic SS welds is typically less than 15%, 

the degradation of mechanical properties occurs primarily by the presence of grain boundary or 

phase boundary carbides, which initiate localized intergranular cracking. Based on fracture 

toughness data for thermally aged SMA and SA welds, the J-R curve for fully embrittled welds 

at 288°C was defined in NUREG/CR-6428 (1996) (Eq. 46 of this report). The curve is about 

40% lower than that for the lower-bound curve for unaged welds. However, the curve was 

based on very limited information. Furthermore, the possible effect of reactor coolant 

environment is not considered in these lower-bound curves. The lower-bound fracture 

toughness J-R curve can be further reduced by 40% in LWR environments relative to that in air. 

6.1.2 IASCC Crack Growth Rates

Susceptibility to IASCC increases with neutron fluence, corrosion potential, and water 

conductivity and decreases with decrease in the corrosion potential of the coolant environment. 

However, low corrosion potential does not provide immunity to IASCC if the fluence is high 

enough. The key 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) that 

influence IASCC had been identified and their effects on IASCC susceptibility of austenitic SSs 

were evaluated in NUREG/CR-7027. The discussions of microstructure, microchemistry, 

radiation hardening, yield strength, Si segregation, and stacking fault energy, are presented in 

this report as background.

In NUREG/CR-7027, the IASCC CGR data for irradiated wrought and CASSs and their weld 

metal and HAZ materials in LWR environments were compiled and compared with the K versus 

CGR disposition curves proposed in NRC NUREG-0313 for nonirradiated sensitized austenitic 

SSs and the EPRI curve for austenitic SSs. In the NWC BWR environment, the IASCC CGRs 

of austenitic SSs irradiated to 5 dpa could be a factor of 10 higher and those irradiated to 10–

12 dpa could be a factor of 25 higher than the NUREG-0313 disposition curve for nonirradiated 

SSs in 8-ppm water. A significant decrease in CGRs is observed in HWC BWR environment 

relative to the NWC BWR environment for SSs irradiated to less than 3 dpa; little or no reduction 
119

for some SSs irradiated to 3–4 dpa; and no decrease for all SSs irradiated to 12–14.5 dpa. The 

CGRs for some of the materials irradiated to 3–4 dpa are above the EPRI disposition curve for 

the HWC BWR environment. In PWR water at 320°C, the CGR data for austenitic SSs

irradiated to 3 dpa are up to a factor of 12 above the NUREG-0313 curve for nonirradiated 

materials in high-purity water with 0.2 ppm DO. The CGRs for SSs irradiated to 11–25 dpa are 

nearly two orders of magnitude above the same curve. For the same material and irradiation 

condition, the CGR in PWR water increase with increasing temperature, the data were 

normalized using activation energy of 100 kJ/mol.

The IASCC CGRs in NWC BWR environments for SA or SMA weld HAZ materials are 

comparable to or higher than for wrought austenitic SSs. Materials irradiated up to about 

2.5 dpa are a factor of up to 20 higher than the CGRs predicted by the NUREG-0313 curve for 

nonirradiated sensitized austenitic SSs in high-purity water with 8 ppm DO. Furthermore, the 

CGRs for these materials in HWC BWR environments are mostly below the 0.2-ppm DO 

NUREG-0313 curve. The available IASCC CGRs for thermally aged and unaged Type 316L 

weld metal at 289°C in high-purity water with 300–2000 ppb DO are bounded by the NUREG0313 curve for nonirradiated sensitized austenitic SSs in 8-ppm DO water. Similar data for 

irradiated Type 316L weld metal are not available. The CGRs for CF-8M CASS material 

irradiated to 0.08 dpa in high-purity water with <50 ppb DO at 320°C are below the NUREG0313 curve in 0.2-ppm DO water.

For irradiated austenitic SSs, the IASCC CGR data in NWC BWR environment at 288°C (550°F)

and K = 20 MPa m1/2 indicate that for neutron doses below about 0.42 dpa, the CGRs are below

the 8 ppm DO NUREG-0313 curve, and for neutron doses above 0.42 dpa the CGRs are 

bounded by Eq. 49. Similar data for CASS materials and austenitic SS welds and HAZ 

materials also indicate that in NWC BWR environment, the IASCC CGRs for nonirradiated 

materials or materials irradiated below 0.42 dpa are bounded by the NUREG-0313 curve for 

nonirradiated sensitized austenitic SSs in high-purity water with 8 ppm DO (i.e., are also 

bounded by Eq. 49).

For irradiated austenitic SSs, the IASCC CGR data in HWC BWR environment at 288°C or 

PWR primary water at 320°C and K = 20 MPa m1/2 indicate that for neutron doses below 

2.2 dpa, the CGRs are below the 0.2 ppm DO NUREG-0313 curve; for neutron doses above 

2.2 dpa, the CGRs are bounded by Eq. 50. Similar data for CASS materials, austenitic SS 

welds, and HAZ materials in HWC BWR environment at 288°C or PWR primary water chemistry 

indicate that the CGRs for nonirradiated CASS materials or materials irradiated up to 2.2 dpa 

are below the NUREG-0313 curve in 0.2-ppm DO water. However, CGR data are not available 

for CASS materials and austenitic SS welds irradiated above 2.2 dpa.

6.1.3 Neutron Embrittlement

The fracture toughness of unaged and nonirradiated wrought and cast austenitic SSs fall in 

Category III (JIc >150 kJ/m2), and failure occurs after stable crack extension at stresses well 

above the yield stress. Neutron irradiation can degrade fracture toughness of these materials to 

the level of Category II (JIc of 30–150 kJ/m2) or Category I (JIc < 30 kJ/m2), where failure may 

occur without general yielding. In such instances, EPFM is generally used for analysis of 

structural integrity, which involves the J-R curve approach.

Plots of JIc versus neutron dose are typically used for developing screening criteria for neutron 

embrittlement. JIc is a measure of fracture toughness at instability without significant stable 

crack extension. It is determined from the intersection of the best-fit power-law J-R curve with 
120

the 0.2-offset line defined as J = 2σf Δa. However, this yields a higher value of JIc for 

Category III materials such as austenitic SSs because of their relatively high toughness, 

ductility, and strain hardening ability. For austenitic SSs, a slope of 4 better defines the blunting 

line, and has been used to determine JIc in most investigations.

Based on data on fast reactor irradiated austenitic SSs, the effects of neutron exposure (in dpa) 

on the fracture toughness JIc have been 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 

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]). The fracture toughness trend for the LWR irradiated austenitic SSs is similar to 

that observed for fast reactor irradiations. Most of the fracture toughness JIc values for SSs 

irradiated at 288–316°C (550–601°F) fall within the scatter band observed for fast reactor data, 

although the LWR irradiations were at lower temperatures. However, some SSs irradiated 

above 4 dpa at LWR temperatures show very poor fracture toughness. For a Type 304 SS 

irradiated to 4.5–5.3 dpa, nine out of 10 specimens showed no ductile crack extension and the 

KIc values were 52.5–67.5 MPa m1/2 (47.7–61.4 ksi in.1/2). The lowest fracture toughness KIc or 

KJc values were in the range of 36.8–40.3 MPa m1/2 (33.5–36.6 ksi in.1/2). 

Specimen orientation has a strong effect on the fracture toughness of irradiated austenitic SSs. 

Types 304 and 304L materials irradiated to 4.7–12.0 dpa in BWRs showed lower fracture 

toughness in the T-L orientation than in the L-T orientation. The toughness values for T-L 

orientation were lower than the limiting fracture toughness KIc of 55 MPa m1/2 (50 ksi in.1/2) 

proposed by industry for flaw tolerance evaluation for austenitic SSs irradiated above 4.5 dpa.

The effects of various parameters such as material type, heat treatment, test and irradiation 

temperature, and neutron energy spectrum, flux, and dose are discussed. 

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 CASS material and austenitic SS weld metal is lower than that of HAZ material, which, in 

turn, is lower than that of solution-annealed materials. The fracture toughness JIc of austenitic 

SSs irradiated to less than 5 dpa decreases with increasing test temperature. For steels 

irradiated to more than 12 dpa, test temperature has little effect on fracture toughness. 

The threshold dose and the rate of the decrease in fracture toughness JIc with neutron dose 

seem to vary for different grades of materials. The decrease in toughness with increasing 

fluence occurs earliest in Type 304L SS, followed by Type 304 SS and then Type 316 SS. 

However, it is not clear whether the threshold for irradiation effects on fracture toughness is 

earlier or the decrease in fracture toughness occurs faster. The average JIc of Type 304 SS 

decreases from more than 150 kJ/m2 (857 in.-lb/in.2) at 1 dpa to 12–24 kJ/m2 (69–137 in.-lb/in.2) 

at about 2.2 dpa. 

A fracture toughness trend curve, represented by either the JIc value or coefficient C of the 

power-law J-R curve, has been developed that bounds the existing neutron embrittlement data. 

The trend curve describes (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. The lower bound curve 

represents a threshold dose of about 0.3 dpa for neutron embrittlement, a minimum fracture 

toughness JIc of ≈118 kJ/m2 for neutron doses below the threshold value, a saturation threshold 
121

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 disposition curve proposed by EPRI for 

BWRs is not bounding for the existing data for BWR-irradiated austenitic SSs.

The lower bound trend curve is consistent with the MRP lower bound model proposed for 

PWRs. The MRP curve also predicts a saturation fracture toughness KIc of 38 MPa m1/2. 

However, the MRP curve predicts fracture toughness values that, for some austenitic SS welds 

or weld HAZ, may be higher than the minimum toughness of the materials in the nonirradiated 

condition. 

The existing fracture toughness data have been used to develop the lower-bound curve for the 

decrease in the coefficient C of the power-law J-R curve with neutron dose. The lower-bound 

trend curve for the change in coefficient C of the power-law J-R curve versus neutron dose may 

be used to analyze material behavior for loading beyond JIc. A median value of 0.37 is 

assumed for exponent n of the J-R curve. The existing data indicate that the lower-bound curve 

for coefficient C based developed in this study may not be adequate for austenitic SS welds. 

Furthermore, the effects of LWR coolant environments may also decrease fracture toughness of 

wrought and cast SSs and their welds. 

As discussed above, the available fracture toughness data on austenitic SS welds and 

associated HAZ materials as well as CASS materials indicate that some of the data for 

austenitic SS welds, particularly Type 316 weld, may not be bounded by the lower-bound 

fracture toughness versus neutron dose curve proposed in NUREG/CR-7027. Furthermore,

limited J-R curve data in LWR coolant environments indicate that the reactor coolant 

environment may further decrease fracture toughness of these materials relative to that in air. 

Therefore, the effects of neutron embrittlement and the potential effect of coolant environment 

on fracture toughness of CASS materials and austenitic SS welds need to be established for 

performing flaw tolerance evaluations for these materials. 

6.1.4 Combined Effects of Thermal and Neutron Embrittlement 

In NRC staff’s evaluation of License Renewal Issue No. 98-0030, discussed in the letter from C. 

I. Grimes, NRC License Renewal and Standardization Branch, to Douglas J. Walters, Nuclear 

Energy Institute, the staff accepted the industry proposal of using fracture toughness 

J2.5 = 255 kJ/m2 to determine susceptibility of CASS materials to loss of fracture toughness due 

to thermal embrittlement. However, at the time when these screening criteria were proposed, 

the combined effects of thermal and neutron embrittlement of CASS materials and austenitic SS 

welds were not well established. Although both the thermal aging embrittlement and the 

neutron irradiation embrittlement of ferrite have been well characterized, the combined effect of 

thermal and irradiation embrittlement is not well established. The limited existing fracture 

toughness data indicate that the concurrent exposure to elevated temperatures and high 

neutron fluence levels could result in a combined effect wherein the service-degraded fracture 

toughness is reduced from the levels predicted independently for either of the two mechanisms

alone. However, the existing fracture toughness data are inadequate to evaluate the combined 

effects of thermal and neutron embrittlement on the threshold dose for embrittlement. 
122

6.2 Information Gaps

6.2.1 Thermal Embrittlement of Duplex Stainless Steels

The updated expressions presented in this report for estimating mechanical properties of CASS 

materials are based on data that contained little or no fracture toughness J-R curve results for 

thermally aged CF-3 and CF-8 materials with more than 25% ferrite. The use of the 

methodology proposed in this report for CF-3 and CF-8 materials with more than 25% ferrite 

needs to be validated. The fracture toughness lower-bound J-R curve for fully embrittled SMA 

and SA welds was defined in NUREG/CR-6428. However, the curve was based on very limited 

information, particularly for Type 316 welds. Additional data on austenitic SSs welds aged at 

reactor operating temperature are needed to better establish the lower-bound J-R curve for 

thermally aged welds, including the effect of reactor coolant environment.

6.2.2 IASCC Crack Growth Rates 

As discussed in NUREG/CR-7027, additional data on PWR materials irradiated between 15 and 

40 dpa at temperatures of 300–350°C are needed to better understand the IASCC susceptibility 

of austenitic SSs in PWR environments, and to understand the reason for high CGRs and large 

data scatter for SSs irradiated to 7–18 dpa and relatively low CGRs in SSs irradiated to

37.5 dpa. 

A review of the existing IASCC susceptibility data for CASS materials and austenitic SS welds 

indicates that the data are too limited to accurately establish the IASCC CGRs of these 

materials in LWR environments. Additional data on irradiated CASS materials and austenitic 

SS weld welds, including weld HAZ, are needed, particularly in low-DO environments. 

6.2.3 Neutron Embrittlement

Most of the data for CASS materials and austenitic SS welds shown in Fig. 62b for the fracture 

toughness coefficient C versus neutron dose trend curve, are from irradiations in fast reactors 

and at temperatures of 370–427°C (698–800°F). The values for coefficient C for LWR irradiated 

materials are likely to be lower. Furthermore, there are few or no fracture toughness data on 

LWR irradiated austenitic SS and CASS materials for dose levels of 1.0–10.0 dpa. Additional 

data on LWR irradiated materials should be obtained to accurately define the lower bound trend 

curve for coefficient C. 

Some of the data obtained in LWR environments falls below the lower-bound fracture toughness 

curve. Therefore, the potential effects of the LWR coolant environment also need to be 

investigated. Furthermore, the contribution of additional precipitate phases, voids, and cavities 

on fracture toughness needs to be investigated.

6.2.4 Combined Effects of Thermal and Neutron Embrittlement 

Additional fracture toughness J-R curve data on the combined effects of thermal and neutron 

embrittlement of austenitic SS welds and CASS materials, including the potential effects of LWR 

coolant environments, are needed to accurately define the lower-bound fracture properties of 

these materials in LWR environments.
123

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NUREG/CR-7185 Effect of Thermal Aging and Neutron Irradiation on Crack Growth Rate and Fracture 

Toughness of Cast Stainless Steels and Austenitic Stainless Steel Welds