Fatigue-crack propagation and fracture toughness behavior of cast stainless steels

Engineering Fracture Mechanics Printed in Great Britain.

Vol.

29, No. 4, pp. 423-434,

0013-7944/88 s3.00+ .oa Pergamon Press pk.

1988

FATIGUE-CRACK PROPAGATION AND FRACTURE TOUGHNESS BEHAVIOR OF CAST STAINLESS STEELS L. A. JAMES and W. J.MILLS Fellow Engineers, Westinghouse Hanford Company, Richland, WA 99352, U.S.A. Abstract-The fatigue-crack growth and fracture toughness behavior of two cast stainless steels, ASME SA 351 Grades CF8 and CWM. was investigated over a wide range of temperatures. The effect of a number of material-related paramekrs such as ferrite level, irack orientation, and heat-to-heat and alloy-to-alloy variations was studied. In addition, the crack growth and J-integral toughness of cast CF8 subjected to neutron irradiation was characterized.

INTRODUCTION CAST stainless steels are employed in a number of structural applications in both nuclear and petrochemical plants including piping, core support structures, valve bodies, pump housing and impellers, etc. Some of these components will be subjected to loadings in service that are cyclic in nature, and extension of flaws is possible if such flaws exist in the appropriate size, orientation, and location. For this reason, it is important to know both the fatigue-crack propagation (FCP) and fracture toughness (FT) behavior of the cast steels in order to assess the structural integrity of the components. Two of the more widely used cast stainless steels are ASME SA 351 Grades CF8 and CF8M, the cast equivalents of Types 304 and 316 stainless steel, respectively. Since these alloys are often employed in elevated temperature applications, the objective of this paper is to report the FCP and J-integral fracture toughness behavior over a range of temperatures.

EXPERIMENTAL

PROCEDURES

Two heats of each cast alloy, CF8 and CF8M, were employed in this study. These heats are identified in Table 1. Chemical compositions and mechanical properties are given in Tables 2 and 3, respectively. The microstructure for the castings consisted of an austenitic matrix with delta ferrite decorating the substructural boundaries (Fig. 1). Compact Type specimens meeting the requirements of ASTM E647-83 for FCP testing and E399-81 for FT testing were obtained from the various castings as shown in Fig. 2. Ferrite levels varied from casting to casting, from location to location within a given casting, and indeed from location to location within a given specimen. For example, six specimens from heat 3 taken from location “A” (see Fig. 1) had an average ferrite number (Magne gage) of 10.0, while seven specimens from location “B” had an average ferrite number of 11.5 (fourteen to seventeen ferrite measurements per specimen). Variations within a given specimen could also be quite large. For example, seventeen ferrite measurements were made at different locations on Specimen 2591 (from Heat 2). The range of measurements was 14.17-17.27 with a mean of 15.41 and a standard deviation of 0.89. Such differences represent real material variability since measurement variability was quite low (nineteen measurements at a single location had a mean of 8.11, a-range of 8.05-8.16, and a standard deviation of 0.04). Hence, with this degree of variability and inhomogenity within the

Table 1. Heat identification Heat ident. 1 2 3 4

Producer/heat

Alloy SA-35 1 Grade SA-351 Grade SA-351 Grade SA-351 Grade

CF8 CF8 CF8M CWM

ESC0/76560-1 ESC0/55368 ESC0/60330 ESC0/60069 423

no.

Condition As-cast As-cast As-cast As-cast

L. A. JAMES AND W. J. MILLS

424 Table Element

Heat

C

composition

1

Heat

0.08 0.54 0.029 0.007 0.089 9.06 19.69 0.24
Mn P S Si Ni Cr Mo V cu Sn Nb Ti co Al B N tNot

2. Chemical

(per cent by weight)

2

0.05 0.72 0.027 0.008 1.08 8.98 19.92

Heat 3

Heat 4

0.07 0.82 0.031 0.008 1.21 9.68 18.77 2.47

0.087 0.72 0.023 0.006 1.16 9.47 18.68 2.32 0.04 0.20 t co.01 co.01 0.09 0.02
:

t

: :

:

t

: t 0.005

: t

determined.

castings, it is difficult to draw conclusions on the influence of ferrite level upon either the FCP behavior or the toughness. ASTM Compact Type specimens were employed for both types of tests. Several specimen sizes were used: in general the width ‘ W’ was 50.8 mm, except for Heat 1 where W = 38.1 mm and the irradiated specimens where W = 29.21 mm. J-integral specimens had thickness (B) equal to W/2; FCP specimens had thicknesses one-half those of the JI, specimens. A few specimens were tested in the irradiated condition. The FCP specimens were irradiated to a level of 16.2 displacements per atom (dpa), and the toughness specimens were irradiated to 19 dpa. The details of the irradiation are given in ref. [l]. The cyclic frequency of all FCP tests was 0.67 Hz, except at room temperature where frequency is not expected to be an important variable. A sinusoidal waveform was used throughout, and the stress ratio (R = K,iJKm,,) was 0.05 for all tests. Fatigue-crack growth rates (da/dN) were calculated using the ‘secant method’, and the stress intensity factor range (AK) was calculated using the standard relationship of ASTM E647-83. A ‘flow stress criterion’[2] was employed to ensure that excessive plasticity in the untracked ligament did not invalidate the results. Toughness tests were conducted at four temperatures: 24”, 371”, 427” and 482°C. The multiple-specimen &-curve technique was used to determine fracture toughness. Analysis procedures differed slightly from those described in ASTM E813-81, because the ASTM test methods and size requirements are generally not applicable to high ductility materials, such as Table

3. Mechanical

properties

Test temp. (“C)

0.2% Offset yield (MPa)

Ultimate (MPa)

Total elong. (X)

Uniform elong. (%)

Reduction in area (%)

1 1 1 1

24 316 427 538

287 178 175 158

550 399 391 331

47.0 33.5 28.0 25.3

42.3 31.7 24.9 23.2

74.8 50.4 55.7 57.7

2 2 2

24 427 538

204 149 169

453 396 306

57.0 41.0 43.1

36.7 34.1 22.5

73.1 47.4 80.3

3 3 3 3

24-I 24t; 427t 538t

308 275 168 156

586 564 507 406

42.1 56.9 48.7 40.4

36.3 46.1 36.4 31.2

67.1 63.1 61.3 59.3

4 4 4

24t 427t 53st

263 163 140

590 451 328

56.5 38.6 39.2

50.6 33.0 29.7

56.3 45.1 65.0

Heat ident.

tFrom location *From location

‘A’ (see Fig. 1). Test conducted ‘B’ (see Fig. 1). Test conducted

by V. K. Sikka, ORNL. by V. K. Sikka, ORNL.

Fatigue

Fig. 1. Typical microstructure microstructure with delta ferrite

and fracture

for stainless located along

of cast stainless

steels

steel casting illustrating a duplex austenitic/ferritic the cores of dendritic branches. (a) Cast CF8. (b) Cast CWM.

425

426

L. A. JAMES

AND W. J. MILLS

Fig. 10. SEM Fractographs for CF8. (a) Unirradiated. Evidence of well-defined microvoid cence. (b) Irradiated to 19 dpa. Channel fracture with small microvoids located on the fracture The microvoids were nucleated by premature failure of delta ferrite particles.

coalesfacets.

Fatigue and fracture of cast stainless steels

427

Fig. 2. Orientation of the specimens.

cast stainless steel alloys. Accordingly, initiation toughness values were termed J, rather than Jr,. During each test, load-line displacements were measured by a high-temperature LVDT technique[3] and recorded continuously on an X-Y recorder as a function of load. The value of J was then determined by the following equation (per ASTM E813-81). J

=

G

(1+ a)

I36 (1 + C?)

where: A = area under load vs load-line displacement curve 6 = unbroken ligament size, (Y= [(2a/b)* + 2(2a/b) + 2]- (2a/b + l), a = crack length. The JR-curves were constructed by plotting values of J as a function of crack extension, Au, and the initiation J, value was then taken to be that value where a least-squares regression line through the crack extension data points intersected the crack blunting line. For unirradiated materials, the blunting line for low strength, high strain-hardening materials was employed[4]:

where: uf = flow strength = $(v,,,,+ u”,,), uys = 0.2% offset yield strength, mUuts = ultimate tensile strength. Neutron irradiation significantly increased strength and diminished strain-hardening capacity, so the conventional ASTM E813-81 blunting line J = 2cr(Aa)

was used to determine postirradiation fracture toughness. The variances of J, and dJR/du (S: and S$, respectively) were determined from statistical analysis of the JR curve data using the procedures outlined in ref. [5]. Values of J, f S1 and dJR/du f S, are reported for each JR curve. Tearing moduli, T, were calculated from the following equation[6]: ePll29:4-c

L. A. JAMES AND W. J. MILLS

428

where E is Young’s modulus. The fracture surface appearance of unirradiated fracture toughness specimens was characterized by direct fractographic examination on an SEM operated at an accelerating potential of 25 kV. To examine the fracture morphology of irradiated specimens, gold-coated celluloseacetate replicas were prepared and studied. RESULTS Fatigue-crack

AND DISCUSSION

propagation

Fatigue-crack growth tests were conducted at five different temperatures and the results are shown in Figs 3-7. The figures are arranged such that results for CF8 are shown on the left, while those for CF8M are shown on the right. Linear least-squares regression lines are plotted through each set of results. A number of observations are apparent when viewing Figs 3-7 and will be discussed below. Comparing the results for CF8 at a given temperature with those for CFSM at the same temperature does not reveal a significant difference between these grades over the entire range of temperatures surveyed. Likewise, there is not a consistent heat-to-heat variation between the two different heats of a given alloy. Considerably greater scatter in FCP rates is observed in the present results on these two cast Types 304 and 316[7]. In fact, the stainless steels than in their wrought counterparts: specimen-to-specimen variability appears to be as large as any potential alloy-to-alloy variability or heat-to-heat variability. Part of this variability may be due to the very uneven crack trajectories that were generally observed. Many of the results for Heat 1 were reported previously[S], and it was shown that the duplex austenite/ferrite microstructure contributed to the very uneven crack trajectories. The cracks generally tended to avoid, and propagate around, the islands of ferrite (an example of this is shown in Figs 7 and 8 of ref. [8]. This behavior is believed to be at least partially responsible for the relatively large degree of scatter in the present results.

1

21’C

,mw

0 SPEC. 2564. HEAT 3 0 SPEC. 25m. HEAT 3

A SPEC.

10-J

fd

10

Fig. 3. Fatigue-crack

1

I

,

I

(b’

100/10

STRESS INTENSITY FACTOR RANGE, AK, MPqfiii

ZBUI. HEAT 6

I

I

I

I

10-S 100

“IDLIla,

growth behavior of two cast stainless steels tested in air at 24°C.

Fatigue and fracture of cast stainless steels

g

316% ,6wF,

316’C ,(yIIBF,

0 SPEC. 241, HEAT I 0 SPEC. 244. HEAT, A SPEC. Z62. HEAT 2

0 SPEC. 2615. HEAT 3 0 SPEC. 2646, “EAT 3 A SPEC. m, HEAT 4

t

_

t (6)

10-5

429

,

10

,

Fig. 4. Fatigue-crack

INTENSITY

d

-1 ‘b)

I

,

,

,

10-6

100/10 STRESS

z F

100

FACTOR

RANGE,

bK. MPmJiii

IeDLIuII

growth behavior of two cast stainless steels tested in air at 316°C.

Duplicate specimens from Heats 1 and 3 were tested at most temperatures, and in these cases, one specimen was of orientation (or location) ‘A’ and one of ‘B’ (see Fig. 2). The orientations for Heat 1 are explicitly listed in ref. [8]. The Heat 1 duplicates represent two different orientations, and the Heat 3 duplicates represent two different locations (and hence ferrite levels), yet again consistent differences are not observed. One specimen of CF8 (Specimen 572 in Fig. 7a) was aged for 12,000 h at 538°C prior to testing. Although the FCP rates in the aged specimen appear to be slightly higher than its unaged counterparts, the difference is actually about the same as the general specimen-to-

3 2 -

10-4

u‘ t

a Ii 5 -

/

E 10-5

A 0427-C SPEC. IlDoDFl 2635. m. 2661. HEAT 3 4

,’

2 9 5 W F 2

1 10-6!;a’

*IRRADIATED TO 16.2 d,,.

STRESS

Fig. 5. Fatigue-crack

INTENSITY

,

(b)

100/10 FACTOR

I RANGE.

AK, MPmfiii

t

*

I

1

lo-6

100

“IOLIuI1

growth behavior of two cast stainless steels tested in air at 427°C.

430

L. A. JAMES AND W. J. MILLS

, STRESS

Fig. 6. Fatigue-crack

III,

INTENSITY

lb1 100110 FACTOR

I

,,,,0-6 100

RANGE,

AK. MPafiii

q

mtlP,&l4

growth behavior of two cast stainless steels tested in air at 482°C.

specimen variation. Hence, it appears that prior thermal aging has little or no influence on FCP behavior at 538°C. Two specimens which had been irradiated to 16.2 dpa were tested at 427°C. The results, plotted in Fig. 5(a) show that significantly higher FCP rates are observed in the irradiated material at the lower values of AK. The material in this condition had been subjected to considerable radiation-induced strengthening (see ref. [l] for the irradiated tensile properties). In view of the apparent lack of a consistent trend for heat-to-heat or alloy-to-alloy variations, crack orientation effects, or thermal aging effects, the data for CF8 and CF8M were combined to produce a single least-squares regression equation for each temperature. The data

10 STRESS

Fig. 7. Fatigue-crack

INTENSITY

100llO FACTOR

RANGE,

AK, MPafi

100 *rDi*o(Iu‘

growth behavior of two cast stainless steels tested in air at 538°C.

Fatigue and fracture of cast stainless steels

431

Table 4. Summary of crack growth equation constants? C (AK)”

g=

n

Comments

3.8243 4.2220 3.2756 2.0008 3.1265 2.5669

Excludes irradiated material Irradiated to 16.2 dpa

c

Temp. 24°C 316°C 427°C 427°C 482°C 538°C

3.3357 1.3127 3.6754 4.5837 4.6513 6.9699

tUnits: da/dN =

x lo-‘0 x 10-r” x lo-“’ x lo-’ x 1O-9 x 10-s

Includes aged material

mm/cycle,AK = MPa&.

on irradiated material were regressed separately. The resulting regression analyses, fitted to the form

$=

C(AK)”

are given in Table 4, and it will be noted that, in general, FCP rates tend to increase with increasing temperature. Similar tests have been conducted on cast stainless steels in air at temperatures in the range 288-32O”C[9-111, and these results agree well with the present results at 316°C (Fig. 4). Results may also be found in the literature for cast stainless steels in an elevated temperature water environment [9-l 21. I

I

I

I

I

1

I

CF8M HEAT 3

HEAT 4. 24% J, = 1397 * 164 kJ/mZ dJR/da = 597 f 99 MPa T= 494 \

A37vc 00 427oc v492oc

l /

HEAT 4 +

24%

n 427%

J, = 593 dJR/da = 397 T= 999

f f

113 kJ/m2 79 MPa

HEAT 3.371492% J, = 929 * 99 kJlm2 dJ&da = 390 f 53 MPa T = 597

ipJ

0

= 40‘ (Aa)

1.0

2.0 CRACK EXTENSION,

3.0 mm

Fig. 8. JR curves for cast CWM. The open square symbol represents orientation (C) (see Fig. 2). All other specimens were of the (A) or (B) orientations.

432

L. A. JAMES AND W. J. MILLS

Fracture toughness

The fracture toughness behavior for the four castings is summarized in Figs 8 and 9. Examination of these JR curves revealed a number of trends that are discussed below. The fracture toughness response for CF8M Heat 3 was insensitive to specimen orientation, location within the casting and test temperature (371, 427 and 482°C). As a result, data for the three test temperatures and two orientations were combined into a single &-curve regression. For Heat 4, J, and dJn/da values at 24°C were found to be much higher than their elevated temperature counterparts. Although the JR-curve slope was substantially greater at room temperature, the tearing modulus was approximately 25% lower because the increase in dJR/da was offset by a larger increase in flow strength. At elevated temperatures, CFSM Heat 3 exhibited the highest J, (829 kJ/m*); CF8M Heat 4 and CF8 Heat 2 displayed intermediate initiation toughnesses (583 and 576 kJ/m*, respectively); and CF8 Heat 1 had the lowest toughness (416 kJ/m*). JR curve slopes for the two CF8M castings (360 and 397 MPa) were slightly higher than those for the CF8 material (259 and 300 MPa). The fracture resistance of all castings was sufficiently high to preclude unstable fracture in most engineering components. Tearing in these materials requires crack lengths of the order of tens of centimeters and stresses well above the yield strength level. As a result, conventional design analysis methods (e.g. ASME Code stress and strain limits) adequately guard against premature fracture, so sophisticated elastic-plastic fracture toughness evaluations are not generally required. Only in special cases, such as quantifying safety margins for critical structures containing either real or hypothetical defects, would ductile fracture mechanics assessments be required. J, and tearing modulus values for the cast materials were within the range of values exhibited by wrought SS alloys[l3] (JC = 178 - 1635 kJ/m*; T = 272-676). This finding indicates that the presence of delta ferrite in the castings did not degrade fracture properties. However, the fracture characteristics of this second phase can be affected markedly at cryogenic temperatures and after long-term thermal exposure. Delta ferrite undergoes a ductile-to-brittle CF8 _

-

I

TESTED 0

l

HEAT HEAT

0

HEAT2

I

I

1

I

1 1 -

I 0

AT 427% IRRADIATED

_

(19 dpal

HEAT 2. .27-C J, = S7S * S4 kJ/m2

HEAT 1.427=X

T=

SSS

IRRADIATED 119 dpal J, * 20 kJ/m?dJR/da = E MPa

0

1.0 CRACK

2.0 EXTENSION.

3.0

mm

Fig. 9. JR curves for unirradiated and irradiated cast CF8. The data point with a vertical slash was not included in the JR regression because it was located near the crack blunting line.

Fatigue and fracture of cast stainless steels

433

transition temperature phenomenon, so its ductility at low temperatures is essentially nil. This degraded plastic straining capacity results in a substantial reduction in the overall fracture resistance of the casting[l4, 151. Long-term exposure at 300 to 500°C induces formation of three different precipitates (G-phase, Type X and chromium-rich a’) within the delta ferrite[l6, 171, which reduce its plastic straining capacity. The diminished fracture resistance of the ferrite phase has been shown to degrade the notch toughness[9,16] and fracture toughness[9, 181. At aging temperatures above 500°C other delta ferrite transformation products, M&Z6 carbides and sigma phase, are expected to control fracture resistance. Neutron irradiation of CF8 to 19 dpa resulted in an order of magnitude reduction in J, and dJn/da (Fig. 9). Moreover, the irradiation-induced strengthening coupled with the decrease in dJn/da caused a 300-fold degradation in tearing modulus. An equivalent plane strain fracture toughness (K,,) of 58 MPa& was computed for the irradiated casting using the following equation: K,, = JEJ, where E is elastic modulus. With the increase flaw sensitivity after irradiation, fracture mechanics analysis and effective NDE procedures are required to verify protection against premature failure for highly stressed components. Fractographic examination revealed that the reduction in fracture properties was associated with a fracture-mechanism transition from ductile microvoid coalescence to channel fracture (Fig. 10). The deep, well-defined dimples in the unirradiated condition demonstrated that the fracture process involved extensive homogeneous plastic deformation. After irradiation, the fracture surface took on a highly faceted crystallographic appearance, indicative of channel fracture[l9,20]. This mechanism results when all dislocation activity is channeled through narrow deformation bands. In stainless steel alloys, lead dislocations sweep out fine defects introduced by irradiation displacement damage, thereby creating defect-free zones that are weaker than the surrounding areas. Subsequent dislocation activity is then channeled through these planar slip bands, and ultimately shear cracks initiate and propagate along them. The planar slip and accompanying channel fracture resulted in the poor postirradiation fracture resistance. The delta ferrite particles initiated the microvoids found on the channel fracture facets. These microvoids, which formed during the early stages of the fracture process, served as stress concentration sites and prematurely initiated channel fracture. This caused the postirradiation fracture resistance of the casting to fall below that for wrought SS alloys (J, = 28 to 31 kJ/m*) after comparable neutron exposure[21].

CONCLUSIONS The fatigue-crack propagation and fracture toughness behavior for CF8 and CFSM was evaluated at room and elevated temperatures, and the results are summarized below: (1) In general, FCP rates increased with increasing temperatures in an air environment. Somewhat greater scatter in da/dN is observed in the cast material relative to the behavior in wrought stainless steel. This increased scatter is believed to be related to very uneven crack trajectories resulting from the duplex austenite/ferrite microstructure. (2) The increased scatter in FCP rates and specimien-to-specimen variability masked any possible effects of ferrite level, crack orientation, heat-to-heat variability or alloy-to-alloy variability. (3) Thermal aging (12,000 h at 538°C prior to testing) did not have an appreciable effect upon FCP behavior at 538°C. On the other hand, neutron irradiation to 16.2 dpa did result in an increase in crack growth rates at 427°C especially at the lower values of K. (4) The exceptionally high fracture resistance of the castings (Jc = 416-829 kJ/m* and T = 597-711 at 371 to 482°C; J, = 1397 kJ/m* and T = 484 at 24°C) demonstrated that fracture control is not a primary design consideration for this class of alloys in the unirradiated

434

L. A. JAMES AND W. J. MILLS

and unaged condition. Conventional design methods, such as the stress and strain limits provided by the ASME Code, are generally sufficient to preclude premature failure. (5) Neutron irradiation of the CF8 casting to 19 dpa resulted in an order of magnitude reduction in J, and two orders of magnitude reduction in T. To guard against brittle fracture after significant neutron exposure, fracture mechanics evaluations based on the supposition of a detectable crack in the most severely stressed location should be performed. (6) The severe degradation in toughness after neutron irradiation was associated with a fracture mechanism transition from microvoid coalescence to channel fracture. Acknowledgement+This paper is based on work performed under U.S. Department of Energy Contract DE-AC06 FF02170 with the Westinghouse Hanford Company, a subsidiary of Westinghouse Electric Corporation. The authors wish to acknowledge the careful experimental work of D. J. Criswell and W. D. Themar.

REFERENCES 111L. A. James, The effect of fast-neutron irradiation on the fatigue-crack

growth behavior of several austenitic stainless steels and weldments. Nucl. Technol. 74, 84-92 (1986). PI L. A. James, Specimen size considerations in fatigue-crack growth rate testing. Fatigue Crack Growth Measurement and Data Analysis, ASTM STP 738, 45-57 (198 1). [31 W. J. Mills, L. A. James and J. A. Williams, A technique for measuring load-line displacements of compact ductile fracture toughness specimens at elevated temperatures. J. Tesr. Eual. 5, 446-451 (1977). [41 W. J. Mills, On the relationship between stretch zone formation and the J-integral for high strain-hardening materials. J. Test. Eual. 9, 56-62 (1981). El W. J. Mills, Fracture Toughness of Stainless Steel Welds. HEDL-SA-3319, Westinghouse Hanford Company, Richland, WA (November 1985). if51P. C. Paris, H. Tada, A. Zahoor and H. Ernst, The theory of instability of the tearing mode for elastic-plastic crack growth. Elastic-Plastic Fracture, ASTM STP 668, 5-36 (1979). 171 L. A. James, Effect of heat-to-heat and melt practice variations upon fatigue crack growth in two austenitic steels. Prooerties of Ausreniric Stainless Steels and Their Weld Metals. ASTM STP 679. 3-16 (1979). PI L. k. Jam&, Fatigue-crack propagation in a cast stainless stekl. Nucl. Techno[.‘26, 46-53 (i975). [91 E. I. Landerman and W. H. Bamford, Fracture toughness and fatigue characteristics of centrifugally cast type 316 stainless steel pipe after simulated thermal service condition. Ductility and Toughness Considerations in Elevated Temperature Service, MPC-8, pp. 99-127. American Society of Mechanical Engineers, New York (1978). [lOI I. L. Bernard, G. Slama, C. Amzallag and P. Rabbe, Influence of PWR environment on fatigue crack growth behavior of stainless steels. Time and Load Dependent Degradation of Pressure Boundary Materials, IWG-RRPC79/2, pp. 27-35. International Atomic Energy Agency, Vienna (1979). Ull C. Amzallag, G. Baudry and J. L. Bernard, Effects of PWR environment on the fatigue crack growth of different stainless steels and inconel type alloy. Proceedings of the International Atomic Energy Agency Specialists’ Meeting on Subcritical Crack Growth, Report NUREG/CP-0044, pp. 263-294. U.S. Nuclear Regulatory Commission (1983). [12] W. H. Cullen, R. E. Taylor, K. Torronen and M. Kemppainen, The Temperature Dependence of Fatigue Crack Growth Rates of A351 CFSA Cast Stainless Steel in LWR Environment. Report NUREG/CR-3546, U.S. Nuclear Regulatory Commission (1984). [13] W. J. Mills, Heat-to-Heat Variations in the Fracture Toughness of Austenitic Stainless Steels. HEDL-SA-3343, Westinghouse Hanford Company, Richland, WA (1986). [14] E. L. Brown, T. A. Whipple and R. L. Tobler, Fracture toughness of CF8 castings at four Kelvin. Mefall. Trans. 14A, 1179-l 188 (1983). [15] T. A. Whipple and E. L. Brown, Deformation and fracture of stainless steel castings and weldments at 4 K. Materials Studies for Magnetic Fusion Energy Application at Low Temperatures-VI. Report No. NBSIR 83- 1690, pp. 213-233. National Bureau of Standards, Boulder, Colorado (May 1983). [16] A. Trautwein and W. Gysel, Influence of long-term aging of CF8 and CF8M cast steel at temperatures between 300 and 500°C on impact toughness and structural properties. Stainless Steel Castings, ASTM STP 756, 165-189 (1982). [17] 0. K. Chopra and H. M. Chung, Long-Term Embrittlement of Cast Duplex Stainless Steels in LWR Systems: Annual Report, October 1984-September 1985. Report No. NUREG/CR-4503, Argonne National Laboratory, Argonne, Illinois (January 1986). [18] S. W. Tagart, Jr, T. V. Marston, R. E. Nickel1 and D. M. Norris, Structural Mechanics Program: Progress in 1981. EPRI Report NP-2705-SR, Electric Power Research Institute, Palo Alto, CA (October 1982). [191 _ - C. W. Hunter, R. L. Fish and J. J. Holmes, Channel fracture in irradiated EBR-II type . . 304 stainless steel. Trans. Am. Nucl. Sot. 15, 254-255 (1972). r201 R. L. Fish. Notch effect on the tensile orooerties of fast-reactor-irradiated tvoe . . ,. 304 stainless steel. Nucl. Technol. 31, 85-95’(1976). [21] W. J. Mills, Fracture Toughness of Irradiated Stainless Steel Alloys. HEDL-SA-3471, Westinghouse Hanford Company, Richland, WA (1986). L

-

(Received

13 March 1987)