7. G.V. Scarich, G.R. Chanani, V.C. Petersen, and D.M. Weaver: Reference 4, pp. 103-14.

8. E647-81, "Standard Test Method for Constant-Load-Amplitude Fatigue Crack Growth Rates above 10 -~ m/cycle", American Society for Testing and Materials, 1981, pp. 765-76.

9. R.R. Boyer and R. Bajoraitis: Air Force Materials Laboratory Report, 1978, AFML-TR-78-131.

10. J.C. Chesnutt, A.W. Thompson, and J.C. Williams: Air Force Materials Laboratory Report, 1978, AFML-TR-78-68.

11. M. Creager and A.W. Sommer: Air Force Materials Laboratory Report, 1977, AFML-TR-77-193.

12. R.E. Peebles and C. A. Kelto: Reference 4, pp. 47-58. 13. G.R. Yoder, L. A. Cooley, and T. W. Crooker: Journal of Engineering

Materials Technology, 1979, vol. 101, pp. 86-90. 14. E.J. Dulis, V.K. Chandhok, F.H. Froes, and L.P. Clark: Pro-

ceedings of the lOth National SAMPE Technical Conference, Kiamesha Lake, NY, published by SAMPE, Azusa, CA, 1978, vol. 10, pp. 316-29.

15. I. Weiss, E H. Froes, and D. Eylon: unpublished research, 1982. 16. G. Welsch, I. Weiss, E H. Froes, and D. Eylon: unpublished research,

1982.

Microfracture Model for Hydrogen Embrittlement of Austenitic Steels

R.E. STOLTZ, N. R. MOODY, and M. W. PERRA

Microfracture models which relate microscopic fracture features to measured thresholds for crack advance have been recently proposed by Gerberich and co-workers. 1-4 These models have treated both static and fatigue thresholds for titanium alloys and martensitic steels. The models are predicated on the observation of a partially cracked (semi- cohesive) zone which, at a quasi-static cracking threshold, separates fully failed and unfailed regions of a cracked body. The ligaments in the partially cracked region provide tractions across the crack plane which reduce the local stress intensity and plastic zone in the unfailed material. This model has been used with success in titanium to account for the area (volume) fraction of ligaments in the zone as a function of temperature 2 and for the variation in crack ad- vance threshold with temperature and hydrogen content. 3

This communication describes the results of fracture surface analysis of selected austenitic steels which ex- hibit cracking in high pressure hydrogen gas. In the alloys studied, evidence of a partially cracked region was observed in samples which exhibited a crack arrest threshold in hydro- gen. By modifying the Gerberich approach, the length of this partially cracked region was related to the alloy yield strength and the arrest threshold stress intensity.

Experimental details are described fully elsewhere. 5"6"7 Briefly, bolt loaded wedge-open-load samples were exposed to 200 MPa hydrogen at 298 K. Crack advance was moni- tored by measuring load decreases in situ. Once a crack arrest threshold had been established the samples were frac- tured in air and the threshold region of the fracture surface was examined.

R. E. STOLTZ, Group Head, is with Surface Metallurgy Group, Exxon Corporate Research Laboratory, Linden, NJ 07036. N.R. MOODY, Member Technical Staff, and M.W. PERRA, Supervisor, Materials Development Division, are both with Sandia National Laboratories, Livermore, CA 94550.

Manuscript submitted December 6, 1982.

Table I, taken from Reference 7, gives results of crack arrest threshold vs yield strength for alloys with nominally clean grain boundaries (see References 5 and 6 for detailed discussion of grain boundary impurity effects in these materials). As seen from the table, alloys with yield strengths above 779 MPa exhibited hydrogen-induced cracking while those with strengths below 717 MPa did not. The two materials indicated by the asterisk (*) were studied in detail. Table II gives the composition and heat treat- ment of these alloys while Figure 1 shows fracture surface features in the threshold region. Hydrogen-assisted crack growth in these materials is intergranular while overload fracture following crack arrest is nominally by ductile void coalescence. The lines on the micrograph separate the par- tially cracked region from the fully cracked and uncracked regions. The measured widths, Ira,a,, of the partially cracked regions are given in Table II. These widths were measured from a series of low magnification micrographs and are the average of the local variation in widths shown at high magnification in Figures l(a) and l(b).

In order to relate the measured size of the partially cracked region to the yield strength and crack arrest threshold stress intensity of Table I, two assumptions are made. First, inter- granular fracture is assumed to occur if the resolved normal stress on the boundary reaches a critical value, o~. The local stress necessary for fracture will be higher than 07 if the

Table I. Crack Arrest Thresholds for Austenites in 200 MPa Hydrogen (Reference 7)

Crack Arrest Yield Strength, Threshold,

Alloy MPa MPa-m "2

304L (A) 372 (NCP 50)* 304L (B) 593 (NCP 110)

316 689 (NCP 132) Nitronic 50 717 (NCP 132)

(Fe-22Cr- 13Ni-5Mn) A286 779 94

Nitronic 40 827 99 (Fe-21Cr-6Ni-9Mn)

JBK-75 (A) 855 116 (modified A286)

316 903 99 *INCOLOY 903 (A) 917 70

JBK-75 (B) 923 66 (modified A286)

*INCOLOY 903 (B) 1055 30 *NCP = no crack propagation at indicated stress intensity level.

Table II. Composition, Heat 1~eatment, and Partially Cracked Zone Sizes for INCOLOY 903 in 200 MPa Hydrogen

Composition, Wt Pct: Fe-37.8Ni-15.2Co-3.1Nb- 1.3Ti-0.17Si-0.15Mn-0.04C

Condition l . . . . . /zm /~c,/xm

(A) Solutionize: 1213 K/ 1 hour 90 79 Age: 993 K/16 hours

(B) Solutionize: 1213 K/ 1 hour 20 18 Age: 993 K/ 8 hours +

893 K/ 8 hours

1528--VOLUME 14A, JULY 1983 METALLURGICAL TRANSACTIONS A

(b) (a)

Fig. 1--Fractographs of INCOLOY 903 (a) Condition A, (b) Condition B, at crack arrest threshold. The solid lines bound the partially cracked region.

boundary is inclined to the principal stress axis. This critical stress approach is supported by experiments 8-H which show that both temper and hydrogen embrittlement can be rational- ized by a local critical stress model. Second, it is assumed that the transition between cracking and no cracking occurs at yield strengths between 779 and 717 MPa (Table I) be- cause below 779 MPa, the resolved normal stress for frac- ture is not achieved. The particular transition yield strength

is, of course, dependent on alloy composition, grain bound- ary chemistry, temperature, and hydrogen fugacity.

In order to estimate the local critical stress for fracture, one must consider the stress elevation due to the crack. Continuum calculations of stress elevations ahead of cracks, which include both the effects of work hardening and plastic constraint, show that local stresses can reach a maximum of from three to five times the yield strength depending on

METALLURGICAL TRANSACTIONS A VOLUME 14A, JULY 1983-- 1529

work hardening rate. The austenitic alloys tested here ex- hibit a relatively high work hardening coefficient, n, of between 0.1 and 0.2. From the work of McMeeking, la the maximum stress attainable at these work hardening values and with a yield strain of 1/300, is between 3.1 and 4.1 times the yield stress. A stress elevation factor of 3.6 was thus chosen as an intermediate value. With the assumption that a minimum yield strength of 779 MPa is necessary for hydrogen-induced cracking in gaseous hydrogen, a o-7 of 2800 MPa is estimated for the alloys listed in Table II.

Figure 2 schematically shows a partially cracked region of the sample with a superimposed stress distribution. The origin of the stress distribution is chosen at the end of the fully cracked region. The axes are given in normalized units: (ryy is the stress normal to the crack plane, (r0 the yield strength, and K the applied stress intensity. The stress distri- bution is that given in References 13 and 14 where n is the work hardening coefficient. While the stress distribution was originally determined for a continuous, uncracked ma- terial, 13'~4 it is assumed here that, on average, the variation in stress with distance is similar for material with isolated flaws. This assumption was also invoked in the earlier appli- cations of the microfracture model for the thresholds. 1-4 Location (A) at the end of the partially cracked region is the distance at which a stress of o-7 is reached. Using n = 0.15, o-~ = 2800 MPa and the values for yields strength and K at crack arrest threshold, the length of the partially cracked zone was calculated. Results of calculations for the two alloys studied is given as l:alc in Table II. Reasonable agreement is observed between calculated and measured zone sizes.

The assumption that a cracking-no cracking transition yield strength can be related to the attainment of a critical local stress for fracture requires further discussion. Speci- men thickness limits the initial applied stress intensities in softer materials due to through-thickness yielding and plane stress effects. It is possible that in tests with thicker samples where plane strain conditions can be maintained to even

PARTIALLY CRACKED !CRACKED !UNCRACKED

I ,,

t ' \ a.5 ; I q._~n I,' X ~ 0.3 [ 0.4 | n + 1

I o" K 2 K 2 I O - - ~ 3.0 ~ [ 0.1 + (70 .1 + X (7

/ " ' , ' , ( L J

oo t 2.5

2 | I I i I "00 0.01 0.02 0.03 0.04

Fig. 2 - - S c h e m a t i c o f part ial ly c racked region and normal ized stress distri- bution. Equat ion for stress distr ibution af ter Refs. 13 and 14.

higher applied stress intensities, hydrogen cracking would occur in lower strength alloys. However, preliminary tests with side grooving of samples, which tends to compensate for thickness limitations on plane strain t5 show that the 717 to 779 MPa transition for crack growth is still observed.~6

A critical yield strength below which austenitic steels appear to be resistant to hydrogen cracking of 779 MPa is surprisingly similar to values observed for martensitic steels in severe hydrogen environments. Tests of a clean mar- tensitic low alloy steel in H2S show a strong upturn in threshold as a function of decreasing strength which indi- cates a critical strength of - 8 0 0 MPa below which the material is immune. ~7 Similarly, 4340 steels under stress corrosion conditions appear resistant below 830 MPa. 18 The common features between the anstenites tested in this work and the martensitic steels are that the fracture path is inter- granular, and the grain boundaries are substantially free from segregated impurity effects. Furthermore, the actual hy- drogen fugacity used in the tests of the austenites, 200 MPa, and the effective fugacity in H2S or stress corroding environ- ments is extremely high. It is tempting to speculate that the critical stress for sustained load intergranular fracture of iron based alloys in the presence of very high hydrogen fuga- cities is less a function of crystal structure than the local structure of the grain boundary which may be similar in both bcc and fcc crystals. (Reference 19 reviews current models for grain boundary structures.) This is a point for further theoretical and experimental consideration.

The calculation for the size of the partially cracked region is dependent only on an assumed value for 0-7 and a knowl- edge of or0 and K at threshold. No characteristic length of the microstructure is involved in the calculation. However, as the yield strength increases and the K~ is observed to de- crease, 7 calculated values for l begin to reach the order of the grain size. (Indeed, the 20/xm calculated for alloy (B) in Table II is approximately two thirds the average grain size.) In order for macroscopic crack advance to occur, it is reasonable to assume that o-~ must be exceeded over a length equal to a grain facet. This requirement puts a lower limit to the threshold stress intensity for polycrystalline materials as the yield strength is increased. For martensitic steels with yield Strengths above 1400 MPa this limit appears to be in the range 10 to 20 MPa-m~n. 4'2~ For the austenites of Table I, assuming a minimum value for l of 10/~m, a minimum threshold of 20 MPa-m ~:2 would be predicted for alloys above 1100 MPa yield strength. Experiments with higher strength austenites are in progress to verify this prediction. These observations are similar to the conclusions reached by Kameda on the relation between blunt notch and sharp crack effects in temper embrittlement of martensitic steels. ~ In those experiments, increases in observed fracture thresholds for sharp crack tests over those calculated from blunt notch results were related to the need to exceed the fracture stress over a reasonable microcrack nucleation distance.

Finally, the fracture model of Figure 2 suggests that in- creases in grain size should raise this limiting value thresh- old stress intensity for high strength materials. This would be true only if no segregation effects occur which would tend to lower the threshold as grain size increases. Gerberich and Wright 4 concluded that even for "clean" 4340 steels some impurity effects would be operative and thus larger grain sizes are detrimental. (The larger the grain size, the greater the concentration of grain boundary segregants.) In

1530--VOLUME 14A, JULY 1983 METALLURGICAL TRANSACTIONS A

austenites where severe temper embrittlement is not ob- served, segregation-free grain boundaries may be possible, and thus increases in residual fracture resistance in austenitic alloys could be achieved through increases in grain size.

In summary, a microfracture model based on a critical local stress for intergranular fracture has been used to ex- plain hydrogen embrittlement effects in austenitic steels. The model predicts variations in the size of a partially cracked region which is observed at crack arrest threshold. A critical stress for fracture in 200 MPa hydrogen gas at 298 K of ~2800 MPa is consistent with the data. Increases in grain size are predicted to raise the lower limiting threshold value for high strength alloys, if grain boundary segregation effects are eliminated.

This work was supported in part by DOE Contract No. DE-AC04-76 DP00789.

REFERENCES

1. W.W. Gerberich and N.R. Moody: in Fatigue Mechanisms, J.T. Fong, ed., ASTM STP 675, ASTM, Philadelphia, PA, 1979, p. 292.

2. N.R. Moody and W. W. Gerberich: Metall. Trans. A, 1980, vol. llA, pp. 973-81.

3. N.R. Moody and W. W. Gerberich: Metall. Trans. A, 1982, vol. 13A, pp. 1055-61.

4. W.W. Gerberich and A.G. Wright: Environmental Degradation of Engineering Materials in Hydrogen, M.R. Louthan, R.P. McNitt, and R. D. Sisson, eds., V. P. I. Press, Blacksburg, VA, 1981, p. 183.

5. R.E. Stoltz and A.J. West: Hydrogen Effects in Metals, I.M. Bernstein and A. W. Thompson, eds., TMS-AIME, Warrendale, PA, 1981, p. 541.

6. M.W. Perra and R. E. Stoltz: ibid., p. 645. 7. M.W. Perra: Environmental Degradation of Engineering Materials in

Hydrogen, M.R. Loutban, R.P. McNitt, and R.D. Sisson, eds., V. P. I. Press, Blacksburg, VA, 1981, p. 321.

8. J. Kameda and C. J. McMahon, Jr.: Metall. Trans. A, 1980, vol. llA, p. 91.

9. J. Kameda and C. J. McMahon, Jr.: Metall. Trans. A, 1981, vol. 12A, p. 31.

10. M. Jokl, J. Kameda, and C.J. McMahon, Jr.: J. Mat. Sci., 1980, vol. 14, p. 375.

11. J. Kameda: Metall. Trans. A, 1981, vol. 12A, p. 2039. 12. R.M. McMeeking: J. Mech. Phys. Solids, 1977, vol. 25, p. 357. 13. D.M. Tracey: J. Eng. Mater. Tech., Series H, 1976, vol. 98, p. 148. 14. K.-H. Schwalbe: J. Eng. Mater. Tech., Series H, 1977, vol. 99,

p. 186. 15. C.F. Shih, H. G. deLorenzi, and W. R. Andrews: Int. J. Fract., 1977,

vol. 13, p. 544. 16. M.W. Perra: unpublished data, Sandia National Laboratories,

Livermore, CA, 1981. 17. R. Viswanathan and S. J. Hudak, Jr.: Metall. Trans. A, 1977, vol. 8A,

p. 1633. 18. W.W. Gerberich and Y.T. Chen: Int. J. Fract., 1973, vol. 9,

p. 369-71. 19. R.W. BaUuffi, A. Brokman, and A.H. King: Acta Met., 1982,

vol. 30, p. 1453. 20. K.N. Akhurst and T.J. Baker: Metall. Trans. A, 1981, vol. 12A,

p. 1059.

METALLURGICAL TRANSACTIONS A VOLUME 14A, JULY 1983--1531