Phase relations in the hydrogen-austenite system

Phase relations in the hydrogen-austenite system

Acta metall, mater. Vol. 41, No. 7, pp. 2235-2241, 1993 Printed in Great Britain.All rights reserved 0956-7151/93$6.00+ 0.00 Copyright © 1993Pergamon...

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Acta metall, mater. Vol. 41, No. 7, pp. 2235-2241, 1993 Printed in Great Britain.All rights reserved

0956-7151/93$6.00+ 0.00 Copyright © 1993PergamonPress Ltd

PHASE RELATIONS IN THE H Y D R O G E N - A U S T E N I T E SYSTEM D. G. ULMER~" and C. J. ALTSTETTER Department of Materials Science and Engineering, University of Illinois at Urbana-Champaign, Urbana, IL 61801, U.S.A. (Received 18 August 1992)

Abstract--X-ray diffraction was used to determine phase relationships and lattice strain as a function of hydrogen concentration and temperature in 310S and 304 austenitic stainless steels. Uniform concentrations of hydrogen were introduced into 20/tm thick foils by cathodic polarization at 35, 55 and 75°C. Hydrogen remained in solid solution in the stable 310S austenite lattice up to concentrations of approximately 16 at.% at 35°C at which point a hydrogen-rich f.c.c, phase (~,*) with a 6% volume increase and a hydrogen-rich h.c.p, phase (e *) formed in the hydrogen-saturated austenite lattice. The ~,:~,* solvus curve was constructed for the hydrogen-310S stainless steel system. In the metastable 304 alloy the e* phase formed instead of V* at hydrogen concentrations above approximately 8 at.%. There was no evidence of formation of ~t'-(b.c.c.) martensite in either alloy. The partial molar volume change due to hydrogen in the 310S alloy, as determined by the X-ray diffraction technique, was measured to be At~/fl = 0.200 + 0.005 independent of the charging temperature.

INTRODUCTION Mechanisms for hydrogen embrittlement of the austenitic stainless steels have been debated for more than 30 years. One reason a mechanism has not been agreed upon is the complicated and controversial role of phase transformation in the embrittlement process. Some compositions of austenitic stainless steels are unstable with respect to formation of b.c.c. (~') and h.c.p. (E) martensite phases during plastic deformation. Hydrogen in interstitial solid solution in the austenite produces a hydrostatic lattice distortion, and at large and inhomogeneous concentrations the lattice strains can exceed the elastic limit. Thus, even with no externally applied load new dislocations can be generated and ~' and E phases can also be formed. This has raised some confusion in the literature because it has not been determined whether the product phases form due to the dissolved hydrogen per se or due to the non-uniformity in the hydrogen concentration, i.e. due to stress-induced transformation. A number of investigators have reported the formation of a metastable f.c.c., hydride-like phase (~ *) [1-4] and E*-martensite, a hydrogen-rich derivative of E [5, 6]. Too often, studies aimed at characterizing hydrogen embrittlement effects in metals have failed to quantify the concentration and distribution of hydrogen in the mechanical test specimens. Highfugacity, low-temperature charging techniques can result in severe near-surface concentration gradients. Embrittlement attributed to intrinsic effects of hydrogen may actually be an artifact of the technique used tPresent address: Rocketdyne Division, Rockwell International, 6633 Canoga Avenue, Canoga Park, CA 91303, U.S.A.

to introduce hydrogen and the consequent plastic deformation and phase transformation. As part of a larger study of hydrogen embrittlement mechanisms, we have attempted to quantify the role of phase transformation in hydrogen embrittlement of austenitic stainless steels by first determining the conditions under which the phase transformation products can exist. In the present study phase transformation in the absence of external stress was characterized as a function of uniform hydrogen concentration and temperature in alloys 310S and 304, two different compositions of austenitic stainless steel alloy with different austenite stabilities. The stress-free terminal solubility curves in hydrogen-metal alloys were determined by the so-called parametric method using X-ray diffraction [7]. In the single phase region, lattice parameter change vs hydrogen concentration was measured in homogeneous alloys to determine the partial molar volume or the relative volume expansion of the austenite lattice due to the solution of interstitial hydrogen atoms. Elastic distortion of the crystal lattice by interstitial hydrogen is an important quantity to measure, because it is a measure of the strength of the elastic interaction between hydrogen and internal stress fields, e.g. dislocations and crack tips, and thus may play an important role in understanding the hydrogen embrittlement of metals. Elsewhere, mechanical properties of the same 310S and 304 material were reported as a function of hydrogen concentration, and correlation of phase transformation with embrittlement was discussed [8]. Early work [9-11] attributed susceptibility of the austenitic stainless steels to a tendency for ~'martensite formation. It was known at the time that

2235

2236

ULMER and ALTSTETTER:

PHASE RELATIONS IN HYDROGEN-AUSTENITE

martensitic steels, particularly high strength steels, could be severely embrittled by hydrogen. The metastable austenitic stainless steels ( < 1 4 w t % N i ) ; that is, those which were susceptible to deformationinduced ~t' martensite formation, were found to be embrittled when tested in a high-pressure (70 MPa) gaseous hydrogen environment at ambient temperature [9, 10] or following pre-charging in which hydrogen was introduced eleetrolytically in a molten salt bath at elevated temperature [11]. On the other hand, stable austenitic alloys containing 14-35wt%Ni, which were not susceptible to =' martensite formation, displayed no significant tendency to embrittlement. The amounts of hydrogen introduced in the austenitic stainless steels by high-pressure gas or molten salt bath charging were relatively low, typically on the order of 50 wt ppm (0.25 at%). At higher hydrogen concentrations it was found that stable austenitic stainless steel could be embrittled even if the material was fully austenitic after straining to fracture [6, 12-15]. It is generally accepted that the b.c.c. =' martensite is not necessary for embrittlement. It has been proposed that the critical role of stress-induced c<' martensite in changing crack growth kinetics is due to its effect on hydrogen transport rather than on the crack advance mechanism [16]. Inoue et al. [17] have presented the strongest evidence for hydrogen embrittlement of austenitic stainless steels due to hydrogen-induced formation of E-martensite. Transmission electron microscopy of thinned areas of tensile specimens of SUS 304 and SUS 310 following cathodic charging concurrently with tensile testing showed that propagation of secondary cracks could occur along the interface between the austenite and E-martensite in both alloys. Unfortunately, the conditions under which the E- or E*-martensite formed in each alloy were not well defined. Local hydrogen concentrations in the thinned areas were not quantified, and the local stress state due to mechanical straining and hydrogen concentration gradients was not known. 30.0

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current

density,

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Fig. 1. Relations between applied current density and of hydrogen introduced into 310S stainless steel specimens charged at 75°C in 0.5 M H2SO4 containing 0.25 g/I NaAsO2.

amount

Detection of the 7" phase in mechanical test specimens has been difficult because of the unstable nature of this hydride-like phase, which reverts to 7 upon outgassing of hydrogen. Although no direct evidence exists, investigators have made arguments for the involvement of y* in the hydrogen embrittlement or stress corrosion cracking (SCC) of the austenitic stainless steels, based on observations of the habit planes of surface cracks formed after hydride formation and decomposition and SCC fracture facets. The y* phase has been observed to form in austenitic stainless steels as a result of high-fugacity, cathodic charging [1-4]. The question is whether or not the hydrogen concentrations/stress state existing under these severe charging conditions can exist at a crack tip, for example, in metals in gaseous hydrogen or stress corrosion service environments in which embrittlement is observed. EXPERIMENTAL

Specimen preparation Commercially produced 310S and 304 austenitic stainless steels, received as cold-rolled foil 51 and 38 # m in thickness, respectively, were annealed for 15min at II00°C in a vacuum of approximately 10 -4 Pa, followed by a water quench. Annealed pieces were electrolytically thinned at room temperature in a stirred 40vol.%HNO3 + 60vol.%HESO 4 solution to a final thickness of 18-22pm. A relatively low anodic current density of 10mA/cm 2 produced a controlled material removal rate of approximately 0.25 # m per min, which enabled specimen thickness uniformity to be maintained within +0.5/~m. The grain sizes, including annealing twin boundaries, were such that there were, on the average, 2-4 grains and 3-5 grains across the final thickness of the 310S and 304 specimens, respectively. Alloy compositions supplied by the manufacturers are shown in Table 1.

Cathodic charging

25.0 +

-o ~

Table 1. Chemicalcompositionof the austeniticalloys(wt%) Alloy C Mn P S Si Cr Ni Mo Cu Fe 3105 0.018 1.77 0.022 0.001 0.67 25.37 19.37 0.40 0.17 Bal. 304 0.51 1.33 0.23 0.012 0.44 18.55 9.51 0.20 0.24 Bal.

Electrolytic charging was used to introduce the large amounts of hydrogen necessary to embrittle the stable austenitic stainless steels. Thin foils, relatively low current densities, elevated charging temperatures and long charging times were used to minimize concentration gradients and plastic strains that can result from the high-fugacity electrolytic charging technique. Stainless steel pieces 2 0 # m thick were cathodically charged in a 0.5 M H2SO 4 + 0.25 g/l NaAsO2 solution at temperatures of 35, 55 and 75°C. Charging times were 60, 18 and 6 h, respectively. These were the times calculated to reach 98 + % of the surface concentration at the center of the specimen using the mathematical solution for diffusion in

ULMER and ALTSTETTER: PHASE RELATIONS IN HYDROGEN-AUSTENITE a plane sheet for equal and constant surface concentrations [18] and the values for hydrogen diffusivity in austenitic stainless steel [19]. At first a constant potential was applied to the specimen using a Luggin probe/saturated calomel electrode positioned a few millimeters from the sample surface. The current density decreased rapidly, reaching a stable value within an hour or so. This stable region was maintained to the end of the charging process by switching from the potentiostatic to the galvanostatic mode. It was found that this electrolytic charging procedure gave the most homogenous distribution of hydrogen within the specimen [20]. Two platinum foil anodes, 50 × 50 mm, were positioned approximately 3 cm on either side of the 25 x 38 mm stainless steel specimen. All electrical connections were made by spot welding 0.25mm diameter platinum wire to the anodes and cathode. Solution temperature in the cell was regulated by circulating the electrolyte through a heated reservoir that was maintained within + 1°C with a temperature controller. A 2-3 s ultrasonic burst was applied to the electrolytic cell every 30 s to clear the specimen surface of any H 2 bubbles which might have interfered with the charging process, and nitrogen gas was continuously bubbled through the electrolyte from a porous frit to remove dissolved oxygen. The specimen was supported by attaching the corners to a mylar frame with acid resistant paint. If careful attention was paid to cleanliness, reproducible results were obtained. All glassware was thoroughly cleaned, washed and rinsed in fresh 5 vol.%H2SO4 solution and deionized water. Prior to each experiment the platinum anodes were cleaned by electrolytic pickling for 10min in concentrated nitric acid, rinsed with deionized water, cleaned by cathodic polarization at 3 A for 5 min in 5 vol.%H2SO4 and rinsed again in deionized water prior to charging [21]. The hydrogen concentration dissolved in the austenite lattice as a result of cathodic polarization was found to increase linearly with increasing current density, shown for example in Fig. 1 for 310S charged at 75°C. The scatter in the data represent the range in hydrogen contents measured for individual specimens cut from a single 2.5 x 3.8 cm foil. The scatter probably indicates the precision of the hydrogen analysis technique rather than true variation of hydrogen concentration in the foils. Note that the data deviate from linearity at current densities above about 1.0 mA/cm 2, which coincided with the onset of H 2 bubble formation on the specimen surfaces. Note also that very low current densities, on the order of a 100/~A/cm2, were able to introduce rather large ( > 1 at.%) hydrogen concentrations. Cathodic current densities of this magnitude are typical of those which can exist at stress-corrosion crack tips even under small applied anodic potentials [22]. Following charging, specimens were removed from the solution, rinsed with deionized water and placed

2237

in liquid N 2 within 15 s to prevent hydrogen loss until ready for X-ray diffraction studies. Following X-ray diffraction measurements, hydrogen concentrations were measured by vacuum hot-extraction and volumetric measurement of the hydrogen gas. The hydrogen analyzer was calibrated regularly using a NIST standard Ti-H alloy. X - r a y diffraction

X-ray diffraction is a powerful technique for nearsurface determination of volume change, phase transformation and concentration/strain gradients [23] produced by diffusion of hydrogen into metals. The combination of relatively high H 2 solubility and low H 2 diffusivity make this technique particularly useful when studying the austenitic stainless steels because of the relatively large lattice strains and because the X-ray penetration distance for Cu-K, radiation is on the order of the effective diffusion distance of hydrogen. Atomic volume

The angular position of X-ray diffraction peaks, 20, are related to interplanar spacing, d, and lattice parameter, a, for cubic materials by the simple Bragg relation 2a

2 = 2d sin 0

sin 0

(1)

4 h 2 + k2.-k l 2

where 2 is the wavelength of the radiation and hkl are the Miller indices for each reflection. Hydrogen occupies the octahedral interstitial sites in f.c.c. austenitic stainless steels and produces a spherical distortion of the lattice. Peisl [24] derived the following linear relationship between the lattice strain, AV/V, due to interstitial hydrogen and the hydrogen/ metal atomic ratio, Cu A V _ a~ - a 3 = CHAV/f~

V

a3

(2)

where an and a0 are the lattice parameters of the hydrogen-charged and hydrogen-free austenite, respectively, Av is the volume change of the lattice per hydrogen atom and f~ is the mean atomic volume of a metal atom. If Vegard's law is followed, a plot of relative volume change, A V / V , as determined from X-ray diffraction, vs hydrogen concentration will result in a linear relationship with the slope of the line equal to the quantity Av/l), which is also referred to as the partial molar volume of hydrogen. In this work filtered Cu-K, radiation was used, and values of an and a0 were calculated for up to five austenite reflections; namely 111, 220, 220, 331 and 331, when possible, in order to minimize uncertainty in an or a0 due to systematic errors in 20 caused by misalignment of the instrument and/or displacement of the specimen from the diffractometer axis. Precise values for an and a0 were obtained by plotting measured values of aH or a 0 for each reflection as a function of cos20 and extrapolating to cos20 = 0 using

2238

ULMER and ALTSTETTER:

PHASE RELATIONS IN HYDROGEN-AUSTENITE

a least squares fit [7]. It is estimated that lattice parameter values obtained using this extrapolation technique were accurate to within Aa = _0.002/~. All X-ray diffraction work was conducted at - 8 0 ° C to minimize hydrogen loss from the surface. Samples were cooled to - 8 0 ° C within a minute by a steady stream of N 2 gas at 3 psig passed through a copper tube immersed in a flask containing liquid N:. An annealed, hydrogen-free specimen was scanned before each hydrogen charged specimen was examined in order to check alignment and to establish the baseline value, a0.

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Two techniques utilizing X-ray diffraction were used to locate phase field boundaries. The disappearing-phase method is based on determination of the composition at which the second phase disappears from a series of two phase alloys. The accuracy of this method is limited by the sensitivity of X-ray diffraction in detecting small amounts of the second phase. The parametric method, the principal method used here, is based on observations of the solid solution itself. This method depends on the fact that the lattice parameter of a solid solution changes with alloy composition up to the saturation limit and then remains constant beyond that point [7].

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RESULTS AND DISCUSSION

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Alloy 310S Portions of the X-ray diffraction patterns for 20/tm thick 310S austenitic stainless steel foils are shown as a function of hydrogen content following cathodic charging for 6 h at 75°C, Fig. 2(a) and for 60 h at 35°C, Fig. 2(b). The maximum information was obtained from the diffraction profiles in the vicinity of the 111 reflection, shown in Fig. 2. The shift of the 111 peak to lower 20 values with increasing hydrogen content is due to dilatation of the lattice by interstitial hydrogen. Remarkably, the 310S austenite lattice was able to retain as much as 25 at.% hydrogen in interstitial solid solution at 75°C. Above 25at.%, the austenite partially transformed to a hydrogen-rich h.c.p, phase (E*) which has been identified by Holzworth and Louthan [25] and Rigsbee [5] to be the same martensitic phase that can be formed by cold working the less stable austenitic stainless steels such as type 304. At lower temperatures, the 310S austenite had a tendency to partially transform to a hydrogen-rich f.c.c, phase (y*) which occupied a volume approximately 6% greater than the expanded f.c.c, matrix (Ye). At 35°C the ?e-7* phase separation occurred at hydrogen concentrations above approximately 16 at.%, shown clearly in Fig. 2(b). The y* phase has been reported by a number of investigators [1-4]. It has been suggested that the y* phase is a metastable hydride similar to that formed in pure nickel [26] and iron-nickel alloys [27] at high hydrogen concentrations. Indeed, the y* phase exhibited

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28 Fig. 2. Shift and broadening of 111 X-ray diffraction peak in 310S with increasing hydrogen content. Cathodic charging at (a) 75°C, and (b) 35°C.

ULMER and ALTSTETTER: PHASE RELATIONS IN HYDROGEN-AUSTENITE some characteristics typical of a metastable hydride. In the 310S austenite, the V* transformation was found to be fully reversible, reverting first to 7~ and then to ~, on outgassing. When the ~ and ),* coexisted, the positions of the ~, and y* X-ray diffraction peaks remained stationary with increasing hydrogen concentration, but the intensity of the V* peak grew at the expense of the V=matrix reflection. This type of behavior is typical of a miscibility gap system, such as observed in the N i - H [26] or P d - H [24] systems. The ~ * peak, on the other hand, continued to shift to lower 20 values with increasing amounts of hydrogen. In the 310S alloy the y*-hydride phase transformation was always accompanied by ~*-martensite formation. It is possible that at these temperatures the E*-martensite formed by faulting of the austenite to accommodate the lattice strains induced by the ?*, which occupied a 6% larger volume than the solid solution. On the other hand, at 75°C, at which temperature the ),* phase did not form, it is possible that the E*-martensite is the thermodynamically more stable phase in 310S. Large broadening of the diffraction peaks was observed, Fig. 2. This broadening was likely caused by random lattice atom displacements

(a)

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Fig. 4. 310S-H pseudo binary phase diagram showing 7e-? * solvus in temperature range from 35 to 75°C. Solvus shown by solid line was constructed by parametric method using X-ray diffraction. Open symbols indicate 2:e solid solution. Filled symbols indicate coexistence of 2:e and 2:* and E*. Half-filled symbol indicates coexistence of 2:eand E*. due to hydrogen, with contributions from dislocations and stacking faults. The linear relationships between the relative volume change A V / V of the 310S austenite solid solution charged at 75 and 35°C, and the atomic fraction of hydrogen, 6",, are shown in Fig. 3(a) and (b), respectively. The solid lines are least squares fits to the austenite solid solution data. From equation (2) the slope of these lines was equal to the partial molar volume, Avlf~, of hydrogen in 310S austenitic stainless steel and was found to be independent of cathodic charging temperature. An average value of the partial molar volume in the 310S was calculated to be Av/l] = 0.200 + 0.005 which is approximately 20% lower than the value predicted by Baranowski et al. [28] for hydrogen in f.c.c, metals. Assuming that there was not significant time for hydrogen segregation or rearrangement on cooling to - 80°C for the X-ray diffraction measurement, the temperature independence of Av/f~ in the 310S austenitic stainless steel suggests that the amount of hydrogen partitioned during charging to trapping sites such as grain boundaries, second phase particles, etc. was an insignificant fraction of the total amount of hydrogen in the specimen, i.e. segregation effects were not significant in the austenitic stainless steels at these large hydrogen concentration levels. The diffraction peaks in Fig. 2 used to illustrate peak shift and phase separation were also used to calculate lattice strains of the solid solution, and the concentrations are indicated in Fig. 3. The solid points in Fig. 3 indicate muitiphase compositions, and the dashed line in Fig. 3(b) indicates that the concentrations of hydrogen in the phases in the two-phase region are essentially constant.

o.oo 0.0

0.1

0.2

0.3

0.4

h y d r ~ e n l m e t a l atom ratio, ~ Fig. 3. Lattice strains in 310S austenite, 2:e,charged at (a) 75°C, and (b) 35°C. Concentrations shown in Fig. 2 are indicated. Filled symbols indicate coexistence of 2:, and 2:* and/or E*.

310S Miscibility gap The parametric method was used to define the stress-free solvus line between the ?e and Ye+ ?* phase fields for the hydrogen-310S austenitic stainless steel binary system. This boundary is shown by the solid

2240

ULMER and ALTSTETTER:

PHASE RELATIONS IN HYDROGEN-AUSTENITE

curve drawn in Fig. 4. The open circles indicate the hydrogen concentration levels at 35, 55 and 75°C at which only a single solid solution was observed. The filled circles indicate the hydrogen concentration levels at which the ),*-hydride coexisted with the ),-austenite. At 35°C, the phase separation occurred at approximately 16at.%, and at 55°C, the phase separation occurred at approximately 18 at.%. The half-filled circle indicates that at 75°C the ),*-hydride phase did not form; the E*-martensite phase appears to be the more thermodynamically stable phase at that temperature. The dashed curve bordering the right side of the two phase field was estimated assuming that Av/f~ of hydrogen in the ),*-hydride was equal to that in the austenite. This was done by extending the partial molar volume curve in Fig. 3(b) until it intersected the relative volume change of the ),*-hydride determined by X-ray diffraction (not shown). For example, at 35°C this would correspond to approximately 33 at.% hydrogen, CH = 0.5, or the formula M2H. This is the concentration expected when half of the octahedral interstitial sites are filled. This result is a good indication that )'* may be an ordered hydride phase. The value of Ca = 0.5 is somewhat less than that found in Ni(CH = 0.7) but about the same as that in a Ni-10 Cu alloy [30]. Alloy 304

Electrolytic charging of the 304 austenitic stainless steel alloy was conducted at 55°C only. The lattice dilatation due to hydrogen is shown in Fig. 5. The point of phase separation in the 304 alloy was not well established. It was found that hydrogen remained in interstitial solid solution up to concentrations between 6 and 10 at.%, above which the 304 partially transformed to e*-martensite. There was some indication that the y* phase formed at hydrogen concentrations above approximately 13 at.%, but identification of diffraction peaks became extremely difficult because of the rapid broadening of the more easily faulted 304 austenite lattice. No evidence of ct'-martensite formation was found within the limits

at.% hydrogen >

0.0

5.0

16.0

15.0

20.0

25.0

I

I

I

i

i

0.08

304 stainless steel Charging temperature:

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CONCLUSIONS 1. The hydrogen solubility limit as measured by X-ray diffraction in the high-nickel (20 wt%Ni) 310S austenitic stainless steel was measured from 16 at.% to as high as 25 at.%H in the temperature range from 35 to 75°C. Under identical electrolytic charging conditions the hydrogen solubility in the low-nickel (10wt%Ni) 304 alloy extended to only approximately 8 at.%H at 55°C. 2. In the single phase region, the relative volume expansion of the 310S austenite lattice due to solution of one interstitial hydrogen atom was measured to be 20% of that for the addition of one metal atom. 3. Once the solubility limit was exceeded, the )'*-hydride and e*-martensite phases formed in the 310S austenite solid solution, while the E*-martensite was the phase formed in the 304 matrix. 4. The hydrogen/metal atom ratio in 7" was estimated to be 0.5, indicating that )'* is an ordered hydride phase. Acknowledgements--The authors would like to thank the Allegheny-Ludlum Steel Company and the TeledyneRodney Metals Company for supplying the stainless steel materials for this work. The authors would also like to thank Dr Peter Flynn and the personnel of the Center for Microanalysis of Materials for their assistance. This work was supported by the Department of Energy under Contract No. EY-76-C-02-1198 through the Materials Research Laboratory at the University of Illinois.

REFERENCES

55°C

0.02

of detection of the X-ray diffraction technique ( ~ 1 vol.%). Three important differences in behavior were observed between the 310S and 304 alloys. First, the maximum hydrogen solubility without phase transformation was found to increase with increasing nickel content, 20 wt%Ni (310S) vs 10 wt%Ni (304), likely a reflection of the higher stacking fault energy and increased austenite stability of the higher Ni alloys [29]. Second, once the hydrogen solubility limits were exceeded, the E*-martensite was the phase that formed in the 304 alloy while both the )'*-hydride and E*-martensite formed in the 310S alloy. Third, the partial molar volume of hydrogen in 304 was somewhat lower (0.19) than in 310S (0.20).

0.4

h y d r ~ e n / m e t a l atom ratio, H/M Fig. 5. Lattice strains in 304 austenite, )'e,charged at 55°C. Filled symbols indicate coexistence of ?= and E*.

1. H. Mathias, Y. Katz and S. Nadiv, Metal Sci. 12, 129 (1978). 2. N. Narita, C. J. Altstetter and H. K. Birnbaum, Metall. Trans. 13A, 1355 (1982). 3. K. Kamachi, Trans. ISIJ 18, 485 (1978). 4. A. Szummer and A. Janko, Corrosion 35, 461 (1979). 5. J. M. Rigsbee, Metallography 11, 493 (1978). 6. M. L. Holzworth, Corrosion 25, 107 (1969). 7. B. D. Cullity, Elements o f X-Ray Diffraction, 2nd edn, p. 379. Addison-Wesley, Reading (1978). 8. D. G. Ulmer and C. J. Altstetter, Acta metall, mater. 39, 1237 (1991). 9. R. M. Vennett and G. S. Ansell, Trans. Am. Soc. Metals 60, 242 (1967). 10. R. B. Benson Jr, R. K. Dann and L. W. Roberts Jr, Trans. T M S - A I M E 242, 2199 (1968).

ULMER and ALTSTETTER:

PHASE RELATIONS IN HYDROGEN-AUSTENITE

11. R. Lagneborg, J. Iron Steel Inst. 207, 363 (1969). 12. M. B. Whiteman and A. R. Troiano, Corrosion 21, 53 (1965). 13. J. Kolts, Stress-Corrosion--New Approaches, ASTM STP 610, pp. 366-380. American Society for Testing and Materials (1976). 14. H. Hanninen and T. Hakkarainen, Metall. Trans. 10A, 1196 (1979). 15. J. H. Huang and C. J. Altstetter, Metall. Trans. 22A, 2605 (1991). 16. T. P. Perng and C. J. Altstetter, Metall. Trans. 18A, 123 (1987). 17. A. Inoue, Y. Hosoya and T. Masumoto, Trans. ISIJ 19, 170 (1979). 18. J. Crank, The Mathematics of Diffusion, p. 45. Clarendon Press, Oxford (1956). 19. T. P. Perng and C. J. Altstetter, Acta metall. 34, 1771 (1986). 20. M. Jeanneret, Master's thesis, Univ. of Illinois at Urbana-Champaign (1988). 21. R. J. Cyle, A. Atrens, N. F. Fiore, J. J. Bellina and M. Jolles, Environment-Sensitive Fracture of Engineering

22. 23. 24. 25. 26. 27. 28. 29. 30.

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