Tensile flow and fracture behaviour of austenitic stainless steels after thermal aging in a hydrogen atmosphere

Materials Science and Engineering, 67 (1984) 91-107

91

Tensile Flow and Fracture Behaviour of Austenitic Stainless Steels after Thermal Aging in a Hydrogen Atmosphere Y. ROSENTHAL, M. MARK-MARKOWITCH and A. STERN Nuclear Research Centre, Negev, P.O. Box 9001, Beer-Sheva (Israel) D. ELIEZER Department of Materials Engineering, Ben Gurion University of the Negev, P.O. Box 653, Beer-Sheva 84 105 (Israel) (Received May 16, 1983; in revised form February 23, 1984)

SUMMARY

1. INTRODUCTION

A hydrogen content o f about 300 at.ppm introduced by concurrent sensitization and precharging from the gas phase at 600 °C o f type 3 0 4 L and 3 1 6 L austenitic stainless steels induced a noticeable degradation of the mechanical properties, changes in fracture mode or morphology and an increase in the a m o u n t o f deformation martensite. The various changes in properties were revealed on posttreatment tensile testing in air at ambient temperature, scanning electron microscopy fractography and X-ray diffractometry. The degradation o f the mechanical properties was described quantitatively by a detailed examination o f the entire range of uniform deformation. The relevant effects were (a) ductility losses o f the order o f 30%, (b) flow stress increases (hydrogen hardening) o f the order o f from 7% (type 3 0 4 L steel) to 10% (type 3 1 6 L steel) at moderate and large strains and (c) a decrease in the strain-hardening capacity o f the order o f 30% at moderate strains. The yield stress and ultimate tensile strength o f both steels were little affected by hydrogenation and were thus shown to be o f doubtful value as criteria o f hydrogen effects when examined o u t of the context o f overall plastic behaviour. Strain-induced martensite is suggested to be the dominant factor in the behaviour o f the type 3 0 4 L steel. The large susceptibility to hydrogen effects shown by the type 3 1 6 L steel is discussed in terms o f nickel content, surface and geometry effects and also void growth.

Blanchard and Troiano [1], Whiteman and Troiano [2], Holzworth [3], Kolts [4], Eliezer et al. [5], H~mninen and Hakkarainen [6], Narita et al. [7] and others have shown over the last two decades that considerable hydrogen embrittlement or hydrogen-induced degradation of properties may be produced in any austenitic stainless steel, however stable with respect to deformation martensite, provided that two conditions were met: (1) a supply of environmental hydrogen at high input fugacities; (2) a build-up of large concentrations of internal hydrogen across a significant part of the specimen's cross section rather than in some shallow subsurface layer. These results, while essentially correct, apparently led to the still popular view that high charging fugacities or internal hydrogen contents are a prerequisite to considerable hydrogen effects in austenitic stainless steels (see for example ref. 8). Thus, relatively little interest was left over for the investigation of hydrogen effects in austenitic stainless steels at low input fugacities or low internal hydrogen contents. Asano and Otsuka [9] and Oriani [10, 11] regarded charging damage induced by high input fugacities as a serious source of the inconsistency and confusion in much of the existing literature on hydrogen embrittlement in various materials. According to Burke et al. [ 12], the intrinsic hydrogen-metal interactions are likely to be obscured by extraneous effects when austenitic stainless steels

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92

are subjected to high fugacity cathodic charging. Such effects have actually been observed and are as follows: expansion of the austenite lattice [7, 13, 14] and build-up of internal stresses in a subsurface layer [5, 7, 15, 16] where the hydrogen content may amount to tens of atomic per cent [7, 8], surface and subsurface cracking [5, 6, 16-18], partial martensite transformation above Ms (the temperature at which the transformation of austenite to martensite starts during cooling) without any externally applied stress [7, 14, 15, 1719], increase in the dislocation and stacking fault density with increasing hydrogen content [15], dislocation motion without an externally applied stress [ 15], microvoid expansion [20] and, possibly, formation of unstable hydrides or hydride-like phases [13, 21-24]. It is worthwhile to note that relatively "mild" cathodic precharging of type 304L and 310 austenitic stainless steels did not prevent some transformation from 7 to e in both steels [12]. In contrast, Thompson and Buck [25] have shown that thermal precharging of type 304L steel (from the hydrogen gas phase) to a low internal hydrogen content (400-450 at.ppm) induced no charging damage and martensite formation. The absence of extraneous charging effects should facilitate investigations of intrinsic hydrogen effects on the properties of austenitic stainless steels with homogeneous (albeit low) internal hydrogen contents. The few results in the literature pertinent to the effects of low to moderate internal hydrogen contents or low input fugacities in austenitic stainless steels are quite intriguing. Bressanelli and Moskowitz [26] observed a noticeable loss of tensile ductility when sheet specimens of type 301 austenitic stainless steel that had been thermally precharged to an internal hydrogen content of about 160 at.ppm were tested in air at room temperature; they found a transverse instead of a shear fracture, accompanied by a loss of total elongation of nearly 32%. The 0.2 yield stress and the ultimate tensile strength apparently were not affected. Hoffmann and Rauls [27] tested DIN (Deutsche Industrie Normen) austenitic stainless steels X5CrNiTi 1811, X5CrNiVMo 1910 and X5CrNi 1811 in hydrogen gas at a~pressure* of 15 MPa and room temperature; they measured losses of reduc-

tion in area of about 5%, 22% and 33% in the respective steels in comparison with tests conducted in air. Walter and Chandler [29] observed a 21% loss of total elongation on testing type 302 steel under a hydrogen gas pressure as low as 0.1 MPa; the ductility losses amounted to approximately 29% for total elongation and 17% for reduction in area as the hydrogen pressure was raised to 0.7 MPa. Hydrogen effects on mechanical properties have sometimes been reported to be larger in thermally precharged austenitic stainless steels than in the same materials when tested in high pressure hydrogen gas. Thus, Vandervoort [30] found no ductility losses or other changes in the mechanical properties of type 21-6-9 austenitic stainless steel specimens exposed to static loading for 300 h in hydrogen gas at 69 MPa and room temperature. Louthan et al. [31], however, reported that the same material, when thermally precharged in hydrogen gas at 69 MPa and 200 °C for about 226 h and subsequently tested in air at room temperature, showed significant ductility losses, approximately 55% for total elongation and 58% for reduction in area. The internal hydrogen content was estimated to be roughly 2300 at.ppm. Further, Thompson [32] reported that type 21-6-9 steel thermally precharged to about 4000 at.ppm H and subsequently tested in air at room temperature showed a loss of reduction in area of approximately 48%; uncharged material pulled in tension in hydrogen gas at 69 MPa and room temperature showed a negligible loss of ductility, as reported in ref. 30. Results qualitatively similar to the above were obtained by Thompson [32, 33] with type 309S and A-286 austenitic stainless steels. Annealed type 304L steel, either thermally precharged with hydrogen (approximately 1500 at.ppm [31] or 4000 at.ppm [32]) and tested in air, or uncharged and tested in hydrogen gas at 69 MPa showed about the same loss of reduction in area, 5~%-59%. Tlae loss of total elongation, however, was 55% for precharged material as against only 35% for uncharged specimens tested in high pressure hydrogen [31]. Thus, while the ductility dependence on the strain rate sensitivity (reduction in area) was * T h e pressures " e q u i v a l e n t " to fugacities typical of cathodic charging are o f t h e o r d e r of 105-107 MPa

[281.

93

equally affected by the two different hydrogen exposures, the dependence on strain hardening (elongation) was affected more by the internal hydrogen than by the environmental hydrogen. The results of Donovan and coworkers [31, 34] revealed a strong influence of plastic strain on the rate of tritium release during the tensile testing of thermally precharged austenitic stainless steels. Few attempts, however, were made using this approach, i.e. a detailed investigation of the flow and strain hardening behaviour of austenitic stainless steels in conditions of hydrogen embrittlement or which produce a hydrogen-induced degradation of properties [12]. Capeletti and Louthan [35] proposed an empirical correlation of yield strength with fracture strain for several austenitic stainless steels tested in tension in hydrogen gas at 69 MPa and room temperature. While such correlation may be useful as an engineering rule of thumb, it has no clear mechanistic foundation. Caskey [36] observed that deuterium thermally precharged in type 304L steel (about 2900 at.ppm) increased the grain size dependence of the flow stress but that this dependence was unaffected in hydrogen-free specimens pulled in hydrogen gas at 69 MPa and room temperature. The significance of flow and strainhardening factors in the hydrogen embrittlement or hydrogen-induced degradation of properties is further emphasized by the few results published so far with respect to type 316 steel. According to Louthan and coworkers [31, 35], type 316 steel pulled in tension in hydrogen gas at 69 MPa and room temperature showed neither ductility losses nor any significant changes in yield strength or ultimate tensile strength (5.6% increase). Similarly, Rosenthal et al. [37] detected no changes in yield strength, ultimate tensile strength and uniform elongation in type 316L specimens precharged in hydrogen gas at 0.5 MPa and 600 °C and subsequently pulled in tension in air at room temperature. However, they observed small but definite hardening effects and a decrease in the strain-hardening exponent over the strain range 5 × 10 -35 × 10 -~- (Fig. 1), as well as surface cracking in the necked region. A hydrogen content of about 270 at.ppm was measured. The results of Caskey [36] and Rosenthal e t al. [37] suggest that a loss of ductility,

Mpe)

50E

~ ~ ~.~.~ ~ 0 ~

20(]

!l,,,i

5

10- 2

I

2

I I il,,,I 5

10-1

,

2

I L Ill 5

£p

Fig. 1. log-log stress-strain (Hollomon) curves of type 316L steel precharged in hydrogen (0.5 MPa; 600 °C) ( ) and thermally aged in argon (0.5 MPa; 600 °C) ( - - -) [37].

however important in engineering applications, may not always be the sole or main criterion of hydrogen embrittlement or hydrogen-induced degradation of properties. A quantitative investigation of the uniform flow and strain-hardening stages preceding plastic instability and fracture of hydrogenated material might be more revealing as to intrinsic hydrogen effects than the few discrete parameters usually reported, i.e. fracture strain, yield stress and ultimate tensile strength. Further, if very little is known about the possible effects of low internal hydrogen contents on the properties of austenitic stainless steels in the annealed (or annealed plus coldworked) condition, even less is known with respect to the response of these steels when they are mildly hydrogenated in the thermally aged condition. The thermal aging of austenitic stainless steels in the temperature range 600-800 °C leads to the precipitation of (Fe, Cr)23C61 mostly as grain boundaries and to a depletion of chromium around the precipitates. This phenomenon is often referred to as "sensitization" because of an increased susceptibility to intergranular corrosion. The few studies of hydrogen embrittlement or hydrogen-induced degradation of properties in sensitized austenitic stainless steels have been concerned almost exclusively with type 304 [38-42] and type 304L [41] steels.

94 The main result of these studies is that these steels fail intergranularly on exposure to environmental hydrogen. The explanation suggested by Briant [43] was that chromium and carbon depletion at grain boundaries allowed the depleted regions to transform to martensite on deformation, thus providing an easy fracture path along the grain boundaries. According to the few results published so far, sensitization apparently has little effect on the resistance to hydrogen of austenitic stainless steels with a higher degree of alloying, particularly those containing nickel. For steels with nickel contents in the range 8-14 wt.%, there is a marked improvement in the resistance to property degradation by internal hydrogen [44, 45] ; there is also an increase in austenite stability [45]. This would explain Thompson's results [ 33 ] for sensitized and unsensitized type 309S steel (14 wt.% Ni); in the presence of a b o u t 425 at.ppm of internal hydrogen introduced by thermal precharging, (a) similar small losses of percentage reduction in area and the same t y p e of dimpled ductile fracture were obtained for both material conditions and (b) no strain-induced martensite was detected in either condition. Briant [46] obtained similar results with type 316 steel (12.8 wt.% Ni); sensitization had very little effect on the resistance to hydrogen effects.

2. PURPOSE OF PRESENT STUDY An investigation of the effects of low internal hydrogen contents (of the order of a few hundreds of atomic parts per million) on the following properties of selected austenitic stainless steels in the thermally aged (sensitized) condition: (1) plastic flow, strainhardening, ductility and fracture characteristics; (2) microstructural stability, i.e. extent of strain-induced martensite formation and possible changes in Md (the temperature at which martensite forms from strained austenite). Md for our particle grade of t y p e 304L steel (with 0.04 wt.% N) was calculated using Angel's formula [47 ], with the result that Md(304L) ~ 340 K. Angel's formula yielded for our t y p e 316L steel (with 0.16 wt.% N) that Md(316L) ~ 290 K. In order to ensure a homogeneous distribution of hydrogen within

the specimens with no surface damage and consequently a bulk response of the steels, thermal precharging from the hydrogen gas phase was chosen as the hydrogenation technique.

3. EXPERIMENTAL DETAILS

3.1. Materials and specimens The two commercial-grade austenitic stainless steels were received as sheets 0.2 mm thick in the bright-annealed condition. Bright annealing of both type 304L and type 316L sheets was performed at the manufacturer's plant according to current industrial practice: 1 9 0 0 - 2 0 0 0 °F ( 1 0 4 0 - 1 0 9 5 °C), in an atmosphere of dissociated ammonia. Residual hydrogen content was below 20 at.ppm for both steels. Table 1 shows the chemical analysis of the steels, while Table 2 summarizes their tensile and microstructural characteristics. Tensile specimens with a nominal gauge length of 50 mm, a gauge width of 6.25 mm and a shoulder width of 25.4 mm were carefully milled on a Tensilkut machine (Sieburg Industries), with the tensile axis parallel to the rolling direction of the original sheet.

3.2. Thermal precharging of hydrogen Hydrogen gas precharging of tensile specimens was conducted at 600 °C (873 K) in vacuum-tight austenitic stainless steel retorts equipped with thermocouple ports and inletoutlet vacuum valves. The retorts containing the specimens were p u m p e d off, filled with high purity hydrogen gas and then heated using resistance-wound furnaces with the temperature controlled to within + 1 °C. The gauge hydrogen pressure in the retorts was kept constant at 0.5 MPa for an exposure duration of 170 h. Control specimens were heated in an argon atmosphere, in temperaturepressure-time conditions identical with those of the hydrogen precharging. After completion of the 600 °C exposures to hydrogen or argon, the furnaces were moved away from the retorts and the retorts were left to cool in air to ambient temperature. According to the earlier techniques, hot extraction was carried out at 900 °C with argon as the carrier gas; hydrogen was determined as such by gas chromatography. In the improved technique,

95 TABLE 1 Chemical analysis of type 304L and 316L austenitic stainless steels Steel

A m o u n t (wt.%) o f following elements

304L 316L

C

Cr

Ni

Mo

Mn

Si

Fe

0.036 -+0.001 0.029 + 0.006

18.5 17.7

7.8 10.2

-1.6

1.5 1.4

0.8 0.6

Balance Balance

TABLE 2 Tensile and microstructural properties of type 304L and 316L austenitic stainless steels in the bright-annealed condition Steel

0.2% yield stress (MPa)

Ultimate tensile strength (MPa)

Uniform elongation (L o = 25 ram) (%)

Viekers hardness (200 g, 10 s)

Grain size (pm)

A m o u n t (vol.%) o f strain-induced martensite a

304L 316L

284.3+11.2 307.7-+15.3

672.7+10.0 641.3-+6.0

72.7+6.2 55.0-+5.2

310-+18 327-+10

56 (ASTM si ze1 2 ) 79 (ASTM size 11)

17.6 0

aX-ray diffraction analysis of fractured specimens.

hot extraction was still performed at 900 °C b u t with oxygen as the carrier gas; the sample was fully oxidized and hydrogen was determined as H20. The present technique is considered to ensure a more complete recovery of the hydrogen. All specimens, hydrogen precharged or controls, were stored under liquid nitrogen until tensile tested. The period of storage in liquid nitrogen varied from a few hours up to a maximum of 24 h. X-ray diffraction tests of previously cooled 304L and 316L specimens, either charged or uncharged (aged in argon), revealed no martensite {less than 1 vol.%). In the initial stage of the reported work a few comparative tensile tests were conducted with charged and control specimens that had n o t been stored in liquid nitrogen; the results were very similar to those obtained with cooled specimens. It is quite possible that storage in liquid nitrogen is not really necessary for specimens charged to low hydrogen concentrations. 3.3. Tensile testing

All tensile testing was conducted in air at room temperature on an Instron 1195 machine at a cross-head speed of 1 mm min -1, i.e. at a strain rate of a b o u t 0.5 X 10 -3 s-1. The exten-

sion was measured with strain gauge extensometers over an effective gauge length of 25 mm. Magnification techniques ensured a high resolution of load-extension readings, particularly at low and moderate strains. Four to six specimens were tested for each investigated condition. The load-extension readings were fed to a PDP-11 computer programmed to reduce such data to stress-strain and strainhardening parameters, and also complete flow curves. Hollomon's empirical power equation was employed as a best-fit function: O" ~- O'OCpln

where o is the true flow stress, o0 the strength coefficient, %1 the true plastic strain and n the strain-hardening exponent. With accurate high resolution load-extension readings the Hollomon equation yields, for some austenitic stainless steels, log-log flow curves consisting of three linear stages characterized by different n and ao values. Such three-stage strainhardening parameters have been used, e.g. by Ch6ne et al. [48], to describe hydrogen effects in austenitic stainless steels. Whether these stages actually represent specific micromechanistic processes is still a debatable issue. However, they can reveal genuine changes in macroscopic strain hardening provided that

96

the double-logarithmic o versus ep~ curves are plotted from a common strain origin when comparing, for example, the hydrogenated and hydrogen-free conditions of a given material. 3.4. Microstructure and fractography After completion of the thermal aging treatments, samples of the two austenitic stainless steels were subjected to grain size and low load hardness measurements and also to thin foil transmission electron microscopy (TEM). The main purpose of the TEM examination was to reveal and characterize car-

bide precipitation at grain boundaries as continuous or discontinuous. Some of the fractured specimens were subjected to X-ray diffraction in order to determine the volume fraction of strain-induced martensite. The Cu Kc~ radiation and standard procedures were employed. Because of the size of the X-rayed area of gauge face, the results were n o t characteristic of the maxim u m strain {fracture strain) in the specimens. Further, the results for the hydrogen-charged specimens and the control specimens are not directly comparable at a given distance from the fracture separation because at any loca-

TABLE 3 M i c r o s t r u c t u r a l p r o p e r t i e s o f t y p e 3 0 4 L a n d 3 1 6 L a u s t e n i t i c stainless steels in t h e t h e r m a l l y aged c o n d i t i o n

Steel, condition

Aging environment

Hydrogen content

Grain size (pro)

Vickers hardness (200 g, 10 s)

Amount (vol.%) o f strain-induced martensite a 23.9 16.5

(at.ppm) 3 0 4 L , t h e r m a l l y aged at 6 0 0 °C

H2 Ar

300-325 --

1 1 2 ( A S T M size 10) 1 1 2 ( A S T M size 10)

2 6 4 + 11 361 + 21

3 1 6 L , t h e r m a l l y aged at 6 0 0 °C

H2 Ar

300-325 --

104 ( A S T M size 10) 104 ( A S T M size 10)

3 4 3 -+ 4 341 + 7

9.8 0

aX-ray d i f f r a c t i o n analysis o f f r a c t u r e d s p e c i m e n s .

TABLE 4 E f f e c t s o f h y d r o g e n o n t h e tensile p r o p e r t i e s o f t h e r m a l l y aged 3 0 4 L a u s t e n i t i c stainless steel

Parameter (units)

S (MPa) for ep = 0 . 0 5 % S (MPa) for ep -- 0.1%

S(MPa) f o r e p = 0 . 2 % S (MPa) for S (MPa) for S (MPa) for S (MPa) for S (MPa) for S (MPa) for S (MPa) for S (MPa) for S (MPa) f o r S (MPa) for U T S (MPa) Sf (MPa) eu (%) ef (%)

ep ep ep ep ep ep ep ep ep ep

-- 0.5% =0.6% = 1% = 2% = 3.5% = 5% = 10% = 20% = 30% = 40%

.~

Value o f parameter after exposure to H 2 at 600 °C

Value o f parameter after exposure to Ar at 600 °C

H2-induced property change

2 6 6 . 5 -+ 2.8 2 6 9 . 3 -+ 2.5 273.2-+2.4 2 8 4 . 8 -+ 2.1 288.5 + 2.8 3 1 2 . 0 -+ 2.3 344.5 + 4.6 3 7 6 . 4 + 6.3 4 2 0 . 0 + 8.4 4 8 0 . 3 -+ 16.2 588.8 + 24.6 6 1 7 . 3 -+ 1.3 6 9 5 . 2 + 25.4 6 9 6 . 6 + 39.8 6 3 9 . 5 +- 35.9 4 8 . 2 3 + 2.24 5 1 . 1 7 + 6.48

2 7 4 . 4 +- 8.8 2 7 6 . 3 + 9.0 280.4+9.0 2 9 0 . 3 + 11.5 2 9 3 . 0 + 12.6 3 0 5 . 4 + 12.0 3 3 1 . 5 + 14.8 3 2 6 . 6 + 15.4 3 9 0 . 3 + 17.2 4 5 6 . 8 + 15.0 547.1 + 13.2 607.1 + 9.5 647.5 + 11.7 7 1 7 . 5 + 9.1 6 9 6 . 1 +- 14.5 7 6 . 0 3 +- 0.80 7 8 . 5 0 + 1.50

- 2 . 9 + 0.2 --2.5 -+ 0.2 - 2 . 6 -+0.2 - - 1 . 9 -+ 0.2 - 1 . 5 -+ 0.1 2.2 -+ 0.2 3.9 + 0.4 3.8 + 0.3 7.6 -+ 0.7 5.1 -+ 0.3 7.6 + 0.5 10.6 + 0 . 3 7.4 + 0.4 -2.9 a --8.1 a - - 3 6 . 6 + 2.1 --34.8 + 5.1

A l l (%)

S, e n g i n e e r i n g stress; ep, e n g i n e e r i n g strain; UTS, u l t i m a t e tensile s t r e n g t h ; Sf, f r a c t u r e stress; eu, u n i f o r m elongat i o n ; e~, t o t a l e l o n g a t i o n . a S t a n d a r d d e v i a t i o n irrelevant, as U T S a n d Sf are n o t flow stresses.

97 TABLE 5 E f f e c t s o f h y d r o g e n o n t h e tensile p r o p e r t i e s o f t h e r m a l l y aged 316L austenitic stainless steel

Parameter (units)

S (MPa) for S (MPa) for S (MPa) for S (MPa) for S (MPa) for S (MPa) for S (MPa) for S (MPa) for S (MPa) for S (MPa) for S (MPa) for S (MPa) for S (MPa) for UTS (MPa) Sf (MPa) e u (%) ef (%)

ep

=

e p ~-

ep = ep = ep ~ ep = ep= ep = ep = ep = ep = ep = ep =

0.05% 0.1% 0.2% 0.5% 0.6% 1% 2% 3.5% 5% 10% 20% 30% 40%

Value o f parameter after exposure to H 2 at 600 °C

Value o f parameter after exposure to Ar at 600°C

H2-induced property change

328.8 329.5 331.0 337.0 343.5 359.3 392.8 426.2 475.0 539.0 642.7 676.0 695.0 690.7 656.6 43.01 51.24

325.8 +- 3.5 326.3 +- 3.9 326.8 + 3.6 322.3 +- 3.8 336.6 -+ 3.1 343.6 -+ 4.1 368.8 -+ 5.9 400.3 -+ 7.2 430.5 -+ 7.9 492.2 -+ 5.3 578.8-+ 3.0 626.7 + 5.4 656.0 -+ 4.5 684.7 +- 20.0 657.0 + 11.3 62.21 + 3.76 63.14 +- 4.02

O.9+0.O 1.0+0.0 1.3+0.0 4.6+0.1 3.0 + 0.1 4.6+0.1 6.5 + 0.2 6.5+0.2 10.3+0.4 9.5+0.2 11.0+0.2 7.9-+0.2 5.9+0.1 0.9 a

+ 4.4 + 4.5 +- 6.2 + 6.7 + 6.4 -+ 7.5 ± 2.2 +- 2.1 +- 10.8 + 4.2 + 7.6 -+ 8.3 -+ 1.7 ± 12.2 ± 7.8 -+ 4.53 -+ 0.47

AII (%)

-0.1 a

-30.9 -+5.1 -18.8+2.4

S, engineering stress; ep, engineering strain; UTS, u l t i m a t e tensile s t r e n g t h ; Sf, f r a c t u r e stress; eu, u n i f o r m elongat i o n ; el, t o t a l e l o n g a t i o n . a S t a n d a r d deviation irrelevant, as UTS and Sf are n o t flow stresses.

tion the local strain was higher in the control specimens (see Tables 4 and 5). Thus, the martensite content data (see Table 3) should be regarded as for the purpose of orientation only. Scanning electron microscopy (SEM) fractography was carried o u t in a search for any hydrogen effects on the fracture mode or morphology of the two austenitic stainless steels. The gauge faces were also examined adjacent to the fracture separation, in order to reveal any hydrogen-induced surface effects.

4. R E S U L T S

4.1. Effects of hydrogen precharging on microstructure Some effects of hydrogen precharging on the microstructure of the type 304L and 316L austenitic stainless steels are summarized in Table 3. Carbide precipitation on grain boundaries as a result of thermal aging is illustrated by the TEM pictures in Fig. 2 (type 304L steel) and Fig. 3 (type 316L steel).

A comparison of the data in Tables 2 and 3 shows that thermal aging, whether in hydrogen or argon, brought about some general changes in the microstructural properties of both steels: (a) grain growth, amounting to 100% in type 304L steel and 32% in type 316L steel; (b) discontinuous carbide precipitation on grain boundaries, with large carbide particles at triple boundary joints in type 304L steel (Fig. 2) {the density o f the grain boundary carbides appears to be rather similar in b o t h 304L and 316L steels); (c) hardening, with the exception of t y p e 304L steel aged in hydrogen. Thermal aging (precharging) in hydrogen, however, led to two specific effects: (1) the occurrence of strain-induced martensite in the relatively stable 316L steel and a marked increase in the amount of strain-induced martensite in the 304L steel; (2) noticeable "hydrogen softening" in the 304L steel, amounting to 27 + 3% and 15 + 2% with respect to the control (aged in argon) and assupplied conditions respectively. To the best of our knowledge, those two effects have not yet been observed in austenitic stainless steels precharged to low internal hydrogen contents.

98 induced change" parameter All% =

Fig. 2. Transmission electron micrograph of type 304L steel aged at 600 °C in hydrogen gas.

Fig. 3. Transmission electron micrograph of type 316L steel aged at 600 °C in hydrogen gas.

4.2. Effects o f hydrogen precharging on tensile properties Tables 4 and 5 summarize the mean values of tensile properties of thermally aged (600 °C) 304L and 316L steels respectively. The properties of hydrogenated specimens are compared with those of control specimens (aged in argon) in terms of a "hydrogen-

100(II~ -- HAt) liar

As well as the mechanical parameters usually reported in the literature dealing with hydrogen effects in metals, namely total (fracture) elongation e~ (or reduction in area for bar specimens), 0.2% yield stress and ultimate tensile strength, Tables 4 and 5 also display the values of uniform elongation eu and selected engineering flow stresses. These additional parameters serve to emphasize the idea that hydrogen-induced or hydrogen-assisted fracture is the final event of a continuous complex strain history. When this idea is overlooked and solely the "conventional" parameters are invoked as criteria of hydrogen effects, an incomplete or even misleading picture might emerge. Thus, as judged from the loss Aef in total elongation, both the metastable 304L steel and the stable 316L steel are markedly embrittled on thermal aging in hydrogen gas. However, since the ductility loss (--18.8%) of the 316L steel is roughly half the ductility loss (--34.8%) of the 304L steel, the more stable steel appears to be somewhat superior to the metastable steel, in the given conditions of hydrogenation and testing. In contrast, as judged from the loss Aeu in uniform elongation the two steels appear to be embrittled to nearly the same degree: --30.9% for the 316L steel as against --36.6% for the 304L steel. The somewhat larger localized (necking) strain e~ -- eu of type 316L steel of about 8% as against approximately 3% for type 304L steel seems to be associated with a different fracture mechanism, as discussed later. Further, hydrogenation apparently had little effect on the 0.2% yield stress and ultimate tensile strength of either steel, if these parameters are examined out of the context of the respective flow curves. The very small differences between the 0.2% yield strength and ultimate tensile strength of hydrogenated specimens and the corresponding values of control specimens could readily have been dismissed as experimental errors and data scatter. The implicit and misleading conclusion would have been that hydrogenation in our experiments had little effect on the "overall" strength of type 304L and

99 316L austenitic stainless steels. The picture is significantly altered, however, if we examine the whole strain history bounded by the events of gross yielding and plastic instability (onset of necking). For this purpose we propose a new criterion of hydrogen effects (already implicit in the definition of the hydrogen-induced property change All in Tables 4 and 5), i.e. the "hydrogen sensitivity of flow stress", which is simply a plot of the hydrogen-induced change ~s% =

100(SH -- SAt) SAt

in engineering flow stress v e r s u s the logarithm of the engineering plastic strain ep1% (Fig. 4). The plot of As v e r s u s log epl in Fig. 4 is drawn up to epl = 40% and thus does not include the hydrogen-induced change AUTS in ultimate tensile strength; the values of ultimate tensile strength are not proper flow stresses as (1) they occur at different strains and (2) they mark the transition from a uniaxial state of stress to a triaxial one. The flow stress of both 304L and 316L steels shows significant hydrogen effects over a large part of the uniform strain range. For the 316L steel the effect is one of general "hydrogen hardening", from macro-yielding up to ep~ 40% (and b e y o n d (see Table 5)). Type 304L steel exhibits slight "hydrogen softening" over the entire gross yielding range, up to about 0.8% strain; the effect then reverts to one of hydrogen hardening, quite parallel to the situation with the 316L steel. Beyond ep~ = 40%, however, As drops steeply towards a negative value of AUTS (see Table 4), giving

rise again to slight softening. It is interesting to note that the relative magnitude of the hydrogen softening effect in type 304L steel was considerably larger on the multiaxial loading typical of indentation testing (Table 3) than in uniaxial tensile testing. Thus, thermal precharging to low internal hydrogen contents did affect the overall strength of both steels, although the largest effects occurred at flow stresses s5%-S4o% which are n o t among the mechanical parameters usually reported in the literature. Any hydrogen effects on the strength of a metal ultimately must originate in some interaction between the dissolved hydrogen and the mechanisms which govern the plastic deformation of the metal. Some qualitative information relevant to such interactions may be obtained from strain-hardening displays, e.g. comparative Hollomon plots (Figs. 5 and 6). Further, Table 6 displays the values of the strain-hardening exponent n in the Hollomon equation and the true plastic strain Con marking the onset of stages 2 and 3 in three-stage Hollomon plots. We observe in Figs. 5 and 6 that both steels have retained after thermal aging the threestage strain-hardening feature previously mentioned in Section 3.3. Hydrogenation induced two specific effects: (a) an overall decrease in ~" (MPe I()) !

I0

/

To

i

n u

AS(:)

,/

l//

/// /;/'

40

30

°

.

~ / / / ' /

-2

A "~'-I I to-'

-4

O.1 1.0 10.0 Fig. 4, Hydrogensensitivity of flow stressfor thermally aged 316L

(~) and 304L (o) steels.

, ,,

I io-'

's

I io-'

I

5

6rl Fig. 5. H o l l o m o n f l o w c u r v e s o f t y p e 3 0 4 L s t e e l t h e r m a l l y aged in h y d r o g e n ( ) a n d in a r g o n ( - - - ) a t 6 0 0 °C.

100 0'~ (MPa 10)

conclusion of gross yielding. Thus, a low internal hydrogen content, on the one hand, seems to facilitate the onset of slip mechanisms and yet, on the other hand, reduces the strain-hardening capacity of both steels. The latter effect cannot be explained as one of dynamic recovery because stage 2 of the Hollomon plots roughly corresponds to the range of fast hydrogen hardening in the plots of hydrogen sensitivity of flow stress (Fig. 6).

|| / ,/

;e |0

n

4.3. Fractography

S

10 -3

fJ

10-2

S

I0-;

S

Fig. 6. H o l l o m o n f l o w c u r v e s o f t y p e 3 1 6 L s t e e l t h e r m a l l y aged in h y d r o g e n ( ) a n d in a r g o n ( - - - ) at 6 0 0 °C.

the strain-hardening exponent n and (b) an onset of strain-hardening stages 2 and 3 at smaller eon strains. Both effects reach their largest magnitude with the onset of stage 2 at around epl = 5 × l 0 -3, which also marks the

After thermal aging in argon, the fractures of both 304L and 316L steels were typically ductile, with a dimple size nearly uniform across the whole fracture face. The fractures of hydrogenated 304L and 316L steels are characterized by the SEM micrographs in Figs. 7 and 8 respectively. Each of these micrographs is representative of three broken specimens. Figures 7(a) and 8(a) were photographed normally to the fracture surfaces, while Figs. 7(b) and 8(b) were photographed normally to the gauge face, adjacent to the fracture separation. The fracture surface of hydrogenated 304L steel (Fig. 7(a)) exhibits a markedly brittle morphology, with extensive interface separation. The exposed interfaces are mostly grain

TABLE 6 S t r a i n - h a r d e n i n g p a r a m e t e r s o f t h e r m a l l y aged 3 0 4 L a n d 3 1 6 L a u s t e n i t i c s t a i n l e s s s t e e l s

Parameter (units)

Steel

Aging environment

Stage 1

Stage 2

Stage 3

0.07 0.07 0.05 0.05

0.17 0.24 0.14 0.22

0.47 0.50 0.38 0.43

n n n n

304L 304L 316L 316L

An - n- H --- n A t X 1 0 0 (%) nAr

304L

--2.9 a

--31.3 a

--6.6 a

An = n i 4 - n A t X 1 0 0 (%) nAx

316L

--2.2 a

--35.5 a

--11.8 a

eon eon

304L 304L 316L

H2 Ar H2

----

4.5 X 10 -3 1.7 X 10 -2 5.4 X 10 -3

6.3 × 10 -2 9 . 6 x 10 -2 5.2 X 10 -2

eon

316L

Ar

--

1.6 X 10- 2

9.0 X 10 -2

eon

H2 Ar H2 Ar

Value of parameter in following stages

+ 0.01 + 0.02 + 0.01 -+ 0 . 0 2

+ + + +

0.01 0.02 0.01 0.05

+ + + +

0.01 0.04 0.01 0.01

a C o m p u t e d f r o m t h r e e - d i g i t v a l u e s o f n, in o r d e r t o r e s o l v e t h e m i n u t e y e t real h y d r o g e n - i n d u c e d d e c r e a s e in n o v e r s t a g e 1.

101

..........

(b)

Fig. 7. Fracture morphology of type 304L steel thermally aged in hydrogen: (a) fracture surface; (b) gauge face.

boundaries; only a few seem to be previous twin boundaries. Some of the exposed interfaces exhibit striations. A large number of exposed crystallite edges seem to suggest preferential crack initiation at triple boundary joints. A few sites of transgranular separation are visible as well. A striking feature associated with the failure of hydrogenated 304L specimen is the absence of surface

I

Fig. 8. Fracture morphology of type 316L steel thermally aged in hydrogen: (a) fracture face; (b) gauge face.

cracking on the gauge face in Fig. 7(b); instead, extensive surface rumpling is apparent, suggesting both strain-induced martensite and deformation twinning. The fracture of hydrogenated 316L steel shows a very different picture. Failure here was associated with a transverse crack (Fig. 8(a)) parallel to the width of the specimen gauge. The crack is located at nearly the halfthickness of the specimen and seems to be the result of a localized shear band in the specimen's neck. Strings of interconnected voids

102

Fig. 9. Shallow cracks on the gauge face of type 316L bar specimens [ 37 ], adjacent to the fracture separation.

are apparent along most of the crack's length. Also a few secondary cracks are associated with strings of interconnected voids. Narrow areas of apparently transgranular separation border the crack on both sides, and yet most of the fracture surface is of the dimpled rupture type, with a mean dimple size about 2.4 times larger than that of control (aged in argon) 316L specimens. In contrast with type 304L specimens the gauge face of hydrogenated 316L specimens shows extensive surface cracking (Fig. 8(b)), seemingly along grain boundaries. In many instances the cracks seem to have been initiated at triple boundary joints. No surface distortion is apparent. It is interesting to note that very similar surface cracking was observed in earlier work [37] on type 316L bar specimens, sensitized and hydrogenated in the same conditions as in the present work (Fig. 9).

5. DISCUSSION

5.1. Initial critical hydrogen concentration The present study has revealed three interesting and rather unexpected effects of low internal hydrogen contents: (1) significant degradation of tensile properties for both steels; (2) enhanced formation of straininduced martensite in both austenitic stainless steels, suggesting an increase in Md; (3) a comparatively high susceptibility of type 316L steel to hydrogen effects.

In general, for hydrogen-induced degradation of properties to occur, some critical hydrogen concentration must be established at microstructural sites, where potential failure is favoured by an appropriate state of stress and strain. In principle, this threshold hydrogen concentration can be reached locally, at suitable trapping sites, even in specimens having a bulk hydrogen concentration less than the critical value. Grain boundaries are known to be strong traps for hydrogen, characterized by interaction energies in the range 0.3-0.65 eV, depending on local crystallography and chemistry and on the state of stress imposed on the matrix (see for example ref. 49). When precipitates such as carbides are associated with grain boundaries, the interaction energy may increase up to 0.8-1.0 eV. It can be assumed that diffusion with trapping determined the hydrogen microdistribution in the specimens subjected to concurrent sensitization and charging in our experiments. Thus, the twofold result of the sensitization and charging process would have been chromium depletion and hydrogen enrichment at the carbide-matrix interface. This enrichment, whether at grain boundaries (type 304L steel) or at additional microsites (type 316L steel), associated with the states of stressstrain imposed by tensile testing obviously proved sufficient to induce the significant effects reported in Section 4. 5.2. Hydrogen-enhanced formation of straininduced martensite Narita and Birnbaum [50] and Narita et al. [ 7 ] have shown very recently that large contents of hydrogen in solid solution increased the Ms and Md values for the transformation from 7 to e in both unstable (type 304) and stable (type 310) austenitic stainless steels in the unsensitized condition. On hydrogen release, e.g. during plastic deformation, e transformed further to a only in the unstable steel while, in the type 310 steel, e reverted to 7 and no c~ was detected. The very high hydrogen concentrations (tens of atomic per cent) in the surface layer of the specimens led to the build-up of crystallographic shears, equivalent to about 5% linear strain [7], and consequently to austenite destabilization without any externally applied stress. Eliezer and coworkers [51] have reported very recently

103

that cathodic precharging without any externally applied stress induced the formation of e martensite in type 316L austenitic stainless steel. In the present work the crystallographic shears required for austenite destabilization in the 304L and 316L steels had to be contributed jointly by an externally applied stress and the locally available hydrogen content. While the contribution from the external stress must have been the major one througho u t the bulk of the specimens, the preliminary sensitization of the specimens might have enhanced locally, at carbide-matrix interfaces, the contribution of the trapped hydrogen content. Further hydrogen enrichment at the carbide-matrix interface by dislocation transport may be assumed to have contributed to the build-up of crystallographic shears. Thus the effects of low hydrogen contents may also be rationalized in terms of the general sequence of microstructural events proposed by Narita et al. [7] for hydrogen embrittlement or hydrogen-induced degradation of properties. In both extreme cases (high and low hydrogen contents) the salient

hydrogen content

point is austenite instability, induced either throughout the lattice or at specific microsites b y suitable combinations of hydrogen content and strain (stress). The overall analogy is illustrated in Fig. 10. The behaviour of type 304L steel in our experiments can be rationalized satisfactorily in terms of the model suggested above and thus confirms the ideas of Briant [40-43, 46] with respect to the role of sensitization in increasing the susceptibility of unstable (low weight percentage of nickel) austenitic stainless steels to hydrogen. The relatively large amount of strain-induced martensite (24 wt.%) which most probably formed preferentially at the interfaces of grain boundary carbides, may be assumed to have played the major role in the hydrogen-induced degradation of properties of the 304L steel; as a result of the large diffusivity and possibly the low solubility of hydrogen in the b.c.c, phase, c~ martensite acted as the easy path for short-range hydrogen transport to crack initiation sites and for intergranular crack propagation as well. Direct TEM evidence a b o u t the role of hydrogen-induced ~ martensite as the easy path for crack propagation in type 304L steel

I

3, phase destabilization j

•

large hydrogen effects

(a)

local chemical 3' phase destabilization local hydrogen enrichment

lattice

shear locally enhanced

3, phase destabilization

moderate although significant hydrogen effects

low lattice hydrogen content I external ] stress

(b) Fig. 10. Schematic relationship between the amount of hydrogen in solution and hydrogen embrittlement or hydrogen-induced degradation of properties in austenitic stainless steels: (a) high hydrogen content, no external stress; (b) low hydrogen content, external stress.

104 was obtained very recently by Minkovitz and Eliezer [ 52 ]. The absence of cracking on the gauge surface of type 304L specimens in the present work as opposed to the surface cracks observed by Briant [40, 41] on type 304 and 304L specimens can be explained in terms of the different hydrogenation methods employed. The surface cracking in Briant's work might have been a consequence of lattice damage initially induced in a surface layer by cathodic charging. Such a condition does not occur on thermal precharging at moderate hydrogen pressures, as was done in our experiments. In general, surface cracking is a stress relief effect induced by large differences between the states of stress in the outer and inner layers of the specimen. In the absence of early surface flaws, considerable stress gradients between the outer and inner layers might have built up during the tensile testing o f the 304L specimens used in the present work. However, profuse martensite formation and deformation twinning would be expected to have facilitated stress relief at the free surface; the extensive surface rumpling observed (Fig. 7(b)) seems to confirm this assumption.

5.3. Hydrogen-induced degradation of properties of type 316L austenitic stainless steel The susceptibility of sensitized 316L steel to hydrogen appears to involve a mechanism considerably more complex than that in the susceptibility of type 304L steels. According to the scarce information published so far, bar specimens of type 316 steel are little affected by hydrogen in either the annealed [45] or the sensitized condition [46], even when some strain-induced ~ martensite forms in the microstructure. Briant [46] observed noticeable hydrogen-induced degradation of properties of the more stable austenitic stainless steels, such as type 316 or 316L (and 310) steels, only when thin sheet specimens were hydrogenated. These steels exhibited surface cracking regardless of the magnitude of the hydrogen-induced degradation effects. The above-mentioned effects were confirmed by the results of the present authors with type 316L bar specimens [37] and thin sheet specimens respectively: (a) shallow surface cracking of a mixed intergranular and trans-

granular character was observed with both bar and thin sheet specimens (Figs. 9 and 8 respectively); (b) ductility losses were zero with the bar specimens but significant with the thin sheet specimens. All the above results indicate the considerable influence that the specimen geometry, i.e. the state of strain and stress, has on the susceptibility to hydrogen effects. In thin sheet specimens (large width-to-thickness ratio) with large shoulders, the lateral contraction of the gauge is partially constrained such that e 2 ~ e 3 (where e2 is the lateral strain and e3 is the through-thickness strain); as a consequence a state of plane strain is superimposed on the axis-symmetric state (see for example ref. 53). As discussed by Stoltz [54] in the context of environmental hydrogen embrittlement at low hydrogen pressures, the state of plane strain is more sensitive to hydrogeninduced surface flaws than is the axis-symmetric state, i.e. for a given material condition and hydrogen content the thin sheet specimens will show a larger hydrogen-induced degradation of properties than will bar or plate specimens. The results of the present authors now suggest a similar response of specimen geometry to internal hydrogen effects. Work is in progress at the present authors' laboratory with the goal of establishing quantitative relationships between the specimen geometry and the magnitude of effects induced by low uniform contents of internal hydrogen. Strain-induced ~ martensite clearly was not the dominant factor in the significant hydrogen-induced degradation effects observed with the thin sheet specimens in the present study. While the ~ phase in this material may be assumed to have formed according to the general mechanism illustrated in Fig. 10(b), the low volume fraction (9 vol.%) apparently was not able to provide for a continuous failure path along grain boundaries or other interfaces. It is interesting to mention in this c o n t e x t Briant's observation [46] that the density of grain boundary carbides was always equivalent in the sensitized 304 steel to that in the sensitized 316 steel. No attempt was made to measure the density of grain boundary carbides in the present work; nonetheless, the representative TEM micrographs of hydrogenated 304L and 316L (Figs. 2 and 3 respectively) seem to suggest that the car-

105

bide density was of the same order of magnitude in both steels. The fact that chromium depletion and the assumed hydrogen enrichment at the interface of grain boundary carbides did not provide for extensive a phase formation and ensuing grain boundary failure in the 316L steel is due to the higher alloy content (nickel plus molybdenum) of the 316L steel when compared with the 304L steel [45, 46]. The shallow surface cracking of hydrogenated specimens observed by Briant [46] and the present authors seems to be typical of the more stable austenitic stainless steels and independent of the surface condition. The latter point is emphasized by the following facts: (a) Briant [46] employed cathodic precharging (inherent surface damage), while thermal precharging at a moderate hydrogen pressure (no surface damage) was employed in both the present work and the earlier work [37] ; (b) similar surface cracking was observed with both bright-annealed and asmachined 316L steel specimens (Figs. 8(b) and 9 respectively). The white particles on the surface of type 316L steel specimens were shown to be surface artifacts; mild ion sputtering or mild chemical attack easily removed the particles. Furthermore, metallographic examination of through-thickness sections did not reveal such particles. The present authors believe that the formation of the particles on thermal aging in hydrogen is associated with the surface chemistry of our particular grade of type 316L steel. The absence of rumpling on the surface of our 316L steel specimens suggests that the available volume fraction of strain-induced a martensite, which is much lower than in the type 304L steel, was not sufficient to allow for surface stress relief. As plastic deformation of hydrogenated specimens leads to an increased hydrogen concentration near the external surface due to dislocation sweeping [34, 55] and the room temperature diffusivity of hydrogen in the f.c.c, lattice is about 104 times lower than in the b.c.c, lattice, a local supersaturation of the lattice with hydrogen at the crack tip would be facilitated, thus leading to surface stress relief by shallow cracking. The large loss of uniform elongation (31%) observed in the present work with hydrogenated thin sheet 316L specimens inherently

was associated with an early onset of macroscopic plastic instability. It is likely that the surface cracks discussed in the preceding might have played the part of surface flaws inducing plastic instability, i.e. necking of the thin sheet specimens. The fact that surface cracking did not induce early necking and ensuing loss of uniform elongation in the 316L bar specimens [37] is apparently due to the size or geometry effect discussed in the first part of this section. The fractographic observations pertinent to thin sheet 316L steel (Section 4.3}, namely the increased dimple size and the transverse crack with features of void linking (Fig. 8(a)), suggest an essentially ductile mode of fracture occurring by hydrogen-assisted microvoid growth combined with shear band localization, as discussed by Thompson [20]. With the onset of necking, microvoids growing at particle interfaces become efficient traps for the hydrogen swept by dislocations [49]; hydrogen atoms recombine at the fresh free surfaces, leading to a build-up of hydrogen pressure inside the microvoids. This pressure aids accelerated microvoid growth and coalescence [ 56 ]. The fracture of type 316L bar specimens [37], which was fully ductile and featured a decreased dimple size, i.e. dominant microvoid nucleation, was probably associated with the size or geometry effect already discussed in this section, namely a larger degree of uniaxial strain hardening and increased triaxiality on necking.

6. C O N C L U S I O N S

(1) A low internal hydrogen content, of the order of 300 at.ppm, was shown to induce significant degradation of tensile properties in thermally aged (sensitized) type 304L and 316L austenitic stainless steels, namely ductility losses, hydrogen hardening (flow stress increase), decreased strain-hardening capacity and changes in fracture mode (type 304L steel) or morphology (type 316L steel). (2) Hydrogenation increased the amount of strain-induced a martensite in both steels. (3) Hydrogen hardening and the changes in strain-hardening capacity mainly occurred over an intermediate range of strains; the yield stress and ultimate tensile strength of

106

b o t h steels w e r e little a f f e c t e d b y h y d r o g e n a t i o n and t h u s w e r e s h o w n to be d o u b t f u l or even misleading criteria of h y d r o g e n e f f e c t s w h e n e v a l u a t e d o u t o f t h e c o n t e x t o f overall deformation behaviour. Uniform elongation was s h o w n t o be an i m p o r t a n t criterion of h y d r o g e n effects, as well as the e l o n g a t i o n t o fracture. (4) T h e h y d r o g e n - a s s i s t e d failure o f t h e 3 0 4 L steel was a n a l y s e d in t e r m s of h y d r o g e n t r a p p i n g at carbide interfaces, e n h a n c e d form a t i o n o f ~ m a r t e n s i t e a n d i n t e r f a c e separation. (5) T h e surprisingly large h y d r o g e n e f f e c t s in t y p e 3 1 6 L steel were discussed in t e r m s o f a size e f f e c t o p e r a t i v e w i t h t h i n sheet specim e n s , surface c r a c k i n g as a m o d e o f surface stress relief a n d failure b y m i c r o v o i d g r o w t h a n d c o a l e s c e n c e c o m b i n e d w i t h shear b a n d localization. (6) T h e r m a l p r e c h a r g i n g t o low a n d u n i f o r m c o n t e n t s o f i n t e r n a l h y d r o g e n , w i t h no surface d a m a g e , a n d a " s c a n " o f the entire plastic r a n g e a p p e a r t o be e x p e r i m e n t a l techniques p a r t i c u l a r l y useful in the s t u d y o f intrinsic h y d r o g e n e f f e c t s in a u s t e n i t i c stainless steels.

ACKNOWLEDGMENTS

T h e a u t h o r s t h a n k f u l l y a p p r e c i a t e the c o m p e t e n t assistance o f Mr. R. F r e n k e l a n d Mr. A. Magen, b o t h o f t h e N u c l e a r R e s e a r c h Centre, Negev. T h e h e l p o f Dr. G. K i m m e l a n d his t e a m at t h e N u c l e a r R e s e a r c h Centre, Negev, is g r a t e f u l l y a c k n o w l e d g e d . T h e T E M w o r k o f Mr. E. M i n k o v i t z , Ben G u r i o n University, is g r e a t l y a p p r e c i a t e d . This w o r k was in partial f u l f i l m e n t o f t h e r e q u i r e m e n t s f o r an M.Sc. Thesis b y Y.R.

REFERENCES

1 P. A. Blanchard and A. R. Troiano, M$m. Sci. Rev. M$t., 57 (1960) 409. 2 M. B. Whiteman and A. R. Troiano, Corrosion, 21 (1965) 53. 3 M. L. Holzworth, Corrosion, 25 (1969) 107. 4 J. Kolts, in H. Craig, Jr. (ed.), Stress Corrosion -New Approaches, A S T M Spec. Tech. Publ. 610,

1976, p. 366. 5 D. Eliezer, D. G. Chakrapani, C. J. Altstetter and E. N. laugh, Metall. Trans. A, 10 (1979) 935.

6 H. H/inninen and T. Hakkarainen, Metall. Trans. A, i 0 (1979) 1196. 7 N. Narita, C. J. Altstetter and H. K. Birnbaum, Metall. Trans. A, 13 (1982) 1355. 8 K. Farrell and M. B. Lewis, Scr. Metall., 15 (1981) 661. 9 S. Asano and R. Otsuka, Ser. Metall., 10 (1976) 1015. 10 R. A. Oriani, Annu, Rev. Mater. Sci., 8 (1978) 327. 11 R. A. Oriani, in M. R. Louthan, Jr., R. P. McNitt and R. D. Sisson, Jr. (eds.), Environmental Degradation o f Engineering Materials in Hydrogen, Virginia Polytechnic Institute, Blacksburg,

VA, 1981, p. 3. 12 J. Burke, A. Jickels, P. Maulik and M. L. Mehta, in A. W. Thompson and J. H. Bernstein (eds.), Effect o f Hydrogen on Behavior o f Materials,

AIME, New York, 1976, p. 102. 13 D. A. Vaughan, D. J. Phalen, C. L. Peterson and W. K. Boyd, Corrosion, 19 (1963) 315t. 14 H. Mathias, Y. Katz and S. Nadir, Met. ScL, 12 (1978) 129. 15 M. L. Holzworth and M. R. Louthan, Jr., Corrosion, 24 (1968) 110. 16 H. Okada, Y. Hosoi and S. Abe, Corrosion, 26 (1970) 183. 17 P. Maulik and J. Burke, Scr. Metall., 9 (1975) 17. 18 H. H~inninen, S. P. Hannula and S. Tahtinen, in Environmental Degradation o f Engineering Materials in Hydrogen, Virginia Polytechnic In-

stitute, Blacksburg, VA, 1981, p. 347. 19 M. Rigsbee, Metallography, 11 (1978) 493. 20 A. W. Thompson, in A. W. Thompson and J. M. Bernstein (eds.), Effect o f Hydrogen on Behavior o f Materials, AIME, New York, 1976, p. 467. 21 D. A. Vaughan and D. J. Phalen, Met. Eng. Q., 5 (1965) 39. 22 K. Kamachi, J. Iron Steel Inst. Jpn., 18 (1978) 485. 23 A. Szummer and A. Janko, Corrosion, 35 {1979) 461. 24 C.-M. Hsiao and W.-Y. Chu, in J. M. Bernstein and A. W. Thompson (eds.), Hydrogen Effects in Metals, AIME, New York, 1981, p. 255. 25 A. W. Thompson and O. Buck, Metall. Trans. A, 7 (1976) 329. 26 J. P. Bressanelli and A. Moskowitz, Trans. Am. Soc. Met., 58 (1965) 499. 27 W. Hoffman and W. Rauls, Weld. J., 44 (1965) 225S. 28 A. Kumnick and H. Johnson, Metall. Trans. A, 6 (1975) 1087. 29 R.J. Walter and W. T. Chandler, Mater. ScL Eng., 8 (1971) 90. 30 R. R. Vandervoort, Met. Eng. Q., 12 (1972) 10. 31 M. R. Louthan, Jr., G. R. Caskey, Jr., J. A. Donovan and D. E. Rawl, Jr., Mater. ScL Eng., 10 (1972) 357. 32 A. W. Thompson, in J. M. Bernstein and A. W. Thompson (eds.), Hydrogen in Metals, American Society for Metals, Metals Park, OH, 1974, p. 91. 33 A. W. Thompson, Mater. Sci. Eng., 14 (1974) 253.

107 34 J. A. Donovan, Metall. Trans. A, 7 (1976) 1677. 35 T. L. Capeletti and M. R. Louthan, Jr., J. Eng. Mater. Technol., 99 (1977) 153. 36 G. R. Caskey, Jr., Scr. Metall., 15 (1981) 1183. 37 Y. Rosenthal, M. Mark-Markowitch, A. Stern and D. Eliezer, Scr. Metall., 15 ( 1 9 8 1 ) 8 6 1 . 38 J. P. Fidelle, L. R. Allemand, C. R o u x and M. Rapin, in J. P. Fidelle and M. Rapin (eds.), L'Hydrog~ne dans les M~taux, Centre d'Etudes Commissariat ~ l'Energie Atomique, Bruy~res-leChatel, 1969, p. 131. 39 J. P. Fidelle, R. Bernardi, R. Broudeur, C. Roux and M. Rapin, in Hydrogen Embrittlement Testing, ASTM Spec. Tech. Publ. 543, 1974, p. 221. 40 C. L. Briant, MetaU. Trans. A, 9 (1978) 731. 41 C. L. Briant, Metall. Trans. A, 10 (1979) 181. 42 C. L. Briant, in Environmental Degradation of Engineering Materials in Hydrogen, Virginia Polytechnic Institute, Blacksburg, VA, 1981, 335. 43 C. L. Briant, Scr. Metall., 12 (1978) 541. 44 R. Lagneborg, J. Iron Steel Inst., London, 207 (1969) 363. 45 G. R. Caskey, Jr., in Environmental Degradation of Engineering Materials in Hydrogen, Virginia Polytechnic Institute, Blacksburg, VA, 1981, p. 283.

46 C. L. Briant, in J. M. Bernstein and A. W. Thompson (eds.), Hydrogen Effects in Metals, AIME, New York, 1981, p. 527. 47 T. Angel, J. Iron Steel Inst., London, 177 (1954) 165. 48 J. Chine, M. Aucouturier, R. Arnould-Laurent, P. Tison and J. P. Fidelle, in J. M. Bernstein and A. W. Thompson (eds.), Hydrogen Effects in Metals, AIME, New York, 1981, p. 583. 49 A. W. Thompson, in Environmental Degradation of Engineering Materials, Virginia Polytechnic Institute, Blacksburg, VA, 1978, p. 3. 50 N. Narita and H. K. Birnbaum, Scr. MetaU., 14 (1980) 1355. 51 E. Minkovitz, M. Talianker and D. Eliezer, J. Mater. Sci., 16 (1981) 3506. 52 E. Minkovitz and D. Eliezer, J. Mater. Sei. Lett., 1 (1982) 192. 53 A. K. Chakrabarti and J. W. Spretnak, Scr. Metall., 8 (1974) 743. 54 R. E. Stoltz, Metall. Trans. A, 12 (1981) 543. 55 J. K. Tien, S. V. Nair and R. R. Jensen, in J. M. Bernstein and A. W. Thompson (eds.), Hydrogen Effects in Metals, AIME, New York, 1981, p. 37. 56 A. W. Thompson and J. M. Bernstein, Adv. Corros. Sci. Technol., 7 (1979) 53.