Hydrogen embrittlement of 316L type stainless steel

Hydrogen embrittlement of 316L type stainless steel

Materials Science and Engineering A272 (1999) 279 – 283 www.elsevier.com/locate/msea Hydrogen embrittlement of 316L type stainless steel E. Herms, J...

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Materials Science and Engineering A272 (1999) 279 – 283 www.elsevier.com/locate/msea

Hydrogen embrittlement of 316L type stainless steel E. Herms, J.M. Olive *, M. Puiggali Laboratoire de Me´canique Physique, UPRESA CNRS 5867, Uni6ersite´ Bordeaux I, 351 Cours de la Libe´ration, 33405 Talence, France Received 27 November 1998; received in revised form 11 June 1999

Abstract Hydrogen embrittlement tests on type 316L stainless steel are performed including cathodic charging during slow strain rate tests. Brittle multiple cracking is observed and relationships between crack growth rate and diffusion are analysed. The influence of hydrogen on the morphology of ductile fracture is found after fractographic examination. Two aspects of ductile failure are observed in accordance with the hydrogen content of the sample; a reduced density of microvoids for higher hydrogen contents and brittle secondary cracking in addition to ductile fracture surfaces for lower hydrogen contents. © 1999 Elsevier Science S.A. All rights reserved. Keywords: Hydrogen embrittlement; Slow strain rate tests; Austenitic stainless steel; Fractography; Ductile fracture

1. Introduction Hydrogen has a deleterious effect on the mechanical properties of many materials such as ferritic steels, nickel based alloys, h.c.p. alloys. Austenitic structures and in particular austenitic stainless steels are less susceptible to hydrogen embrittlement because of low diffusivity and high solubility of hydrogen in the f.c.c. structures. Brittle fracture associated with hydrogen embrittlement (HE) is, however, observed for stainless steels in severe environmental conditions such as cathodic charging [1,2]. Depending on test conditions and material sensibility to HE, several fracture modes are observed for Fe–Ni– Cr alloys: intergranular fracture, quasi-cleavage, twin boundary separation and dimpled failure. In the case of ductile failure associated with HE, an influence of hydrogen on the morphology of microvoids (decrease of density and size) is observed [3]. However, this influence depends on the material and hydrogen charging conditions. A decrease of microvoid sizes due to hydrogen was reported for nickel based alloys [4,5], an opposite effect was observed for 304L [6] and Fe–Cr– Mn austenitic steel [7]. * Corresponding author. Tel.: + 33-5-56846222; fax: +33-556846964. E-mail address: [email protected] (J.M. Olive)

The aim of this study is to identify the different degradation processes involved in HE by cathodic charging fOR 316L type stainless steel.

2. Experimental conditions The hydrogen embrittlement tests were performed on annealed AISI 316L stainless steel (composition in wt%: C: 0.02; Cr: 16.92; Ni: 12.07; Mo: 2.62; Si: 0.39; Mn: 1.68; S: 0.018). Specimens were wires of 1.5 mm in diameter and 200 mm in length. After solution heat treatment: annealing 1050°C during 30 min and quenching, the materials was fully austenitic with an average grain size of 60 mm. Specimen were subsequently polished by abrasive paper grade 1200. Cathodic hydrogen charging at room temperature was employed using a 1N H2SO4 solution containing 0.25 g l − 1 As2O3 as a hydrogen recombination poison. A current density of 100 mA cm − 2 was applied between the specimen and a platinum anode. A 24 h precharging was applied to increase the severity of the test. Slow strain rate tests (SSRT) were performed during cathodic polarisation on cathodically precharged 24 h specimen and for three strain rates ranging between 7× 10 − 7 and 5× 10 − 6 s − 1. Scanning electron microscopy analyses of specimens and fracture surfaces were realised on each sample after tests.

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Table 1 Result of hydrogen embrittlement during slow strain rate test of AISI 316L AISI 316L

o; = 7×10−7 s−1

o; = 1×10−6 s−1

o; =5×10−6 s−1

Strain to fracture Test duration (h) Average crack propagation rate (10−10 ms−1) Crack depth (10−6 m) Depth theoretically covered by H2 (SSRT duration+24 h precharging) (10−6 m)

20% 80 5 140 50

13% 22 15 120 30

15% 9 28 90 27

3. Results and discussion

3.1. Hydrogen absorption During the cathodic precharging of 24 h duration, brittle multiple cracking induced by hydrogen diffusion develops. The average crack depth does not exceed 3 mm although the average crack width is 20 mm. Based on the hydrogen diffusion coefficient D0 in 316 stainless steel [2], about 4×10 − 12 cm2s − 1 at room temperature, the distance–time relation deduced from Fick’s second law, x 2 =Dt, gives 23 mm for t =24 h. The thickness of metal where the critical concentration for cracking is reached during precharging, is about 0.13 times the distance predicted by diffusion. Cathodic polarisation during slow strain rate test results in a severe reduction of ductility for AISI 316L stainless steel (Table 1). Strain to fracture that is 55% for 316L stainless steel tested in air decreases to 15% under cathodic polarisation. This reduction of strain is due both to brittle cracking (Fig. 1) and a modification of ductile failure as will be described in the next section. Table 1 shows that the average crack propagation rate increases with strain rate. The average crack propagation rate is estimated from the maximum crack depth observed on the fracture surface and the test duration, this latter quantity being equivalent to the propagation time since crack initiation occurs during precharging. The crack growth rate (Vp in ms − 1) can be approximated as a function of strain rate (o; in s − 1) by:



Vp=7.3× 10 − 10

o; o; 0

(Table 1) is estimated from hydrogen diffusion coefficient. The crack depth is at least three times larger than the zone affected only by pure diffusion. From the ratio thickness of metal where the critical concentration for cracking is reached to the diffusion zone thickness evaluated during the precharging as being 0.13, one can predict that the entire specimen will contain hydrogen at the end of the SSRT. This effect of plasticity is typically associated with hydrogen transport by dislocations.

3.2. Fracture mode In the adopted conditions, the hydrogen content is larger near the surface of the sample than in the bulk. Brittle mixed mode of fracture is observed in the periphery of the fracture surface (average crack depth of 100 mm) and ductile fracture mode in the center of fracture surface in accordance with the lower hydrogen

0.73

(1)

with o; 0 = 7×10 − 7 s − 1, the lowest value of strain rate used in this study. This means that the critical hydrogen concentration for decohesion near the crack tip is reached faster for higher strain rates, in the range 7 ×10 − 7 s − 1 – 5× 10 − 6 s − 1. Hydrogen diffusion is well known to be activated by plastic deformation. Eq. (1) indirectly shows that this activation is time dependent. This result is consistent with a transport of hydrogen by dislocations. The depth of the hydrogen affected zone reached after the total test duration (precharging+ SSRT)

Fig. 1. Brittle transgranular and intergranular multiple cracking on specimen surface. Cathodic charging on type 316L steel during slow strain rate test, cathodic polarisation 100 mA cm − 2, strain rate 10 − 6 s − 1, strain to fracture 13%.

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Fig. 2. Schematic presentation of typical fracture surface obtained after hydrogen embrittlement test, cathodic charging on type 316 L slow strain rate test.

content reached in these areas (Fig. 2). Dimple failure is commonly observed at any temperature for f.c.c. metals if the overload is the main cause of fracture, contrary to b.c.c. metals which show brittle to ductile transitions at low temperatures. Ductile fracture corresponds to the growth of microvoids and their coalescence in regions of strain discontinuities (inclusions, ...). The fracture surface appears essentially composed of two parts: a brittle crown area near the sample surface followed by ductile part in the center (Figs. 2 and 3). Three aspects on the ductile part of the fracture surface appear from the vicinity of the brittle zone to the center of the fracture surface:

Fig. 3. Transition between brittle and ductile fracture. Cathodic charging of type 316L steel during slow strain rate test, cathodic polarisation 100 mA cm − 2, strain rate 10 − 6 s − 1, strain to fracture 13%.

Fig. 4. Ductile fracture near brittle zone. Cathodic charging of type 316L steel during slow strain rate test, cathodic polarisation 100 mA cm − 2, strain rate 10 − 6 s − 1, strain to fracture 13%.

1. Near the brittle zone (till 100 mm) small microvoids (B 1 mm diameter) of low spatial density are observed (Figs. 3 and 4). This particular pattern corresponds to a decrease in ductility caused by hydrogen which is present in lower concentrations than near the surface of the specimen where brittle fracture is observed. 2. Near the aforementioned area (till 350 mm), the ductile fracture surface is characterised by the presence of larger microvoids (3 mm) similar in size to those on the fracture surface of specimens tested in air but with signs of brittle secondary cracking (Fig. 5). These microvoids are characterised by not really well defined tear ridges in comparison with microvoids usually observed in this material after mechanical failure in air. Fig. 6 shows the common aspects of a ductile fracture surface obtained on the same material after mechanical failure in air: microvoids are well defined and without secondary cracks. 3. The third zone is an apparently unaffected zone where ductile failure is similar as in fracture surfaces of the same steel in air. No hydrogen effect can be observed in this part of the sample. The hydrogen effects near the center of the sample can not be explained only by hydrogen diffusion. Therefore secondary brittle cracking in the ductile zone is in accordance with hydrogen transport by dislocations. Similar observations were reported after hydrogen embrittlement tests on 304 stainless steel in hydrogen sulphide (H2S) environment in SSRT conditions [8]. Secondary cracking in the ductile central area of the

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Fig. 5. Ductile zone near the center of sample with secondary brittle cracks. Cathodic charging of type 316L steel during slow strain rate test, cathodic polarisation 100 mA cm − 2, strain rate 10 − 6 s − 1, strain to fracture 13%.

sample is also observed. The fracture surfaces of the 304 samples show brittle crown areas in the periphery and ductile central areas with secondary cracking. This latter lower influence of hydrogen is related to a lower content of hydrogen in the specimen center than near the specimen surface because of the low diffusivity of hydrogen in f.c.c. structures.

Fig. 6. Ductile fracture obtained on 316L stainless steel after tensile test in air. Strain rate 10 − 2 s − 1, strain to fracture 55%.

Ductility reduction is observed in hydrogen embrittlement conditions for many alloys. The fracture mode varies with test conditions and alloy composition [3]: usually ductile failure occurs for a nickel content ranging between 12 and 30% (stainless steels), and the mode changes to intergranular fracture for a nickel content higher than 30 or 35%. But intergranular fracture appears for sensitised austenitic stainless steels and transgranular fracture (cleavage, quasi-cleavage) occurs after severe hydrogen charging or in samples containing high amounts of hydrogen [9–11]. In general, moderate hydrogen embrittlement can be seen on fracture surface thanks to a reduction of density and size of microvoids. This aspect has been examined on 316 stainless steel [12], nickel based alloys [5] and superalloys [4,13], or even after impact tests on 304 stainless steel [3] in gaseous hydrogen. The same effect is also observed when hydrogen is introduced in an Fe–Cr–Mn austenitic stainless steel under neutron irradiation [7]. Ductile failure on hydrogen charged samples can be explained by the presence of relatively low amounts of hydrogen: a high content of hydrogen is necessary to introduce brittle fracture in austenitic alloys. Metallurgical factors play also a role. A transition of fracture mode on pure nickel is observed depending on grain size [14]. The presence of hydrogen changes the fracture mode from transgranular ductile rupture to brittle intergranular fracture only when the grain size is higher than 10 mm. Ductile failure is commonly associated with nucleation and growth of microvoids around particles of a second phase in metal when a critical value of stress is exceeded. The influence of hydrogen on ductile fracture is attributed to an accumulation of hydrogen at matrix–particles interfaces, which lowers the interfacial strength [15]. On the other hand, in our case after this first affected zone (reduction of density and size of microvoids), the second zone near the center of the fracture surface shows two aspects different from classical dimple fracture: (i) brittle secondary cracking apparently intergranular and transgranular, (ii) presence of non typical microvoids (shape of microvoids is modified because of not well defined tear ridges). These two aspects, present on fracture surfaces between the first and second zones (except brittle cracks on periphery), show the role of hydrogen content on the nature of fracture in hydrogen embrittled 316L stainless steel. The hydrogen content decreases from the surface of the sample to the center owing to the low diffusivity of hydrogen in the austenitic phase. The presence of a not fully ductile fracture surface far away (350 mm) from the brittle areas is a significant effect of the hydrogen transport by dislocations. Pure diffusion does indeed not provide sufficient hydrogen penetration far from the surface of the sample because of the low diffusivity of hydrogen in AISI 316L.

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These characteristics of hydrogen embrittlement of 316L stainless steel show relationships between hydrogen and plasticity in accordance with hydrogen–plasticity interaction models in hydrogen embrittlement and stress corrosion cracking [16].

[2]

[3] [4]

4. Conclusion Hydrogen embrittlement tests consisting of cathodic charging during SSRT were performed on annealed AISI 316L stainless steel. In these conditions 316L stainless steel suffers from a severe ductility reduction associated with brittle multiple cracking. In addition to brittle fracture, an influence of hydrogen in the material on the ductile fracture surface is also detected and analysed. In relation with the hydrogen content of the material, hydrogen embrittlement can occur not only in the form of brittle cracks. Ductile fracture can occur along with ductility reduction. The role of hydrogen in the development of fracture is an influence on density, nature and size of microvoids present on dimpled fracture surface.

[5] [6] [7] [8] [9] [10] [11] [12] [13]

[14] [15]

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