Effect of Hydrogen on the Tensile Behavior of Austenitic Stainless Steels 316L Produced by Laser-Powder Bed Fusion

Abstract

Hydrogen was doped in austenitic stainless steel (ASS) 316L tensile samples produced by the laser-powder bed fusion (L-PBF) technique. For this aim, an electrochemical method was conducted under a high current density of 100 mA/cm² for three days to examine its sustainability under extreme hydrogen environments at ambient temperatures. The chemical composition of the starting powders contained a high amount of Ni, approximately 12.9 wt.%, as a strong austenite stabilizer. The tensile tests disclosed that hydrogen charging caused a minor reduction in the elongation to failure (approximately 3.5% on average) and ultimate tensile strength (UTS; approximately 2.1% on average) of the samples, using a low strain rate of 1.2 × 10⁻⁴ s⁻¹. It was also found that an increase in the strain rate from 1.2 × 10⁻⁴ s⁻¹ to 4.8 × 10⁻⁴ s⁻¹ led to a reduction of approximately 3.6% on average for the elongation to failure and 1.7% on average for UTS in the pre-charged samples. No trace of martensite was detected in the X-ray diffraction (XRD) analysis of the fractured samples thanks to the high Ni content, which caused a minor reduction in UTS × uniform elongation (UE) (GPa%) after the H charging. Considerable surface tearing was observed for the pre-charged sample after the tensile deformation. Additionally, some cracks were observed to be independent of the melt pool boundaries, indicating that such boundaries cannot necessarily act as a suitable area for the crack propagation.

1. Introduction

Additive manufacturing processes selectively add materials layer-by-layer to build 3D models. Among the most popular processes for the additive manufacturing of metals is laser-powder bed fusion (L-PBF) [1]. During this process, a high-energy laser beam irradiates the powder layer, and energy is absorbed by powder particles through bulk-coupling and powder-coupling, generating an extremely high temperature and a rapid cooling rate within the molten pool [2].

Austenitic stainless steels (ASSs) can be used in a wide range of applications due to their properties, such as a good combination of strength and ductility and corrosion resistance [3]. The 316L grade is a type of ASS, which has attracted the attention of engineers and researchers in both academic and industrial areas [4]. Numerous studies have shown that L-PBF of 316L ASS results in a microstructure different from that of traditional 316L ASS, which can cause improved mechanical responses [5,6]. The mentioned unique properties of ASSs, including 316L ASS, have made them a promising material for application in the hydrogen industry, especially in the production of gaseous or liquid hydrogen storage used in transportation vehicles [7]. However, it is well understood that the hydrogen embrittlement (HE) is a critical issue for this group of materials when used in highly deformed conditions [8,9,10,11,12,13,14]. This has stimulated researchers to study the HE of ASSs and to overcome the problem.

Recently, some researchers [15,16,17] have reported the response of laser-powder bed fusioned (L-PBFed) 304L and 316L in mild and extreme hydrogen environments. Baek et al. [16] reported the high resistance of L-PBFed 304L for hydrogen services, in which no phase transformation from austenite to martensite was detected, despite a meta-stable microstructure of 304L. They provided the tensile test results under the crosshead speed of 0.0012 mm/s, which corresponds to an initial strain rate of 10⁻⁴/s and a reduction in elongation from approximately 64 to 59% that occurred in a hydrogen atmosphere. However, a fracture surface with small and uniform dimples was observed for their samples in both the air and hydrogen atmospheres, indicating no sign of HE. In another study, Lee et al. [17] systematically investigated the hydrogen effects on the tensile ductility of 304L fabricated by L-PBF. The hydrogen was charged for one and five days using current densities of 20 and 50 mA/cm², respectively. A slight effect on ductility was reported in their study, with a small amount of α′ martensite (approximately 3 vol.%) at the strain level of 60%, suggesting that the high local mechanical stability of austenite could be responsible for the HE resistance. In addition, they found a negligible effect of the solidification cell structure (cellular structure emanates from the Marangoni effect and elemental segregation during the rapid cooling during the L-PBF process [18]) on austenite stability, as well as on H trapping behavior. It was reported in their study that cellular structure may not provide additional H trapping sites and thus a H content increment, while an increase in the dislocation density of the L-PBFed sample could be more responsible as the trapping sites, and HE susceptibility might not be increased by increasing the H content. Finally, they found that in the absence of deformation-induced martensite in meta-stable stainless steels (SSs), a slow crack growth can be expected during the tensile deformation. In another recent study, Park et al. [6] studied the H uptake and its influence on L-PBFed ASS 316L using a low current density of 5 mA/cm² for 6 and 24 h nanoindentation measurement. They found that the cellular structure of L-PBFed ASSs containing a high density of dislocations might not enhance the H absorption capacity and H-induced mechanical degradation.

Based on these recent findings, it seems that L-PBFed ASSs may have suitable HE resistance. However, the tensile behaviour of L-PBFed ASS 316L with a high Ni content for the stabilization of the austenite phase in an extreme hydrogen environment is lacking. In this study, ASS 316L powders that contain 12.9 wt.% Ni were used to examine the HE resistance of L-PBFed ASSs under extreme hydrogen environments. Shibayama et al. [19] reported that the crack growth rate under the hydrogen charging at a higher current density was higher than that at a lower current density, which denoted that a higher hydrogen fugacity promoted the faster growth of cracking. For this purpose, in this study, a high current density, which may lead to a higher amount of hydrogen after charging [20], was used for three days before the tensile test.

2. Materials and Methods

The gas-atomized Fe-Cr-Ni ASS pre-alloyed powder with a Gaussian size distribution ranging from 15 to 50 μm was used in this work. A cylindrical bar with a height of 60 mm and a diameter of 10 mm was vertically printed by an EOSINT M280 3D (EOS, Freiburg, Germany) printing system, with a discontinuous Yb-fiber laser and maximum power capacity of 200 W under argon atmosphere. The chemical composition of the L-PBFed bars was determined by an optical emission spectrometer (Oxford Instruments—PMI-Master Smart, UK). The combined parameter energy density (ED) is frequently used to evaluate printing parameters. In the present work, the value of ED was determined to be 100 J/mm³. The chemical composition of the L-PBFed 316L alloy is listed in the upper part of

Table 1. Chemical compositions of bulk L-PBFed austenitic stainless steel (ASS) 316L in conjunction with conditions of different experiments in this study.

	Chemical Compositions
Material	C	Mn	P	S	Cr	Mo	Ni	Si	Cu	Fe
L-PBFed SS316L	0.025	0.76	0.009	0.010	17.30	2.30	12.90	0.642	0.026	Bal.
	Condition of Experiments
Sample		Strain Rate (S−1)
Uncharged, UC-1.2		1.2 × 10−4
Charged, C-1.2		1.2 × 10−4
Charged, C-4.8		4.8 × 10−4

.

The L-PBF sample bars were machined to produce cylindrical dog-bone specimens with a gauge length of 16 mm and a diameter of 4 ± 0.1 mm, according to the ASTM E8/16a standard. Hydrogen was pre-charged into the tensile samples by the electrochemical method (Figure 1). The samples were immersed in a distilled water solution containing 3 wt.% NaCl and 0.3 wt.% NH₄SCN. The efficiency of this solution in the H permeation was reported several times in the literature [19,21,22,23,24,25,26,27,28,29]. Ammonium thiocyanate (NH₄SCN) is widely used as a reagent (hydrogen recombination poison) for promoting hydrogen absorption by steels, and during hydrogen charging, the solution was de-aerated with N₂ [22,30]. Samples were connected to a cathode, and a Pt coil was connected to the anode. To ensure absorption of a high amount of hydrogen into the samples, a high current density of 100 mA/cm² was applied for three days at room temperature. Three tensile tests per condition were performed at room temperature just 20 min after charging (to ensure minimum outgassing) using a Universal Testing Machine (STM-20, Santaam, Tehran, Iran). The lower part of Table 1 summarizes the used samples in this study with or without the hydrogen charging process tensile tested at two different strain rates.

After the tensile test, the microstructures of the samples and fractured surfaces were investigated using an optical microscope (OM) (3d imaging: OPTIKA B-383MET-Ponteranica, Italy) and an scanning electron microscope (SEM) (MIRA3 TESCAN, Czech Republic, SEM HV: 15.0 kV, window Type: Moxtek AP1.3). For this aim, the samples were mechanically ground with fine SiC papers (grit number up to 2000) and subsequently polished with 3 µm diamond suspension and finally with 0.05 µm silica suspension. Finally, to confirm the austenite stability, XRD analysis was performed using an Explorer (GNR-Analytical Instruments Group, Italy, CuKα, λ = 1.5406 Å, 40 kV, 30 mA, and step size 0.01).

3. Results

Figure 2 shows the microstructure of the L-PBFed 316L before the hydrogen charging. It is worth noting that the laser scanning direction was rotated by 67° after each layer, leading to differently oriented melt pools, as depicted. Many grains include melt pool boundaries (MPBs), suggesting the epitaxial growth of new grains from re-melted zones [31]. The presence of MPBs and such epitaxial growth, which forms a texture, is among the specific characteristics of the L-PBFed materials that can directly affect their mechanical properties along different axes.

In Figure 3, the tensile curves—in conjunction with tensile properties, such as tensile yield strength (YS), ultimate tensile strength (UTS), and elongation (EL), of specimens—are illustrated, from which the following results can be extracted. It is well understood that L-PBF causes the formation of cell structures and a high density of dislocations, which results in the strengthening of ASSs [32]. The charging of hydrogen caused lower values of elongation and strength, which means that the overall performance of these steels with a high Ni content is suitable after the intensive pre-charging of H.

Figure 4 shows the cross-sectional microstructures of fractured tensile samples. The following results from Figure 4 can be summarized as follows: (1) In the case of samples without hydrogen charging (Figure 4a,b), there were some fine cracks from which the fracture could be initiated. These fine cracks are marked by the white arrows in Figure 4a. (2) After hydrogen charging (Figure 4c–f), the cracks were larger, as indicated by the white arrows in Figure 4c,d. The cracks were propagated along the MPBs (dashed white lines in Figure 4e). No sign of MPB-induced cracking can be observed in Figure 4f, in contrast with Figure 4b,e, indicating that MPBs do not necessarily act as suitable zones for crack propagation, and there might be multiple sources of crack initiation and propagation in H-charged L-PBFed SSs. Moreover, the deformation mechanisms of all samples were controlled by slip and twinning, due to the formation of large amounts of mechanical twins inside the grains.

The SEM micrographs of tensile fractured surfaces in conjunction with energy-dispersive X-ray spectroscopy (EDS) analysis are shown in Figure 5 and Figure 6. The fracture surfaces included both the cleavages and dimples as the characteristics of brittle and ductile fracture modes, respectively. However, the ductile mode was the dominant fracture mechanism. In addition, there were some porosities in the fracture surface of the L-PBF samples (white arrow in Figure 5b). Moreover, the EDS analysis of points A and B in Figure 5b showed that in point A, beside the voids, a segregation of alloying elements occurred. There was a similar trend in the case of points A and B in Figure 6 for C-4.8 sample. The segregation of elements in L-PBFed stainless steel 316L has been reported many times [33,34]. A high solidification rate and lack of equal mobility of elements during solidification are the main reasons for obtaining such elemental segregation that may cause the formation of a dislocation wall and can improve the mechanical properties of such steels [33,34].

4. Discussion

In this study, hydrogen charging caused a decrease in the ductility of UTS of 316L ASS based on tensile properties and fracture modes, illustrated in Figure 3, Figure 5 and Figure 6, which agrees with previous reports [15,16,17]; in which a suitable performance of L-PBFed microstructures against H was reported due to the presence of the cellular structure. Baek et al. [16]—in the case of 304L—concluded that in the presence of hydrogen, the strain-induced plasticity transformation of austenite to martensite is postponed in extensive strains, which results in larger elongations compared to rolled 304L. According to the microstructures of deformed samples and XRD analysis results (see Figure 4 and Figure 7, respectively), the martensite phase was not observed. This can be explained by the high content of Ni as a strong austenite stabilizer, where the Cr amount as a ferrite stabilizer is not more than 17.3 wt.%. A slight reduction in UTS*UE after hydrogen charging might be explained by the HELP mechanism [35], although no considerable embrittlement was observed for the C-1.2 sample. However, increasing the strain rate caused a decrease in the average amount of UTS*UE. It was already explained by [20] that a higher current density can cause damage on the surface of the ASSs, and, therefore, from Figure 3, it seems that the surface damage caused by a higher strain rate was more effective on the overall tensile behaviour of L-PBFed ASS. Figure 8 shows the presence of more surface cracks (surface tearing and damage) on the longitudinal direction of the broken sample. It should be noted that such surface cracks are independent of the manufacturing process, and the L-PBF prior to the H charging may not be responsible for their formation. Such surface cracks in the longitudinal direction of failed samples were also detected and reported in [25] for the casted/annealed samples.

Irrespective of the surface damage caused by H charging that is independent of the manufacturing process, the origin of the lower UTS*UE of hydrogen-doped 316L ASS can be explained by the following observations: The MPBs can act as nucleation and propagation sites for hydrogen-assisted cracking (see Figure 4). Thus, their presence is among the possible reasons for the lower UTS*UE. It should be noted that no sign of H-assisted cracks for grain boundary–twin boundary–slip line–mechanical twin interactions was found in another study [17] for SSs. During the tensile loading, a high density of dislocations formed by the high cooling rate of solidification [36] could be responsible for H trapping, and hence, they might have a negative effect on the average UTS*UE amount [37,38]. In addition, the cellular structure with a high density of dislocation around the cell wall’s structure can absorb fewer hydrogen atoms, and thus, they might not be responsible for the H trapping [17,20]. From all the findings of the current study and the literature, it can be concluded that the high density of dislocations after L-PBF for H trapping—which might affect the HELP mechanism as well as the presence of MPBs—is the possible cause of the slight L-PBF-induced reduction in UTS*UE after H charging on L-PBFed SS 316L. In addition to L-PBF-induced causes, the surface damage after electrochemical H charging is another important cause of the UTS*UE reduction.

5. Conclusions

The tensile behavior, microstructures, and fracture surfaces of 316L ASSs produced by the L-PBF process were compared before and after electrochemical hydrogen charging. For this purpose, an extreme hydrogen environment including a high current density of 100 mA/cm² was applied for three days. The hydrogen charging caused a minor loss in the ductility and UTS of the samples. Approximately 3.5% in elongation to failure and 2.1% in the UTS of the samples were obtained using a low strain rate of 1.2 × 10⁻⁴ s⁻¹. It was also found that an increase in the strain rate from 1.2 × 10⁻⁴ s⁻¹ to 4.8 × 10⁻⁴ s⁻¹ led to a reduction of approximately 3.6% on average for the elongation to failure and 1.7% on average for the UTS in the pre-charged samples, with a higher detection of surface cracks (caused by the electrochemical H charging process) on the broken samples. It was found that the MPBs, which are among the specific features of L-PBFed structures, did not necessarily cause the crack propagation on the pre-charged surfaces. However, the total reduction in UTS*UE (GPa%) for the samples after being exposed to the extreme hydrogen environment was minor and based on the XRD results of the fractured sample; the usage of a high Ni content to stabilize the austenite structure should be mentioned as an important reason for this minor reduction.