Hydrogen resistance of a 1 GPa strong equiatomic CoCrNi medium entropy alloy

Journal Pre-proof Hydrogen resistance of a 1 GPa strong equiatomic CoCrNi medium entropy alloy Chandrahaasan K. Soundararajan, Hong Luo, Dierk Raabe, Zhiming Li

PII:

S0010-938X(19)31616-6

DOI:

https://doi.org/10.1016/j.corsci.2020.108510

Reference:

CS 108510

To appear in:

Corrosion Science

Received Date:

9 August 2019

Revised Date:

10 January 2020

Accepted Date:

27 January 2020

Please cite this article as: Soundararajan CK, Luo H, Raabe D, Li Z, Hydrogen resistance of a 1 GPa strong equiatomic CoCrNi medium entropy alloy, Corrosion Science (2020), doi: https://doi.org/10.1016/j.corsci.2020.108510

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Hydrogen resistance of a 1 GPa strong equiatomic CoCrNi medium entropy alloy Chandrahaasan K. Soundararajan1, Hong Luo2,1*, Dierk Raabe1, Zhiming Li1,3*

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1, Max-Planck-Institut für Eisenforschung, Max-Planck-Straße 1, 40237 Düsseldorf, Germany 2, Institute for Advanced Materials and Technology, University of Science and Technology Beijing, Beijing 100083, China 3, School of Materials Science and Engineering, Central South University, Changsha 410083, China Correspondence to: H. Luo ([email protected]), Z. Li ([email protected])

The influence of hydrogen on the deformation behavior in an equiatomic CoCrNi

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medium entropy alloy (MEA) is investigated.

Hydrogen influences the deformation microstructure of the MEA.

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The primary hydrogen-induced failure mode of the alloy is intergranular cracking.

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Two main competing mechanisms are found in the CoCrNi MEA when exposed to

Abstract

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hydrogen.

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In this work, we study the influence of hydrogen on the deformation behavior and microstructure evolution in an equiatomic CoCrNi medium entropy alloy (MEA) with an ultimate tensile strength of ~1 GPa. Upon deformation, hydrogen-charged samples exhibit enhanced dislocation activity and nanotwinning. Hydrogen shows both positive and negative effects on the deformation behavior of the CoCrNi MEA. More specifically, it weakens grain boundaries during loading, leading to intergranular cracking. Also, it promotes the formation of twins which enhance the material’s resistance to crack

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propagation. The underlying mechanisms responsible for the hydrogen resistance of the CoCrNi MEA are discussed in detail.

Keywords: Hydrogen; Medium entropy alloy; Deformation; Microstructure; Twinning

1. Introduction High entropy alloys (HEAs) containing multi-principal elements have opened a new field in

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alloy design [1-4]. The versatility of these alloys in adjusting compositions over wide ranges and thereby tuning intrinsic features such as solid solution hardening and the stacking fault energy

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for achieving enhanced mechanical properties are attractive features when designing improved materials [5-7].

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Hydrogen in metals is a long-standing research topic owing to its embrittling effects which can initiate catastrophic failure, a phenomenon known as hydrogen-embrittlement [8-11]. Most

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engineering metallic alloys, particularly those with strength levels beyond 700 MPa, are at least to some degree susceptible to this detrimental phenomenon, yet, possibly due to quite different

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underlying mechanisms in the individual materials. Recently, several studies have dealt with hydrogen embrittlement effects in HEAs [12-16]. The equiatomic CoCrFeMnNi HEA shows relatively high resistance to hydrogen embrittlement when exposed to hydrogen concentrations

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of approximately 80 wt. ppm [12, 17], but it still undergoes embrittlement at higher concentrations [13]. Although addition of interstitial carbon to this material leads to an increase in strength [18, 19], deformation in the presence of hydrogen results in the initiation of microcracks around carbides and hence C-doped material shows higher susceptibility to hydrogen embrittlement compared to the C-free reference alloy [20]. To date, several mechanisms have

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been proposed to explain such embrittlement phenomena in metals, namely, hydrogen-enhanced localized plasticity (HELP) [21-24], hydrogen-enhanced decohesion (HEDE) [25-27], as well as hydrogen-stabilized superabundant vacancy enrichment [28, 29], hydride formation [30, 31] and microvoid coalescence [22, 32]. In the CoCrFeMnNi alloy family and its lower entropy variants, the so called medium entropy alloys (MEA), specifically the equiatomic ternary CoCrNi alloy shows a higher strength-ductility combination than many other related binary, ternary, and quaternary variants [33, 34]. Also, its 2

fracture toughness is superior to that of the widely studied equiatomic CoCrFeMnNi HEA [35]. Though the mechanical properties and deformation mechanisms of the equiatomic CoCrNi MEA were recently revealed in more detail, nothing is known about its deformation behavior in presence of hydrogen. Yet, hydrogen effects in high and medium entropy alloys are of high interest in this alloy class as both, detrimental and strain hardening effects were observed and improved hydrogen-resistant alloys are urgently needed for future hydrogen-exposed transportation and industry applications. Here we thus study the influence of hydrogen in the

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equiatomic CoCrNi MEA by using tensile testing with hydrogen pre-charging. The mechanisms of hydrogen-induced cracking and the evolution of the deformation microstructure are probed via the correlative use of electron backscatter diffraction (EBSD) and electron channeling contrast

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imaging (ECCI).

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2. Materials, Experimental Methods and Probing Procedures 2.1. Alloy Preparation

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The equiatomic CoCrNi MEA was synthesized from pure metals (>99% purity) by melting and casting using a vacuum induction furnace. The as-cast ingot had a dimension of 65 mm × 25

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mm × 80 mm, and the chemical composition measured by wet-chemical analysis is shown in Table. 1. The ingot was hot rolled at 1050 °C to a thickness reduction of 50% and subsequently homogenized at 1200 °C for 2 hours in Ar atmosphere. To obtain a fine-grained microstructure,

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the homogenized alloy plate was cold rolled to a thickness reduction of 60% and annealed at 800 °C for 30 minutes in Ar atmosphere followed by water quenching. Table. 1: Chemical composition of the CoCrNi alloy (in wt. %) according to wet-chemical analysis.

Co

Cr

Ni

Fe

O

C

S

wt. %

35.21

30.70

34.00

0.057

0.012

0.0073

0.0013

33.76

33.36

32.73

0.057

0.042

0.034

0.0023

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Elements

at. %

2.2. Hydrogen charging and mechanical characterization Prior to hydrogen charging, the sample surfaces were ground with silicon carbide papers from 600 to 2500 granulation. Hydrogen was then introduced into the sample through electrochemical

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charging. We used 0.5 M H2SO4 solution along with 5 g/L ammonium thiocyanate addition (hydrogen recombination agent) as an electrolyte. A platinum sheet was used as a counter electrode. Charging was conducted at 25 °C for two different conditions. One set of samples were pre-charged at 25 mA/cm2 for 1 day, and another set was pre-charged at 25 mA/cm2 for 5 days. Each set consisted of three samples to ascertain reproducibility during tensile testing. After charging, the samples were rinsed in distilled water and subsequently dried before the tensile test. The surfaces of the charged samples did not show any signs of macro-cracks or corrosion products. Tensile testing was immediately started for all samples using a Kammrath & Weiss

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tensile stage at an initial strain rate of 10-4 s-1. Flat samples with thickness of 1 mm, gauge length of 10 mm and gauge width of 3 mm were used for the tensile testing [36]. The Digital Image

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Correlation (DIC) technique was employed to determine the local plastic strain during tensile deformation [37, 38]. The amount of hydrogen in each charging condition was measured by a

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custom-designed UHV based Thermal Desorption Analysis instrument in conjunction with a mass spectrometer detector set-up (TDA-MS) from 25 °C to 800 °C, and the corresponding

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heating rate was 26 °C min−1. The dimension of the samples used for TDS experiments was 15 mm × 15 mm × 1 mm. The total hydrogen content was determined by measuring the cumulative

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desorbed hydrogen from 25 °C to 800 °C. The amount of hydrogen-charged after 1 day and after 5 days, respectively, was 42.73 wt. ppm and 52.78 wt. ppm, showing that the hydrogen uptake is not much increasing in this material even when charging it for 5 days (increase by only ~10.05

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wt. ppm). The relatively small increase in total hydrogen concentration in 5 days charged sample is due to the relatively low coefficient of hydrogen diffusion in the material and the formation of passive layer on the surface. The chromium content in the material is high (33.36 at. %), which promotes the formation of the passive layer, thereby inhibiting hydrogen diffusion into the

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material.

2.3. Microstructural characterization After tensile testing, the fracture surfaces were examined using scanning electron microscopy

(SEM). Then the samples were ground and polished in the ND-RD plane for microstructure investigation. With the aid of the DIC data, electron backscattered diffraction (EBSD) mapping was conducted over different regions of the sample to record the microstructure at different local strains. All EBSD measurements were conducted in a JEOL JSM 6500F SEM, equipped with the 4

TSL OIM software for collecting and indexing the Kikuchi patterns, operated at an acceleration voltage of 15 kV and a beam current of 3 nA. The annealing twin fraction was evaluated using TSL OIM data analysis software with multiple EBSD maps. The Kernel average misorientation (KAM) was calculated with respect to the first nearest neighbors, serving as an approximate measure for local lattice curvature and the associated geometrically necessary dislocation content [39, 40]. Electron channeling contrast imaging (ECCI) analysis of the deformation microstructures was performed on a Zeiss-Merlin instrument at an acceleration voltage of 30 kV

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and a beam current of 3 nA. ECC images were captured on the same regions probed before by EBSD to correlate the microstructural features with the corresponding crystallographic orientations. X-ray diffraction (XRD) measurements were performed using an X-ray BRUKER-

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a step size (Δ2θ) of 0.03° and count time of 30 s per step.

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X2 instrument equipped with Co Kα (λ = 1.788965 Å) radiation operated at 40 kV and 30 mA at

3. Results

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3.1. Recrystallized microstructure

Fig. 1 shows the initial microstructure of the CoCrNi MEA after cold rolling and annealing at

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800 °C for 30 minutes. The inverse pole figure (IPF) and phase maps of the same sample region are shown in Figs. 1a and 1b, respectively. The alloy exhibits a single-phase face-centered cubic (FCC) structure. The EBSD (Fig. 1a-b) and ECC results (Fig. 1c) show a fully recrystallized

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microstructure along with a high fraction of annealing twins (~56%) in the CoCrNi MEA. The grain size of the microstructure is in the range of 1~10 µm. The XRD result in Fig. 1d confirms

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the single FCC structure.

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Fig. 1. Recrystallization microstructure of the CoCrNi MEA after cold rolling and annealing at 800 °C for 30 minutes. a) Inverse pore figure (IPF) map; b) Phase map overlapped with high angle grain boundaries

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>15° (HAGB) and twin boundaries (TB); c) ECC image; d) XRD pattern.

3.2. Tensile properties of the medium entropy alloy with and without hydrogen Fig. 2 shows the tensile test results for H-charged and uncharged specimens. The strain hardening curves are shown in Fig. S1. Increase in yield strength (YS) of about 4% and 8% for 1 day (42.73 wt. ppm H) and 5 days (52.78 wt. ppm) charged samples was observed, respectively.

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Marginal increase (~3%) in the ultimate tensile strength (UTS) for both 1 day (42.73 wt. ppm H) and 5 days (52.78 wt. ppm) charged samples was observed. Interestingly, there is very little or no significant difference in fracture strain (58%) between the 1 day (42.73 wt. ppm H) charged and uncharged samples. The hydrogen enhanced plasticity was not observed in this MEA probably due to the relatively high hydrogen concentration compared to that reported in a previous work [12]. The hydrogen charged samples show significant stress drops at engineering strains higher than ~15% (Fig. 2a). This is associated with the formation of surface cracks under high strain 6

with high concentration of hydrogen and the slightly earlier necking of the charged samples. The overall differences in yield strength, ultimate tensile strength, and fracture elongation values for

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hydrogen pre-charged and uncharged samples lie within 10% (Fig. 2b).

Fig. 2. a) Typical engineering stress-strain curves for uncharged and hydrogen-charged samples; b)

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Quantification of important values averaged over several engineering stress-strain curves (The I bar indicates standard deviation). The total diffusible hydrogen concentrations (HTotal) accumulated in

results are indicated in (a).

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samples charged for 1 and 5 days, respectively, were measured using thermal desorption analysis and the

3.3. Fracture morphology after tensile testing

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An overview of the fracture surfaces of uncharged, as well as of the 1 (42.73 wt. ppm H) and 5 days (52.78 wt. ppm) charged samples is provided in Figs. 3a, b and c, respectively. The fractography shown in Fig. 3 corresponds to the samples with stress-strain curves shown Fig. 2a. The fracture surface at the core and edge of the sample without hydrogen charging (Fig. 3a, 3d) reveals microvoid coalescence, indicating ductile fracture. In the 1 day (42.73 wt. ppm H)

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hydrogen-charged sample, three distinct zones, i.e., brittle, transition, and ductile regions were identified (Fig. 3e). This type of subdivision is related to the hydrogen concentration gradient after electrochemical charging. At the surface of the sample, with its high hydrogen concentration, brittle fracture prevails, characterized by intergranular cracking (Fig. 3e). According to the hydrogen diffusion coefficient of FCC materials [41], the hydrogen concentration near the surface is very high, which can be more than 3 times higher than the average hydrogen concentration in the sample. At the core of the sample, with its low hydrogen 7

concentration, ductile fracture dominates as revealed by multiple microvoids and their coalescence (Fig. 3e). A mixed fracture mode is observed in the transient region between surface and core. In the 5 days (52.78 wt. ppm) hydrogen-charged sample the same trend is found as in the 1 day (42.73 wt. ppm H) charged sample, indicating that prolonged hydrogen exposure does not fundamentally alter the failure modes and their associated microstructure characteristics. The only difference is that the extension of the brittle region in the 5 days (52.78 wt. ppm) charged sample (~66 µm) (Fig. 3f) is larger than that in the 1 day (42.73 wt. ppm H) charged sample (~37

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µm), suggesting that prolonged charging affects a larger portion of the sample, hence creating more hydrogen embrittled regions in the samples. This observation points at the important role of hydrogen transport in the current alloy. Furthermore, the sizes of dimples at the core of the

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samples were slightly decreased with an increase in charging time, which can be explained as a

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mark of shear deformation localization caused by hydrogen [42, 43].

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Fig. 3. Fractured surfaces of the CoCrNi MEA after tensile testing under different charging conditions. a) and d) Uncharged sample; b) and e) 1 day (42.73 wt. ppm H) hydrogen-charged sample; c) and f) 5 days (52.78 wt. ppm) hydrogen-charged sample.

3.4. Deformation microstructure For the uncharged sample, with the aid of the DIC mapping, EBSD maps were recorded at different sample regions corresponding to different local strain levels (5%, 20%, and 60%). Figs. 8

4a, 4c and 4e show the KAM map for local strains of 5%, 20% and 60% respectively. At low strains up to 5%, local orientation gradients evolve at grain boundaries (Fig. 4a). The average KAM value increases from 0.34° for 5% local strain to a value up to 1.84° for 60% local strain. At 60% local strain Kikuchi pattern recognition quality has decayed as indicated by undetected orientation spots (black pixels) in Fig. 4e-f. Previous work reported that above 50% true strain, the equiatomic CoCrNi MEA develops a nanotwin-hexagonal close-packed (HCP) lamellae structure [44]. However, these features are hard to detect in the current case by transmission

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Kikuchi diffraction (TKD) or diffraction contrast imaging using STEM [45], due to the very small size (several nanometers) of these features and the small sample size utilized by these techniques. From the phase maps shown in Figs. 4b, 4d, and 4f, no obvious deformation-induced

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HCP structures were observed, probably due to the limited EBSD resolution.

Fig. 4. EBSD maps of the deformation microstructure of uncharged samples at different local strain levels after tensile testing (ND-RD plane). (a-b) 5%; (c-d) 20%; (e-f) 60%. (a, c, e) KAM maps correspond to 5%, 20% and 60% local strain regions, respectively. (b, d, f) Phase maps correspond to 5%, 20% and 60% local strain regions, respectively. “TD” refers to tensile direction, “HAGB” refers to high angle grain

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boundaries (>15° misorientation).

Correlative ECCI-EBSD was done on previously recorded regions to reveal the individual

deformation mechanisms of the CoCrNi MEA with and without hydrogen charging. Fig. 5 shows the ECC images of uncharged samples at different strain levels. At 5% local strain planar dislocation slip bands emerge (Fig. 5b). Also, short stacking faults form within the grains. At 20% local strain, besides partial dislocations and stacking faults, deformation-induced

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nanotwins appear (Fig. 5c). At 60% local strain, a large density of deformation-induced

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nanotwins and shear bands were found (Fig. 5d).

Fig. 5. ECC images of the CoCrNi MEA captured at different local strain levels of tensile samples in ND-

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RD plane without hydrogen charging. (a) The un-deformed initial microstructure of the recrystallized alloy; (b) Deformation structure at 5% local strain showing dislocation shear and short stacking faults; (c) Deformation structure at 20% local strain with stacking faults and nanotwins; (d) Deformation structure at 60% local strain showing a high density of nanotwins and shear bands (marked by yellow arrows).

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Fig. 6 compares the ECC images at the regions corresponding to 5% and 60% local strains

for the uncharged and hydrogen-charged samples. The uncharged sample (Fig. 6a) shows a low dislocation density at 5% local strain. In contrast, the 5 days (52.78 wt. ppm) hydrogen-charged sample (Fig. 6b) shows a high density of dislocation slip bands, indicating that solute hydrogen enhances strain localization and slip planarity [24, 46]. Several long stacking faults (~1 µm length) were also observed in the charged samples at 5% local strain, similar to what has been reported about the hydrogen-charged FCC phase in duplex stainless steels [47]. No voids were 10

observed at the intersections of dislocation slip bands, different from what has been seen in Nibase alloys [48]. Fig. 6c and 6d show the microstructures of the uncharged and 5 days (52.78 wt. ppm) hydrogen-charged CoCrNi MEA at 60% local strain, respectively. Deformation-induced nanotwins were observed in both samples, however, the hydrogen-charged samples exhibit a higher density of nanotwins, and the spacing between the twins has reduced from ~260 nm in the

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uncharged state to ~200 nm in the 5 days (52.78 wt. ppm) hydrogen-charged state.

Fig. 6. ECC images of the samples at different local strain levels (not in the cracking regions) with and without hydrogen charging. (a) 5% local strain without hydrogen charging; (b) 5% local strain with 5

days (52.78 wt. ppm) hydrogen charging (image recorded at the central region of the sample); (c) 60% local strain without hydrogen charging; (d) 60% local strain with 5 days (52.78 wt. ppm) hydrogen charging (image recorded at the central region of the sample). DSBs, NTs, SFs refer to dislocation slip bands, nanotwins and stacking faults, respectively.

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3.5. Hydrogen-induced cracking The 1 day (42.73 wt. ppm H) and 5 days (52.78 wt. ppm) hydrogen-charged samples show similar cracking behavior; therefore only the results for the 5 days (52.78 wt. ppm) charged sample are presented here. Fig. 7 shows EBSD maps and ECC images recorded at the cracked region of the hydrogen-charged sample after the tensile test. As shown in the IPF map (Fig. 7a), the material exhibits intergranular cracking in the presence of hydrogen, a similar trend as reported for several other FCC HEAs [12-14]. The KAM map (Fig. 7b) taken at the cracked

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regions (e.g., the region ‘e’) shows a relatively low value (average KAM value: 0.48°), suggesting that these regions did not experience large plastic deformation. Interestingly, ECC

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images of the sample region near the crack reveal a large number density of nanotwins (Fig. 7d), suggesting that hydrogen is likely to promote nanotwinning in this MEA. However, the KAM

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measure is not sufficiently highly resolving for mapping this effect as misorientation pattern.

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Fig. 7. Combined EBSD and ECCI analysis of a crack in the surface regions of a 5 days (52.78 wt. ppm) hydrogen-charged sample after tensile testing (ND-RD plane). The observed region was in the surface area with relatively high hydrogen concentrations. (a) IPF map; (b) KAM map considering the first nearest-neighbor orientation points; (c) Phase and boundary map; (d) ECC image reveals the microstructure around the crack; (e) Enlarged ECC image of the marked region in (b) and (d). “HAGB” refers to high angle grain boundaries (>15° misorientation).

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Fig. 8 shows the ECC images of a primary crack together with several secondary cracks in the 5 days (52.78 wt. ppm) hydrogen-charged sample after tensile deformation. Besides the main crack, secondary crack propagation was observed. Fig. 8b shows that such a secondary crack progresses along the nanotwin boundaries, promoting further crack propagation. In another case (Fig. 8c), the secondary crack interacts with nanotwins at an oblique angle, impeding crack propagation. Coherent twin boundaries are usually assumed to have high resistance to hydrogeninduced crack propagation [49] due to their high lattice match and low hydrogen solubility.

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However, in some cases, they were reported to act as initiation and propagation sites for cracks [50]. Our present results suggest that nanotwins in the current MEA can have both, promoting

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and impeding effects on crack propagation, depending on their angle of interaction.

Fig. 8. (a) ECC image of the main crack and secondary cracks in the 5 days (52.78 wt. ppm) hydrogencharged CoCrNi MEA after tensile testing; The observed region was in the surface area with relatively

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high hydrogen concentration. (b) A secondary crack meets with the nanotwins and propagates along the nanotwin boundaries; (c) Another secondary crack gets deflected by nanotwins. “NTBs” refers to Nanotwin boundaries.

4. Discussion 4.1. Mechanism of hydrogen-induced intergranular cracking

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The elongation to fracture value was slightly reduced from ~58% for the uncharged alloy to ~54% for 5 days (52.78 wt. ppm) hydrogen-charged sample, suggesting that hydrogen does not notably embrittle the CoCrNi MEA. This means that the CoCrNi MEA with its ultimate tensile strength of ~1 GPa has a much higher resistance to hydrogen embrittlement than many conventional alloys such as the Ni-based superalloy 718 [51], High-Mn steels [52] and DP steels [53], which were evaluated using the similar hydrogen charging conditions and slow strain rate tensile testing. The damage mode in the CoCrNi MEA is hydrogen-induced intergranular

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cracking (Fig. 7). Although this effect was also reported for many other materials, e.g., Ni [54], Fe [55], and HEAs [14, 15, 20], the underlying mechanisms are discussed controversially. Several mechanisms, i.e., hydrogen-enhanced localized plasticity (HELP), hydrogen-enhanced

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decohesion (HEDE), and adsorption-induced dislocation emission (AIDE) have been proposed to explain intergranular cracking [56]. Martin et al. [54] characterized deformation microstructures

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near the fracture surface under the presence of hydrogen. They concluded that hydrogen-induced intergranular cracking is driven by hydrogen-enhanced dislocation mobility. In contrast, Harris et

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al. [57] pointed out that intergranular cracking still occurs at 77 K. At this temperature interactions between mobile hydrogen and dislocations are frozen in as the hydrogen can no

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longer follow, decorate and screen the dislocations. Other authors have suggested that the hydrogen together with other embrittling elements (e.g., Sulphur and Phosphorus) tends to decorate grain boundaries, promoting decohesion and thus causing intergranular failure [58].

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However, the contents of embrittling elements in the current MEA are extremely low (Table 1), and their effects on the mechanical properties of the alloy can be ignored. In the present work, based on the KAM map shown in Fig. 7b, besides the cracked regions with low strain localization, high strain localization occurs at grain boundaries of the deformed specimens compared to the interior of the grains. This suggests that intergranular failure is in the current

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case indeed likely to be attributed to the stress-driven enrichment of hydrogen at grain boundaries [59]. It is however not entirely clear from the present results whether the interfacial cracks are initiated alone by decohesion or if their formation is also supported by hydrogenenhanced slip. 4.2. Hydrogen effect on deformation mechanisms and strain hardening The stacking fault energy (SFE) plays an essential role in altering the deformation mode of FCC metals and alloys [60]. The SFE of CoCrNi is reported as ~22 mJ/m2 [61]. Due to this low 14

value, the recrystallization microstructure of the CoCrNi MEA contains a large number of annealing twins. Planar slip and stacking faults prevail at low strains (~5%) while twinning starts at higher local strain (~60%). These are typical deformation characteristics of FCC metals with low SFE, and consistent with previously reported results [45]. The dynamic Hall-Petch effect associated with nanotwinning is mainly driving the high strain hardening and ductility of the CoCrNi MEA [62, 63]. More specifically, the deformation-induced nanotwinning acts as a barrier against dislocation motion, reducing their mean free path and increasing the work

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hardening rate. Comparison of the deformation behavior of hydrogen-charged and uncharged samples at 5% local strain shows the significant role of hydrogen in that context: The HELP effect promotes

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formation of highly localized and coarse slip bands at the surface of the charged samples [64], and it was also observed in the present MEA as shown in Fig. 6b. In addition the hydrogen-

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charged samples exhibit a higher density of nanotwins when compared to uncharged samples (Fig. 6d) at the same strain (60%), suggesting that the hydrogen reduces the SFE of the MEA

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[12, 65]. This effect was also shown in simulations for the case of pure Ni [66]. These two effects, viz., the reduction in SFE and the HELP-enhanced dislocation activity due to the

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presence of hydrogen, reduce the activation energy for twin formation, promoting a higher nucleation rate of mechanical twins during mechanical loading [67]. The marginal increase in yield strength of hydrogen-charged samples can be ascribed to the

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interstitial strengthening. Hydrogen exists in the lattice as an interstitial solute, prefers to occupy octahedral sites and thus imposes symmetric cubic distortion fields in the lattice [70]. During plastic deformation, hydrogen interstitial solute atmospheres can exert a drag force on moving dislocations, which may slightly increase the flow stress. Progressive increase in yield strength with higher charging time is attributed to the larger intrusion depth of hydrogen into the matrix.

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The enhanced twinning activity owing to hydrogen at the later deformation stages leads to an increase in the ultimate strength observed for the charged samples through the dynamic HallPetch effect.

5. Conclusions

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The resistance of the 1 GPa strong CoCrNi MEA against hydrogen embrittlement was investigated using tensile testing of hydrogen pre-charged samples and combined EBSD and ECCI analysis of the deformation microstructures. The main conclusions are: 1. The equiatomic CoCrNi MEA shows relatively high resistance to hydrogen embrittlement compared to conventional alloys, which were evaluated by electrochemical hydrogen charging and the similar strain rate tensile testing. 2. Hydrogen influences the deformation microstructure of the MEA. As observed in the grain

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interiors, at low local strain (~5%), the dislocation tangle is clearly observed in the hydrogencharged sample, which is attributed to an increased dislocation density compared to the uncharged sample. At high local strain (e.g., ~60%), the number density of nanotwins formed

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under the presence of hydrogen is higher than that in samples devoid of hydrogen.

3. The primary hydrogen-induced failure mode of this alloy is intergranular cracking. Based on

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the KAM results, the stress concentration at grain boundaries were clearly observed in the presence of hydrogen, which may act as the cracking initiation sites.

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4. Two main competing mechanisms are found in the CoCrNi MEA when exposed to hydrogen, namely, hydrogen-induced intergranular cracking as weakening mechanism and hydrogen-

mechanism.

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promoted mechanical nanotwinning causing dynamic strain hardening as a strengthening

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Author statement

We have made substantial contributions to the conception or design of the work; or the acquisition, analysis, or interpretation of data for the work; AND

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We have drafted the work or revised it critically for important intellectual content; AND We have approved the final version to be published; AND We agree to be accountable for all aspects of the work in ensuring that questions related to the accuracy or integrity of any part of the work are appropriately investigated and resolved. All persons who have made substantial contributions to the work reported in the manuscript, including those who provided editing and writing assistance but who are not authors, are named in the Acknowledgments section of the manuscript and have given their written permission to be 16

named. If the manuscript does not include Acknowledgments, it is because the authors have not received substantial contributions from nonauthors.

Data Availability The raw/processed data required to reproduce these findings cannot be shared at this time as the data also forms part of an ongoing study.

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Conflict of Interest

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We declare that we have no financial and personal relationships with other people or

organizations that can inappropriately influence our work, there is no professional or other

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personal interest of any nature or kind in any product, service and/or company that could be

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construed as influencing the position presented in, or the review of, the manuscript entitled.

Acknowledgments:

The kind support of F. Schlüter, B. Breitbach, M. Nellessen and K. Angenendt at the MaxPlanck-Institut für Eisenforschung is gratefully acknowledged. H. Luo would like to

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acknowledge the final support from Supported by PetroChina Innovation Foundation (2019D5007-0308); Z.L. would like to acknowledge the final support from the Hunan Special Funding for the Construction of Innovative Province (Grant No. 2019RS1001) and the Innovation-Driven Project of Central South University (Grant No. 2020CX023).

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