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Corrosion Science, Vol. 216, p. 111097, 2023
https://doi.org/10.1016/j.corsci.2023.111097
High strength and high ductility of a severely deformed high-entropy alloy
in the presence of hydrogen
Abbas Mohammadi1,2, Payam Edalati3, Makoto Arita4, Jae Wung Bae5,
Hyoung Seop Kim6,7 and Kaveh Edalati1,*
1 WPI, International Institute for Carbon-Neutral Energy Research (WPI-I2CNER), Kyushu
University, Fukuoka 819-0395, Japan
2 Department of Materials Science and Engineering, The Ohio State University, Columbus,
OH, 43210, USA
3 Department of Materials and Metallurgical Engineering, Amirkabir University of
Technology, Tehran, Iran
4 Department of Materials Science and Engineering, Faculty of Engineering, Kyushu
University, Fukuoka 819-0395, Japan
5 Department of Metallurgical Engineering, Pukyong National University, Busan, 48513,
Republic of Korea
6 Graduate Institute of Ferrous & Energy Materials Technology, Pohang University of Science
and Technology, Pohang, 37673, Republic of Korea
7 Institute for Convergence Research and Education in Advanced Technology, Yonsei
University, Seoul, 03722, South Korea
Abstract
High-strength materials usually exhibit low ductility, particularly in the presence of hydrogen
embrittlement phenomena. In this study, three strategies are combined to achieve excellent
strength-plasticity combinations in the presence of hydrogen: (i) selecting an FCC high-entropy
alloy with slow hydrogen lattice diffusion, (ii) adding aluminum to the alloy to hinder surface-
to-bulk hydrogen diffusion, and (iii) introducing low-mobility lattice defects like nanotwins
and Lomer-Cottrell locks by severe plastic deformation to suppress hydrogen-enhanced
localized plasticity and stress concentration. The Al0.1CrFeCoNi alloy severely deformed by
high-pressure torsion exhibits an ultrahigh yield strength of 1.96 GPa and a high elongation to
failure of 10%.
Keywords: Hydrogen compatibility; Multiprincipal-element alloys; High-entropy alloys
(HEAs); Ultrafine-grained (UFG); Severe plastic deformation (SPD); High-pressure torsion
(HPT)
*Corresponding author (E-mail: kaveh.edalati@kyudai.jp; Tel: +81-92-802-6744)
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1. Introduction
Achieving high strength is a major goal in many engineering applications, but high-
strength materials usually show low ductility due to limited dislocation activity in their
microstructure. To overcome the undesirable strength-ductility trade-off relationship, some
strategies were developed within the past few decades: grain boundary engineering by severe
plastic deformation (SPD) [1-3], introducing nanotwins [4-6], lattice softening [7-9],
generating bimodal [10,11], gradient [12], heterostructured [13] and lamellar [14] materials,
engineering precipitates [15-17], and controlling transformation/twinning-induced plasticity
(TRIP/TWIP) [18-20]. Although these strategies are effective under atmospheric conditions,
achieving high strength and high ductility in the presence of hydrogen is still a significant
challenge due to the occurrence of the hydrogen embrittlement phenomenon.
In general, austenitic steels with the FCC crystal structure are well-known for the slow
diffusion of hydrogen in their lattice and outstanding resistance to hydrogen embrittlement
[21,22]. Modified austenitic steels with the addition of aluminum can show even higher
resistance to hydrogen embrittlement, due to reduced hydrogen diffusion through the surface
because of the formation of a protective Al2O3 layer [23-25]. However, a general shortcoming
of these steels is their low yield strength. Due to the demand to achieve a higher yield strength
in the presence of the hydrogen atmosphere, high entropy alloys (HEAs) with the FCC crystal
structure were recently introduced for such applications [22,26].
High-entropy materials are a new class of materials containing at least five principal
elements [27,28]. These materials, which are divided into two groups of HEAs [29] and high-
entropy ceramics [30], have an entropy of mixing of 1.5R or higher (R: gas constant). The
HEAs show good mechanical and functional properties due to the four core features resulting
from the cocktail effect (individual elements introduce their own property to the alloy), valence
electron effect, lattice strain effect (because of variations in the atomic size of elements) and
slow diffusion effect [31,32]. In general, these alloys exhibit rather good solid solution
hardening usually in the form of FCC or BCC crystal structures, and good thermal stability due
to their low Gibbs free energy [32,33]. Although high solid solution strengthening, slow
hydrogen diffusivity, high resistance to hydrogen-assisted cracking, and high performance
under different loading and chemical conditions make the FCC HEAs potential materials with
high hydrogen embrittlement resistance [22,26], these FCC HEAs still show low yield strength
levels for many practical applications.
Grain refinement and crystal defect generation are promising strategies to increase yield
strength while maintaining resistance to the hydrogen embrittlement [34,35]. For example,
decreasing the grain size to 4 µm increases strength while keeping high hydrogen
embrittlement resistance in 304 austenitic stainless steel [34]. Reducing the grain size to ~600
nm increases strength and suppresses hydrogen-assisted failure in TWIP austenitic steel [35].
The improvement in strength and hydrogen embrittlement resistance was also reported in HEA
CrMnFeCoNi alloy when grains are refined to 1.9 µm [36]. The resistance to hydrogen
embrittlement in fine-grained FCC alloys stems from (i) the reduction of hydrogen content per
unit grain boundary area (preventing intergranular decohesion), and (ii) lower local stress
concentration at grain boundaries in a fine-grained microstructure due to the reduction in the
number of dislocation pile-ups. Although grain refinement to the submicrometer and
nanometer levels can significantly enhance the yield strength due to the Hall-Petch hardening
mechanisms, limited works have been done on hydrogen embrittlement of such ultrafine (UFG)
microstructures [37] due to the expected effects of large grain boundary diffusion and
hydrogen-enhanced decohesion (HEDE) [38].
Severe plastic deformation (SPD) is a well-known strategy to enhance the yield strength
of metallic materials through grain refinement to the submicrometer and nanometer levels [3].
In addition to grain refinement, SPD processes can enhance yield strength by engineering
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different microstructural defects such as dislocations, twins, and stacking faults [3]. Such
defects can be effectively introduced even in HEAs by the application of SPD [39,40].
Although these defects can be potential sites for hydrogen trapping [41,42] or hydrogen-
assisted cracking through hydrogen-enhanced localized plasticity (HELP) [43], their presence
at an optimized level is expected to lead to high yield strength and high ductility in the presence
of hydrogen.
In this study, to achieve a good combination of high strength and high ductility in the
presence of hydrogen, three strategies are combined and some of the best strength-plasticity
combinations are reported. (i) An FCC high-entropy alloy is selected to have slow diffusion of
hydrogen through the lattice. (ii) Some amounts of aluminum are added to the alloy to form a
protecting oxide surface layer and reduce hydrogen diffusion through the surface. (iii) The
grain size and crystal lattice defects are controlled by SPD processing to delocalize the stress
concentration. The Al-contained HEA, Al0.1CrFeCoNi, which was reported to have good
mechanical properties under atmospheric conditions [44,45], was selected and its
microstructure was controlled by SPD processing through the high-pressure torsion (HPT)
method.
2. Experimental procedures
The material in this study was selected by the addition of aluminum to the medium-
entropy alloy CrFeCoNi, which was reported to show good hydrogen embrittlement resistance
[46]. In addition to the formation of an oxide layer, aluminum is expected to enhance the lattice
distortion and solution hardening due to its different atomic radius compared to chromium,
iron, cobalt and nickel [47]. However, the addition of aluminum can lead to the formation of a
BCC phase [47] which is not desirable for hydrogen embrittlement resistance [21]. To quantify
the composition of the material, valence electron concentration (VEC) calculations were used
in this study. Earlier studies suggested that the FCC and BCC phases coexist at 6.87 ≤ VEC <
8.0, while a single FCC phase becomes stable only at VEC ≥. 8.0 [48]. The VEC values of
preliminarily selected alloys in this study, Al0.1CrFeCoNi and Al0.3CrFeCoNi, were 8.12 and
7.88, respectively. Therefore, the Al0.1CrFeCoNi alloy with a single FCC phase was selected.
Another reason for the selection of Al0.1CrFeCoNi was its lower stacking fault energy
compared to Al0.3CrFeCoNi (6-21 mJ/m2 versus 51 mJ/m2 [49]), which makes another positive
feature for good hydrogen embrittlement resistance [21].
The HEA ingot with a nominal composition of Al0.1CrFeCoNi was prepared by vacuum
induction melting under an argon atmosphere. Raw elements of aluminum, chromium, iron,
cobalt, and nickel were used with purity levels above 99.5 wt%. Disc-shaped samples with 10
mm in diameter and 1 mm in thickness were extracted from the ingot using wire-cut electrical
discharge machining (EDM, Brother, HS-300). The extracted discs were homogenized at 1273
K for 1 h and quenched in iced water. The damaged layers caused by the EDM and
homogenization processes were removed from all samples by mechanical polishing using
sandpapers. The microstructural features and distribution of elements in Al0.1CrFeCoNi after
homogenization were examined by using scanning electron microscopy (SEM, JEOL JSM-
7900F) equipped with an electron backscatter diffraction (EBSD) detector and energy-
dispersive X-ray spectroscopy (EDS) under an acceleration voltage of 15 kV.
To conduct the HPT process, a pressure of P = 6 GPa was applied to the polished disc
samples between two anvils of an HPT machine (Riken Enterprise, High-Speed-H.P.T.60ton),
and the strain was introduced by rotating the lower anvil for N = 1/16, 1/8, 1/4, 1/2, 1, 5, and
10 turns with a rotation speed of 1 rpm at room temperature (T = 300 K). The principles of the
HPT method are discussed in detail in Ref. [50]. The microstructural characterization and
mechanical behavior of HPT-processed samples were examined by various techniques, as
explained below.
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First, Vickers microhardness (Mitutoyo, HM-103) was measured on the upper surface
of the discs after mechanically polished to the mirror-like condition on both sides of the discs.
This measurement was done under an applied load of 5 N for 15 s at distances of 0.15, 1, 2, 3,
and 4 mm from the disc center and in four radial directions. The average of four measurements
at each position was calculated as the microhardness of the corresponding position.
Additionally, shear strain corresponding to hardness measurements at each radial distance was
calculated as γ = 2πrN/h (r: distance from the disc center, N: turn numbers, and h: disc
thickness) [50].
Second, to examine the crystal structure for samples before and after the HPT process,
X-ray diffraction (XRD, Rigaku SmartLab) patterns were obtained. The XRD analysis was
operated using the Cu Kα irradiation at an accelerating voltage of 45 kV, and an electrode
current of 40 mA. The scanning range was 20⁰ to 100⁰ at a scanning step of 0.05⁰ and a scanning
speed of 2⁰ per minute.
Third, to examine microstructure, transmission electron microscopy (TEM, JEOL JEM-
ARM200F) was conducted under an acceleration voltage of 200 kV using bright/dark-field
images, selected-area electron diffraction (SAED) analysis, and high-resolution images with
corresponding fast Furrier transform (FFT) diffractograms. For TEM, foils were fabricated by
cutting 3 mm diameter discs from 2-5 mm away from the center of HPT-processed samples
using an EDM machine. The thickness of 3 mm discs was first reduced to 100 μm by
mechanical polishing and then to a very low thickness transparent to electrons by polishing
with an electrochemical polisher in a solution of 5 vol% perchloric acid, 25 vol% butanol, and
70 vol% methanol at 12 V and 263 K.
Fourth, to examine the surface depth profiles and understand the oxidation states of
elements, X-ray photoelectron spectroscopy (XPS, Physical Electronics, PHI 5600) was
conducted on the homogenized alloy after polishing it to mirror-like surfaces using M Kα
radiation. The surface was removed layer by layer by etching via argon ion sputtering with a
current of 1.5 μA and an acceleration voltage of 3 kV for time periods of 0-130 s before XPS.
This sputtering rate roughly corresponds to 1 nm/min surface removal (calibrated using Ga2O3).
Fifth, to examine tensile properties, flat dog-bone specimens were cut from samples
before and after the HPT process. Hydrogen charging was electrochemically conducted for 72
h at room temperature in 3% NaCl containing 3 g/L of NH4SCN aqueous solution at a cathodic
current density of 50 mA.cm−2 where the platinum wire was used as a counter electrode. Tensile
tests were conducted using a machine (Made by Arita, Kydai) at a low initial strain rate of
5.6×10−4 s−1 (constant crosshead speed of 0.05 mm.min-1) at room temperature for hydrogen-
charged and unchanged specimens. The tensile tests were carried out quickly within 2 minutes
after hydrogen charging.
Sixth, the diffusible amounts of hydrogen for a few representative samples were
measured using thermal desorption spectroscopy (TDS, ESCO EMD-WA1000S). TDS was
performed immediately after the tensile test in the temperature range of 308 K to 773 K at a
heating rate of 400 Kh-1.
Seventh, to characterize the fracture surface of tensile specimens, SEM using an
acceleration voltage of 15 kV was used.
3. Results
Fig. 1 shows the microstructure of the initial homogenized sample before HPT
processing including (a) SEM band contrast image, and corresponding (b) EBSD crystal
orientation map, (c) EBSD phase map, (d-h) EDS elemental maps, and (i) EDS spectrum and
corresponding atomic fractions of elements. Microstructural characterization of the initial
sample reveals that the average grain size of the HEA is ~25 µm and it has a single FCC phase
with a reasonably homogeneous distribution of elements (except for some slight heterogeneity
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for aluminum and cobalt), in good agreement with earlier reports [44,45]. Some dark points
are visible in Fig. 1a, which correspond to micro-pores formed during vacuum induction
melting.
Figure 1. Microstructure of homogenized high-entropy alloy Al0.1CrFeCoNi before HPT
processing. (a) SEM band contrast image, (b) EBSD crystal orientation mapping, (c) EBSD
phase mapping, (d-h) EDS elemental mappings, and (i) EDS spectrum and corresponding
atomic fractions of elements.
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Figure 2. Saturation of hardness of high-entropy alloy Al0.1CrFeCoNi to 520 Hv after HPT
processing. Hardness versus (a) number of turns and (b) shear strain before and after HPT
processing with different turns.
Fig. 2 shows the variations of hardness with (a) distance from the disc center and (b)
shear strain for samples processed with the different numbers of HPT turns. Hardness enhances
by increasing the number of turns and increasing the distance from the disc center and reaches
the steady-state level of ~520 Hv after only N = 1 turn. This level of hardness is ~3.3 times
higher than the hardness of 155 Hv for the initial homogenized sample. The saturation of
hardness to the steady-state level of 520 Hv occurs at the early stages of straining which is
rather fast. This fast saturation of hardness was also observed in other HEAs [39,40] including
the Cantor alloy [51]. It is well established that the saturation of hardness is due to a balance
between hardening phenomena such as grain refinement [51] and dislocation generation [52],
and softening phenomena such as recovery [53], recrystallization [53] and grain boundary
migration [53].
Fig. 3 shows the XRD profiles of the HEA before and after HPT processing, confirming
that all samples consist of a single FCC phase. Despite the absence of phase transformation,
peak broadening occurs after HPT processing. This peak broadening, which was also reported
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after HPT processing of various HEAs [39,40,45], is due to the lattice strain, dislocation
generation, and crystallite size refinement.
Figure 3. Stability of FCC structure in high-entropy alloy Al0.1CrFeCoNi after severe plastic
deformation. XRD profiles before and after HPT processing with different turns.
Fig. 4 shows TEM bright-field images (left side), dark-field images (center), and
corresponding SAED patterns (right side) for (a) the initial sample and for samples processed
by HPT for (b) 1/16, (c) 1/4, (d) 1/2, (e) 5 and (f) 10 turns. These microstructural
characterizations reveal that the initial sample includes coarse grains with sharp high-angle
grain boundaries in good agreement with SEM observations in Fig. 1. The application of shear
strain via the HPT process leads to the formation of dislocation structures after N = 1/16 and
1/4 turns and to significant grain refinement to the nanometer level even after N = 1/2 turns.
These microstructural evolutions lead to changes in the shape of SAED patterns from a dotted
form to a ring form. Fig. 5, in which the average size measured for more than 100 grains using
the dark-field images is plotted versus the number of HPT turns, shows the effect of HPT
processing on grain refinement more clearly. The average grain size reduces with increasing
the number of HPT turns and reaches a saturation level of 29 nm after N = 10 turns. Such a
grain size is smaller than those achieved in pure metals or conventional alloys after HPT
processing because of more significant solute-defect interactions in HEAs during straining
[50,52].
Fig. 6 shows the high-resolution TEM images for the sample processed by HPT for N
= 1/4 turns, i.e. at the early stages of straining. There are many dislocations in the sample, as
indicated by τ in Figs. 6a and 6d. Examination of the microstructure in higher magnifications
confirms that while some isolated dislocations are present in the microstructure such as the one
in Fig. 6b, some dislocations pairs as the ones shown in Fig. 6e stay at close distances to each
other due to the formation of Lomer-Cottrell locks. In addition to dislocations, a number of
stacking faults are also observed in the sample, as shown in Fig. 6c. The formation of
dislocations and stacking faults is usually the dominant mechanism at the early stages of
straining by HPT [53].
Fig. 7 shows the high-resolution TEM images after processing a sample with 10 HPT
turns, where the microstructure saturates to the steady state. A change of deformation
mechanism from sliding of dislocations to the creation of nanotwins occurs at the steady state,
as shown in Figs. 7a and Fig. 7b. In addition to coherent twin boundaries, some semi-coherent
grain boundaries are also observed at the steady state, as shown in Figs. 7c and 7d. The
formation of nanotwins can be desirable for the purpose of this study because the introduction
of twins was suggested as a strategy to enhance the combination of strength and ductility under
atmospheric conditions [4-6] and in the presence of hydrogen [26].
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Figure 4. Grain refinement of high-entropy alloy Al0.1CrFeCoNi by HPT processing. TEM
bright-field images (left side), dark-field images (center) and corresponding SAED patterns
(right side) for (a) initial homogenized condition, and for HPT-processed samples after (b)
1/16, (c) 1/4, (d) 1/2, (e) 5 and (f) 10 turns. Dark-field images were taken with diffracted beams
indicated by arrows in SAED patterns.
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Figure 5. Saturation of grain size of high-entropy alloy Al0.1CrFeCoNi to steady-state level of
29 nm after HPT processing. Variations of grain size versus number of turns for samples before
and after HPT processing.
Fig. 8 shows the XPS spectra of (a) Al 2s, (b) Cr 2p3/2, 1/2, (c) Fe 2p3/2, 1/2, (d) Co 2p3/2,
1/2, (e) Ni 2p3/2, 1/2 and (f) O 1s after argon sputtering for various periods of time. Except for
nickel, other elements have an oxidized state on the surface. The oxidized states of elements
fully transform to the metallic states after sputtering for 90 s, suggesting that the thickness of
the oxide layer is about 1.5 nm. While aluminum and chromium keep their oxidized state all
through the 1.5 nm thickness, iron and cobalt are oxidized almost up to the middle of the oxide
layer. This Al-containing oxide layer is expected to suppress hydrogen diffusion through the
surface to the bulk [23,24].
Fig. 9a shows the tensile stress versus strain curves of uncharged and hydrogen-charged
specimens for the initial homogenized sample and for HPT-processed samples after N = 1/16,
1/8, 1/4, 1/2, 1, 5, and 10 turns. Examination of these curves indicates several important points.
First, the HPT process drastically increases the yield strength to reach a maximum level of 1.96
GPa. This tendency, which was also observed during the hardness measurement, should be
mainly due to the grain size refinement to 29 nm. Second, in all samples, the yield strength of
the sample in the presence of hydrogen is the same or even higher than the yield stress in the
absence of hydrogen. Third, the ductility decreases in the presence of hydrogen, but all samples
still exhibit good ductility. These good ductility levels contrast with the nanograined Cantor
HEA which breaks in the elastic region without showing any ductility after hydrogen charging
[51], but agree with the reported hydrogen embrittlement resistance of fine-grained austenitic
steels [34,35,37].
Fig. 9b shows the hydrogen desorption spectra measured by TDS for the initial
specimen and for samples treated by HPT for N = 1/2 and 10 turns. The hydrogen contents of
the initial specimen treated by HPT for N = 1/2 and 10 turns are 5.4, 6.7, and 7.9 wt.ppm,
respectively. The hydrogen content increases by only 30% with increasing the applied shear
strain by HPT, despite the significant grain refinement and increasing the volume fraction of
grain boundaries after HPT processing. Here, it should be noted that the hydrogen content for
the Cantor alloy using the same procedure increases from 3 wt.ppm for the initial specimen to
27 wt.ppm for the sample processed by HPT with N = 10 turns [51]. Although the grain size
for the Cantor alloy after HPT processing (25 nm) was close to the current alloy (29 nm), the
differences in their hydrogen uptake are high (27 wt.pp compared to 7.9 wt.ppm), If it is
roughly assumed that the hydrogen diffusion through grain boundaries and hydrogen trapping
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in defects are similar for the current alloy and for the Cantor alloy (due to similarities in the
compositions and crystal structures), the low hydrogen content in Al0.1CrFeCoNi can be
attributed to sluggish diffusion through the surface due to the presence of an Al-containing
oxide layer [23-25].
Figure 6. Formation of dislocations, stacking faults and Lomer-Cottrell locks in high-entropy
alloy Al0.1CrFeCoNi processed with HPT for 1/4 turns. (a, d) High-resolution TEM images and
corresponding FFT diffractograms, (b, c) lattice image of area indicated in (a), and (e) lattice
image of area indicated in (d) and reconstructed by inverse FFT.
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Figure 7. Formation of twins and high-angle grain boundaries in high-entropy alloy
Al0.1CrFeCoNi processed with HPT for 10 turns. (a, b) High-resolution TEM images and
corresponding FFT diffractograms, (b) lattice image of area indicated in (a) and reconstructed
by inverse FFT, and (d) lattice image of area indicated in (c).
Fig. 10 shows the fracture surfaces of the uncharged (a-c) initial homogenized sample
and samples after HPT processing for (d-f) 1/2 and (g-i) 10 turns. The fracture surfaces of the
three samples are covered by dimples, indicating the ductile fracture of the uncharged
specimens. Significant necking can be observed for the initial specimen with clear shear lips
which indicates the high plastic strain for this specimen, while plasticity and necking become
less significant after HPT processing. The reduction of plasticity after SPD is a natural
consequence of grain refinement and limited dislocation activity [1-3].
Fig. 11 shows the fracture surfaces of the hydrogen-charged (a-c) initial homogenized
sample and samples after HPT processing for (d-f) 1/2 and (g-i) 10 turns. A brittle fracture is
observed in the vicinity of the surface of three specimens, where the concentration of hydrogen
is high, and this leads to the initiation of hydrogen-assisted cracking. Intergranular cracking
with clear slip traces can be observed on the fracture surface of the initial homogenized
specimen, indicating the possible occurrence of the HELP mechanism [21]. In contrast, slip
traces are eliminated after HPT processing, indicating that the HELP mechanisms may not be
the main hydrogen-assisted failure mechanism in the nanograined HEA [38].
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Figure 8. Formation of Al-containing oxide layer on surface of high-entropy alloy
Al0.1CrFeCoNi . XPS spectra of (a) Al 2s, (b) Cr 2p3/2, 1/2, (c) Fe 2p3/2, 1/2, (d) Co 2p3/2, 1/2, (e)
Ni 2p3/2, 1/2 and (f) O 1s after argon sputtering for etching times of 0-130 s.
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Figure 9. High yield strength of high-entropy alloy Al0.1CrFeCoNi processed by HPT with and
without hydrogen charging. (a) Tensile stress-strain curves in presence and absence of
hydrogen for samples before and after HPT processing. (b) TDS profiles and hydrogen content
for initial sample and for samples processed by HPT for 1/2 and 10 turns.
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Figure 10. Indication of plastic fracture in high-entropy alloy Al0.1CrFeCoNi in absence of
hydrogen. Fracture surfaces examined in different magnifications for (a-c) initial homogenized
sample, and for samples processed by HPT for (d-f) 1/2 and (g-i) 10 turns. Images at center
and right are magnified views of areas indicated in images on left.
Figure 11. Indication of plastic fracture in high-entropy alloy Al0.1CrFeCoNi in presence of
hydrogen. Fracture surfaces examined in different magnifications for (a-c) initial homogenized
sample, and for samples processed by HPT for (d-f) 1/2 and (g-i) 10 turns. Images at center
and right are magnified views of areas indicated in images on left.
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4. Discussion
The results of this study illustrate that simultaneous engineering of crystal structure,
composition, and microstructural defects is a successful strategy to achieve an outstanding
combination of high yield strength and low susceptibility to hydrogen embrittlement in HEAs.
Although high-entropy materials are known for their high potential for different applications
[26-33], two questions naturally arise from this study. First, what are the mechanisms
underlying such high yield strength and high plasticity in the presence of hydrogen? Second,
how good is the mechanical property of the developed materials compared to the best
hydrogen-resistant alloys?
Regarding the first question, it is noted that hydrogen can diffuse and trap at lattice
defects such as dislocations [54,55] and grain boundaries [56], and thus the diffusion of
hydrogen and its interaction with defects are determining factors for the strength and ductility
of materials in the presence of hydrogen. The FCC structure of current HEA leads to slow
lattice diffusion [21,31], while the presence of aluminum in the alloy can provide a protective
oxide layer on the surface, as shown in Fig. 8 using XPS depth profiles, to reduce the diffusion
of hydrogen from the surface to the bulk [23-25]. Despite gradual changes in the microstructure
from coarse-grained to nanograined, all alloys show good resistance to hydrogen
embrittlement. The coarse-grained homogenized sample shows high ductility and less
susceptibility to hydrogen embrittlement, but its yield strength is low for practical applications.
The failure of this material should be due to dislocation pile-ups on the slip plane and their
interaction with grain boundaries [25,36]. The interaction of slip and grain boundaries
introduces stress concentration resulting in hydrogen segregation and hydrogen-assisted
cracking along the grain boundaries through the HELP mechanism [21,38]. After processing
by HPT at the early stages of straining, dislocations including low-mobility Lomer-Cottrell
locks, formed by the interaction of two full dislocations, appear. These immobile dislocation
locks can act as dislocation barriers leading to stress delocalization. They contribute to
mitigating dislocation pile-ups at grain boundaries resulting in decreasing the stress
concentration and lowering stress concentration-assisted hydrogen segregation and cracking.
After increasing the applied shear strain via the HPT process to a steady state, nanograins
containing nanotwins and some semi-coherent boundaries are formed, while the density of
dislocations is reduced. In these nanograined samples, the pile-up of dislocations is not easy
because grain boundaries can act as dislocation sinks [1] and this is the main reason for the
reduced ductility of these nanograined materials. The existence of numerous coherent nanotwin
boundaries and some semi-coherent boundaries is effective in reducing the movement of
dislocations [4-6] and suppressing the HELP mechanism through-dislocation-twin interactions
[26]. Here, it should be noted that the effect of defects on the mobility of dislocations is not the
only factor that determines their contribution to suppressing hydrogen embrittlement. The
effect of defects on hydrogen trapping (i.e. hydrogen trapping energy) [57,58], hydrogen
diffusion [59,60] and crack initiation [21,31] are other factors that should be taken to account
to analyze their effect on hydrogen embrittlement.
Regarding the second question, it should be noted that different strategies were
successfully applied to achieve high strength and high ductility in the absence of hydrogen
including grain boundary engineering [1-3], introducing nanotwins [4-6], lattice softening [7-
9], generating bimodal [10,11], gradient [12], heterostructured [13] and lamellar [14] materials,
engineering precipitates [15-17], and controlling TWIP and TRIP effects [18-20]. However, it
is challenging to achieve high strength and high ductility in the presence of hydrogen due to
the interaction of hydrogen and lattice defects as well as because of the hydrogen effects on
decohesion (DEDE mechanism). As attempted in this study, selecting appropriate materials
with slow hydrogen diffusion through the surface and bulk is important in achieving good
mechanical properties under hydrogen. Table A1 of the Appendix gives yield stress, elongation
16
to failure, and hydrogen content for some of the reported materials with high hydrogen
embrittlement resistance [35,36,51,61-77], and Fig. 12 compares the combination of yield
strength and elongation of these materials with those achieved in the current study. There is a
clear trade-off relation between yield stress and ductility, as shown in many studies conducted
by tensile tests under atmospheric conditions [1-20]. However, the main difference is that the
HEA samples processed by HPT in this study clearly break the trade-off relation and show
some of the best combinations of strength and ductility in the presence of hydrogen. Taken
altogether, Fig. 12 confirms that the strategies employed in this study are quite effective to
develop high-strength alloys with high resistance to hydrogen embrittlement for hydrogen-
related applications.
Figure 12. Development of high strength and high ductility in presence of hydrogen in high-
entropy alloy Al0.1CrFeCoNi processed by HPT. Yield strength versus elongation to failure
achieved in this study in presence of hydrogen compared to various alloys reported in literature
and given in Table A1. High-entropy alloys are Cantor alloys from Cr-Mn-Fe-Co-Ni system
and medium-entropy alloys are from either Cr-Fe-Co-Ni or Cr-Mn-Fe-Co systems.
5. Conclusions
In order to achieve a good combination of strength and ductility in the presence of
hydrogen, severe plastic deformation using the high-pressure torsion (HPT) technique was
applied to an Al-containing high entropy alloy, Al0.1CrFeCoNi, which was expected to have
slow hydrogen diffusion through the surface and lattice. The following conclusions were
achieved.
1. When low strain is applied by HPT, the mechanism of plastic deformation is based on
dislocation activities and Lomer-Cottrell lock formation; however, when strain increases
by rising the number of HPT turns, the grain size reduces to 29 nm and nanotwins and some
semi-coherent boundaries appear.
2. The yield strength increases and plasticity decreases with increasing the applied strain by
HPT, while these changes become more significant in the presence of hydrogen.
3. Due to the effect of Lomer-Cottrell locks, nanotwin boundaries, and semi-coherent
boundaries by HPT processing, the movement of dislocations is slowed down in the alloy.
Accordingly, the pile-up of dislocations at grain boundaries and stress concentration are
suppressed, leading to good hydrogen embrittlement resistance of the alloy.
4. The current alloy shows some of the best combinations of high yield strength and good
plastic elongation (e.g., 1.96 GPa and 10%) reported so far in metallic materials in the
17
presence of hydrogen, suggesting that selecting a HEA with slow surface and lattice
diffusion and engineering its lattice defects by severe plastic deformation is a potential
approach to deal with the hydrogen embrittlement issue.
Appendix
Table A1 summarizes yield strength, elongation to failure and hydrogen content
reported for some major alloys in the literature.
Table A1. Yield strength, elongation to failure and hydrogen content for high-entropy alloy
Al0.1CrFeCoNi processed by HPT in this study compared with different alloys reported in
literature.
Alloy
0.2% Offset Yield
Stress (MPa)
Elongation to
Failure (%)
Hydrogen
(wt.ppm)
Ref.
304 Steel
245
28
63
[61]
304 Steel
200
50
[62]
304 Steel
740
21
[62]
304 Steel
1090
6
[62]
316L Steel
230
47
36.5
[61]
Fe-31Mn-3Al-3Si (wt%)
410
63
[35]
Fe-31Mn-3Al-3Si (wt%)
680
54
[35]
Fe-31Mn-3Al-3Si (wt%)
220
96
[35]
Fe-15Mn-2.5Al-2.5Si-0.7C (wt%)
510
15
[63]
Fe-18Mn-3Si-0.6C (wt%)
520
51
[63]
Fe-18Mn-1.5Si-0.6C (wt%)
400
57
[63]
Fe-23Mn-0.5C (wt%)
1000
5.5
[43]
Fe-22Mn-0.6C (wt%)
540
52
[65]
Fe-18Mn-0.6C (wt%)
730
35
[64]
Fe-18Mn-0.6C (wt%)
350
81
[66]
Fe-18Mn-0.6C (wt%)
420
30
[67]
Fe-17Mn-0.8C (wt%)
400
70
[68]
Fe-12.5Mn-1.1C (wt%)
400
38
[69]
Inconel 718
500
20
[70]
CrFeCoNi
600
50
129
[71]
Cr10Mn30Fe50Co10
190
25
16.3
[72]
Cr10 Mn30Fe49.5Co10C0.5
320
31
19.1
[73]
Cr10Mn30 Fe48.5Co10Ni1.0C0.5
365
22
16.6
[72]
Cr22 Mn13 Fe21Co22 Ni22
580
20
75.4
[71]
Cr22Mn13Fe21Co22Ni22
260
23
67.1
[71]
CrMnFeCoNi
300
21
113
[73]
CrMnFeCoNi
350
24
[73]
CrMnFeCoNi
220
10
178
[73]
CrMnFeCoNi
220
14
[73]
CrMnFeCoNi
470
44
4.1
[74]
CrMnFeCoNi
470
21
13.5
[74]
CrMnFeCoNi
300
51
54.3
[71]
CrMnFeCoNi
270
42
76.5
[61]
CrMnFeCoNi
870
4
54.5
[75]
CrMnFeCoNi
720
5
54.7
[75]
CrMnFeCoNi
550
21
52.7
[75]
CrMnFeCoNi
400
30
52.7
[75]
CrMnFeCoNi
220
42
63.2
[75]
CrMnFeCoNi
300
20
113
[36]
CrMnFeCoNi
620
16
129
[36]
CrMnFeCoNi
920
11
125
[36]
CrMnFeCoNi
500
33
5.3
[76]
CrMnFeCoNi
700
15
6.1
[76]
CrMnFeCoNi
220
73
3
[47]
CrMnFeCoNi
1090
9
[47]
CrMnFeCoNi
1230
9
10.1
[47]
CrMnFeCoNi
1270
3
[47]
Al0.1CrFeCoNi
250
87
5.4
This study
Al0.1CrFeCoNi
1300
25
This study
Al0.1CrFeCoNi
1360
15
This study
Al0.1CrFeCoNi
1520
12.2
This study
Al0.1CrFeCoNi
1770
11.1
6.7
This study
Al0.1CrFeCoNi
1960
10
This study
Al0.1CrFeCoNi
1950
3.9
This study
Al0.1CrFeCoNi
1900
2.5
7.9
This study
18
Acknowledgments
This work was supported in part by the MEXT, Japan through Grants-in-Aid for
Scientific Research on Innovative Areas (JP19H05176 & JP21H00150) and Challenging
Research Exploratory (JP22K18737).
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