The Mechanism of the High Resistance to Hydrogen-Induced Strength Loss in Ultra-High Strength High-Entropy Alloy

Abstract

The resistance of the Al_0.5Cr_0.9FeNi_2.5V_0.2 high-entropy alloy (HEA) to hydrogen embrittlement (HE) was investigated by a slow strain rate test (SSRT), and the fracture surface was examined through a scanning electron microscope. Compared with other high-strength steels, Al_0.5Cr_0.9FeNi_2.5V_0.2 showed insignificant strength loss after hydrogen charging. The fracture surface of the hydrogen-charged specimens mainly consisted of dimples, and no intergranular morphology was observed. The coupling effect of the dispersed nano-structured precipitates and high-density dislocations in Al_0.5Cr_0.9FeNi_2.5V_0.2 improves the resistance to hydrogen-induced strength loss.

1. Introduction

Due to the presence of hydrogen, metals often experience loss of ductility and strength. This phenomenon, caused by hydrogen in metals, is known well as hydrogen embrittlement (HE) and has been widely reported and investigated [1,2,3]. Much research has been conducted to explain the phenomenon of HE, and there are two mechanisms which are the most accepted; these are hydrogen-enhanced localized plasticity (HELP) and hydrogen-enhanced decohesion (HEDE) [4,5,6,7]. The HELP theory proposes that hydrogen atoms in metals enhance the mobility of dislocations and cause localized plastic deformation. At the same time, the HEDE mechanism postulates that hydrogen atoms embrittle metals by reducing the cohesive forces between matrix atoms. The synergistic interplay of the HELP and HEDE mechanisms is also investigated in numerous papers, through computation simulations or modeling [8,9], as well as experimentally [10,11], for different metals. This synergistic effect of the HELP and HEDE mechanisms depends on the adopted experimental conditions, such as the stress state and hydrogen concentration [12]. It was also found that steels with a higher strength often suffer more severe HE [13,14], which can cause the premature failure of the steels in service. HE is gradually becoming an unneglectable problem with the further development of high-strength steels. In addition, hydrogen diffusion and accumulation in high-strength steels generally occur at grain boundaries under stress. Hydrogen at the grain boundary would reduce its strength and cause the intergranular cracking of the steel [15,16], and steels experience considerable strength loss in a such a case. Defects such as precipitates, dislocations, and vacancies are preferential hydrogen trapping sites inside grains, making hydrogen diffusion towards the grain boundary difficult and lowering the hydrogen concentration at the grain boundaries [17,18,19,20,21]. Therefore, clarifying the distribution of hydrogen is crucial to the elucidation of the hydrogen embrittlement performance of steels.

Precipitates and dislocations in steels can improve the strength of steels and can also act as hydrogen traps [22,23,24,25], which limit the diffusion and enrichment of hydrogen at the grain boundaries. However, high-strength steels with different microstructures usually possess different resistances to hydrogen embrittlement. Li et al. [26] introduced large numbers of dislocations in screw-thread steel by pre-straining, and the strained steel showed a high susceptibility to HE. Strained screw-thread steel showed significant strength loss, and there was an obvious intergranular fracture feature in the fracture surface of the hydrogen-charged specimens. The distribution of precipitates in the martensite matrix for high-strength maraging steels also gave a similar performance with regard to HE. Wang et al. [27] studied the HE susceptibility of TM210 maraging steels with different precipitate densities. They found that all the specimens exhibited significant strength loss and intergranular fracture morphology on the fracture surface of the hydrogen-charged specimens, whatever the density of the precipitate was. Other studies about the HE performance of maraging steel steels also found similar results [28,29,30,31]. The literature above suggested that neither large amounts of dislocations nor precipitate structures in steels could improve its HE resistance.

High-entropy alloys (HEAs) are formed by five or more equal, or approximately equal, metals. Recently, HEAs have been widely investigated due to their excellent mechanical properties, corrosion resistance, and high temperature resistance [32,33,34,35]. Liang et al. [36] developed precipitation-strengthened HEAs (Al_0.5Cr_0.9FeNi_2.5V_0.2), which were composed of a near-equiatomic disordered face-centered cubic (FCC) matrix and high-content ordered L1₂ nano-precipitates with dense dislocations, to produce an ultra-high strength while maintaining good ductility and work-hardening capacity. Comparing the mechanical properties of this HEA to other high-strength metallic materials, HEAs exhibit a good combination of strength and ductility. Currently, the mechanical performance of this kind of HEA in the presence of hydrogen remains to be clarified.

In this study, the hydrogen embrittlement behaviour of Al_0.5Cr_0.9FeNi_2.5V_0.2 high-strength HEA [36] was investigated and was compared with other similar high-strength steels. The slow strain rate tensile test (SSRT) was used to study the hydrogen embrittlement susceptibility of Al_0.5Cr_0.9FeNi_2.5V_0.2, and the fracture morphology of uncharged and hydrogen-charged specimens was also observed in order to analyse the crack nucleation position. This study reveals that the coupling effect of the precipitates and dislocations improves the resistance to hydrogen-induced strength loss. This study also aimed to provide a conception that improves the hydrogen embrittlement resistance of high-strength steel.

2. Materials and Methods

An HEA Al_0.5Cr_0.9FeNi_2.5V_0.2 as-cast alloy was processed by arc-melting pure metals (Al, Cr, Fe, Ni, and V, with a purity greater than 9.9%) in a high-purity argon atmosphere. The as-cast alloy was subsequently solution treated at 1200 °C for 24 h and then water quenched. The quenched alloy was cold-rolled at room temperature with a total thickness reduction of 72%. Finally, the rolled alloy was aged at 600 °C for 1 h and cooled in air. Homogenized and coarse equiaxed FCC grains formed after the solution treatment. After the cold-rolling and ageing, a large number of nano-sized precipitates and high-density dislocations were dispersed inside the fine recrystallization grains. An annealed 22MnB5 product sheet was heated to 930 °C and then quickly moved to the mold. This high-temperature billet was stamped and formed in the mold, and quenching was completed simultaneously. The starting temperature of the hot forming was 830 °C, and the cooling rate was greater than 27 °C/s. A fully lath martensite structure was obtained, and the prior austenite grain size was about 19 μm.

To assess HE susceptibility of Al_0.5Cr_0.9FeNi_2.5V_0.2, SSRTs were conducted. The dog-bone-shaped tensile specimen with a gauge section of 10 × 3.2 × 0.8 mm³ was used. The tensile specimen was polished up with 2000-grit SiC sandpaper, washed with deionized water, and then pre-charged with hydrogen. High-strength martensite steel 22MnB5 (wt%: 0.37 C, 0.23 Si, 1.12 Mn, 0.012 S, 0.008 P, 0.064 Al, 0.029 Ti, 0.11 Cr, 0.003 B), which shows a similar strength to Al_0.5Cr_0.9FeNi_2.5V_0.2, was also charged by hydrogen and tensile tested as a reference. Hydrogen was introduced into the tensile specimens through a pre-charging method, using a cathodic electrochemical charging technique in a 0.05 mol/L H₂SO₄ +1 g/L CH₄N₂S (thiourea) solution with a constant current density of 4 mA/cm², and the charging time was 18 h. Thiourea was added as a poisoning agent to the solution to improve the efficiency of the hydrogen charging. SSRT was performed immediately after the end of the hydrogen charging. Tensile tests were performed with strain rates of 1 × 10⁻⁴ s⁻¹ and 1 × 10⁻⁶ s⁻¹ for the uncharged and the hydrogen-charged specimens, respectively. A hydrogen analyser G4 PHOENIX (Bruker Corporation, Billerica, MA, USA) was used to measure the total hydrogen content of the specimen immediately after the tensile test. The fracture morphologies of the uncharged and hydrogen-charged samples were observed by a sigma 300 scanning electron microscope (SEM, Carl Zeiss AG, Oberkochen, Germany). Each experiment was repeated three times to ensure the reliability of the experimental data. All the above experiments were carried out at room temperature.

3. Results

Figure 1 shows the engineering stress-strain curves of Al_0.5Cr_0.9FeNi_2.5V_0.2 obtained by SSRT under hydrogen-uncharged and hydrogen-charged conditions.

Table 1. Mechanical properties of Al_0.5Cr_0.9FeNi_2.5V_0.2 specimens with and without hydrogen charging.

Sample	Yield Strength (σs)/MPa	Tensile Strength (σb)/MPa	Total Elongation (δ)/%	Hydrogen Content (CH)/wppm
Hydrogen-Uncharged	1630.8	2073.6	6.0	0.2
Hydrogen-Charged	1554.7	1691.7	0.4	15.3

summarizes the mechanical performances of the hydrogen-uncharged and hydrogen-charged specimens according to Figure 1. The engineering stress-strain curve of the hydrogen-uncharged specimen shows that the tensile strength of Al_0.5Cr_0.9FeNi_2.5V_0.2 exceeds 2000 MPa, while the total elongation is maintained around 6% in the hydrogen-uncharged condition, showing an excellent combination of mechanical properties. The high-content ductile Ni₃Al-type ordered nano-precipitates dispersed in the substrate contribute to its high strength. However, the total elongation significantly decreases after hydrogen charging while the tensile strength merely decreases slightly. Because of the dramatic decrease in total elongation, the relative tensile strength loss is chosen to calculate the HE susceptibility, i.e.,

$$I \left({\sigma }_b\right)=\frac{{\sigma }_{b, air-}{\sigma }_{b, H}}{{\sigma }_{b, air}}\times 100\%,$$

(1)

where σ_b,air and σ_b,H are the tensile strengths for the uncharged and hydrogen-charged specimens, respectively. Based on this equation, Al_0.5Cr_0.9FeNi_2.5V_0.2 HEA has a tensile strength loss of 18.4%, while there is no distinct change in yield strength. After the tensile test, the hydrogen contents of the fractured HEA specimen are measured to be 15.3 wppm and 0.2 wppm for the hydrogen-charged and the uncharged specimens, respectively.

Figure 2 shows the engineering stress-strain curves of 22MnB5 steel obtained by SSRT under hydrogen-uncharged and hydrogen-charged conditions; the relative mechanical properties of the two kinds of specimens are listed in

Table 2. Mechanical properties of 22MnB5 steel specimens with and without hydrogen charging.

Sample	Yield Strength (σs)/MPa	Tensile Strength (σb)/MPa	Total Elongation (δ)/%	Hydrogen Content (CH)/wppm
Hydrogen-Uncharged	1179.4	1582.5	9.0	0.1
Hydrogen-Charged	-	732.8	0.2	1.8

. As with Al_0.5Cr_0.9FeNi_2.5V_0.2, 22MnB5 steel also possesses high strength (~1600 MPa). After hydrogen charging, 22MnB5 steel also experiences a dramatic total elongation loss, which is considerable compared with Al_0.5Cr_0.9FeNi_2.5V_0.2. Unlike Al_0.5Cr_0.9FeNi_2.5V_0.2, it can be seen from Figure 2 that hydrogen-charged 22MnB5 steel also displays a significant strength loss, which is up to 53.7%. The total hydrogen content for uncharged 22MnB5 steel is 0.1 wppm. Although the total hydrogen content of Al_0.5Cr_0.9FeNi_2.5V_0.2 after the tensile test is 15.3 wppm, while for 22MnB5 steel it is only 1.8 wppm, the strength loss of Al_0.5Cr_0.9FeNi_2.5V_0.2 is even smaller. Consequently, Al_0.5Cr_0.9FeNi_2.5V_0.2 showed better resistance to hydrogen-induced strength loss than 22MnB5 steel, which has a similar strength.

The strength losses of high-strength steels (1500~2000 MPa) after hydrogen charging that have been investigated by other researchers are listed in

Table 3. Summary of strength loss, hydrogen charging, SSRT conditions, and fracture modes of other high-strength steels after hydrogen charging from the literature. Note that all the specimens were electrochemically hydrogen charged, and all the SSRTs were conducted at room temperature.

Steels	Hydrogen Charging and SSRT Conditions						Tensile Stress (σb)/MPa	Tensile Strength Loss/%	Fracture Mode	Ref.
	Hydrogen-Charging Method	Charging Solution	Current Density (mA/cm2) or Cathodic Potential (V)	Time (h)	Strain Rate (s−1)	Hydrogen Content (wppm)
Screw-thread steel	Pre-charging	0.5 mol/L H2SO4 + 1 g/L CH4N2S	0.3 mA/cm2	24	2 × 10−5	23.0	~1500	~65.0	Intergranular and quasi-cleavage	[21]
18Ni (300) maraging steel	Pre-charging and dynamic hydrogen charging	0.6 M NaCl	−1.2 VSCE (0.32 mA/cm2)	24	1 × 10−6	-	1950	68.6	Intergranular and quasi-cleavage	[23]
TM210 maraging steel	Dynamic hydrogen charging	0.2 mol/L NaOH + 0.22 g/L CH4N2S	0.5 mA/cm2	-	1 × 10−7	1.16	1970	63.6	Intergranular	[22]
18Ni (300) maraging steel	Pre-charging and dynamic hydrogen charging	0.6 M NaCl	−1.2 VSCE (0.32 mA/cm2)	24	1 × 10−6	-	2180	61.5	Intergranular and ductile transgranular	[24]
Maraging steel (Marval 18)	Pre-charging	3% NaCl + 0.3% NH4SCN	0.2 mA/cm2	160	1 × 10−7 ~ 1 × 10−4	4.7	2030	~52.5	Intergranular	[25]
18Ni (300) maraging steel	Pre-charging	0.6 M NaCl	−0.42 VSCE	24	1 × 10−5	-	1873	50.4	Intergranular and quasi-cleavage	[26]

. After hydrogen charging, the tensile strength loss of Al_0.5Cr_0.9FeNi_2.5V_0.2 is merely 18%, which is quite small compared to other high-strength steels. Moreover, intergranular fracture failure is a dominant morphology on the fracture surface of the conventional high-strength steels after hydrogen charging. A comparison of Al_0.5Cr_0.9FeNi_2.5V_0.2 with the literature data also demonstrates that Al_0.5Cr_0.9FeNi_2.5V_0.2 exposed to H-rich environments experiences less strength loss despite its higher hydrogen content, thereby showing better HE resistance. The intergranular morphology shown on the fracture surfaces of conventional high-strength steels indicates that the hydrogen uptake during hydrogen charging is distributed at the grain boundary. It is well known that the accumulation of hydrogen at the grain boundary would degrade the grain boundary strength [15,16], resulting in intergranular fracture. Intergranular fracture and premature failure of steels are often accompanied by dramatic strength loss. Therefore, the good resistance to hydrogen-induced strength loss of Al_0.5Cr_0.9FeNi_2.5V_0.2 may lie in the distribution of the hydrogen. Possessing good mechanical properties and better resistance to hydrogen-induced strength loss, Al_0.5Cr_0.9FeNi_2.5V_0.2 will have good application prospects.

The fracture morphology of the uncharged and the hydrogen-charged Al_0.5Cr_0.9FeNi_2.5V_0.2 is shown in Figure 3. The fracture surface of the uncharged specimen shows complete and evident dimple morphology, regardless of the center region (Figure 3b) or the edge region (Figure 3c), revealing a microvoid coalescence (MVC)-type fracture. The distribution of dimples on the fracture surface also indicates a ductile fracture mode for the uncharged specimen. Meanwhile, the edge area around the uncharged fracture surface shows a shear lip (surrounded by dotted lines), as shown in Figure 3a. For the hydrogen-charged specimen, a semi-shear lip morphology (surrounded by dotted lines) is also observed (Figure 3d). The region close to the crack initiation site also predominantly shows a feature of shallow dimples (Figure 3e). As the crack propagates, the morphology of the fracture surface that is along the crack does not change. The region away from the crack initiation site still shows shallow dimples, as shown in Figure 3f. In a similar way to the uncharged specimens, the hydrogen-charged specimens show a ductile fracture feature, indicating a ductile fracture mode. The fracture mode of Al_0.5Cr_0.9FeNi_2.5V_0.2 does not change, with or without hydrogen charging. The hydrogen-charged HEA shows shallower dimples than the uncharged specimens, due to the enhanced mobility of dislocation (the HELP mechanism). Dislocation glide is accelerated in the presence of hydrogen, and the initiation and coalescence of microvoids are promoted, leading to the formation of shallower dimples. In addition, the cohesive force is also reduced by hydrogen (the HEDE mechanism), resulting in premature fracture of the specimens under lower stress. Due to the characteristics of the dimples on the fracture surface of the hydrogen-charged specimens, it can be inferred that the HELP mechanism plays a major role during the tensile process.

The fracture morphology of the uncharged and hydrogen-charged 22MnB5 steel is shown in Figure 4. The specimen without hydrogen charging shows obvious necking at lower magnification (Figure 4a). The center region (Figure 4b) and the edge region (Figure 4c) on the fracture surface of the uncharged 22MnB5 steel both show typical dimples, regardless of their locations, revealing a ductile fracture mode. After hydrogen charging, 22MnB5 steel displays obvious brittle fracture characteristics. Both the center region (Figure 4e) and the edge region (Figure 4f) on the fracture surface of the hydrogen-charged 22MnB5 steel exhibit intergranular fracture morphology without any dimples. This indicates that the hydrogen in 22MnB5 steel is more likely to be enriched at the prior austenite grain boundaries rather than inside the grain. According to the HEDE mechanism, the enrichment of hydrogen at the grain boundary reduces the strength of the prior austenite grain boundary and then causes the intergranular fracture. Moreover, secondary cracks along the grain boundaries are observed on the fracture surface. From these results, the fracture mode has changed from the ductile dimple fracture to the entire brittle intergranular fracture after hydrogen charging.

4. Discussion

High-strength alloys have been widely used due to their excellent mechanical property combinations. However, alloys with higher strength are more likely to suffer severe HE. Thus, it is of great importance for high-strength alloys to improve their resistance to HE. Additionally, maintaining high strength in the presence of hydrogen is required for high-strength alloys in their real service environments. Hydrogen atoms are prone to accumulation at grain boundaries and cause severe strength loss to alloys. According to the present SSRT experiments of two kinds of materials and the relative literature, the Al_0.5Cr_0.9FeNi_2.5V_0.2 studied in this experiment showed less strength loss than the other conventional high-strength steels when hydrogen was introduced into the materials under the experimental conditions in this work. The fracture surface of the hydrogen-charged specimen shows the same ductile fracture mode as the uncharged specimen and thereby shows better resistance to HE. The Al_0.5Cr_0.9FeNi_2.5V_0.2 investigated in this study consists of a near-equiatomic disordered face-centered cubic (FCC) substrate with high-content ductile Ni₃Al-type ordered nano-precipitates and achieves ultra-high strength while maintaining good ductility. It is well elucidated that a large number of nano-sized precipitates and high-density dislocations are dispersed inside the grains [36]. The interaction between the dislocations and precipitates contributes to the excellent resistance to hydrogen-induced strength loss. It is well known that the FCC structure possesses better solubility but lower hydrogen diffusivity than that of the body-centered cubic (BCC) structure [37,38]. Due to its FCC substrate, Al_0.5Cr_0.9FeNi_2.5V_0.2 can dissolve a considerable amount of hydrogen while the hydrogen cannot easily become enriched at the grain boundary through bulk diffusion, and the hydrogen can only be effectively transported with mobile dislocations. On the one hand, the hydrogen is transported with mobile dislocations under stress. When the mobile dislocations carrying hydrogen atoms meet dispersed nano-precipitates, the dislocation will most likely lose the hydrogen atoms to the precipitates due to their high hydrogen-binding energy [3,25]. Once trapped by the precipitates, it is hard for the hydrogen to escape, and the local hydrogen concentration at the grain boundaries becomes relatively low, making the occurrence of intergranular fracture less possible. On the other hand, the presence of precipitates could hinder the movement of the dislocations [39], reducing the mean free path of the dislocations. In addition, a large number of dislocations are generated inside the substrate during the aging process [36], which further reduces the mean free path of the dislocations [40]. In this case, the enrichment of hydrogen is caused by the dislocations, and the dislocations can also act as hydrogen traps, which can also trap the hydrogen inside the grain and further reduce the concentration of hydrogen at the grain boundary [41,42]. All in all, hydrogen can only efficiently transport with hydrogen in the FCC structure of Al_0.5Cr_0.9FeNi_2.5V_0.2 under stress, dispersing hydrogen inside the grain and becoming trapped by irreversible hydrogen traps. Meanwhile, when the dislocations become fixed by the nano-precipitates, the dislocations can also trap hydrogen inside the grain, which further prevents the accumulation of hydrogen at the grain boundary. Thus, the possibility of hydrogen enrichment at the grain boundary is reduced through this coupling effect between the precipitates and dislocations. The ductile fracture morphology of Figure 3e,f, which mainly consists of dimples without intergranular fracture morphology, verified this hypothesis well. To assess the HE behaviour of Al_0.5Cr_0.9FeNi_2.5V_0.2, martensite steel 22MnB5 with a similar strength was also tested as a reference because high-strength steels are susceptible to HE. In this work, 22MnB5 steel showed dramatic hydrogen-induced strength loss. The intergranular fracture morphology on the whole fracture surfaces of the hydrogen-charged specimens reveals that hydrogen accumulates at the grain boundaries. Moreover, hydrogen-charged HEA shows shallower dimples than the uncharged specimens, indicating that HELP is the dominant mechanism during the tensile process and that the HEDE mechanism plays a less prominent part, while for the hydrogen-charged 22MnB5, the HEDE mechanism controls the fracture process.

The coupling effect of the precipitates and dislocations inside the grain makes it difficult to accumulate hydrogen atoms at the grain boundaries, reducing the possibility of cracks initiating at the grain boundaries and propagating along the grain boundaries. Table 3 shows that other conventional high-strength steels all exhibit dramatic hydrogen-induced strength loss and intergranular morphology on the fracture surface of the hydrogen-charged specimen, the microstructure of which consists of either a large number of dislocations inside the grain [26] or single precipitate structures [27,28,29,30,43]. This indicates that the single dislocation structure or the precipitate structure do not benefit the improvement of HE resistance. In contrast with the abovementioned literature, the ultra-high-strength Al_0.5Cr_0.9FeNi_2.5V_0.2 studied in this work greatly reduced the hydrogen-induced strength loss through the coupling effect of the precipitates and dislocations. Thus, the coupling effect of the well-dispersed precipitates and dislocations not only ensures the excellent mechanical properties of the alloy, but also significantly improves its resistance against hydrogen-induced strength loss.

5. Conclusions

To obtain the resistance of Al_0.5Cr_0.9FeNi_2.5V_0.2 to hydrogen embrittlement, a steel of a similar strength level, 22MnB5, was also tested for comparison. After hydrogen charging in a 0.05 mol/L H₂SO₄ + 1 g/L CH₄N₂S solution at 4 mA/cm² for 18 h, an SSRT was carried out to determine the hydrogen embrittlement susceptibility of the two steels, and their fracture morphologies were observed. The abovementioned results lead to the following conclusions:

(1)

Compared to high-strength martensite steel 22MnB5 with a similar strength level, Al_0.5Cr_0.9FeNi_2.5V_0.2 shows less strength loss after hydrogen charging and exhibits better resistance against hydrogen-induced strength loss.

(2)

The fracture surfaces of uncharged and hydrogen-charged Al_0.5Cr_0.9FeNi_2.5V_0.2 are characterized by dimples and no intergranular morphology is observed, mainly due to the HELP mechanism. The fracture mode of 22MnB5 steel changes from the dimple characteristic to the entire intergranular fracture feature after hydrogen charging, and the HEDE mechanism plays the predominant role in the fracture process.

(3)

Dispersed nano-structured precipitates and high-density dislocations hindered the aggregation of hydrogen, and this coupling effect effectively improves the resistance of Al_0.5Cr_0.9FeNi_2.5V_0.2 to hydrogen-induced strength loss.