1 Introduction

As a class of new materials, high-entropy alloys (HEAs) have properties that traditional materials do not have due to their unique design concepts [1,2,3], such as high strength, high ductility, high hardness, superplasticity, superconductivity, and other excellent properties (as shown in Fig. 1). The part of HEAs [4] represents a four-dimensional (4D) quinary phase diagram and the Co–Cr–Fe–Mn–Ni phase diagram. For the high strength part, the above one [5] shows the strain distribution after the tensile test, while the below one [6] provides the elemental distribution map and selected area electron diffraction (SAED) pattern about (CoCrFeNi)94Ti2Al4 HEA after 2 turn high-pressure torsion (HPT). The part of high hardness [7] shows the {111} X-ray diffraction (XRD) pole figures of sample after 1/4 and 1 turn HPT. The superconductivity part [8] reveals the pressure dependence of electronic density of states and electrical resistivity of the as-cast TiZrHfNb HEA under different pressures. The part of superplasticity [9] shows the electron backscatter diffraction (EBSD)-inverse pole figure (IPF) map and the corresponding phase map of annealed sample. Therefore, HEAs have attracted extensive attention from researchers [10,11,12,13,14]. At the same time, HEA's multi-major element concept provides a great impetus for the development of new alloys for advanced applications [4, 15]. The progress of high-pressure experimental technology has directly promoted the research in the field of HEAs [16]. At present, the main high-pressure methods used for studying the properties of HEAs are HPT technology and diamond anvil cells (DAC). In this paper, the mechanical properties of HEAs under high pressure are reviewed, and the potential properties of HEAs under high pressure are prospected.

Fig. 1
figure 1

reproduced with permission from Ref. [4]. Copyright 2017 Elsevier. The above part of high strength is reproduced with permission from Ref. [5]. Copyright 2021 Elsevier. The below part of high strength is reproduced with permission from Ref. [6]. Copyright 2019 Elsevier. The part of high hardness is reproduced with permission from Ref. [7]. Copyright 2019 Elsevier. The part of superconductivity is reproduced with permission from Ref. [8]. Copyright 2022 Elsevier. The part of superplasticity is reproduced with permission from Ref. [9]. Copyright 2020 Springer Nature

Some excellent properties of HEAs. The part of HEAs is

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2 HEAs

The earliest research on HEAs can be traced back to the proposal of crystal multi-major element alloys (MPEAS) in 1970 [11]. After a series of developments, the HEAs were systematically introduced in several works published in 2004 [3, 17,18,19,20]. HEA is an equiatomic and multi-element system, and there are many definitions of HEAs [21]. Two more representative definitions were based on component and entropy.

Defined by component, the HEA is an alloy consisting of five or more metal elements in tantamount or approximately tantamount quantities. The concentration of major elements should be between 5 and 35 at.%, and trace elements can be contained, but the concentration of minor elements should be less than 5 at.% [11, 20, 22]. Trace elements can change the properties of matrix HEAs, so different trace elements can produce different properties, thus expanding the HEA's system. However, this combination-based definition is only specified and stated in the dimension of elements, without mentioning the effect of entropy, which is not conducive to further understanding of the properties of HEAs.

From the viewpoint of entropy, the HEA is defined as the alloy whose mixed configuration entropy \((\Delta \)Smix) of the n-element equiatomic alloy changes from elemental state to random solution state (ideal state or regular state) [22]. \(\Delta \)Smix can be calculated according to Boltzmann's hypothesis

$$\Delta S_{\text{mix}}\hspace{0.17em}=\hspace{0.17em}{R}\mathrm{ln }n,$$
(1)

where R is the gas constant (8.31 JK−1·mol−1) and \(n\) is the number of elements for equimolar alloys. The calculation result of Eq. (1) is shown in Fig. 2a.

Fig. 2
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a Variation of the \(\Delta \)Smix/R with the number of elements in equiatomic alloy; b the classified alloys according to the mixing entropy in the random solution states

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According to the mixing entropy in the random solution state [22], alloy can be roughly divided into the following categories, (i) low entropy alloys, (ii) medium entropy alloys, and (iii) high-entropy alloys, as shown in Fig. 2b.

2.1 Four core effects

For the sake of further understanding the properties of HEAs, Yeh [22] summarized four core effects of HEAs, which favored the excellent properties of HEAs: (1) high-entropy effects; (2) sluggish diffusion; (3) severe lattice distortion; and (4) cocktail effects.

2.1.1 High-entropy effect

For the HEAs, high-entropy effect is the most basic characteristic. In simple terms, because of the high mixing entropy of the HEAs, it does not form a complex and fragile microstructure inside, but only consists of several solid solution phases or even a single phase. It can also be understood that the high entropy favors the stable high-entropy phase, namely the solid solution phase, instead of the intermetallic phase [22,23,24,25,26].

2.1.2 Sluggish diffusion effect

The sluggish diffusion effect is a significant reason why HEAs have many excellent properties [27], which leads to excellent high-temperature strength [28, 29]. In short, the diffusion and phase transition dynamics of HEAs are slower than those of traditional alloys [25, 30].

2.1.3 Severe lattice-distortion effect

Because the size of different elements in HEAs is imparity in many cases, this size difference will bring about serious lattice distortion and will also affect the properties of HEAs [22, 25]. As shown in Fig. 3, the different dimensions of the elements cause severe lattice distortion.

Fig. 3
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Schematic illustration of lattice distortion

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2.1.4 Cocktail effect

Ranganathan [31] was the initial person to propose the cocktail effect of alloys. The original intention was “pleasant mixture” to describe the characteristics of trace elements in alloys that regulated their properties. Later, it meant a synergistic mixture that the final result was unpredictable and larger than the sum of its parts. The cocktail effect of HEAs is that their properties are not only the superposition of the properties of each element, but also related to the interaction between elements [20, 23].

3 High-pressure technology

With the rapid development of high-pressure technology, it provides new possibilities for material research [32, 33]. At present, the main methods to realize high pressure in laboratory are HPT technology and diamond anvil cells.

In 1935, Bridgman [34] first proposed HPT technology. The working principle is shown in Fig. 4a. HPT is a severe plastic deformation (SPD) technique in which samples are subjected to torsional shear strain under high hydrostatic pressure [35]. By placing the sample between two anvils, when the external force drives the relative rotation of the two anvils, the friction between the sample and the anvil drives the rotation of the sample, causing the shear deformation of the sample. Many materials can withstand large shear strains without failure through HPT treatments [34, 36, 37].

Fig. 4
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a Schematic diagram of high-pressure torsion and b diamond anvil cells

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In 1955, Bundy [38] synthesized diamond with high-pressure and high-temperature technology, making it possible for high-pressure technology to be widely used in material research. The first diamond anvil cell was made in 1958 with the tools at hand [39]. The diamond anvil structure is shown in Fig. 4b.

The DAC is composed of two opposite diamond anvils. By turning external screws, the sample is squeezed to generate hydrostatic/non-hydrostatic pressure [16, 33, 40]. In 1977, Buras et al. [41] first combined synchrotron radiation XRD method with DAC, which promoted the development of high-pressure technology.

4 Mechanical properties of HEAs under HPT

4.1 Change in hardness

High-pressure treatment can refine the grain size of HEAs and improve their hardness. For example, it can affect a sample by changing the hardness and softness of its phase. Stepanov et al. [42] made the AlNbTiVZr0.5 HEA disk rotate 0.25, 0.5, 1, and 5 rotations at room temperature, 8 GPa pressure, and 0.5 r·min−1, respectively. The nano-hardness of single phase indicated that the increase of hardness was due to the structural refinement of B2 phase. At the same time, the Laves phase became softer during HPT processing. As shown in Fig. 5a, the microhardness increased with the increasing number of turns and the increasing distance from the center of the disk until it reached the saturation state. After five HPT rotations, the microhardness of the sample increased from 550 HV to a maximum of 665 HV.

Fig. 5
figure 5

Reproduced with permission from Ref. [42]. Copyright 2018 Elsevier. b Two-dimensional contour map of hardness of samples treated with different HPT; c IPF maps of samples with monotonic 100°, cyclic 100°, monotonic 360°, and cyclic 360° HPT-treated. Reproduced with permission from Ref. [43]. Copyright 2019 Springer Nature. d The Vickers hardness of nanocrystals CoCrFeNiMn after different HPT treatment as a function of the distance from the center of the disk; e simulation diagram of nanocrystalline CoCrFeNiMn torsional stress with disk center distance after compression and 1/5 HPT treatment. The distribution of torsional stress on the surface of semicircular disk was shown in different colors. Reproduced with permission from Ref. [44]. Copyright 2015 Springer Nature. f The Vickers microhardness evolution of nanocrystalline CoCrFeMnNi as a function of center distance after 1/4, 1/2, 1, and 2 turns of HPT treatment. Reproduced with permission from Ref. [45]. Copyright 2017 Elsevier

a The relationship between the microhardness of AlNbTiVZr0.5 alloy and the distance from the center of disk after different HPT turns (N).

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Meanwhile, the change of strain path also affects the hardness of HEAs. Sonkusare et al. [43] performed monochromatic and cyclic HPT on a CoCuFeMnNi HEA disk at room temperature and 5 GPa. Five forward and reverse cyclic single-step and multi-step torsions at 100° and 360° were performed, respectively. It was shown that the microhardness of the HPT-treated sample was significantly higher than that of the annealed sample. Figure 5b visually shows the change in hardness for the entire quarter of the disk subjected to HPT. The periphery of the 100° cyclic HPT sample exhibited the highest hardness due to the change of strain path, and significant grain refinement occurred with the evolution of complex dislocation substructure. The low hardness of HPT alloy with 360° cycle was caused by the unstable dislocation substructure which promoted grain boundary migration and grain growth. In Fig. 5c, the crystal orientation or IPF map near the center of the HPT-processed disk showed that there was a large color gradient inside the grains, indicating that there was an orientation gradient due to the dislocation activity.

HPT can be used to improve the hardness of nanocrystalline HEAs. Lee et al. [44] performed 1/4, 1/2, 1, and 2 turns of nanocrystalline CoCrFeNiMn HEA disks at room temperature, 6.0 GPa pressure, and 1 r·min−1. The Vickers hardness test of the samples after treatment is shown in Fig. 5d. HPT treatment significantly improved the hardness of the samples, and the grain refinement degree was positively correlated. The nanocrystalline structure was also obvious in the early stages of HPT processing due to the severe torsional stresses which strengthen the atomic diffusion. At the same time, the finite-element analysis simulation of the radial variation of 1/5 HPT-treated samples from the center of the disk was carried out. In Fig. 5e, finite-element simulation estimated the variation of torsional stress τt from the center of the disk along the radial direction. The simulated rotation angle was 70°, which was close to 1/5 HPT. Different stresses were represented on the disk in different colors. As the distance from the center increases, τt increased significantly and exhibited maximum stress of 3.26 GPa at the edge.

Heczel et al. [45] conducted HPT tests at room temperature, 3.0 GPa pressure, and 1 r·min−1 on nanocrystalline CoCrFeMnNi with 1/4, 1/2, 1, and 2 turns, respectively. The hardness measured along the disk diameter after HPT treatment is shown in Fig. 5f. At 1/4 turn, the hardness gradient along the radius of the disk was large. At 1/2 turn tight, there was mild hardening in the center and periphery. It was also observed that the hardness of the middle area increased greatly from 1/2 turn to 1 turn. It was observed that the hardness changed little during the process of 1–2 turns, and a higher number of turns might produce the completed saturation of hardness. The maximum hardness of the outer disk reached 5.38 GPa after two turns. Basically, the smaller the grain size, the greater the hardness; the greater the change of grain size, the greater the degree of hardness improvement. Not only grain refinement but also twins were taken shape in the grain. The EBSD images (Fig. 6a, b) of the 1 HPT-treated sample showed that twins with an average spacing of 3 μm were formed in the crystal.

Fig. 6
figure 6

Reproduced with permission from Ref. [45]. Copyright 2017 Elsevier. c Vickers microhardness versus distance from disk center; SEM–EDS analysis of d the initial AlCrFeCoNiNb high entropy alloy and e 1/4 turns HPT-treated and f 4 turns HPT-treated samples to check the element distribution map. Reproduced with permission from Ref. [46]. Copyright 2021 Elsevier. g The Vickers microhardness of Al0.1CoCrFeNi at 6 GPa after 1 and 2 cycles of hot-isostatic pressed (HIPed) and HPT treatments as a function of the distance from the center of the disk. Reproduced with permission from Ref. [47]. Copyright 2016 Elsevier

a, b EBSD images of nanocrystalline CoCrFeMnNi sample disk center after 1 turn of HPT, and further magnification of part of the region.

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HPT can also be used to improve the hardness of nanostructured dual-phase HEAs. Edalati et al. [46] treated nanostructured dual-phase AlCrFeCoNiN HEA with N = 1/16, 1/4, 1, and 4 cycles of HPT at 300 K temperature, 6 GPa pressure, and 1 r·min−1 rotation speed, respectively. The measurement of microhardness of samples treated under different conditions (Fig. 6c) showed that after 4 cycles of HPT treatment, the sample reached the ultra-high hardness of 1030 HV. This was caused by nanoparticles with high lattice distortion, interface formation dislocation generation, and average size of 10 nm. Meanwhile, scanning electron microscope–energy-dispersive X-ray spectroscopy (SEM–EDS) analysis was performed on the initial sample, 1/4 and 4 laps of HPT samples to detect the distribution of various elements, and the results were shown in Fig. 6d–f. It should be noted that after HPT treatment, the composition of the phase did not change, but each phase was mixed, and it was obvious with the increase of HPT treatment winding number.

The saturation state of the material can be investigated by the HPT effect due to the balance between hardening and softening of the material. Yu et al. [47] treated Al0.1CoCrFeNi with HPT for 1 and 2 turns at a speed of 1 r·min−1 and a pressure of 6 GPa. Vickers microhardness was measured along the radius of the treated samples. The results are shown in Fig. 6g. Under hot-isostatic pressing, the microhardness increased from 135 HV to about 482 HV, and increased to saturation with the increase of strain (from center to edge). At the same time, the results showed that the edge hardness of the sample was about 482 HV, which reached the balance of softening and hardening, and the material reached the saturation state.

HPT is a rapid synthesis method for preparing HEAs nanocomposites [48]. Taheriniya et al. [49] connected two single-phase equiatomic bcc-HfNbTaTiZr and fcc-CoCrFeMnNi HEAs disks by HPT method. The connection was achieved by rotating 5, 10, 15, and 20 times at an ambient temperature and hydrostatic pressure of 9 GPa, and Vickers microhardness measurements were performed on the treated samples (as shown in Fig. 7a). The results showed that the microstructure of the newly formed nanocomposites was refined and the hardness was increased by 5–20 revolution. Under the same conditions, the hardness of co-deformed double bcc–fcc nanocomposites exceeded that of the two single-phase materials. Figure 7b, c shows high-resolution and high-speed nanoindentation patterns obtained by continuous stiffness measurement (CSM) method from the edge of HPT deformed samples after 10, 15, and 20 revolutions. The hardness distributions of different phases and phase boundaries as a function of deformation were summarized. Figure 7a–c jointly confirms that the significant increase in hardness of mixed HEAs was related to the non-uniform co-deformation of the two-phase lamellar HEA composite.

Fig. 7
figure 7

Reproduced with permission from Ref. [49]. Copyright 2021 Elsevier. d The variation of as-cast Vickers hardness and HPT HEA with annealing temperature. Reproduced with permission from Ref. [50]. Copyright 2015 Elsevier

a Relationship between dual-phase HEAs microhardness and center distance after 5, 10, and 15 HPT treatments. The Vickers microhardness curves of the single-phase fcc and bcc alloys were given; high-resolution and high-speed nanoindentation images of dual-phase HEAs at 10, 15, and 20 revolutions for b cluster diagram of average hardness of two different phases and for c hardness intensity distribution in two phases.

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The combination of HPT and thermal annealing treatment is a method to further improve the hardness of HEAs. Tang et al. [50] treated Al0.3CoCrFeNi samples with HPT at 8 r·min−1 at room temperature and 6.0 GPa, resulting in the steady-state grain size and uniform microstructure of the samples. Then, the untreated samples and HPT samples were annealed at different temperatures, and the microhardness was measured, respectively. As shown in Fig. 7d, the maximum hardness of the HPT sample was about 4 times that of the original hardness of as-cast HEA, and the increase of hardness was mainly caused by the grain refinement to the nanometer range. In addition, during the annealing process, HPT samples formed hard secondary phase (ordered annealing bcc structure rich in Al and Ni), which further improved the hardness.

HPT can also be applied to the synthesis of dense HEAs. Elemental powders of Co, Cr, Fe, Mn, and Ni were mixed in argon with a medium molar ratio by Kilmametov et al. [51]. Then, the mixed powders were treated with HPT, respectively. The microhardness test results are shown in Fig. 8a. Due to the uneven distribution of elements, incomplete phase formation, or low material density, the hardness in the central region of the sample increased abnormally after 1 rotation. After 10 rotations, the maximum hardness of 7.3 GPa was measured at the edge of the sample. At the same time, the uniformity of elements at nanometer scale was analyzed by atom probe tomography (APT) and the oxidation particles were studied. Figure 8b shows the reconstruction of Co, Cr, Fe, Mn, and Ni as well as the distribution of CrO and O ions at the tip. The results showed that the concentration of the main elements along the tip did not fluctuate much and a Mn-enriched grain boundary could be authenticated during the reconstruction process. The distribution of CrO and O could reveal the formation of several nano-sized Cr oxides.

Fig. 8
figure 8

Reproduced with permission from Ref. [51]. Copyright 2019 Elsevier. c Vickers hardness of isochronous (1 h) treated samples at different annealing temperatures; d Vickers hardness of samples treated for unequal time at the same annealing temperature (450 °C). Reproduced with permission from Ref. [52]. Copyright 2015 Elsevier

a Relationship between microhardness and distance from center of CoCrFeMnNi treated with 1, 10, and 100 turns of HPT; b distribution of the atoms in the tip atom. Mn-enriched grain boundary is marked during Mn reconstruction. Both the detected CrO and O ions are displayed. Only 1% of the other ions were extracted from the probe.

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HPT combined with annealing treatment can break through the saturation state of the material and further improve the hardness of the material. Schuh et al. [52] treated CoCrFeMnNi HEA with HPT for 1/4, 1, and 5 turns at 7.8 GPa pressure, 0.2 r·min−1, and room temperature. The microhardness was significantly improved. Then, the samples treated with 5 cycles were annealed at disparate temperatures. The results are shown in Fig. 8c. At 0–500 °C, the microhardness increased significantly even in the saturated state. The 5 HPT-treated samples were then isothermally treated at 450 °C for up to 200 h. As shown in Fig. 8d, the hardness continued to increase from the SPD state to a peak of 910 HV after 100 h of annealing, and then began to decline again. To correlate hardness variations with microstructure, 3-dimensional (3D)-APT treatment was performed on samples annealed at 450 °C, and the results are shown in Fig. 9a, b. The volume fraction increased obviously at 1 h. At 15 h, an additional phase mainly composed of Fe and Co is formed, and the volume fraction of the first two phases further increases. With the extension of annealing time, the composition of Mn, Ni phase remained unchanged and the content of Cr phase increased. After that, the content of each phase continued to change, but due to the high brittleness of the material, the experiment could not be continued, so it needs to be further studied.

Fig. 9
figure 9

Reproduced with permission from Ref. [52]. Copyright 2015 Elsevier. c Engineering stress extension curves of CoCrFeNiMn HEA at 6 GPa, 1 r·min−1 and 1.0 × 10–3 s−1 strain rates at different temperatures. Reproduced with permission from Ref. [55]. Copyright 2017 Elsevier

a Chemical composition distribution on probe image; b Cr, Fe, Ni, Co, and Mn ion maps reconstructed by CoCrFeMnNi 3D-APT after severe plastic deformation.

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4.2 Changes in superplasticity

Superplasticity is the extreme plastic deformation of fine-grained materials (up to 10 microns) at temperatures greater than 0.5 Tm (Tm is the melting temperature) with a tensile elongation of at least 400% [53, 54]. Shahmir et al. [55] applied HPT treatment to CoCrFeNiMn HEA for 5 turns at room temperature, 6.0 GPa pressure, and 1 r·min−1. The grain size was refined to about 10 nm, and grain boundary sliding (GBS) occurred. The maximum total elongation at each temperature was measured, as shown in Fig. 9c. It could be seen that superplastic stretching (> 400%) occurred at 873 and 973 K, which proved the superplasticity of the material. Figure 10a–d shows the microstructure of the grip section and the gauge section of the specimens deformed at 973 K, at a strain rate of 1.0 × 10–3 s−1 and a stretch direction of tensile direction (TD). The arrows indicated the precipitates of the sample tested at 973 K.

Fig. 10
figure 10

Reproduced with permission from Ref. [55]. Copyright 2017 Elsevier. e, f SEM images of the fracture samples after tensile test at 1073 K, 3.3 × 10–3 s−1, 700% elongation. Reproduced with permission from Ref. [56]. Copyright 2019 Elsevier

a SEM images of 973 K samples with initial strain rate 1.0 × 10–3 s−1 and maximum elongation is 610%; b–d position corresponding to B, C and D in figure a, respectively. Arrows indicate precipitation and TD.

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Nguyen et al. [56] treated V10Cr15Mn5Fe35Co10Ni25 HEA with 5 turns of HPT at room temperature, 1 r·min−1, and 6 GPa under applied pressure. The total elongation results are shown in Table 1. The material reached superplasticity and the temperature range for obtaining good elongation was larger. It could also be seen that the total elongation decreased with the decrease of strain rate, which was probably due to dynamic recovery process. The XRD patterns showed that there was a small fraction of σ phase in the samples tested at 973 K and 1073 K, and the formation of σ phase could be used as an obstacle to restrict the growth of fine grains and stabilize fine grains, thus contributing to the promotion of superplastic deformation. At a strain rate of 3.3 × 10–3 s−1, the electron microscopy of 973 K and 1073 K samples from grip to fracture is shown in Fig. 10e, f. At 973 K, the width of the gauge decreased gradually to achieve quasi-uniform deformation, while at 1073 K, the width of the gauge was completely uniform.

Table 1 Elongation results and peak stresses of the HPT-processed V10Cr15Mn5Fe35Co10Ni25 HEA at strain rates and different temperatures.
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Nguyen et al. [9] also achieved ultra-high strain rate superplasticity of nanostructured HEAs using HPT technology. An extremely rare high strain rate superplasticity (HSRS) has been achieved in high strength materials. The 10 mm Al9(CoCrFeMnNi)91 HEA disk was processed by HPT for 5 turns at room temperature, 6 GPa applied pressure, and 1 r·min−1. The transmission electron microscope–electron dispersive X-ray spectroscopy (TEM–EDS) analysis of the sample (Fig. 11a) showed that sigma phase (rich in Cr) was formed at 1073 K, and fcc and B2 phases were removed in the superplastic tests. Newly formed σ phase and B2 phase could prevent significant grain growth and facilitate grain boundary slip (GBS) with high strain rate sensitivity. In B2 and fcc phases, hole interconnections and dislocation activity were observed to be adapted to GBS. Fine grain size was associated with high strain rate sensitivity and GBS adaptation to plastic deformation. In other words, the combination of grain refinement induced by HPT and finite grain growth during thermoplastic deformation resulted in high strain rate superplasticity through GBS to which dislocation activity was suited. This kind of superplasticity under high strain rate greatly saved production time and promoted industrial production. Figure 11b–d shows the microstructure of the samples at different locations. It was found that there was no significant difference in grain size between the deformed and undeformed regions of the samples; thus, it can be concluded that there was almost no dynamic grain growth in the deformed region. The superplasticity test results are shown in Fig. 12a. The superplastic elongation of Al9(CoCrFeMnNi)91 HEA nanostructure was found to be 2000% of the original length at high strain rate (5 × 10–2 s−1).

Fig. 11
figure 11

Reproduced with permission from Ref. [9]. Copyright 2020 Springer Nature

a TEM-EDS images and phase diagrams of the grip, middle, and tip portions of the tested specimens were obtained at 1073 K and 5 × 10–2 s−1; b–d Al9(CoCrFeMnNi)91 HEA scanning transmission electron microscope (STEM) images in the grip, middle, and tip.

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Fig. 12
figure 12

Reproduced with permission from Ref. [9]. Copyright 2020 Springer Nature. b Engineering stress–strain curves of CoCrFeMnNi HEA powder billet after 4 HPT and post-HPT annealed; EBSD-IPF diagrams of the post-HPT annealed samples at c 700 °C for 15 min; d 700 °C for 60 min; e 800 °C for 15 min; f 800 °C for 60 min. Reproduced with permission from Ref. [5]. Copyright 2021 Elsevier. g Engineering stress–strain curves of V10Cr15Mn5Fe35Co10Ni25 HEA under different conditions. Reproduced with permission from Ref. [63]. Copyright 2020 Elsevier

a Engineering stress–strain curves of Al9(CoCrFeMnNi)91 in tensile tests at 873 K, 973 K and 1073 K and at 10–1 ~ 10–2 s−1, respectively.

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4.3 Good combination of ductility and strength

The balance between ductility and strength is a long-lasting problem in the field of materials research, and it is no exception in HEAs [57,58,59]. However, Pan et al. [60] recently solved this problem by controllable introduction of a novel gradient nano-scaled dislocation-cell structures in a stable single-phase fcc-Al0.1CoCrFeNi HEA. Also, we found that HPT and annealing treatment of HEAs could effectively solve this problem [61].

Asghari-Rad et al. [5] treated CoCrFeMnNi samples with 4 cycles of HPT at 5 GPa pressure and 1 r·min−1. The results in Fig. 12b showed that the tensile strength of the specimen was about 1867 MPa after HPT, but the elongation was low. The HPT disks were then annealed at 700 °C and 800 °C for 15 and 60 min in an argon atmosphere. After annealed processing, the toughness of samples was restored, but the tensile strength was decreased. The reason was that with the increase of annealing time and temperature, the dislocation density and grain growth caused by HPT decrease. Figure 12c–f shows EBSD-IPF maps of sample treated with HPT under various annealing conditions. The results showed that the average grain size of the sample reached 844 nm after annealing at 800 °C for 60 min. It increased with the increase of temperature and annealing time. When annealed at 700 °C, the σ particles produced in the samples would reduce the tensile strength. The final sample annealed at 800 °C for 60 min obtained a significant combination of strength and elongation.

Shahmir et al. [62] treated the CoCrFeNiMn HEA with 1, 5, and 10 turns of HPT at room temperature and 6.0 GPa. The samples were then annealed at 473–1173 K for different times. The results are as shown in Table 2. The short-term annealing at 1073 K and 10 min could prevent the grain from growing and obtain high strength (830 MPa) and good plasticity.

Table 2 Microhardness, yield strength (YS), ultimate tensile strength (UTS), and elongation (δ) of CoCrFeNiMn HEA before HPT and after post-deformation annealing (PDA) at 1073–1173 K for 10–60 min.
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Asghari-Rad et al. [63] processed V10Cr15Mn5Fe35Co10Ni25 HEA for five turns of HPT under quasi-constraint conditions, 25 °C, 6 GPa, and 1 r·min−1. The samples were then annealed at 600–800 °C for 2–60 min. As shown in Fig. 12g, annealing at 700 °C for 10 min could obtain the best mechanical properties with high tensile strength and good elongation. Figure 13 shows the EBSD-IPF diagram of V10Cr15Mn5Fe35Co10Ni25 annealed at different temperatures and times. The results showed that the grain growth tended to be single at 700 °C with increasing annealing time. The annealing process at 800 °C for 60 min increased sharply the average grain size to ~ 1.4 μm.

Fig. 13
figure 13

Reproduced with permission from Ref. [63]. Copyright 2020 Elsevier

EBSD-IPF maps of annealed samples after deformation. The color scale on the right corresponds to the inverse pole diagram of [001].

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5 Mechanical properties of HEAs under DAC

DAC is also commonly used in the study of HEAs. Using XRD and DAC techniques, Zhang et al. [64] studied the deformation behavior of MoNbTaVW HEA under compression conditions. Figure 14a shows the function images of azimuthally unrolled diffraction patterns of MoNbTaVW and W with direction angle before compression, under 5 GPa–20 GPa pressure and after pressure relief. At 5 GPa, both W and MoNbTaVW form in the bcc structure, resulting in strong lattice strain and texture. At 20 GPa, the diffraction peak of MoNbTaVW was much sharper than W, indicating that it had a larger lattice strain at high pressure. In the process of pressure relief, no new diffraction peaks were found, which confirmed the phase stability of W and MoNbTaVW. The XRD patterns (Fig. 14b) showed that MoNbTaVW had stronger resistance to high pressure than pure tungsten. It was founded that the active dislocation behavior was the main reason for the high strength of MoNbTaVW.

Fig. 14
figure 14

Reproduced with permission from Ref. [64]. Copyright 2021 Elsevier

a Azimuthal diffraction images of MoNbTaVW (left) and pure W (right) at different pressures. The diffraction of Au is indicated at the top of the figure. Pressure is marked on the right side of each figure; b XRD patterns of the MoNbTaVW/Au mixture during decompression and compression.

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Also using XRD and DAC techniques, Zhang et al. [65] investigated the effect of dilute carbon doping on the mechanical properties of Mo15Nb20Re15Ta30W20 HEA. The results show that high pressure/temperature promotes the formation of Ta3C and Nb3C carbides. The stable Ta3C and Nb3C carbides may further improved the mechanical properties of the Alloy. Table 3 indicates that low carbon concentration had small effect on the hardness of the alloy.

Table 3 Chemical compositions (at. %) and Vickers microhardness of intermediate alloy (alloy) and C0.1-alloy (C-doped).
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Zhang et al. [66] synthesized carbon-doped MoNbTaW alloy through ball milling. The performance of the alloy at 3.1 GPa and different temperatures was studied using the resistor-heating DAC. As shown in Fig. 15a, fcc carbides (TaC and NbC) and bcc (compositionally complex alloys) phases are present in the sample, and it is found that carbides are more easily obtained when high concentrations of C are added. XRD patterns of the samples (Fig. 15b) indicated good stability with increasing temperature from room temperature to 870 °C at 3.1 GPa. The mechanical properties of C-MoNbTaW were expected to be improved due to the emergence of favorable carbide phases (TaC and NbC) of hard components in the composites.

Fig. 15
figure 15

Reproduced with permission from Ref. [66]. Copyright 2021 MDPI. c HRTEM image of quenched Al0.3CoCrFeNi sample from 53 GPa. Reproduced with permission from Ref. [72]. Copyright 2019 AIP Publishing. d XRD patterns of CrMnFeCoNi at different pressures; e XRD patterns of quenched CrMnFeCoNi from different pressures. Reproduced with permission from Ref. [74]. Copyright 2017 Springer Nature

a XRD images of sintered bulk sample (red) and carbon-doped MoNbTaW powder by ball mill method (black); b XRD results of C-MoNbTaW/MgO/Pt mixture with different temperatures at 3.1 GPa.

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Yusenko et al. [67] synthesized the fcc-structured Ir0.26Os0.05Pt0.31Rh0.23Ru0.15 refractory alloy and studied its compression behavior with DAC technique. The alloy was stable up to 49 GPa at room temperature and the compressibility curve could be fitted with a bulk modulus of 300 (22) GPa, which is little bit smaller than the hcp-Ir0.24Os0.21Re0.16Rh0.18Ru0.20 (317(2) GPa) [68]. Ahmad et al. [69] performed XRD experiments on three typical HEAs fcc-Ni20Co20Fe20Mn20Cr20, bcc-Hf25Nb25Zr25Ti25, and hexagonal close-packed (hcp) Re25Ru25Co25Fe25 (at.%) under extreme pressure and temperature conditions. The results showed that the three HEAs exhibited great stability up to 80 GPa. The bulk modulus and its pressure derivative of three HEAs were determined and the bcc-Hf25Nb25Zr25Ti25 had a relatively larger compressibility than the other two alloys. Yu et al. [70] synthesized HoDyYGdTb HEA and performed high-pressure XRD experiments up to 60.1 GPa at room temperature. The calculated bulk modulus (B0 = 45.8 GPa) and atomic volume (V0 = 19.3 cm3·mol−1) agreed well with the values calculated by the “additivity law” [71], which is favorable for mechanical properties.

Wang et al. [72] investigated the effect of Al content on the structure of HEAs AlxCoCrFeNi (X = 0, 0.1, 0.3, 0.75, 1.5) under high pressure. XRD patterns showed that the structure of Al0.3CoCrFeNi remained stable up to 53 GPa. Moreover, high-resolution TEM (HRTEM) image (Fig. 15c) showed that a new nano-twinned fcc structure was formed, which was proposed as a source of high maximum strength of the alloy [73]. Tracy et al. [74] studied the structure properties of CrMnFeCoNi HEA by angle-dispersive XRD measurements at pressures of up to 54.1 GPa. The results are shown in Fig. 15d. A pressure-induced structure transition from fcc phase to a unique hcp phase was observed starting at 14 GPa. The hcp phase could be quenched to ambient pressure, forming fcc-hcp mixtures. As shown in Fig. 15e, the content of hard hcp phase and ductile fcc phase could be controlled effectively, thus providing a distinctive performance with combined high ductility, strength, and toughness.

6 Conclusion and outlooks

As an emerging alloy material in recent years, HEAs have been at the forefront of advanced metal material research. As a powerful means to study the properties of materials, high-pressure technology brings new possibilities to the properties of HEAs. HPT treatment can be applied to many kinds of HEAs, and the hardness of these alloys can be significantly improved. High-density HEAs and high-entropy nanocomposites were synthesized by HPT treatment, and their hardness was significantly improved. In addition, a combination of good ductility and high strength can be effectively achieved by HPT processing combined with thermal annealing. Besides, the DAC technology combined with XRD technology also provides a way for synthesizing novel structure as well as optimizing the mechanical properties of HEAs.

Although sufficient kinds of HEAs have been developed, only a few of them have been studied under high pressure. In the future, with the help of high-pressure technology, more HEAs will be explored extensively under high pressure. Inspired by the combination of high-pressure technology and thermal annealing treatment, potential properties such as creep under high-pressure high/low-temperature conditions, the dynamic response under high strain rate based on HPT technology, and pressure-directed microstructure optimization (e.g., plasticity-assisted redistribution, grain boundary effect, optimal combination of strength, and plasticity) are worth more exploration in future research.