**Effect of Annealing on Microstructure and Tensile Behavior of CoCrNi Medium Entropy Alloy Processed by High-Pressure Torsion**

**Praveen Sathiyamoorthi 1,2, Jae Wung Bae 1,2, Peyman Asghari-Rad 1,2, Jeong Min Park 1,2, Jung Gi Kim <sup>1</sup> and Hyoung Seop Kim 1,2,3,\***


Received: 19 September 2018; Accepted: 2 November 2018; Published: 6 November 2018

**Abstract:** Annealing of severely plastic deformed materials is expected to produce a good combination of strength and ductility, which has been widely demonstrated in conventional materials. In the present study, high-pressure torsion processed CoCrNi medium entropy alloy consisting of a single face-centered cubic (FCC) phase with a grain size of ~50 nm was subjected to different annealing conditions, and its effect on microstructure and mechanical behavior was investigated. The annealing of high-pressure torsion processed CoCrNi alloy exhibits partial recrystallization and near full recrystallization based on the annealing temperature and time. The samples annealed at 700 ◦C for 2 min exhibit very fine grain size, a high fraction of low angle grain boundaries, and high kernel average misorientation value, indicating partially recrystallized microstructure. The samples annealed for a longer duration (>2 min) exhibit relatively larger grain size, a low fraction of low angle grain boundaries, and low kernel average misorientation value, indicating nearly full recrystallized microstructure. The annealed samples with different microstructures significantly influence the uniform elongation, tensile strength, and work hardening rate. The sample annealed at 700 ◦C for 15 min exhibits a remarkable combination of tensile strength (~1090 MPa) and strain to failure (~41%).

**Keywords:** medium entropy alloy; high-pressure torsion; partial recrystallization; tensile strength

#### **1. Introduction**

High entropy alloys (HEAs), also known as compositionally complex alloys, baseless alloys, and concentrated solid solution alloys, exhibit unique and remarkable properties as compared with the conventional alloys [1–3]. The unique properties of HEAs are mainly attributed to the distinct alloy design concept based on multi-principal elements [4,5]. The alloys with multi-principal elements are widely classified into medium entropy alloys (MEAs) and HEAs based on the configurational entropy [5,6]. Among the HEAs, alloys based on the Co–Cr–Fe–Mn–Ni system, Al–Co–Cr–Cu–Fe–Ni system, and refractory elements are widely reported in the literature [1,2,5].

Recently, CoCrNi MEA has received widespread attention from researchers because of its unique properties such as exceptional fracture toughness at cryogenic temperature, high friction stress, superior dynamic shear properties, and annealing-induced hardening at intermediate temperatures (500–600 ◦C) [7–10]. It is widely reported that HEAs with a single face-centered cubic (FCC) phase

exhibit low yield strength [11–13]. Several strategies have been reported to increase the strength of a single FCC phase HEAs by solid solution strengthening, precipitation strengthening, grain boundary strengthening, and partial recrystallization through thermo-mechanical treatments [11,14–16]. Among the several strategies to enhance the mechanical properties, thermo-mechanical treatment is an effective technique to achieve a good combination of strength and ductility by achieving microstructure with different grain size and/or different volume fraction of secondary phases [11,17,18]. The CoCrNi MEA with a single FCC phase, like any other HEA with a single FCC phase, also exhibits lower yield strength [14,19]. Recently, it has been reported that the addition of 3 at% W resulted in a 20% increase in intrinsic strength [14]. In another study, cold rolling followed by annealing of CoCrNi alloy is reported to have an ultra-fine-grained size (~750 nm) with a good combination of tensile strength (964 MPa) and strain to failure (66%) [20]. In this connection, high-pressure torsion (HPT) followed by annealing has been effectively used in conventional alloys and HEAs to improve their mechanical properties [21–23]. Generally, HPT is an effective process in fabricating materials with nano-grains or ultra-fine grained structure by simultaneous application of torsion and high hydrostatic pressure [22]. The HPT processed CoCrNi MEA is reported to have a grain size of ~50 nm with an exceptional ultimate tensile strength of ~2.2 GPa and a considerable strain to failure of 9% [24]. It is also reported that annealing of the HPT processed CoCrNi MEA in the temperature range of 500–600 ◦C has resulted in an increase in hardness due to a decrease in dislocation density and grain boundary relaxation [7]. However, the systematic investigation of post HPT annealing of CoCrNi MEA with respect to microstructure evolution and strength enhancement is limited.

Thus, the present study is aimed at achieving a better combination of strength and ductility in the HPT processed CoCrNi MEA by post-HPT annealing. Consequently, in the present study, the post-HPT annealing effect on the microstructural and mechanical behaviors of CoCrNi MEA is investigated.

#### **2. Experimental Methods**

The CoCrNi alloy for the present study was fabricated using a vacuum induction melting of pure metals followed by homogenization, cold rolling, recrystallization, HPT, and annealing of HPT samples. The detailed procedure about the fabrication and secondary processing route can be found in our earlier publications [7,24]. The microstructure of starting material for HPT consists of a single FCC phase with an average grain size of ~19 μm and profuse annealing twins [24]. After the HPT process, the microstructure consists of a very fine grain size of ~50 nm with a high density of nano-twins and stacking faults [24]. In this study, the HPT processed sample with a single FCC phase was subjected to different heat treatment conditions in the Ar atmosphere to investigate the effect of post-HPT annealing. The post-HPT annealed samples were characterized using X-ray diffraction (XRD, Rigaku D/MAX-2500), field emission scanning electron microscopy (FE-SEM, FEI XL30S) coupled with electron backscatter diffraction (EBSD), Vickers hardness, and tensile testing equipment (Instron 1361). The deformed microstructures of post tensile samples were characterized using a focused ion beam (FIB) coupled with transmission Kikuchi diffraction analysis (TKD, FEI-FIB/TKD).

Tensile samples (dimension: 1.5 mm gauge length and 1 mm gauge width) were cut from the annealed HPT disc using electric discharge machining. Tensile tests were carried out at a quasi-static strain rate of 10−<sup>3</sup> s−<sup>1</sup> and the strain on the sample was measured using digital image correlation (DIC, ARAMIS v6.1, GOM Optical Measuring Techniques). The Vickers hardness measurement was performed using a load of 300 gf and a dwell time of 15 s. For EBSD experiment, the samples were mechanically polished using SiC carbide papers followed by diamond and colloidal polishing. For TKD experiment, a transparent sample near to the fracture surface of deformed tensile samples was lifted out by FIB.

#### **3. Results and Discussion**

Figure 1 shows the XRD patterns of CoCrNi alloy under different processing conditions: HPT and post-HPT annealing. It can be seen that in all the processing conditions, the alloy exhibits a single FCC phase.

**Figure 1.** X-ray diffraction (XRD) patterns of CoCrNi alloy in different processing conditions illustrating the presence of a single face-centered cubic (FCC) phase.

The formation of sigma phase/secondary phases is widely reported for HEAs processed by HPT and annealing. In Al0.5CoCrFeMnNi HEA, the formation of sigma phase is reported when the HPT processed sample is annealed at 800 ◦C for 1 h [25]. Similarly, for the HPT processed CrFe2NiMnV0.25C0.125 HEA, sigma phase has been observed during annealing at 550 ◦C [26]. In CoCrFeMnNi HEA, the formation of sigma and secondary nano-phases have been reported after annealing of the HPT processed sample [27]. However, in the present alloy, there is no evidence of formation of sigma phase or secondary phase in the temperature range of 600–800 ◦C. The formation of sigma phase in this alloy was also not observed when annealed in the temperature range of 500–600 ◦C [7].

Figure 2 shows the inverse pole figure (IPF) map of the post-HPT annealed samples, and the corresponding kernel average misorientation (KAM) map and misorientation angle chart are shown in Figures 3 and 4, respectively.

**Figure 2.** Electron backscatter diffraction (EBSD) inverse pole figure map of post-high-pressure torsion (HPT) annealing samples. The color scale on the right side corresponds to [001] inverse pole figure.

The IPF map indicates an increase in grain size with increasing annealing temperature and time. The samples annealed at 600 ◦C for 60 min and 700 ◦C for 2 min consist of a microstructure with almost similar average grain size (~600–670 nm), while the samples annealed at 700 ◦C over 2 min and 800 ◦C for 60 min show that the average grain size is in the range of approximately 960 nm to 5 μm.

The KAM map is generally used to represent the strain distribution based on the dislocation density. The KAM maps in Figure 3 were estimated up to the third nearest neighbor kernel with a maximum misorientation of 5◦. The KAM map (Figure 3) shows that the strain distribution in samples annealed at 600 ◦C for 60 min and 700 ◦C for 2 min is relatively high as compared with samples at other annealing conditions. This clearly indicates that the strain induced into the material by HPT processing is not completely recovered in samples annealed at 600 ◦C for 60 min and 700 ◦C for 2 min. In general, deformed samples exhibit a larger KAM value (>1◦), while recrystallized samples exhibit a lower KAM value (<1◦) [28]. An almost similar average KAM value of approximately 0.83–0.84 is observed for samples annealed at 600 ◦C for 60 min and 700 ◦C for 2 min, while the average KAM value is in the range of 0.37◦–0.45◦ for samples annealed at 700 ◦C over 2 min and 800 ◦C for 60 min.

The misorientation chart (Figure 4) indicates that the fraction of low angle grain boundaries is higher in samples annealed at 600 ◦C for 60 min and 700 ◦C for 2 min, but decreases significantly with increasing annealing time and temperature. However, a high fraction of 60◦ misorientation angle can be seen in all the conditions, with the fraction increasing with temperature and time. The presence of a high fraction of low angle grain boundaries indicates that the samples are not fully recrystallized, and the high fraction of 60◦ misorientation angle indicates the presence of abundant annealing twins.

**Figure 3.** Kernel average misorientation (KAM) map of post-HPT annealing samples, and the corresponding color scale (right side).

Based on the IPF map, KAM map, and misorientation angle chart, the annealed samples can be divided into two categories: (a) microstructure with fine grain size, high KAM value, and high fraction of low angle grain boundaries (partially recrystallized); and (b) microstructure with relatively coarser grain size, low KAM value, and low fraction of low angle grain boundaries (nearly fully recrystallized). Accordingly, the sample annealed at 600 ◦C for 60 min, and 700 ◦C for 2 min will be referred to as samples with partially recrystallized microstructures. Similarly, the samples annealed at 700 ◦C for 15 min, 700 ◦C for 30 min, 700 ◦C for 60 min, and 800 ◦C for 60 min will be referred to as samples with recrystallized microstructures.

**Figure 4.** Misorientation angle chart of HPT processed CoCrNi alloy after annealing at different temperatures and times.

The mechanical properties of CoCrNi alloy at different processing conditions were assessed by Vickers hardness measurement and uniaxial tensile test. The hardness values under different processing conditions are shown in Figure 5a. The hardness of the HPT processed sample decreases with increasing annealing temperature and time. The annealed HPT samples with partially recrystallized microstructure show the smallest decrease in hardness as compared with the samples with near fully recrystallized microstructure. This can be attributed to the fine grain size and high KAM value in the samples with partially recrystallized microstructure.

**Figure 5.** Mechanical properties of CoCrNi alloy in different processing conditions. (**a**) Hardness and (**b**) engineering stress–strain curve.

The engineering stress–strain curves of CoCrNi alloy at different processing conditions are shown in Figure 5b. The tensile curves clearly indicate that the annealing of the HPT processed sample leads to a reasonably good combination of strength and ductility. The tensile strength of the annealed samples follows a similar trend to that of hardness value. The annealed samples with partially recrystallized microstructure show the lowest yield strength reduction compared with the samples with recrystallized microstructure. In steels, it has been reported that the strength increases with an increase in pre-strain [29]. As KAM indicates the strain distribution based on the dislocation density, the annealed samples with high KAM can be regarded as a material with pre-strain. Thus, high strength in the samples with partially recrystallized microstructure can be attributed to the fine grain size and high KAM value.

The post-HPT annealing has significantly improved the uniform elongation and strain to fracture at the cost of a reduction in strength for samples with recrystallized microstructure. It is interesting to note that only the strain to fracture has been improved significantly and not the uniform elongation for the samples with partially recrystallized microstructure. The uniform elongations in the samples annealed at 600 ◦C for 60 min and 700 ◦C for 2 min are almost similar to each other and are also similar to the HPT processed sample. However, the strain to fracture is significantly higher in the samples annealed at 600 ◦C for 60 min and 700 ◦C for 2 min than the HPT processed sample, with the sample annealed at 600 ◦C for 60 min showing higher elongation to fracture.

In high-grade pipeline steels, it has been reported that the samples with a high percentage of low angle grain boundaries resulted in the smaller difference between yield strength and tensile strength [30]. In the present study, the samples with a high fraction of low angle grain boundaries exhibit a smaller difference between the yield strength and tensile strength. In addition, it has been shown that defects such as dislocation and vacancies can pass through the low angle grain boundaries, unlike the pile-up of defects at high angle grain boundaries [31]. This could be a possible reason for reduced work hardenability, thereby resulting in reduced uniform elongation in materials with a high fraction of low-angle grain boundaries. In addition, low angle grain boundaries are considered to be resistant to intergranular fracture in metallic materials [32]. This could be a possible reason for reasonable post-necking deformation in samples with partially recrystallized microstructure. Thus, the presence of low angle grain boundaries resulted in an increase in strain to fracture without enhancement in uniform elongation.

In conventional alloys, enhancement of uniform elongation in post-deformed annealed samples is observed when the fraction of recrystallized grains exceeds a certain volume fraction. In Cu–Al alloys, a reasonable improvement in uniform elongation is observed when the recrystallized grains are over 80% of the volume fraction [33]. In nanostructured Cu processed by rolling at liquid nitrogen, an increase in uniform elongation is reported when the fraction of secondary recrystallized grains is about 25% [34]. In the present study, an increase in uniform elongation is observed when the fraction of low-angle grain boundaries and the average KAM value is lower than 5% and 0.5◦, respectively.

The shape of the flow curves (Figure 5) is distinct between the samples with partially recrystallized and recrystallized microstructures, clearly indicating the influence of microstructure on the flow curve. In samples with the partially recrystallized microstructure, the flow curve shows softening behavior. The rate of softening in the sample annealed at 600 ◦C for 60 min is different from that of the sample annealed at 700 ◦C for 2 min, with the one at 700 ◦C for 2 min showing pronounced softening. As the average KAM value and the average grain size are almost similar for 700 ◦C for 2 min and 600 ◦C for 60 min samples, the pronounced softening in the one at 700 ◦C for 2 min clearly indicates that the fraction of low angle grain boundaries significantly influences the softening behavior. Thus, the sample annealed at 700 ◦C for 2 min with a higher fraction of low-angle grain boundaries shows more pronounced softening than the 600 ◦C for 2 min sample. In the samples with recrystallized microstructure, the flow curves show strain hardening behavior, and the strain hardening ability is observed to increase with increasing annealing temperature and time.

In order to further understand the flow curve behavior, the work hardening rate (WHR) curves are plotted as a function of true strain for different processing conditions and are shown in Figure 6. In general, the WHR is not discussed with respect to plastic deformation after the plastic instability. However, the WHR after plastic instability is also shown in Figure 6a to understand the enhancement in strain to fracture without a significant enhancement of uniform elongation in samples with partial recrystallization. For the HPT sample, the WHR decreases rapidly and enters into the softening after yielding, and decreases continuously till fracture. This is commonly observed in severely deformed or nano-grained materials because of the limited dislocation activity and ineffectiveness in dislocation generation and accumulation during deformation. For the samples with a partially recrystallized microstructure, the WHR decreases initially and enters the softening region, and with further strain, the WHR curve increases and slowly approaches a peak, and then decreases gradually. The amount of increase in WHR from softening is higher for the sample annealed at 700 ◦C for 2 min than that at 600 ◦C for 60 min. Interestingly, the WHR curve of the 600 ◦C for 60 min sample increases after softening and reaches hardening region (i.e., from a negative value to positive value of WHR) and then

decreases again as the strain approaches the fracture strain. As previously mentioned, the difference softening behavior in the partially recrystallized microstructure samples can be attributed to the difference in the fraction of low angle grain boundaries. The increase in WHR curve in the softening zone can be a possible reason for the increase in strain to failure without an increase in uniform elongation in the annealed samples with partially recrystallized microstructure.

The WHR curves of samples with recrystallized microstructure are shown in Figure 6b. The WHR curves show three distinct stages of WHR. In the first stage (Stage A), the WHR decreases rapidly at the initial stage of plastic deformation (elastic to plastic), and is observed in all the annealed samples. In 700 ◦C for 60 min and 800 ◦C for 60 min samples, the second stage (Stage B) shows a gradual decrease in WHR followed by a slow decrease (Stage C) with the increase in true strain. In the 700 ◦C for 30 min sample, the second stage of WHR is very narrow, and is not significant as in 700 ◦C for 60 min and 800 ◦C for 60 min samples. In the 700 ◦C for 15 min annealed sample, Stage B shows an increase in WHR followed by the slow decrease in WHR (Stage C).

**Figure 6.** Work hardening rate as a function of true strain for samples under different processing conditions.

The difference in WHR stage, especially Stage B, of the annealed samples can be attributed to the difference in grain size and the amount of strain present in the sample before tensile testing [15,35]. The increase in WHR in Stage B has been reported for ultra-fine grained CoCrMnFeNi HEAs, and it is attributed to the yield drop phenomenon observed in ultra-fine grained materials [15]. In commercial pure Ti, it is demonstrated that the increase in WHR in Stage B is observed when the pre-strain reaches 3.5% and is attributed to the presence of high density of dislocations and deformation twinning in pre-strained samples. In the present study, the increase in WHR in Stage B of the 700 ◦C for 15 min sample could be attributed to the fine grain size (~960 nm) and the presence of strain (KAM value ~0.45◦) (Figure 3). As the grain size and KAM value of the 700 ◦C for 30 min samples fall in between the those of the 700 ◦C for 15 min sample and 700 ◦C for 60 min sample, Stage B is observed only for a narrow strain as there could be a balance between increase and gradual decrease in WHR.

Figure 7 shows the deformed microstructure (transmission Kikuchi diffraction inverse pole figure map) near the fracture surfaces of the samples with partially recrystallized microstructure (600 ◦C for 60 min) and recrystallized microstructure (700 ◦C for 30 min). As a result of the large strain developed in the material, the confidence index is low, and there are some unindexed regions in Figure 7. Both the microstructures show the presence of fine annealing twins. It is to be noted that the formation of deformation twins is difficult to form in ultra-fine grained materials as the critical twining stress increases with decreasing grain size [36,37]. Indeed, the formation of deformation twinning is not observed in ultra-fine grained CoCrMnFeNi HEA [15]. The presence of very fine twins in Figure 7 could be annealing twins, and the deformation twins may not have formed as a result of the fine grain size of the annealed samples in the present study. Besides, the WHR curve also did not show a plateau region indicating the formation of twins during deformation.

**Figure 7.** Transmission Kikuchi diffraction inverse pole figure map of post-deformed samples of 600 ◦C for 60 min and 700 ◦C for 30 min annealed samples.

Figure 8 shows the comparison of tensile properties of the present study with the conventional alloys and other high strength HEAs. (Figure 8 is adopted and modified from the literature [38]. The values for HEA wire with nano-twins are taken from the literature [39]). It is clear that the strengths of the CoCrNi in HPT processed state and after annealing conditions are quite high as compared with the other HEAs and other high strength conventional alloys.

**Figure 8.** Comparison of tensile properties of HPT processed and post-HPT annealed CoCrNi alloy with other high strength high entropy alloys (HEAs) and conventional alloys (This figure is adopted and modified from the literature [38]. The values for HEA wire with nano-twins are taken from the literature [39]).

The superior mechanical properties achieved in the present alloy can be primarily attributed to the presence of fine grains, the formation of fine annealing twins, high shear stress, and high friction stress. It is well known that the strength increases with the decrease in grain size based on the Hall–Petch relation [9,40]. The presence of very fine twins can enhance the mechanical properties by reducing the dislocation mean free path and resisting the dislocation slip, a mechanism very similar to the

Hall–Petch relation [33]. The CoCrNi alloy is reported to have very low stacking fault energy and high shear stress [19,41]. In materials with high shear stress and low stacking fault energy, the cross-slip during deformation becomes difficult and thereby enhances the mechanical properties [36,42]. It is reported that the friction stress of CoCrNi alloy is higher than that of many conventional FCC metals because of the fluctuation of Peierls potential for dislocation motion, and it can enhance the strength of the material [9]. Thus, the presence of fine grains, fine annealing twins, high shear stress, high friction stress, and low-angle grain boundaries resulted in superior mechanical properties in the post-HPT annealed CoCrNi samples.

#### **4. Summary**

In the present study, the effect of annealing on the microstructural and mechanical properties of HPT processed CoCrNi HEA was investigated. The results clearly indicate that the initial microstructure plays an important role in the mechanical properties of CoCrNi alloy. The initial microstructure with fine grain size of ~660 nm, high fraction of low angle grain boundaries, and high KAM value shows strain softening followed by strain hardening, whereas only strain hardening is observed for the microstructure with grain size >900 nm, a low fraction of low-angle grain boundaries, and low KAM value. The sample annealed at 700 ◦C for 15 min shows the optimized room temperature mechanical properties with a remarkable combination of strength (1090 MPa) and strain to failure (41%). Thus, the present study demonstrates that the severe plastic deformation followed by annealing is one of the effective ways of enhancing the mechanical properties in CoCrNi MEA.

**Author Contributions:** Conceptualization, P.S.M.; Data curation, P.S.M.; Formal analysis, P.S.M. and J.G.K.; Investigation, P.S.M., J.W.B., P.A.-R., and J.M.P.; Methodology, P.S.M., J.W.B., P.A.-R., and J.M.P.; Supervision, H.S.K.; Writing—original draft, P.S.M.; Writing—review & editing, P.S.M., J.G.K., and H.S.K.

**Funding:** This research was supported by Future Materials Discovery Project through the National Research Foundation of Korea (NRF) funded by the Ministry of Science and ICT (NRF-2016M3D1A1023383). Praveen Sathiyamoorthi is supported by Korea Research Fellowship program through the National Research Foundation of Korea (NRF) funded by the Ministry of Science and ICT (2017H1D3A1A01013666).

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**


© 2018 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

### *Review* **High-Pressure Induced Phase Transitions in High-Entropy Alloys: A Review**

#### **Fei Zhang 1,2, Hongbo Lou 1, Benyuan Cheng 1,3, Zhidan Zeng <sup>1</sup> and Qiaoshi Zeng 1,4,\***


Received: 24 January 2019; Accepted: 26 February 2019; Published: 2 March 2019

**Abstract:** High-entropy alloys (HEAs) as a new class of alloy have been at the cutting edge of advanced metallic materials research in the last decade. With unique chemical and topological structures at the atomic level, HEAs own a combination of extraordinary properties and show potential in widespread applications. However, their phase stability/transition, which is of great scientific and technical importance for materials, has been mainly explored by varying temperature. Recently, pressure as another fundamental and powerful parameter has been introduced to the experimental study of HEAs. Many interesting reversible/irreversible phase transitions that were not expected or otherwise invisible before have been observed by applying high pressure. These recent findings bring new insight into the stability of HEAs, deepens our understanding of HEAs, and open up new avenues towards developing new HEAs. In this paper, we review recent results in various HEAs obtained using in situ static high-pressure synchrotron radiation x-ray techniques and provide some perspectives for future research.

**Keywords:** high pressure; polymorphic transition; high-entropy alloy

#### **1. Introduction**

Developing multicomponent metallic alloys with superior properties has played a vital role in the advancement of human civilizations since the Bronze Age. Conventional metallic alloys are usually based on one or two principle elements, such as Fe-, Al-, Mg-, and TiAl- based alloys. Adding alloying elements into the host lattice of the principle elements forming solid solutions has been the major strategy to optimize the microstructures and properties of alloys. However, in the traditional metallurgy, the development of alloys was restricted by the limited solubility of the alloying element in the solvent lattice, which leaves the areas besides the corner of their multicomponent phase diagrams unexplored. Deviation from the phase diagram corner was believed to readily result in the formation of useless, brittle intermetallic compounds. In 2004, Yeh et al. and Cantor et al.'s discoveries of single solid-solution phase alloys formed with multi-principal elements challenged the traditional metallurgy experience and established an exciting new concept for alloy design. By mixing five or more elements with equimolar or near-equimolar ratios, the system could be stabilized in a single phase solid-solution by their maximized configurational entropy [1,2]. Since then, this new class of so-called high-entropy alloys (HEAs) has attracted considerable attention and research interests in the advanced metallic materials community [3–5]. Numerous new HEAs have been developed

with simple crystal structures such as face-centered cubic (*fcc*) [1,2], body-centered cubic (*bcc*) [1,6,7], and hexagonal close-packing (*hcp*) [8–12]. With unique compositions and atomic structures, HEAs show many interesting properties for potential applications, such as high ductility and strength over a wide temperature range and excellent resistance to both corrosion and wear [13–21]. On the other hand, due to their complex composition with a high chemical disorder, many fundamental questions of the HEAs remains challenging to address [3,22,23].

One of the critical unsolved questions about HEAs is their phase stability. Many pure elements in the periodic table show rich polymorphic phase transitions [24]. By mapping out the phase diagram (transition paths) with varying temperature and pressure, the phase stability of each structure can be clarified. A famous example is iron. Phase transitions between three different prototype polymorphs with *fcc*, *hcp*, and *bcc* structures were extensively studied in iron [25–27]. The *bcc* phase of iron was confirmed to be stable at ambient conditions, while the *fcc* phase is stable at high temperatures and the *hcp* phase is favorable at high pressures. As a combination of multiple elements, HEAs do not simply inherit the structure and properties of their constituent elements as expected with a "linear effect" [28]. Regardless of the various compositions and structures, HEAs are reported to be surprisingly stable over a large temperature range in previous experiments. It is believed that HEAs are thermodynamically stabilized by their high configurational entropy which can largely lower down the Gibbs free energy. Also, the high chemical complexity and packing disorder cause severe local lattice distortion and extremely sluggish atomic diffusion in HEAs, which could further stabilize HEAs kinetically. For instance, the equiatomic CoCrFeMnNi alloy (also named as Cantor's alloy) [2], is a prototype *fcc*-structured HEA. Extensive studies demonstrate that the Cantor's alloy can maintain its *fcc* structure from cryogenic temperatures up to its melting temperature without any polymorphic phase transition [13,29–31].

It has been empirically established that the competition between configuration entropy and enthalpy, the difference between the atomic radius and electronegativity of constituent elements [32–34], and also the overall valence electron concentration [35] are a few key thermodynamic parameters for the formation of HEAs. All these parameters are very susceptible to pressure tuning. Actually, pressure is a very powerful tool to tune the atomic/electronic structure of various materials and has been employed to understand materials and to search for novel materials through rich pressure-induced phase transitions in diverse systems, such as pure elements [24], alloys [36–42], oxides [43–46], and metallic glass [47–49]. Among them, the Ce3Al system is of particular interest, where a pressure-induced intermetallic compound and metallic glass to *fcc* solid solution transitions were discovered due to the significant reduction of the difference between both the atomic radii and electronegativity of Ce and Al during compression [42,49]. For HEAs, their synthesis process is closely associated with the competition between intermetallic compounds, metallic glasses, and simple crystalline solid solutions. Therefore, in contrast to the seeming "ultra-stability" during heating or cooling, HEAs might exhibit rich tunable behavior under high pressure.

Very recently, the structural stability of various HEA systems has been explored using in situ high-pressure synchrotron radiation-based x-ray diffraction techniques. Many interesting polymorphic transitions have been discovered. These results are summarized in Table 1 and are subsequently reviewed in detail. The microstructural and compositional metastability of HEAs was nicely reviewed by Wei et al. [50]. In this paper, we focus on very recent results about phase stability and transitions under high pressure and provide a brief review of the relevant experimental methods, issues, and perspectives for future study.


**Table 1.** A summary of the pressure-induced polymorphic transitions in high-entropy alloys (HEAs) investigated by in situ high-pressure XRD (ME: methanol:ethanol = 4:1 (volume ratio) mixture; MEW: methanol:ethanol:water = 16:3:1 (volume ratio) mixture).

#### **2. Experimental Methods**

Diamond anvil cells (DACs) are the most commonly used device to generate high pressure for the in situ studies of materials. DACs are versatile devices for generating pressures up to hundreds of GPa and for combining a full range of in situ measurements [66]. DACs are composed of two opposing diamond anvils that squeeze materials in between them to generate hydrostatic/non-hydrostatic pressure. Since diamonds are transparent to almost the entire electromagnetic spectrum, various in situ electromagnetic radiation detection approaches can be employed to study the structure and properties of samples inside DACs, such as in situ x-ray and neutron diffraction techniques, x-ray emission/absorption spectroscopy, Raman spectroscopy, and Brillouin scattering. Among them, in situ angular dispersive x-ray diffraction (XRD) based on an intense synchrotron radiation x-ray source is the primary technique in the high-pressure structural study of various materials [67].

The size of the DAC anvil culets is typically small with a diameter of approx. 500 μm to 20 μm depending on the target maximum pressure. The sample chamber is a small hole (with a diameter of approx. 1/3 of the culet size and a height of approx. 50 μm to 30 μm) drilled in the gasket indent which is pre-indented by the two anvils. Therefore, only tiny samples with a typical maximum dimension less than tens of microns can be accommodated in the sample chamber in a DAC. Particularly, the sample thickness should be less than the height of the gasket hole during the entire compression process to avoid bridging the anvils (otherwise, it will be uniaxially compressed with large shear stress). Metals with high shear strength such as T301 stainless steel, Re, and W are usually used as the gasket materials for typical high-pressure experiments. Different pressure transmitting mediums (including solid, liquid, and gas) can be selected based on the target pressure range and the required degree of hydrostaticity or operational convenience in the experiment. The pressure in the sample chamber can be determined either by the pressure–volume (*P–V*) equation of state (EOS) of standard materials (e.g., MgO, NaCl, Au, and Pt) using in situ high-pressure XRD [68] or by the ruby fluorescence peak shift excited by an optical laser or x-ray [69]. The standard material or ruby balls should be loaded close to the sample in the sample chamber to minimize the pressure difference. Due to the small sample volume, a high-brightness synchrotron radiation X-ray beam focused down to tens of microns (<20 μm) is required. To go through the two thick diamond anvils (3–5 mm in total) and to cover a large enough range in *d* space with the limited two-theta opening of DACs, a high-energy x-ray is usually required (>20 keV). Using in situ high-pressure synchrotron radiation XRD diffraction coupled with a DAC (a typical experimental setup is shown in Figure 1), researchers have been able to obtain detailed structural information of samples as a function of pressure, including their unit cell parameters, atomic positions, thermal parameters, and even electron density distributions.

**Figure 1.** A schematic illustration of the in situ high-pressure synchrotron radiation XRD setup using a diamond anvil cell (DAC).

#### **3. Structural Stability and Evolution of HEAs under High Pressure**

Over the last fourteen years, intense effort has been devoted to developing numerous HEAs, which provides hundreds of new alloys for fundamental and applied studies. Recently, the structural stability of some typical HEAs has been investigated using in situ high-pressure synchrotron radiation XRD. Herein, we briefly review these results in separate groups according to their initial crystal structures.

#### *3.1. Fcc-Structured HEAs*

Among the various HEAs, single-phase *fcc*-structured alloy systems tend to show a relatively low yield strength but an excellent ductility and strain hardening capability. The CoCrFeMnNi (Cantor's alloy) as a prototype *fcc*-structured HEA has attracted the most extensive investigation. Cantor's alloy shows high structural stability over a broad temperature range at ambient pressures (from extreme low temperature (~3 K) up to its melting temperature) [13,29–31]. In contrast, Cantor's alloy undergoes unexpected polymorphic transformations from *fcc* to *hcp* phases under applied high pressure [51–54]. Under the best hydrostatic pressure conditions provided by the pressure-medium helium, the *fcc* to *hcp* phase transition starts at approx. 22 GPa but does not fully complete even up to approx. 41 GPa (Figure 2a) [52]. The transition is sluggish and irreversible. The *hcp* phase can be retained during the decompression down to ambient pressure, as shown in Figure 2b. The fabricated *hcp* CoCrFeMnNi HEA can almost maintain its volume fraction during decompression. Therefore, *hcp*-*fcc* dual-phase composites with tunable volume fractions can be readily synthesized by decompression from different maximum pressures between approx. 22 and 41 GPa (Figure 2a) [52].

**Figure 2.** (**a**) The in situ high-pressure XRD patterns of the CoCrFeMnNi HEA under high pressure in a DAC at room temperature [52]: The x-ray wavelength is 0.2952 Å and (**b**) the change of the *hcp* phase volume fraction as a function of pressure during compression (solid symbols) and decompression (open symbols). The volume fractions of the *hcp* phase were calculated based on the peak area changes of the *fcc*-(200) (blue circles) and *hcp*-(101) peaks (red squares) in Figure 2a, which yield consistent results of the volume fractions. During decompression, the volume fraction of the *hcp* almost remains constant, which makes the synthesis of the *hcp*-*fcc* dual-phase composite possible by following a different decompression path, as shown by the dashed arrows [52].

According to previous theoretical simulations, the Gibbs free energy of the *hcp* phase of the CoCrFeMnNi HEA may be smaller than its well-known *fcc* phase at room temperature [70,71]. However, there was no clear experimental evidence to confirm the simulation results. For example, the CoCrFeMnNi HEA samples synthesized by various melt-quenching methods always only show an *fcc* structure, and the irreversibility of the *fcc* to *hcp* phase transition under high pressure also questions the relative stability of the *fcc* and *hcp* phases [51,53,54]. To clarify the phase stability, the synthesized *hcp* phase was further examined using in situ high-temperature XRD measurements at different pressures [52]. During heating at constant pressures, the *hcp* phase transforms back to the *fcc* phase and the critical transition temperatures increase with increasing pressure. Therefore, these results demonstrate that the well-known *fcc* phase of the CoCrFeMnNi HEA is thermodynamically favorable at high temperatures. In contrast, the *hcp* phase is indeed more stable at relatively lower temperatures and higher pressures [52].

The pressure-induced *fcc* to *hcp* polymorphic transition has been observed in Cantor's alloy by independent research groups. However, quite different onset pressures were reported ranging from approx. 7 GPa to even above 49 GPa [51,53–55]. Tracy et al. loaded a CoCrFeMnNi HEA sample (annealed at 1200 ◦C for 24 h) in a DAC with silicone oil as the pressure medium and compressed it up to 54.1 GPa. A sluggish martensitic transformation from the *fcc* to an *hcp* phase was also observed starting at approx. 14 GPa [51]; Huang et al. reported that the *fcc* to *hcp* phase transition occurred at approx. 7.1 GPa in a CoCrFeMnNi HEA sample. The initial sample was processed by a series of heating, cold rolling, and milling and then loaded into a DAC with neon as the pressure medium [54]; Yu et al. prepared a nanograined (approx. 100 nm) CoCrFeMnNi HEA sample by mechanical alloying and a high-pressure sintering process. They compressed the sample with silicone oil as the pressure medium in a DAC up to 31GPa, but no phase transition was observed [56]. Ahmad et al. investigated the structure of Cantor's alloy up to approx. 49 GPa with neon as the pressure-transmitting medium in a DAC. Surprisingly, no obvious phase transition was observed as well [55]. This experimental inconsistency suggests that the polymorphic phase transition in Cantor's alloy may be susceptible to the sample and experimental conditions, such as the different sample grain sizes and different pressure mediums used in the experiments above.

To clarify this speculation, Zhang et al. [52] systematically investigated the effect of the non-hydrostaticity of the pressure environment and the grain size of the samples on their pressure-induced phase transitions. The experiments were carefully designed to study only one factor at each time. To address the effect of the pressure environment, they loaded the same sample with three distinct pressure mediums with different degrees of hydrostaticity, such as helium (the most hydrostatic), amorphous boron (the most non-hydrostatic), and silicone oil (quasi-hydrostatic in-between). According to the in situ high-pressure XRD results (as shown in Figure 3), the onset pressures for the *fcc* to *hcp* transition were estimated to be approx. 22 GPa in helium, approx. 2 to 6 GPa in amorphous boron, and approx. 7 GPa in silicone oil. These results demonstrate that the degree of the pressure medium's hydrostaticity has a positive effect on the onset pressure of the *fcc*-to-*hcp* phase transition in Cantor's alloy [53].

**Figure 3.** The in situ high-pressure XRD patterns of the CoCrFeMnNi HEA sample with helium (**a**), silicone oil (**b**), and amorphous boron (**c**) as the pressure mediums [53].

To study the effect of grain size (an important internal factor) on the pressure-induced phase transition in the CoCrFeMnNi HEA, Zhang et al. [53] loaded two distinct samples into one sample chamber in a symmetric DAC. The two samples were carefully located with equivalent positions to the chamber center to ensure they had identical pressure environments. To highlight the grain size effect, the two selected samples had a huge difference in grain size; one was synthesized by gas-atomization (GA) (approx. 5 μm), and the other was obtained by high-pressure torsion (HPT) (approx. 10 nm). When the two samples were compressed from 0.3 GPa up to 31.4 GPa, both of them showed an *fcc*-to-*hcp* phase transition but with quite different onset pressures (as shown in Figure 4). For the HPT sample with nano-sized grains, the phase transition was observed from approx. 12.3 GPa,

while the GA sample with a bigger grain size had a much lower onset pressure of approx. 6.9 GPa. The underlying mechanism is still not clear and calls for further investigation [53].

**Figure 4.** The in situ high-pressure XRD patterns of the gas-atomization (GA) and high-pressure torsion (HPT) CoCrFeMnNi HEA samples loaded in the same DAC with silicone oil as their pressure medium [53].

The effect of different alloying elements on the pressure-induced polymorphic phase transition was investigated by Zhang et al. [57] in three *fcc*-structured medium-entropy alloys and HEAs (NiCoCr, NiCoCrFe, and NiCoCrFePd). They observed a similar martensitic phase transition from *fcc* to *hcp* in the CoCrFeNi alloy starting at approx. 13.5 GPa, as shown in Figure 5a. This phase transformation was also sluggish and irreversible, as reported in Cantor's alloy. The volume fraction of the *hcp* phase was only about 36% when the pressure reached 39 GPa. However, with different alloying elements, the NiCoCr and NiCoCrFePd alloys exhibited distinct compression behaviors under high pressure, as shown in Figure 5b,c, respectively. Only a small amount (<5 wt.%) of the *hcp* phase emerged at 34.4 GPa in the NiCoCr alloy, and its amount barely changes with increasing pressure. In the NiCoCrFePd system, as the element Mn in the Cantor's alloy is replaced by the Pd, no obvious *hcp* phase emerges up to 74 GPa [57].

**Figure 5.** The XRD profiles of CoCrFeNi (**a**), CoCrNi (**b**), and CoCrFeNiPd (**c**) at high pressures: The diffraction peaks marked by symbol \* are from the Au pressure standard [57].

Many of the HEAs have a minor second phase. Ma et al. [58] investigated an equiatomic CoCrCuFeNiPr HEA sample with dual phases (major disordered-*fcc* and minor ordered-*fcc* phases) using in situ synchrotron radiation high-pressure energy-dispersive XRD (EDXRD). They observed a pressure-induced fast ordering transition from approx. 8 GPa to 16.0 GPa followed by a slow transition up to 106.4 GPa. The initially ordered domain in the CoCrCuFeNiPr HEA was believed to act as embryos. With increasing pressure, the embryos grow into the ordered phase [58].

#### *3.2. Bcc-Structured HEAs*

In addition to the *fcc*-structured HEAs, *bcc*-structured alloys are another major member of the HEA family, which include both the chemically disordered A2 and ordered B2 phases. *Bcc*-structured HEAs often exhibit high yield strength in a very high-temperature regime. Therefore, the stability of the *bcc*-structured HEAs is an exciting topic that has been explored extensively at various conditions, recently also by high pressure.

Ahmad et al. investigated the structural stability of a TiZrHfNb alloy with a disordered *bcc* structure during compression up to 50.8 GPa; no phase transition was found [55]. Yusenko et al. explored another *bcc*-structured Al2CoCrFeNi HEA; it had no phase transition up to 60 GPa as well [59]. Guo et al. studied the structural evolution of a superconducting (TaNb)0.67(HfZrTi)0.33 HEA during compression up to approx. 100 GPa; its *bcc* structure seemed very robust without any detectable structural transition [63]. No phase transition has ever observed in the *bcc*-structured HEAs. Therefore, it seems that the *bcc*-structured HEAs are incredibly stable, and much higher pressure may be needed to induce phase transitions (compared to the *fcc* family).

To lower down the transition pressure of possible polymorphic phase transitions, Cheng et al. [61] employed a creative strategy to focus on relatively less stable compositions. They chose an equiatomic AlCoCrFeNi HEA and monitored its structural evolution during compression up to 42 GPa. The AlCoCrFeNi alloy had an ordered *bcc*-structure (B2 phase) and was reported to sit in the transition zone between the *fcc* and *bcc* phases, as *x* varies in the Al*x*CoCrFeNi HEA system (0 < *x* < 2) [72]. Indeed, they discovered a phase transition from the initial B2 phase to a highly distorted form starting at relatively low pressure of approx. 17.6 GPa, by combining ex situ high-resolution transmission electron microscope (HRTEM) with in situ high-pressure synchrotron radiation XRD data, as shown in Figure 6. Besides the XRD peak splitting, severe peak weakening and broadening occurred during compression, which may have been caused by the significant lattice distortion developed in the sample. Therefore, their work was unable to resolve the atomic structure of the high-pressure phase. Nevertheless, it is the first time that a pressure-induced polymorphism was suggested in a *bcc*-structured HEA [61].

**Figure 6.** The structural evolution of the AlCoCrFeNi HEA as a function of pressure monitored by in situ high-pressure XRD patterns at room temperature (**a**) and the locally enlarged plot of the XRD patterns for each peak upon compression (**b**) and decompression (**c**) to show more details of the peak shape and width: The x-ray wavelength is 0.4959 Å [61].

With slightly lower Al content but still located in the *bcc*-*fcc* transition zone, the Al0.6CoCrFeNi HEA was studied by Wang et al. using in situ synchrotron radiation XRD in a DAC with both silicone oil and helium as the pressure-transmitting medium up to approx. 40 GPa [62]. The Al0.6CoCrFeNi HEA powders were prepared by the GA method. A single *bcc* phase was obtained with a high quenching rate in the GA process. They revealed a *bcc*-to-orthorhombic phase transition, which started at approx. 10.6 GPa and completed at approx. 21.4 GPa. Interestingly, another body-center-tetragonal (*bct*) phase emerged and coexisted with the high-pressure synthesized orthorhombic phase when the pressure was released. These results indicate that the orthorhombic phase may be metastable at ambient conditions but could be partially maintained due to the possible large energy barrier [62]. Moreover, after annealing at 1000 ◦C for 2 h, the initial GA *bcc* Al0.6CoCrFeNi HEA can transform into a more stable *fcc* phase. During compression, the *fcc* phase of the Al0.6CoCrFeNi HEA could completely transform into an *hcp* phase similar to Cantor's alloy. The samples recovered from high-pressure compression were characterized by transmission electron microscopy (TEM) and further confirmed that all of the five polymorphs could stably/metastably exist at ambient conditions (transition path between them is shown in Figure 7). Severe lattice distortion, which is tunable by high pressure or temperature was suggested to play a crucial role in the formation of various polymorphs and the transition between them in the Al0.6CoCrFeNi HEA. These findings suggest that HEAs could behave quite differently from the expectation of a linear combination of its constituent element; they may also exhibit structural flexibility/tunability far beyond that of their solution components [62].

**Figure 7.** A schematic illustration of the atomic structure for five polymorphs observed in the Al0.6CoCrFeNi HEA and the transition paths between them [62].

#### *3.3. Hcp-Structured HEAs*

HEAs commonly form with *fcc* or *bcc* structures. Recently, *hcp*-structured HEAs were observed in high-pressure experiments via an *fcc* to *hcp* polymorphic transition and was obtained in the melt-quenched alloys mainly consisting of heavy *hcp* metals, e.g., the CoOsReRu, CoFeReRu and CoReRuV, Ir0.19Os0.22Re0.21Rh0.20Ru0.19, or rare earth *hcp* elements [8–12]. Ahmad et al. [55] investigated the structural stability of the quarternary equiatomic ReRuCoFe alloy under high pressure. They compressed the sample from 0.9 GPa to 80.4 GPa in a DAC. The unit cell parameters of *a* and *c*, its ratio *a*/*c*, and the sample volume all continuously decreased with increasing pressure, which indicates the *hcp* structure of the ReRuCoFe alloy is stable under compression up to approx. 80 GPa. Yusenko et al. investigated the structural stability of the *hcp* Ir0.19Os0.22Re0.21Rh0.20Ru0.19 HEA at room temperature during compression up to 45 GPa but also observed no phase transition [64].

For the *hcp*-structured rare earth HEAs, Yu et al. reported a series of pressure-induced phase transitions in the HoDyYGdTb HEA by in situ XRD measurements in a DAC using synchrotron radiation x-ray. Four polymorphs were observed following a transition sequence of *hcp*→Sm-type→d*hcp*→d*fcc* during compression up to 60.1 GPa (Figure 8), which resembles the rich pressure-induced polymorphic transitions in its constituent elements [65]. The Sm-type phase firstly appeared when the pressure reached 4.4 GPa. At 13.6 GPa, the *hcp* (102) diffraction peak disappeared, indicating the completion of the *hcp* to Sm-type phase transition. With further increasing pressure to

26.7 GPa, the d*hcp* phase emerged and persisted to 38.3 GPa, and then, the d*hcp*-d*fcc* phase transition occurred at 40.2 GPa.

**Figure 8.** The in situ high-pressure XRD patterns of the HoDyYGdTb HEA during compression [65].

#### **4. Conclusions and Outlooks**

HEAs are the focus of advanced metallic alloy research and have been attracting more and more attention over the last decade. Recent studies on HEAs under high pressure have added another dimension to the exploration of HEAs. The exciting findings in HEAs during compression under high pressure deepen our understanding of HEAs, providing a new avenue towards new HEAs development and helpful guidance for applications at extreme conditions. As a new research direction, growing interest in the structure and properties of HEAs under high pressure is expected to continue. Future outlooks are briefly summarized below:

(1) Synergic effect of pressure–composition–temperature. HEAs open up an almost infinite composition space for alloy design. Hundreds of different HEAs have been developed, but so far, only a few of them have been studied under high pressure. Inspired by the existing high-pressure work, more exciting novel phenomena and new structures are expected with broader exploration in more HEAs. Meanwhile, the elusive composition effect on the phase transitions of HEAs remains to be addressed. Besides, combining high pressure with temperature (from cryogenic temperatures up to the melting temperatures) can further clarify the stability of various HEAs and is worth more effort in future research.

(2) Combing multiple high-pressure techniques for better understanding of HEAs. Over the last few years, the major high-pressure research of HEAs has focused on the crystal structure evolution during compression and decompression. However, we still lack an in-depth understanding of the transformation mechanism. The atomic size ratio, electronegativity difference, valence electron concentration, magnetic states, etc., which are all critical for the formation and transformation of HEAs, have not been systematically explored under high pressure yet. To get more detailed information of the atomic and electronic structure of the multicomponent HEAs under high pressure, besides XRD measurements, more experiments combining other powerful in situ element-sensitive techniques are required, such as in situ high-pressure extended x-ray absorption fine structure (EXAFS), in situ high-pressure x-ray emission spectroscopy (XES), and in situ high-pressure x-ray magnetic circular dichroism (XMCD).

(3) Involving more variables for high-pressure studies of HEAs. Existing studies have shown that the phase transitions of HEAs are sensitive to shear stress. Therefore, high-pressure torsion (HPT) which generates extreme shear deformation under high pressure could be another powerful technique for HEAs structure tuning with well-controlled shear stress and deformation. In addition, the results reviewed in this paper focus on the static compression of HEAs using DACs. The strain rate effect on the high-pressure behaviors of HEAs has not been extensively investigated. The dynamic compression of HEAs with another dimension of an extremely high strain rate is also worth more exploration.

(4) Properties studies of HEAs using large-volume press. With unique compositions and disordered atomic structures, HEAs show many unusual properties. Under high-pressure compression, HEAs with new structure could be synthesized. Meanwhile, the grain size and defects could also be considerably changed, which could affect their properties as well. Guo et al. measured the superconducting behavior of the (TaNb)0.67(HfZrTi)0.33 HEA under high pressure. They surprisingly observed extraordinarily robust superconductivity even up to 190.6 GPa [63]. Although the properties of HEAs under high pressure may be interesting, besides the equation of states (EOS) which can be readily measured by in situ XRD in DACs, the vast other properties have not been well-studied. One critical issue of the DAC samples is the requisite tiny sample size. Fortunately, since the critical pressures reported for the polymorphic transitions in HEAs are mostly around 20 GPa or below, a large-volume press (LVP) with approx. 1000 times larger sample volumes than DACs under similar pressure conditions (typically <25 GPa) [73] could be used to synthesize millimeter or centimeter-sized HEAs readily for various properties characterization. Very recently, Yu et al. used a 10-MN double-stage LVP and compressed Cantor's alloy with a diameter of 1.5 mm and a height of 2 mm to 20 GPa [74]. They synthesized a bulk equiatomic CoCrFeMnNi HEA containing a mixture of *fcc* and *hcp* phases for property characterization. Cantor's alloy recovered from high-pressure treatment (20 GPa) showed a doubled hardness of the as-cast *fcc* samples because of enhanced dislocations, twins, stacking faults, and the *hcp* laths [74].

(5) Theoretical calculations. Recent progress in the experimental discovery of the polymorphic phase transitions in HEAs was first inspired by the finite-temperature ab initio calculation work done by Ma et al., which predicted that the hcp phase in certain magnetic states would be more stable than the *fcc* phase of Cantor's alloy at room temperature [70]. HEAs with complex compositions are a challenge for theoretical simulations, but it is quite encouraging that much exciting work has been successfully done on HEAs [23,75]. In the high-pressure community, many calculation methods have also been successfully established to handle materials under high pressure [76–78]. Theoretical simulations definitely will continue to play a vital part in predicting new phenomena and in interpreting elusive experimental results of HEAs. So far, there is still limited computational calculation works on the high-pressure behaviors of HEAs. However, we believe more exciting works can be expected in HEAs under high pressure by combing advanced experimental tools with simulation methods closely.

The concept of high entropy has been extended into many material systems including high-entropy nitrides, carbides, oxides, and metallic glasses. Therefore, the proposed research above is suitable for other new high-entropy materials as well.

**Funding:** This work was funded by the National Thousand Youth Talents Program in China and the National Natural Science Foundation of China, grant numbers 51871054 and U1530402, and the Fundamental Research Funds for the Central Universities.

**Acknowledgments:** The authors would like to thank Peter K. Liaw for the useful discussions and suggestions and Freyja O'Toole for editing the manuscript.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**


© 2019 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

### *Review* **Welding of High Entropy Alloys—A Review**

#### **Jing Guo 1,\*, Cong Tang 1, Glynn Rothwell 1, Lisa Li 1, Yun-Che Wang 2, Qingxiang Yang <sup>3</sup> and Xuejun Ren 1,\***


Received: 22 February 2019; Accepted: 15 April 2019; Published: 24 April 2019

**Abstract:** High-entropy alloy (HEA) offers great flexibility in materials design with 3–5 principal elements and a range of unique advantages such as good microstructure stability, mechanical strength over a broad range of temperatures and corrosion resistance, etc. Welding of high entropy alloy, as a key joining method, is an important emerging area with significant potential impact to future application-oriented research and technological developments in HEAs. The selection of feasible welding processes with optimized parameters is essential to enhance the applications of HEAs. However, the structure of the welded joints varies with material systems, welding methods and parameters. A systemic understanding of the structures and properties of the weldment is directly relevant to the application of HEAs as well as managing the effect of welding on situations such as corrosion that are known to be a service life limiting factor of welded structures in conditions such as marine environments. In this paper, key recent work on welding of HEAs is reviewed in detail focusing on the research of main HEA systems when applying different welding techniques. The experimental details including sample preparation, sample size (thickness) and welding conditions reflecting energy input are summarized and key issues are highlighted. The microstructures and properties of different welding zones, in particular the fusion zone (FZ) and the heat affected zones (HAZ), formed with different welding methods are compared and presented in details and the structure-property relationships are discussed. The work shows that the weldability of HEAs varies with the HEA composition groups and the welding method employed. Arc and laser welding of AlCoCrFeNi HEAs results in lower hardness in the FZ and HAZ and reduced overall strength. Friction stir welding results in higher hardness in the FZ and achieves comparable/higher strength of the welded joints in tensile tests. The welded HEAs are capable of maintaining a reasonable proportion of the ductility. The key structure changes including element distribution, the volume fraction of face centered cubic (FCC) and body centered cubic (BCC) phase as well as reported changes in the lattice constants are summarized and analyzed. Detailed mechanisms governing the mechanical properties including the grain size-property/hardness relationship in the form of Hall–Petch (H–P) effect for both bulk and welded structure of HEAs are compared. Finally, future challenges and main areas to research are highlighted.

**Keywords:** high-entropy alloys; welding; Hall–Petch (H–P) effect; lattice constants

#### **1. Introduction**

High-entropy alloys (HEAs) are based on the promising alloy design ideas of configurational entropy maximization, which offer great flexibility in developing different material systems [1–6]. Compared with the classic alloys, HEAs possess better performance in areas such as hardness, wear resistance, fatigue and resistance to corrosion and oxidation [3–6] as well as novel physical properties [2,3]. In general, HEA has high intrinsic strength and ductility, and the low diffusion rate of HEAs at high and low temperatures makes their applications in harsh condition possible [1,7,8]. There are many different material systems and processing methods developed in the past decades, which offer great flexibility in materials and processing methods selection for different conditions with a suitable welding process. In the meantime, the complexity of the material systems requires development of detailed understanding of the material behavior in different manufacturing processes including both hot and cold working operations. Joining of HEAs in similar or dissimilar material systems is increasingly important for expanding the applications of HEAs [9], among which welding process of HEAs is challenging due to the complex chemical, physical, and mechanical nature of welding and the need for tailoring the properties and structures of the weldment for complex loading conditions in similar and dissimilar welding.

HEA can be produced by many different routes such as liquid processing approaches (e.g., arc melting, Bridgman solidification, atomization and laser cladding), additive manufacturing, mechanical alloying (e.g., powder metallurgy), mixing elements of the vapor state including sputter deposition, atomic layer deposition and vapor phase deposition [1,10–18]. Welding is a complex process due to the fact that many conditions or factors affect the welding process. The quality of the welded joints can be assessed by the microstructure of the welding zones formed and mechanical structural integrity of the weldment as well as corrosion and fatigue [12]. There are many different welding processes such as arc welding, laser welding, friction stir welding, electron beam welding, electrical resistance spot welding as well as other processes such as electroslag welding, vacuum diffusion welding, etc. [9,12–16]. These welding processes differ significantly in the welding mechanism, working principles of the equipment and efficiency, but each has been found to have wide applications in different industries and application conditions. The energy input and the maximum temperature could be many orders of magnitude different from each of the welding methods, which directly influence the elemental and material behavior such as evaporation, melting and cooling. For example, the heat density ranges from 105–106 W/cm<sup>2</sup> for gas metal arc welding or plasma arc welding to 107–108 W/cm2 for laser or electron beam welding [15,16]. The structural zones formed (Figure 1) after different welding processes could be significantly different in geometry, size, microstructure and compositions, which is dependent on many factors such as the behavior of material systems (elements, evaporation, melting, flow, phase changes, cooling, etc.), the original structures of the workpiece, size of the specimen and the welding parameters, etc. The structure and properties of the welded joints (such as size/shape of the welding zones, the hardness distribution, residual stress, defects, etc.) also affects the performance of the overall joints under different modes of loading and service conditions [17]. These directly influence the suitability of each welding process on joining different HEAs, the application of HEAs and future developments in related areas [1,5,6,18].

In this paper key recent work on welding of HEAs is systematically reviewed. The main works on welding of different HEAs with different welding techniques are presented. The experimental details including samples preparation, sample size (thickness) and welding conditions reflecting the speeds and energy input are summarized. The different welding zones formed with different welding methods and the microstructures are compared. The hardness change of the fusion zone (FZ) and heat affected zones (HAZ) in different welding processes is compared and linked to the material systems. The mechanism of microstructure and property changes in the FZ and HAZ and their influence on the mechanical properties is discussed. Data on the Hall–Petch (H–P) effect (grain size-property/hardness relationship) for both bulk and welded structure of HEAs is compared. The future challenges and main areas to research are highlighted on the weldability of HEAs, welding process selection and potential effects of the microstructure and properties in different welding zones of HEAs on their applications.

**Figure 1.** Welded zones formed by different welding methods. [http://is.gliwice.pl/en/strona-cms/ electron-beam-welding-laboratory].

#### **2. Welding of HEAs in Di**ff**erent Welding Processes**

#### *2.1. Arc Welding of HEAs*

Arc welding is one of the most common welding methods, in which the energy from an arc of electric current between the material and a consumable/non-consumable electrode stick (either) is employed to melt the workpiece(s) with/without a protective atmosphere. Typical arc welding includes shielded metal arc welding (SMAW), gas tungsten arc welding (GTAW) (also known as TIG (Tungsten inert gas)), gas metal arc welding (GMAW), flux-cored arc welding (FCAW), submerged arc welding (SAW) [15,16]. Most of the work on welding HEAs has been focused on gas tungsten arc welding (GTAW) [19,20]. In the welding process, a non-consumable tungsten electrode is used to produce the weld by melting the base metal (Figure 2a), and the weld area is protected from atmospheric contamination by an inert shielding gas such as argon or helium. In the work by Sokkalingam et al [19], two Al0.5CoCrFeNi–HEA plates (homogenized at 1423 K for 24 hours followed by furnace cooling) of a thickness of 2.5 mm were butt-welded with a current of 40 A and voltage of 12 V at a welding speed of 80 mm/min. The FZ and HAZ formed exhibited distinctively different structures from the base metal (BM) (as shown in Figure 2b). The dominant structure in the FZ mainly consisted of fine elongated grains with diameter of 8–12 μm and lengths of 80–120 μm. The HAZ had a distinctive coarse grain structure with grain boundaries delineated by the body centered cubic (BCC) phase. Vickers micro-hardness tests across the three zones revealed an inferior hardness of the FZ and HAZ compared to the BM, particularly at the center of the welding pool (Figure 2c). Tensile tests of welded sample perpendicular to the welding seam showed that both the yield strength (YS) and the ultimate tensile strength (UTS) were reduced after welding, but the welded sample has maintained a high proportion of the ductility (Figure 2d), which is better than samples with other welding processes (detailed in the following sections).

CoCrFeMnNi system is another major HEAs system with significant potential in structural applications [20–23]. Wu et al [20] investigated the welding of CoCrFeMnNi HEAs with low-energy-density, high-heat-input gas tungsten arc (GTA). The raw material of the workpiece was arc-melted and drop-cast into an ingot of 25.4 mm × 12.7 mm × 127 mm, a rolling operation along the longitudinal direction was then performed in air at room temperature to reach a final thickness of 1.6 mm. The sheets were annealed at 900 ◦C for 1h to obtain an equiaxed microstructure. Buttwelding of two pieces of the sheet alloys were performed at a power level of 8.4 V and 75 A with a welding speed of 25.4 mm/min. The FZ of the welded joints exhibited a centerline grain structure of large columnar grains grown from the fusion lines to the centerlines (Figure 3). Tensile tests of the specimen at 293 K and 77 K showed that the effective YS of the GTA weld was close the BM, but the UTS was significantly lower, in particular for the low temperature test. The ductility of the GTA sample was maintained over 50% (Figure 3). The difference in the microstructures of the FZ between the two material systems (i.e., Al0.5CoCrFeNi and CoCrFeMnNi) is apparent, both showed columnar grain structure, but the effect of the welding on the mechanical response is different. The FZ of Al0.5CoCrFeNi–HEA (Figure 2d) showed a drop in YS and UTS, and maintained high proportion of the ductility against the BM; while the drop in YS for CoCrFeMnNi (Figure 3) was relatively limited but there was a clear drop of the UTS, but it had only maintained about half of the ductility. The work also revealed limited depletion

behavior of Mn in the weld zone, which is different from welding of high-Mn stainless steels [24]. Nano-twins (Σ3) common to CoCrFeMnNi HEA and other face-centered cubic (FCC) metals and alloys with low stacking fault energy were also observed [25]. The nano-twins and twin bundles formed during deformation had significantly beneficial effect on mechanical properties such as tensile strength.

**Figure 2.** Structure and hardness of the welded zones and engineering stress-stain curves of the welded samples of Al0.5CoCrFeNi high entropy alloy (HEA) [19]: (**a**) Schematic of gas tungsten arc (GTA) welding process, (**b**) Scanning electron micrographs of base metal (BM)–heat affected zone (HAZ)–fusion zone (FZ) interfaces, (**c**) Microhardness profile on the surface of the welded sample and (**d**) Engineering stress–engineering strain curve for the BM and welded sample.

**Figure 3.** Structure of the welding zones and tensile test results of welded joint of CoCrFeMnNi alloy by GTA [20].

#### *2.2. Laser Welding*/*Laser Beam Welding*

Research on laser beam welding (LBW) of HEAs has been reported in several recent publications [26–30]. In a laser welding process, pieces of metal are melted and joined through a concentrated heat source provided by the laser beam, forming a narrow, deep weld (Figure 4a). LBW has much higher welding rates than arc welding and is more suitable for high volume applications. LBM is based on keyhole or penetration mode welding [17]. Compared to arc welding, the heat density for laser is much higher and the welding point/zones is narrower and the cooling could be much faster [17,18]. Kashaev et al [26] studied LBW of a CoCrFeNiMn-type HEA produced by self-propagating high-temperature synthesis (SHS). The synthesis of the initial CoCrFeNiMn-type alloy was carried out with the use of thermite-type SHS powders containing oxides of the target elements (NiO, Cr2O3, Co3O4, Fe2O3, MnO2 and high purity Al as the metal reducer). The SHS-fabricated alloy was characterized by ∼2 times reduced Mn content in comparison with that of the other principal components and the presence of impurities including Al, C, S, and Si. In the experiments, welded butt-joints were produced using an 8.0kW fiber laser with a fiber optic (300 μm core diameter) and a 300 mm focal length with a welding speed in the range between 3.0 m/min and 6.0 m/min. Detailed metallurgical analysis revealed that the difference in microstructure and grain orientation distribution between the BM, HAZ and the FZ was not significant. As shown in Figure 4b, the FZ was not fully symmetric. Detailed Vickers hardness tests at the top, middle and bottom of cross-section showed a pronounced increase in microhardness from (153±3) HV 0.5 (BM) to (208±6) HV 0.5 (FZ) (Figure 4b). Nam et al [28] recently also reported the change of hardness on LBW of equiatomic Co0.2Cr0.2Fe0.2Mn0.2Ni0.2. The specimen in the work was prepared *via* vacuum induction melting, casting and homogenisation at 1100 ◦C for 24 h followed by air cooling. The tensile test data (Figure 4c) showed that the rolled specimen was much stronger with significant work hardening than the cast specimen. Systematic tests also showed that the stress-strain curves were not very sensitive to the change of welding speed, further details can be found in the paper [28]. Tests on the cross section of the weld of rolled specimen and cast specimen revealed hardness increase in the FZ and HAZ over the BM, but to a distinctively different extent. The hardness increase in the FZ of the rolled HEA specimen was limited while the hardness increase in the FZ of the cast specimen was over 30% (~130 to ~170 HV 0.5) (Figure 4d). Similar levels of hardness increase were also reported in the work on welding of CrMnFeCoNi plates with thickness of 1 or 2 mm [29], in which metallurgical analysis showed the FZ formed a dendritic structure with the dendritic and interdendritic regions relatively rich in Fe and Mn, respectively. The hardness of the FZ was much higher than the BM (~185 HV0.1 *vs.* ~143 HV0.1). Despite the difference in material systems and the condition of the BM, the thickness and welding powers and speeds [26,28,29], all these works showed an increase in the hardnesses of the FZ and HAZ.

The trend of the other reported works on LBW showed a difference in terms of the hardness changes in the FZ [19,27]. In the work by Sokkalingam et al [19], 1 mm thick sheets of Al0.5CoCrFeNi were welded with a power of 1.5 kW and a traverse speed of 600 mm/min. The FZ that was exposed to rapid heating and cooling rate showed clear grain refinement (Figure 5a) with two kinds of dendritic structures: few longer columnar dendrites at the region next to fusion line, at the interloop boundaries and weld center with an average dendritic spacing of 4.8 μm and smaller equiaxed dendrites at in-between areas. Different from the results shown in Figure 4, Vickers microhardness test showed that the FZ became softer than the BM (Figure 5b). Similar hardness drop in LBW has also been reported by Nam et al [27] on Co0.2Cr0.2Fe0.2Mn0.2Ni0.2 (Figure 5c,d). In the work, the HEA slab used in laser butt-welding was homogenized at 1100 ◦C for 24 h and hot-rolled to 3 mm, followed by air cooling, then cold-rolled to 1.5 mm at room temperature (25 ◦C). As shown in Figure 5d the hardness of the FZ is much lower than the BM. The data also clearly showed that the width of the HAZ was affected by the welding speed but the bound of hardness of the FZ was not sensitive to the welding speed.

**Figure 4.** Hardness distribution in laser welded joints of HEAs: (**a**) Schematic diagram showing the laser welding process, (**b**) Microhardness profile of a butt–joint [26], (**c**) Stress-strain curves of the cast and rolled BM at various welding velocities [28] and (**d**) Hardness distribution in the transverse welds of the cast and rolled HEAs at various welding velocities: 6–10 m min−<sup>1</sup> [28].

**Figure 5.** Structure and hardness of laser beam welded Al0.5CoCrFeNi HEAs: (**a**) Welding zones and boundary [30], (**b**) Microhardness profile on the surface [30], (**c**) Structure of the welded joints [27] and (**d**) Hardness distributions in the transverse weld for various welding velocities [27].

#### *2.3. Electron Beam Welding*

In an electron beam welding (EBW) process, a beam of high-velocity electrons is applied to two materials to be melted and joined. The heat density is high and the melted zone is normally thinner than that in laser and arc welding. In the work by Wu et al [31] on weldability of a high entropy CrMnFeCoNi alloy using a controlled test, a face centered cubic (FCC) CrMnFeCoNi alloy was selected. Welds produced by EBW showed no cracking. Tensile tests data showed that the welded joints possessed mechanical properties comparable to those of the BM at both room and cryogenic temperatures. Compared with the BM, deformation twinning was more pronounced in the FZ of the tested alloy. In another work, Wu et al [20] investigated low-heat-input EBW of CoCrFeMnNi HEA. The EB welds were made at a power level of 125 kV/2.2mA, and at a welding speed of 38 mm min<sup>−</sup>1. The work also showed that the welded specimen has maintained the strength and ductility of the BM indicating good weldability of the HEA in this condition (Figure 6). In the work by Nahmany et al [32], electron beam surface re-melting was employed to modify the surface properties of two five-component AlxCrFeCoNi HEAs (x-0.6 and 0.8) prepared by vacuum arc-melting. The effects of electron beam heat on the structure and mechanical properties of deep penetration welding of HEAs were investigated. The quality of weld was found to be dependent on the welding condition (Figure 7), when the welding heat input was increased from 72 J/mm (P2-1 and P3-2) to 108J/mm (P2-3, P3-4), cracks were observed, which is thought to be due to the residual stresses. In addition, the hardness of the FZ also increased relative to the BM, which is different from EBW. The reason for the hardness difference requires further studies but this highlighted the effect of depth and energy input on the quality of the welding process.

**Figure 6.** Structures and properties of electron beam butt–welded joints of CrMnFeCoNi alloy: (**a**) Electron backscattered diffraction (EBSD) maps showing the grain structure of the welded joints on transverse surface and (**b**) Tensile test results of welded joint [20].

**Figure 7.** Structures of electron beam deep penetration welding of AlCrFeCoNi HEAs with different heat input: (**a**) 72 J/mm [32] and (**b**) 108 J/mm [32].

#### *2.4. Friction Stir Welding*

Several works have studied the friction stir welding (FSW) of HEA systems [29,33–37]. Different from arc, laser and electron beam welding, FSW (Figure 8) is a form of solid-state joining, in which two facing workpieces are jointed through the heat generated by friction between the rotating of a non-consumable tool and the workpiece material [38]. The volume of the material affected in the welding process is much wider than other fusion based welding processes. Zhu et al [34] studied the FSW of a typical FCC CoCrFeNiAl0.3 HEA. The sample was made by arc-melting and cast into ingot plates with a dimension of 2 mm × 10 mm × 30 mm. The welding process was performed with speeds of 30 and 50 mm/min while the rotation rate and load force were kept at 400 rpm and 1500 kg, respectively. The tool has a shoulder diameter of 12 mm, probe diameter of 4 mm and probe length of 1.8 mm. The FSW joint consisted of four different regions: the stir zone (SZ), thermomechanically affected zone (TMAZ), HAZ and BM (Figure 8). The SZ exhibited refined grain size arising from recrystallization and it exhibited higher hardness due to grain size refinement. The TMAZ exhibited a mixed microstructure comprising coarse and fine grains due to partial recrystallization. The XRD results indicated that the HEA remained an FCC structure after FSW. The SZ showed a refined equiaxed microstructure due to recrystallization, and the hardness of the FZ was found to be much higher (220 HV) than the BM (180 HV). Another study [35] reported the work on FSW of a ductile Co16Fe28Ni28Cr28 HEA of a low content of Co with a particular focus on microstructural evolution and weld strength in comparison to typical FCC HEAs. The work revealed a similar trend of hardness increase in the SZ compared to the hardness of the BM (~250 HV vs. 150 HV). The work also showed that the grain size decreased slightly while the hardness increased with increasing the welding speed.

**Figure 8.** (**a**) Schematic diagram showing the friction stir welding (FSW) process and (**b**) Microstructure formed in FSW of Co16Fe28Ni28Cr28 alloys [34]: stir zone (SZ), thermomechanically affected zone (TMAZ), heat affected zone (HAZ) and BM.

Jo et al [29] investigated the microstructure and mechanical properties of friction stir welded CrMnFeCoNi HEA. The material was prepared by vacuum induction melting and hot rolling at 1100 ◦C. The dimension of the FSW plates was 55 mm × 60 mm × 2 mm. The pin diameter was 4–5.76 mm, the pin length was 1.85 mm, the shoulder diameter was 12 mm and the tilt angle was 3◦. The FSW was carried out at a welding speed of 150 mm/min and tool rotation speeds of 600 and 70 RPM. Detailed metallurgical analysis showed that FSW refined the grain size in the weld region by a factor of ∼14 when compared with the BM (Figure 9a). The hardness in the weld region was much higher than the BM (~215 HV vs.144 HV) (Figure 9b). The tensile strength and ductility of FSW CrMnFeCoNi were comparable to that of annealed CrMnFeCoNi. This is probably associated with considerable microstructure refinement in the SZ and further recrystallization in FSW.

**Figure 9.** (**a**) Grain size at different distance from weld center positions and (**b**) Vickers hardness in the cross-section of FSW CrMnFeCoNi HEA [29].

Shaysultanov et al [36] studied FSW of a carbon-doped CoCrFeNiMn HEA in butt joints. Along with the principal elements, a small amount (0.9 at.%) of C was added to the alloy produced by SHS. The CoCrFeNiMn alloy was produced using thermite-type SHS, in which a mixture of powders (oxides of the target elements NiO, Cr2O3, Co3O4, Fe2O3, MnO2, pure carbon C, and Al as the metal reducer) was used as the starting material. The as-cast alloy was cold rolled and annealed at 900 ◦C to produce a refined microstructure. The microhardness measurement showed an approximately 40 HV increase in the area of the SZ in comparison with the BM. Tensile tests were performed on samples perpendicular to the welded seam as well as on the sample along the welding direction (Figure 10). The data for both sample conditions showed a clear noticeable rise in strength in comparison with the BM. This can be associated with the microstructure refinement and some increase in the volume fraction of M23C6 carbides [36]. The sample with the loading axis perpendicular to the weld maintained around 50% of the ductility, while the sample fully taken from the seam showed comparable ductility to the BM.

**Figure 10.** Microhardness distribution across the weld seam (**a**) and tensile stress-strain curves of the FSW specimens cut across (**b**) or along (**c**) the weld seam [36].

#### **3. Grain Structure, Element Distribution and Precipitations in Welded Structures of HEAs**

The structures at different levels (grain, precipitation and lattice) are important to the integrity of the weldment. The structure, precipitation and element segregation/redistribution in a welding process are also important for some service conditions such as corrosion, creep etc. [1,12,15,16]. Apart from the difference in the general structure and hardness profiles of welding zones detailed in Section 2, research work on welding of HEAs also revealed significant difference in secondary phase and precipitation associated with different welding processes and material systems. In arc welding processes, such as GTAW, which is like a miniature casting process, the weld metal experiences a rapid cooling rate. The main grain structure is elongated columnar dendrites nucleated from the fusion line and equiaxed grains near the weld centerline [19,20]. Sokkalingham et al [19] has performed detailed XRD analysis on GTAW for Al0.5CoCrFeNi and found that both the volumetric fraction and lattice constant for both FCC and BCC phases change. The volume fraction of FCC increases (from 76 to 97.8%) but the lattice constant of FCC decreased from 0.3568 to 0.2424, and the lattice constant of BBC increased significantly, from 0.2132 to 0.3020. This is an interesting finding, and further research is required on how this is linked to the element segregation and how this change may affect the mechanical behaviour including shear and corrosion [39,40].

Composition distribution and element segregation across the welded joint at macro and micro scales are important to the understanding of the mechanical and corrosion resistance of the welded joint, which may be associated with many different mechanisms, such as evaporation, time and temperature for element diffusion, oxidation and precipitation, etc. [4,19,20,30,41–43]. For CoCrFeMnNi alloy [20], the GTA weld exhibited microsegregation behaviour in the Mn- and Ni-rich interdendritic region and Co-, Cr-, and Fe-rich dendrite cores, but the depletion of Mn was not observed, while for EBW weld, there was clear Mn depletion in the FZ (Figure 11) [20]. As shown in Figure 11b, the average value of the Mn is lower than the nominal value (20%). This was probably caused by evaporation of Mn due to the high power density associated with the EB welding process. For Al0.5CoCrFeNi, the elemental analysis in FZ revealed that the dentrites were rich in Fe and Co, depleted of Al and Ni, and had an even distribution of the Cr element [19]. Similar results have been observed in LBE of Al0.5CoCrFeNi [30]. It is not fully clear if evaporation of Al had played role influencing the structure and properties. However, elemental analysis showed that the concentration of Al element in dentrites and interdentrites of the as-welded (AW) sample is lower than the grains and the grain boundaries of the BM. The dendrites of the AW have a Al concentration of 4.78 %, while the grain of the BM has a concentration of 5.27%; The data also shows that the interdentrites of the AW region consist of 10.69% Al, while the Al concentration in the BM-grain boundary is about14.43%. So, it is possible that evaporation of Al may have occurred during the LBE process, which requires further quantitative studies together with other mechanisms proposed by the authors, such as precipitation of BBC phase (Al-Ni rich phase). In both material systems (i.e., AlCoCrFeNi and CoCrFeMnNi), the primary dendrite arm spacing and the extent of elemental segregation were less in the welds than in the cast ingot as the cooling in welding is much higher than in casting. In LBW, the heat density and the cooling rate is much higher than arc welding, the FZ gives the appearance of a columnar grained microstructure with random crystallographic orientation, and no significant change in microstructure near the FZ-BM boundary was observed [29]. In another work [26], the differences in microstructure and grain orientation distributions between the BM, HAZ and the FZ were not significant for CoCrFeNiMn-type HEA in LBW. For CrMnFeCoNi, the interdendritic and dendritic regions are enriched in Mn and Fe, respectively, similar to that at the FZ-BM boundary. LBW also resulted in precipitation of the nanoscale B2 phase particles for CoCrFeNiMn-type HEA. In the weldment of Al0.5CoCrFeNi, the equiaxed polygonal grains with grain size of 60 μm in homogenized state transformed to longer columnar dendrites at the region nearer to fusion boundary and inter-loop boundaries with an average dendritic spacing of 4.8 μm and fine equiaxed dendrites at the weld center on welding [19]. The work also found that, in the LB weldment, the hardening factor (Al-Ni rich BCC phase) was less in the weld metal than that in the BM, therefore the weld metal exhibited a lower hardness than that of BM.

*Entropy* **2019**, *21*, 431

For FSW, the SZ comprised of refined grains arising from dynamic recrystallization. The HAZ exhibited a columnar structure [34,35]. In another work [29], EDS-line scan showed that in general the dendrites were enriched in Fe and depleted in Ni and Mn. The fluctuation in composition was significant: between ~5 and 15 at%. The cause for this compositional fluctuation in the FZ, which is expected to stay hot longer than the FZ-BM boundary and thereby allow for greater homogenization, is not clear and needs further study. In the work, relatively dark spherical particles were also detected in the High-angle annular dark-field image (HAADF) and all the element maps were found to be rich in Mn, S and O indicating that these were Mn sulfides/oxides. In the work on FSW of carbon-doped CoCrFeNiMn HEA [36], the microstructure of the SZ was shown to contain M23C6 carbides. More importantly, the increase in the volume fraction and size of the carbides (to 7% and 150 ± 62 nm, respectively) after FSW in comparison with that in the BM can be clearly noted. All these may have contributed to the strength and hardness increase of the welded zone in FSW as illustrated in Figures 9 and 10. This suggests that, for FSW, there is no specific post welding processing needed to obtain a hardened joint, i.e., the alloy hardens due to the "natural" heat treatment during welding [36]. This is an advantage of interstitial alloying of HEAs. This indicates that higher strength of the FSW joints can be achieved in similar alloys by tailoring microstructure *via* adjusting the chemical composition and processing parameters even [26,33,44,45]. A typical example is shown in Figure 12, which shows that the FSW can be used for processing HEAs with enhanced strength and ductility in a friction stir processing engineered dual phase HEA [33].

**Figure 11.** Elemental mapping of fusion zone (FZ) for electron beam welding of CoCrFeMnNi: (**a**) Microstructure of the electron beam (EB) weld zone area and electron microprobe analyzer (EMPA) compositional mapping of the marked area in upper figure and (**b**) Compositional profile along the arrow shown in (**a**) [20].

**Figure 12.** Structure and stress strain curves of high entropy alloys (HEAs) with enhanced properties through friction stir welding [45].

#### **4. Discussion and Future Works**

The recent work summarized in Sections 2 and 3 showed that HEAs have good welding characteristics. A broad range of welding methods were found to be applicable to the welding of HEAs and most cases showed no significant defects or cracking. The size of different welding zones varies with the welding method, such as electron beam, laser and arc welding. FSW resulted in a much wider FZ due to the involvement of the rotating pin. In most of the cases, the shape of the welded beam is symmetrical, which represents good weldability for structure and quality control. However, there are reported cases in which some butt-joints showed a non-symmetrical weld shape (Figure 4b), this is most likely due to coarse columnar grains of the base material [26]. Most of the reported works were based on butt-welding, which is one of the simplest forms of welding configuration. More complex setup such as tensile, shear, bending, torsion, and impact, in which the stress condition is more complex and close to real service conditions, need to be considered.The residual stresses also need to be studied systematically. Residual stress is influenced by many factors such as localized heating and cooling, differential volumetric change occurs both at macroscopic and microscopic level [46]. Limited work has been reported on the measurement of residual stresses in the work reviewed, which were mostly based on relatively simple, less constrained samples. In the case of EBM deep penetration welding of AlxCrFeCoNi HEAs, it was shown that residual stresses could cause cracking (Figure 7b). HEAs have greater freedom in designing composition, mechanical/physical properties, such as yield strength, stiffness, thermal conductivity and expansion, all these may offer opportunity to optimize the material design to manage the residual stresses in welding processes.

As detailed in Sections 2 and 3, the works published clearly show that the change of the strength/hardness of the FZ and HAZ varies between the composition systems and the welding process employed. This is a significant difference from other engineering materials such as carbon or stainless steels, for which, the FZs normally have a hardness increase [47,48]. For nonferrous metals such as magnesium, the hardness for the FZ normally becomes higher but the hardness of the HAZ may drop [49]. For HEAs, hardness drop in the FZ was observed for Al0.5CoCrFeNi in laser and arc welding, while the hardness drop for CoCrFeMnNi is relatively limited or slightly increased. Sokkalingam et al [19,30] suggested that the hardness drop for the Al0.5CoCrFeNi in the LB weldment was due to the reduction of the hardening factor (Al-Ni rich BCC phase) in the weld metal than in the base metal. The processing method of the sample also influences the hardness changes. For example, as shown in Figure 4, the hardness change for the cast specimen after welding was much more profound than that for the rolled specimen. Other parameters such as welding speed and power also showed some influence on the hardness change but not as significant as the material groups and welding methods. Another issue, which requires further work, lies in the question how the hardness/strength change of the FZ and HAZ may affect the overall ductility and toughness of the welded joints. Hardness/strength drop is not an ideal situation in terms of strength following general engineering principles, however, this may potentially offer a way to enhance the ductility and toughness of the welded joints. This can be beneficial to some service condition where toughness is more critical than strength or hardness. Preliminary numerical modelling work by the authors (in-process, result not shown) indicates that soft FZs in HEAs may help to spread the load more evenly and reduce stress concentration, which commonly exists in welded joints with a hardened FZ [17], thus improve the overall ductility or toughness of the welded structure. In the work by Wu et al [31] on EBW of CrMnFeCoNi, tensile test results showed that, compared with the BM, deformation twinning was more pronounced in the FZ of the tested alloy. Further combined experimental and numerical modelling work is required to systematically investigate the structure and stress in the FZ, HAZ and the HAZ-BM or HAZ-FZ interface within welded joints of HEAs under controlled material, welding and testing conditions.

For HEAs and welded structures, many factors may contribute to the strength of the materials (such as grain refinement, precipitation, residual stress, volume ratio of different phases) [15,16,22–25,41–43]. Grain size is still an important but not necessarily the dominating factor depending on the HEA system and the welding processes utilized. For FSW, the SZ becomes much harder [19,20,26] than the BM. One major mechanism contributing to the hardness increase is the refinement of the grain, but the trend of the grain size effect does not fully follow the theories or data established for conventional bulk HEAs. For bulk HEA alloys, the Hall–Petch (H–P) trend between grain size and strength is applicable for a range of HEAs based on either tensile tests or hardness data including CoCrFeNi [50–53], CoCrFeNiMn [7,54–56] and AlxCoCrFeNi [57–61]. While for welding, the suitability of H–P relationship varies with the materials and welding conditions. For example, the data for FSW of Co16Cr28Fe28Ni28 [35] and Al0.1CoCrFeNi [34,62] followed the H–P trend well, but for arc welding and laser welding, the hardness and grain size relationship is opposite to the H–P trend [19,30]. These differences reflect that the controlling operative strengthening mechanisms are more complicated than pure grain refinement, and other factors such as phase change, precipitation all may affect the hardness to different extents. A more quantitative data driven methodology is required to clarify or establish the dominating mechanism(s). For example, the combined contribution of grain refinement and precipitation was quantified in [36] through combined Hall-Petch coefficient analysis. The work showed that a decrease in grain size by a factor of two (from 9.2 μm in the initial condition to 4.6 μm after FSW) resulted in an increase in strength by ~55 MPa. The rest of the increase in the yield strength (~145 MPa) was attributed to the precipitation strengthening associated with some increase in the volume fraction of the M23C6 carbides. Such quantitative analysis is not only important for establishing the strengthening mechanism, it is also essential to provide guidelines for developing effective preor post-treatments of welded structures. The published research work has successfully highlighted the key issues influencing the mechanical properties of the welded joints, which is important for expanding the applications of HEAs and welding processes. However, the data available is still limited to draw direct comprehensive conclusion on the property differences between different welding methods. Even though the general heat input is different between various welding techniques, but more

systematic data on the cooling curves of comparable materials systemsare required to be conclusive when quantifying the main strengthening mechanisms. Physical simulation with controlled cooling (such as GleebleTM) will help to quantify the effect cooling history on the microstructure and properties of different welding zones [62]. Establishment of the mechanisms of phase change, grain refinement and precipitation is also directly beneficial to the development of new materials and interruptive control technologies to maintain or improve the beneficial properties of HEAs such as high temperature stability and corrosion resistance of welded structures of HEAs [63–65].

#### **5. Summary**

Recent works on welding of HEAs with various welding methods of different setup and heat input were reviewed in detail focusing on the research on main HEA systems when applying different welding techniques. The structures and properties of the welding zones in particular the FZ and the HAZ formed with different welding methods were compared and presented in details and the structure properties relationships were discussed. The works showed that weldability of HEAs varies with the composition groups and the welding methods employed. Arc and laser welding of AlCoCrFeNi HEAs resulted in lower hardness in the FZ and HAZ and reduced strength. FSW resulted in higher hardness in the FZ and maintained the strength of the welded joints under tensile load. The welded of HEAs are capable to maintaining reasonable proportion of strength and the ductility. The key structural changes including element distribution, the volume fraction of FCC and BCC as well as some reported lattice were summarized and analyzed. The effects of evaporation for high energy welding (such as LBM and EBM) on composition and structure requires further quantitative study. Detailed mechanism(s) governing the mechanical properties including the contribution of the grain size-properties/hardness relationship in the form of Hall–Petch (H–P) effect for both bulk and welded structure of HEAs were discussed. Future research is required to establish the strengthening mechanisms of the welded joints and the effect of the observed hardness changes in the FZ on the strength and toughness of welded structures of HEAs. Such quantitative analysis will provide guidelines for developing effective preor post-treatments of welded structures. Studies of residual stress for different welding processes and pre-post weld treatments, in particular for under complex loading conditions, is required for enhancing the applications of HEAs. It is also beneficial to the development of new materials and interruptive control technologies to maintain or improve the beneficial properties of HEAs such as high temperature stability and corrosion resistance of welded structures of HEAs.

**Author Contributions:** All the authors contributed to the review. Original Draft preparation: J.G., C.T.; Analysis and editing: L.L., G.R., Y.-C.W. and Q.Y.; supervision: G.R., X.R.

**Funding:** This project has received funding from the European Union's Horizon 2020 research and innovation programme under the Marie Skłodowska-Curie grant agreement No: 793114".

**Acknowledgments:** The authors would like to acknowledge the Royal Society for the support through an International Exchange Grant.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**


© 2019 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

### *Article* **Intermediate-Temperature Creep Deformation and Microstructural Evolution of an Equiatomic FCC-Structured CoCrFeNiMn High-Entropy Alloy**

**Chengming Cao 1, Jianxin Fu 1, Tongwei Tong 1, Yuxiao Hao 1, Ping Gu 1, Hai Hao <sup>2</sup> and Liangming Peng 1,\***


Received: 30 November 2018; Accepted: 9 December 2018; Published: 12 December 2018

**Abstract:** The tensile creep behavior of an equiatomic CoCrFeNiMn high-entropy alloy was systematically investigated over an intermediate temperature range (500–600 ◦C) and applied stress (140–400 MPa). The alloy exhibited a stress-dependent transition from a low-stress region (LSR-region I) to a high-stress region (HSR-region II). The LSR was characterized by a stress exponent of 5 to 6 and an average activation energy of 268 kJ mol−1, whereas the HSR showed much higher corresponding values of 8.9–14 and 380 kJ mol−1. Microstructural examinations on the deformed samples revealed remarkable dynamic recrystallization at higher stress levels. Dislocation jogging and tangling configurations were frequently observed in LSR and HSR at 550 and 600 ◦C, respectively. Moreover, dynamic precipitates identified as M23C6 or a Cr-rich σ phase were formed along grain boundaries in HSR. The diffusion-compensated strain rate versus modulus-compensated stress data analysis implied that the creep deformation in both stress regions was dominated by stress-assisted dislocation climb controlled by lattice diffusion. Nevertheless, the abnormally high stress exponents in HSR were ascribed to the coordinative contributions of dynamic recrystallization and dynamic precipitation. Simultaneously, the barriers imposed by these precipitates and severe initial deformation were referred to so as to increase the activation energy for creep deformation.

**Keywords:** high entropy alloy; tensile creep behavior; microstructural evolution; creep mechanism

#### **1. Introduction**

High-entropy alloys (HEAs), as a novel type of material generally consisting of five or more principle elements, have attracted extensive attention in recent years because of their attractive crystallographic and mechanical properties [1–7]. Although it is anticipated that simple solid solution phases are easier to form rather than the intermetallic compounds or complex-ordered phases due to their high configurational entropy [8,9], only a few HEAs have exactly single and stable solid solution structures, among which CoCrFeNiMn alloy is widely concerned [10–14]. It possesses a single face-center-cubic (fcc) solid solution phase even after a series of thermomechanical processing [15,16] and its constituent elements have more sluggish diffusion compared to other conventional alloys [17]. These characterizations contribute to excellent mechanical properties with a satisfactory combination between tensile strength and ductility [18,19]. Apart from previous investigations mainly focusing on its temperature dependence of strength [20,21], which provided a fundamental understanding of

the strengthening mechanism in fcc solid solution HEAs, recent reports have elucidated the thermally activated process, where dislocation lines overcome nanoscale clusters or short-range orders for controlling the deformation rate [22,23]. This intrinsic mechanism indicated effective heat-resistance during high-temperature plastic deformation processes. Owing to the aforementioned sluggish elemental diffusion, it is optimistic that the CoCrFeNiMn alloy will exhibit a promising perspective in high-temperature applications.

When referring to high-temperature performance for most alloy systems, creep resistance is one of the most important standards for service life and safety reliability of engineering structures. In particular, it is necessary to evaluate the stress exponents and activation energies to demonstrate the mechanism operative in the power-law creep deformation, as these two parameters are usually employed to predict the steady-state creep strain. However, only quite limited reports on tensile creep properties of HEAs at elevated temperatures have been available. The investigation on creep behavior of AlxCoCrFeNi (x = 0.15, 0.60) alloys using the stress relaxation method [24] demonstrated a lower creep resistance for higher Al content alloy, which was attributed to the easier dislocation cross-slip due to its higher stacking fault energy. The high-stress exponent of 8.8 and activation energy of 334 kJ mol−<sup>1</sup> were simply explained by the increasing grain boundary in the initial body-center-cubic (bcc) phase. The dislocation configurations and microstructural evolution were not examined to confirm the validity of the proposed creep mechanism. Recently, He et al. [25,26] investigated the high-temperature plastic flow behavior of single-phase FeCoNiMn and precipitation-hardened (FeCoNiCr)94Ti2Al4 HEAs using strain-rate-jump/stress increment tests at different temperatures (750–900 ◦C). In general, the deformation in these two alloys was divided into two regimes, dependent on the applied strain rate or testing temperatures. Specifically, the obtained higher stress exponent (6–9) and activation energy (>600 kJ mol<sup>−</sup>1) in the latter alloy were proposed to be associated with the strong interactions between dislocations and coherent precipitates, resulting in a threshold stress term in the flow constitutive equation [26]. It should be pointed out that the testing temperatures were up to 750–900 ◦C and far beyond the softening temperature (~600 ◦C) of the investigated alloys [19]. From a scientific view, the strain-rate-jump or stress increment test method is inadequate using a single specimen at respective temperatures, as the microstructure during elevated-temperature plastic flow is generally not invariant. Moreover, such plastic-flow behavior is, to some extent, obviously different from creep deformation under a constant applied stress. Consequently, it is still necessary to systematically investigate the intermediate temperature (not higher than 600 ◦C) creep deformation and examine the microstructural evolution for exploring the intrinsic deformation mechanism of HEAs.

In the present study, the tensile creep tests under constant stress were performed on multiple fine-grained CoCrFeNiMn HEA specimens in the temperature range of 500–600 ◦C. Specifically, some tests were interrupted at the steady-state stages and cooled quickly to reserve the crept microstructures for examining the microstructural evolution. The objective was to reveal the inherent relationship between microstructural evolution and creep properties of single-phase CoCrFeNiMn high-entropy alloy.

#### **2. Experimental Procedures**

Ingots with a nominal composition of Co20Cr20Fe20Ni20Mn20 (in atomic percentage) were prepared by vacuum induction melting the constituent elements with at least 99.9 mass % purity under argon atmosphere. To ensure chemical homogeneity, the ingots were re-melted at least five times and then drop-cast into a rectangular steel mold with a dimension of 120 × <sup>65</sup> × 10 mm3. The ingots were homogenized at 1100 ◦C for 24 h in a vacuum, followed by furnace cooling. Cold rolling was subsequently conducted on these ingots with a reduction of 40% in thickness, with several cross-rolling steps to ensure flatness. The rolled sheets were eventually annealed at 900 ◦C for 1 h.

Flat dog-boned specimens with a dimension of 25 mm in gauge length and 5.6 mm × 1.5 mm in cross-section were electric-discharge machined for creep tests. The specimens were mechanically ground to 2000-grid sand paper to remove surface asperities. Tensile creep tests were conducted

at temperatures of 500, 550, and 600 ◦C under applied stresses of 140–400 MPa on a CSS-3905 multi-functional testing machine. The temperature was measured with a temperature accuracy of ±1 ◦C by three thermocouples closely attached to the upper, middle, and lower sections of the specimen, respectively. The creep strain was continuously measured using a Linear Variable Differential Transducer (LVDT) extensometer with a strain resolution of ± 0.1 μm. The acquisition of strain-time data was accomplished by a computer, and data processing was conducted through a computer program. At least seven applied stress levels were chosen to obtain a wide range of creep rates for each testing temperature. Several specimens were subjected to interrupted tests at a steady-state stage and cooled quickly under the load using liquid nitrogen to freeze the structures produced during creep deformation.

The microstructural observation was conducted using optical microscopy (OM) (AxioImager.A1m, Jena, Germany) after etching in an aqueous solution of HCl+H2O2+Cu(NO3)2. The phase constitution was identified by X-ray diffraction (XRD) with Cu-Ka radiation (PANalytical X'pert PRO, Almelo, Netherlands). Thin foils from the interrupted specimens for transmission electron microscope (TEM) observations were twin-jet electropolished in an ethanol solution containing 5% perchloric acid at −25 ◦C and an applied voltage of 35 V. TEM investigations were conducted on a JEOL JEM2100F microscope operated at 200 kV. Fracture surfaces were examined by a scanning electron microscope (SEM) (XL30 ESEM, Philip, Netherlands) equipped with an energy-dispersive spectrometer (EDS)).

#### **3. Results**

#### *3.1. Initial Microstructures*

Figure 1 shows the initial microstructures of the alloy in as-cast and annealed states prior to high-temperature creep deformation. It is evident that the cast alloy exhibits typical dendritic and interdendritic structures, caused by the segregation during the quick freezing process in the mold. After being cold-rolled and recrystallized, equiaxed and homogenous grain structures are obtained with many annealing twins visible inside the grains. The average grain size is evaluated to be approximately 25 μm. Energy-dispersive spectroscopy (EDS) analyses demonstrate that different grains have almost identical elemental compositions as Cr = 19.9, Co = 19.9, Fe = 19.9, Ni = 19.8, and Mn=19.5 with a little loss of Mn due to its slight evaporation at an elevated temperature. As depicted by XRD patterns in Figure 1c, the alloy consists of a single fcc solid solution phase without precipitates or intermetallic compounds in both cast and recrystallized states. Furthermore, almost no peak shift is observed in the two states, with a cell parameter of a = 0.361 nm for the fcc phase. The present results are different from those observed in the Al0.3CoCrFeNi HEA where precipitation of nanometer scale-ordered L12, B2, and sigma phases occurred in the case of different thermo-mechanical processing routes [27].

**Figure 1.** Microstructures of the alloy in different states. (**a**) As-cast, (**b**) thermo-mechanical treatment (cold-rolled and annealed at 900 ◦C for 1 h), and (**c**) X-ray diffraction (XRD) patterns.

#### *3.2. Steady-State Creep Deformation Behavior*

Figure 2 shows the selected creep curves of the alloy at 500–600 ◦C under different stress levels. It can be found that each of the individual curves exhibits a rather short primary creep stage and a relatively long steady-state region where the creep strain increases linearly with time, especially under lower stress levels. Necking and fracture of the specimens do occur, and hence, a tertiary stage of the curves is recorded in cases of high stress, where the creep rate accelerates with time until the final fracture occurs.

**Figure 2.** Selected creep strain versus time at (**a**) 500 ◦C, (**b**) 550 ◦C, and (**c**) 600 ◦C.

It is generally accepted that the steady-state creep rate . *ε* of metallic materials can be correlated with the applied stress *σ* using the well-known power-law equation as follows [28]:

$$\dot{\varepsilon} = \frac{AD\_0 Gb}{kT} (\frac{b}{d})^p \left(\frac{\sigma}{G}\right)^n \exp(-\frac{Q}{RT}) \tag{1}$$

where *A* is a material-dependent constant, *G* the shear modulus, *b* the Burgers vector, *D*<sup>0</sup> the frequency factor, *k* the Boltzmann's constant, *T* the absolute temperature, *d* the grain size of polycrystalline materials, *Q* the activation energy for creep deformation, and *p* and *n* the grain size and stress

exponents, respectively. The stress exponent *n* and activation energy value *Q* are determined according to Equation (1) by plotting the steady-state creep rate against applied stress on a double logarithmic scale and the reciprocal of the absolute temperatures 1/*T* at constant stress levels on a semi-logarithmic scale, as shown in Figures 3 and 4, respectively. It is evident that the data exhibit two distinct regions. In the low-stress region (subsequently denoted as LSR-region I), the stress exponent values vary in the range of 5 to 6, and the average activation energy *Q* takes a value of 268 kJ mol−1. On the contrary, in the high-stress region (subsequently denoted as HSR-region II), the stress exponents are in the range of 8.9–14, and the activation energy is *Q* = 359–410 kJ/mol with an average value of 380 kJ mol−1. The applied stress at which the transition occurred increased from 200 to 350 MPa, with the temperature decreasing from 600 to 500 ◦C.

**Figure 3.** Dependence of creep rates on applied stress showing the transition in stress exponents.

**Figure 4.** Arrhenius plot of steady-state creep rate versus temperature to determine the activation energy for (**a**) low-stress region I and (**b**) high-stress region II.

*Entropy* **2018**, *20*, 960

It should be noted that the average activation energy of 268 kJ mol−<sup>1</sup> in LSR-region I is only slightly lower than the range of activation energies for lattice diffusion of the constituent elements in the alloy (288–317 kJ mol−<sup>1</sup> [17]), and thus can be considered to be comparable to the lattice diffusion of the constituent elements in the alloy. In addition, this value is also quite close to those (284–333 kJ mol<sup>−</sup>1) reported for the steady-state flow behavior of this alloy at higher temperatures and lower stress levels [26]. The two characteristic parameters suggest a dislocation-climb mechanism operative in LSR. Nevertheless, the abnormally high stress exponent in HSR-region II is obviously beyond the values of 3 to 5 responsible for the dislocation glide or climb mechanism [29,30] reported in some Mg [31], Al [32], and Ti [33,34] alloys. The underlying operative mechanism for HSR creep deformation in the present alloy will be addressed in the following section. In order to eliminate the influence of temperature on creep rate, the normalized creep rates . *εkT* exp(*Q*/*RT*)/*G* versus the shear modulus-compensated stresses *σ*/*G* were plotted in Figure 5, where activation energies *Q* were taken to be 268 kJ mol−<sup>1</sup> and 380 kJ mol−<sup>1</sup> for region I and region II, respectively, and the shear modulus *<sup>G</sup>* = <sup>85</sup> − 16/(*e*448/*<sup>T</sup>* − <sup>1</sup>) [35]. It is evident that almost all data points in the two stress regions at different temperatures can be represented by a single line with a respective slope of 5.5 and 10.6 for LSR and HSR. This strongly implies that creep deformation in both regions may be controlled by lattice diffusion of constituent elements in the alloy. However, it should be noted that the correlation between . *εkT* exp(*Q*/*RT*)/*G* and *σ*/*G* in HSR still yields an abnormally high stress exponent.

**Figure 5.** Normalized creep rate versus shear modulus-compensated stress for (**a**) low-stress region I and (**b**) high-stress region II.

#### *3.3. Crept Microstructures*

Figure 6 shows the microstructures of the interrupted samples at different temperatures and stress levels to illustrate the evolution of grain morphologies. Compared to the initial microstructure of the alloy, obvious coarsening can be observed when the specimens are subjected to low stress levels (500 ◦C/200MPa and 600 ◦C/140MPa), as depicted in Figure 6a,d. Similar grain coarsening phenomena are also visible at intermediate stress levels (500 ◦C/320MPa and 600 ◦C/200MPa) near the transition stress indicated in Figure 3. However, the volume fraction of small equiaxed grains increases at the coarse grain boundaries (indicated by arrows in Figure 6b,e), revealing strong evidence of the occurrence of dynamic recovery and recrystallization during creep deformation. Moreover, higher stress levels (500 ◦C/400MPa and 600 ◦C/320MPa) lead to the formation of completely recrystallized and refined microstructures (Figure 6c,f). The resulting average grain sizes are statistically evaluated to be approximately 56, 45, and 22 μm at 500 ◦C, whereas 48, 43, and 31 μm at 600 ◦C under different chosen stress levels.

**Figure 6.** Optical microstructures showing the evolution of grain morphology after being crept at (**a**) 500 ◦C/200MPa, (**b**) 500 ◦C/320MPa, (**c**) 500 ◦C/400MPa, (**d**) 600 ◦C/140MPa, (**e**) 600 ◦C/200MPa, and (**f**) 600 ◦C/320MPa.

The dislocation substructures are shown in Figures 7–9, with selected area electron diffraction (SAED) patterns inserted to identify the precipitations. In general, the dislocation densities in HSR were distinctly higher than those in LSR at three testing temperatures. After being deformed at 500 ◦C/200 MPa (LSR), only a few short straight dislocation segments were displayed inside the grains (Figure 7a), and in particular, several lath-shaped areas were formed with closely-spaced dislocations arranged (Figure 7b). However, the substructures reveal increasing dislocation activity in LSR with increasing temperature, as the dislocations are extensively curved or even looped (Figures 8a and 9a). Apart from the limited number of alignments of pile-ups in Figure 8b, it is noted that cusped configurations are frequently observed inside grains with the bowed segments on either side. Such cusped configurations may be attributed to intrinsic pinning due to the occurrence of jogs along the screw dislocations [34]. In contrast, the dislocations in HSR exhibit a higher density of configurations. Numerous short straight segments are tangled and tend to form cell substructures at 500 ◦C/400 MPa (Figure 7c). However, the dislocations at 550 ◦C/360 MPa show a relatively homogeneous distribution still with high densities of curved dislocations tangling with each other (Figure 8c). The tangled dislocations at 600 ◦C/320 MPa are severely curved, with a large number of loops formed in a wide range of sizes (Figure 9c).

Another striking aspect of microstructural evolution relies on the precipitation of dispersoids during a long-term deformation process at elevated temperatures. On the whole, considerable precipitates of irregular shape were predominantly observed at the grain boundaries in HSR. The sizes of these precipitates were within a range of 50–200 nm and increased with testing temperature. In cases of 500 ◦C/400MPa and 600 ◦C/320MPa (Figures 7d and 9d), quantitative micro-EDS analysis indicates that the chemical compositions (at. %) of the precipitates take almost identical values of Cr = 26.6, Mn = 20.2, Fe = 18.9, Co = 17.5, Ni =16.6 (in at. %) and a minor amount of carbon. The SAED pattern along the [011] zone axis exhibits intense diffraction spots from the fcc matrix, accompanied by weak spots from the precipitates. The crystal structure of these precipitates is identified as fcc M23C6 carbide with a lattice parameter of 1.06 nm, which is consistent with both the general identifications of

grain-boundary precipitates in austenitic steel [36] and the second phase observed in coarse-grained CoCrFeNiMn alloy after being subjected to prolonged exposures at 700 ◦C [37]. The tiny presence of carbon in the alloy may have originated from either the potential contamination of starting materials, or the melting system. Nevertheless, as depicted in Figures 8d and 9b for 550 ◦C/360MPa and 600 ◦C/140MPa, the micro-EDS analysis demonstrates chemical compositions (at %) of the precipitates with Cr = 48.8, Mn = 15.5, Fe = 14.0, Co = 11.5, and Ni = 10.2. The present results are quite consistent with previous reports for the prolonged annealing CrMnFeCoNi system [38,39], where the precipitates were identified as a quinary variant of the binary Cr-Fe σ phase. The interplanar spacings are calculated to be d1/d2/d3 = 0.648/0.402/0.337 nm and show a satisfactory agreement with the values of d1/d2/d3 = 0.648/0.383/0.330 nm for (110)/(111)/(201) lattice planes in the [112] zone axes for the σ-FeCrMo phase [40]. Accordingly, the precipitates in the 50 ◦C/360MPa crept alloy are concluded to be Cr-rich tetragonal σ phase with lattice parameters of a/c = 0.916/0.509 nm. Similarly, the weak spots along the [221] zone axes in the SAED pattern in Figure 9b are also indicated to be very close to the Cr-rich σ phase, with lattice parameters of a/c = 0.895/0.473 nm.

**Figure 7.** Transmission electron microscope (TEM) images of dislocation substructures in the interrupted specimens after being crept at 500 ◦C under (**a**) and (**b**) 200 MPa (low-stress region (LSR)); (**c**) and (**d**) 400 MPa (high-stress region (HSR)).

**Figure 8.** TEM images of dislocation substructures in the interrupted specimens after being crept at 550 ◦C under (**a**) and (**b**) 160 MPa (LSR); (**c**) and (**d**) 360 MPa (HSR).

**Figure 9.** TEM images of dislocation substructures in the interrupted specimens after being crept at 600 ◦C under (**a**) and (**b**) 140 MPa (LSR); (**c**) and (**d**) 320 MPa (HSR).

#### *3.4. Fractographs*

All the fractographs in Figure 10 for HSR from 500 ◦C to 600 ◦C exhibit typical ductile rupture, where both the width and depth of dimples increase with temperature. However, slip striations around the dimples become notable at 600 ◦C. Meanwhile, a small number of particles identified as Mn-containing oxides by EDS are visible on the fracture surfaces, which is attributed to an in-situ oxidation during fracture process [26]. It is worth noting that the crept samples fracture in a ductile transgranular manner at high stress, which is contrary to common knowledge, as large amounts of brittle are generally considered to be destructive to the stability of grain boundaries during the tertiary creep stage [41,42]. The ductile creep failure may be associated with the dislocation interactions during high-temperature deformation, which contribute to the progressive generation of vacancies and nucleation of voids.

#### **4. Discussion**

The observed jog configurations are indicative of a dislocation climb process, and thus the stress-assisted dislocation climb controlled by lattice diffusion may be the operative creep mechanism of the present high-entropy alloy in LSR. Contrary to LSR, creep deformation in HSR exhibits abnormally high stress exponents of 8.9–14 and an average activation energy of 380 kJ mol−1. Combined with the dislocation configurations in Figures 8c and 9c, the same creep mechanism may also be considered rate-controlling. The unusually high stress exponents and activation energy in several traditional alloys have been frequently described using the power-law breakdown equation [43–46]. However, no deviation from the linear relationship on double-logarithmic curves of log . *ε*~log *σ* were observed, and therefore, the power-law breakdown assumption may be inappropriate for the present alloy.

It is documented that dynamic recrystallization during the creep process could cause an increase in the stress exponent [47–49]. Recrystallization is generally accompanied by intense diffusion of solutes, resulting in a refined-grain microstructure and a larger opportunity for grain boundary sliding. Under such circumstances, the creep rate would increase with decreasing grain size [50]. As shown in Figure 6, obvious dynamic recrystallization occurred in the present CoCrFeNiMn alloy in HSR. In fact, low stacking-fault energy of alloys is conducive to promoting recrystallization [51]. Based on Equation (1) where *p* is generally equal to 2 for a dislocation-controlled creep, an increase in the stress exponent Δ*n* due to dynamic recrystallization is expressed as [47]:

$$
\Delta n = n\_{app} - n = \frac{2\ln(d\_1/d\_2)}{\ln(\sigma\_2/\sigma\_1)}\tag{2}
$$

where *napp* is the experimentally measured apparent stress exponent, *n* is the stress exponent prior to the occurrence of recrystallization, *d*<sup>1</sup> and *d*<sup>2</sup> denote the average grains size before and after dynamic recrystallization, and *σ*<sup>1</sup> and *σ*<sup>2</sup> are the corresponding stress levels. Using the relevant evaluated grain sizes, the Δ*n* values were calculated to be 6.5 and 1.4 for 500 ◦C and 600 ◦C, respectively. The corresponding modified stress exponents *n* then fall down to ~7.5. However, this value is still higher than those for traditional solid solution alloys, which implies that the sole recrystallization is insufficient and additional mechanisms should be involved to be responsible for the entire increase in stress exponents.

TEM observations provide clear evidence of dynamic precipitation in HSR. As a result, precipitation hardening may have contributed to the high stress exponents [25]. In general, the precipitates are predominantly formed along the grain boundaries and would exert a boundary obstacle stress *σbo*, which prevent the annihilation of moving dislocations at the boundaries. The stress controlling the dislocation velocity is reduced to an effective value obtained by subtracting the stress necessary for overcoming precipitates from the applied stress. Equation (1) is then replaced by [52,53]:

$$\dot{\varepsilon} = \frac{AD\_0 Gb}{kT} (\frac{b}{d})^p \left(\frac{\sigma - \sigma\_{b0}}{G}\right)^n \exp(-\frac{Q}{RT}) \tag{3}$$

By re-plotting the double-logarithm relationship between . *ε*/ exp(−*Q*/*RT*) and (*σ* − *σbo*)/*G* in HSR, the stress exponents can then be decreased by subtracting the precipitation-hardening portion. Unfortunately, it is difficult to precisely estimate the values of boundary obstacle stress and the decrease in the stress exponent, due to the ambiguous interaction between dynamic precipitation and recrystallization. Nevertheless, combinative effects of precipitation hardening and recrystallization can provide reasonable insights into the intrinsic reason for the abnormal stress exponents and activation energy for creep in CoCrFeNiMn alloy. Moreover, the estimated abnormal activation energy in HSR might also stem partially from precipitation hardening [22]. The nano-sized dispersoids could effectively fasten the grain boundaries, and thus extra activation energy is needed to overcome the barriers for recrystallization and further growth of grains. In addition, the samples experience drastic

plastic deformation in a short time to establish a constant stress state for subsequent creep deformation. As a result, some grains are severely elongated, which, in turn, exerts an extra influence on increasing the barriers for dynamic recovery and recrystallization.

#### **5. Conclusions**

The intermediate-temperature creep behavior and microstructural evolution of CoCrFeNiMn HEA have been studied. Optical micrographs and TEM images of the dislocation substructure after creep deformation revealed the effects of dynamic recrystallization and dynamic precipitation on the deformation mechanisms. The following conclusions can be drawn:


**Author Contributions:** L.P. initiated this research project. C.C. performed the creep tests. T.T. and Y.H. prepared the entropy alloy. J.F. performed the microstructural characterization and thermal-mechanical processing under the supervision of P.G., H.H. and L.P. All authors discussed the results and approved the final manuscript.

**Funding:** This research was funded by the National Natural Science Foundation of China (11572306) and the Fundamental Research Funds for Central Universities (WK2090050040).

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**


© 2018 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

### *Article* **Mechanical Properties and Microstructure of a NiCrFeCoMn High-Entropy Alloy Deformed at High Strain Rates**

#### **Bingfeng Wang 1,2, Xianrui Yao 2, Chu Wang 2, Xiaoyong Zhang 1,\* and Xiaoxia Huang <sup>2</sup>**


Received: 1 October 2018; Accepted: 17 November 2018; Published: 21 November 2018

**Abstract:** The equiatomic NiCrFeCoMn high-entropy alloy prepared by arc melting has a single crystallographic structure. Mechanical properties and microstructure of the NiCrFeCoMn high-entropy alloy deformed at high strain rates (900 s−<sup>1</sup> to 4600 s−1) were investigated. The yield strength of the NiCrFeCoMn high-entropy alloy is sensitive to the change of high strain rates. Serration behaviors were also observed on the flow stress curves of the alloy deformed at the strain rates ranging from 900 s−<sup>1</sup> to 4600 s−1. The Zerilli–Armstrong constitutive equation can be used to predict the flow stress curves of the NiCrFeCoMn high-entropy alloy. Large amounts of deformation bands led to obvious serration behaviors of the NiCrFeCoMn high-entropy alloy under dynamic loading.

**Keywords:** high-entropy alloy; electron microscopy; plasticity methods; plasticity; serration behavior

#### **1. Introduction**

The concept of multi-component high-entropy alloys (HEAs) has been presented in the beginning of this century. Most HEAs usually contain five elements with nearly equal atomic ratios [1]. HEAs have excellent mechanical properties, such as low strength and high plasticity [2–12]. From recent literatures, the hardness of FeCoNiCrMn high-entropy alloy has even increased to 6700 MPa [13]. The Al0.6CoCrFeNi high-entropy alloy also displayed the excellent strength–ductility combination for nanoscale deformation twins induced by dynamic loading and high-density dislocation substructure [14]. However, the NiCrFeCoMn high-entropy alloy has still been regarded as a typical case for its single-phase face-centered cubic (FCC) [15] and relative promising mechanical properties. For example, its yield strength increases with the angle of rotation at high-pressure torsion (HPT) at room temperature at the cost of reduced ductility [16] and can reach up to 350 MPa when reducing its grain size [17,18]. Therefore, the NiCrFeCoMn high-entropy alloy prepared by arc melting can be used to fabricate big parts of industry devices or as the transitional layer between the two types of alloys, e.g., the HEA solder used for welding pure titanium and chromium–nickel–titanium stainless steel [19]. At the present time, dynamic impacts are widely found in aeronautical engineering, the automotive industry, and marine engineering. The as-cast NiCrFeCoMn high-entropy alloy could be applied to many dynamic deformation processes such as penetration, impact cyclic loading, and shock loading. Hence, it is vital to comprehend the dynamic behavior of as-cast NiCrFeCoMn high-entropy alloy under low/high-speed loading and expand its applications.

Many researchers were devoted to investigating the deformation behavior of the high-entropy alloy at different strain rate levels (1 × <sup>10</sup>−<sup>8</sup> <sup>s</sup>−<sup>1</sup> to 10 s−1) [20]. From the recent research, serration

behavior was found in the high-entropy alloy CoCrFeMnNi prepared by powder metallurgy at a low strain rate of 1 × <sup>10</sup>−<sup>3</sup> <sup>s</sup>−<sup>1</sup> and the high strain rates (1 × <sup>10</sup><sup>3</sup> <sup>s</sup>−<sup>1</sup> to 3 × <sup>10</sup><sup>3</sup> <sup>s</sup><sup>−</sup>1) [21]. Serration behavior was supposed to be associated with Cottrell atmosphere interaction with moving dislocations, slip bands, and dynamic strain aging, i.e., the dynamic breakaway/locking of dislocations from/by mobile solute atoms at intermediate temperatures. Until now, only several researches on dynamic behavior of high-entropy alloys under high strain rates (beyond 1 × <sup>10</sup><sup>3</sup> <sup>s</sup><sup>−</sup>1) have been published. Kumar et al. [22] had investigated the strain-rate sensitivity of yield strength for the Al0.1CrFeCoNi high-entropy alloy under high strain rates. Li et al. [23] had also suggested that the Al0.3CoCrFeNi high-entropy alloy exhibits high strain-rate sensitivity. Dirras et al. [24] had found that the Ti20Hf20Zr20Ta20Nb20 high-entropy alloy would be strongly localized under deformation at high strain rates. He et al. [25] had discussed the strain-rate sensitivity effect for the FeCoNiCrMn high-entropy alloy and found a higher strain-rate sensitivity of 0.022 than that of traditional FCC metals. Park et al. [26] had even found that the strain-rate dependency of the yield strength under dynamic conditions would be much higher than that under quasi-static conditions. However, the mechanism of the special phenomenon for the high-entropy alloy deformed at high strain rates (beyond 1 × 103 <sup>s</sup><sup>−</sup>1) is not clear yet. Serration characteristics are ubiquitous in many structural and functional materials such as high-entropy alloy (AlCoCr1.5Fe1.5NiTi0.5, Al0.3CoCrFeNi etc.). It would lead to the instability of mechanical properties. By studying the serration behavior, it may be possible to have an early warning signal for oncoming epileptic seizures and perhaps economic trends.

In the present work, we used the split-Hopkinson pressure bar to investigate the dynamic mechanical behavior of NiCrFeCoMn high-entropy alloy prepared by arc melting. The aims are: (1) To report the mechanical properties and microstructure under dynamic loadings, (2) to obtain the plastic model, and (3) to discuss the microstructural mechanism for the serration behavior when the alloy deforms at high strain rates.

#### **2. Experiments and Procedures**

The material was prepared by arc melting and the chemical composition is given in Table 1. Figure 1 shows the optical micrograph of the initial microstructure of the sample. The obtained specimen exhibits an as-cast dendrite structure and the columnar crystals are distributed evenly. Elemental scanning results in Figure 2 show the distribution of each element evenly in the alloy. From the XRD pattern shown in Figure 3, it can be found that the NiCrFeCoMn high-entropy alloy is composed of a simple FCC solid solution.

**Table 1.** Chemical composition of the NiCrFeCoMn high-entropy alloy.

**Figure 1.** Optical micrograph of the as-cast NiCrFeCoMn alloy.

**Figure 2.** Element planar pattern of the as-cast NiCrFeCoMn alloy.

**Figure 3.** The XRD pattern of the as-cast NiCrFeCoMn alloy.

In this investigation, cylindrical specimens were used for mechanical testing. To ensure uniaxial compressive condition, the end faces of the compressive specimens were ground on each side with SiC paper. During the process of the impact loading and the electrical signal collection, we had adopted advanced waving plastic and anti-jamming techniques. Two kinds of compressive tests were adopted to do mechanical testing at an ambient temperature (298 K) as follows: (1) Quasi-static compressive tests were performed with an INSTRON 3369 machine at the strain rate of 1 × <sup>10</sup>−<sup>3</sup> <sup>s</sup>−1, (2) dynamic compressive tests were performed with a split-Hopkinson pressure bar at the strain rates of 900 s−1, 1700 s−1, and 4600 s−1. The cylindrical specimens had three diameters. Specimens for the strain rate of about 900 s−<sup>1</sup> had a height of 8.4 mm and diameter of 6 mm, for the strain rate of about 1700 s−<sup>1</sup> they had a height of 5.6 mm and diameter of 4 mm, and for the strain rate of about 4600 s−<sup>1</sup> they had a height of 2.8 mm and diameter of 2 mm. The strain rate, the true strain, and the true stress of cylindrical specimens can be obtained by the following equations:

$$\dot{\varepsilon} = -\frac{2\mathbf{C}\_0}{\mathbf{L}\_\mathbf{S}} \varepsilon\_r(t),\tag{1}$$

$$\varepsilon = -\frac{2\mathbf{C}\_0}{\mathbf{L}\_\\$} \int\_0^t \varepsilon\_r(t)dt,\tag{2}$$

$$
\varepsilon = -\frac{2\mathcal{C}\_0}{\mathcal{L}\_\mathbb{S}} \int\_0^t \varepsilon\_\mathbb{r}(t)dt,\tag{3}
$$

where E0 and C0 are elastic modulus and elastic wave speed in a split-Hopkinson pressure bar, A0 is the cross-sectional area of the bar, As and Ls are the cross-sectional area and the length of the cylindrical specimens, and *εr*(*t*) and *εt*(*t*) are the experimentally measured strain of incident and transmitted stress pulse on the split-Hopkinson pressure bars respectively.

Cylindrical specimens were sectioned to two halves along the impacting axis by line cutting. Afterwards, the half sections were polished with the etchant of 50 mL hydrochloric acid, 50 mL water, and 10 g CuSO4·5H2O. Samples were cured and ground using SiC papers. Polishing steps were employed by using diamond pastes and polishing cloths. A mixture solution of 90% acetic acid and 10% perchloric acid at room temperature and at an applied voltage of 27 V for 15 s was then used for electro-polishing. Optical microscopy was performed with a POLYVAR-MET microscope. The crystallographic structure was identified by X-ray diffraction (Rigaku D/MAX-2500 X-ray diffractometer, Rigaku Corporation, Tokyo, Japan) using a Cu target at an operating voltage of 40 kV and current of 250 mA. Electron backscatter diffraction (EBSD) patterns were collected using a ZEISS EVOMA10 scanning electron microscope (SEM, Carl Zeiss SMT Ltd., Cambridge, UK) equipped with a detector and operated at an accelerating voltage of 20 kV. The working distance for SEM is about 10 mm. Transmission electron microscopy (TEM, Royal Philips, Amsterdam, Netherlands) observations were carried out with a Tecnai G2 T20 ST transmission electron microscope operated at 200 kV.

#### **3. Results**

Figure 4 presents the true stress–strain curves of the NiCrFeCoMn high-entropy alloy. It can be seen that the yield strength of the NiCrFeCoMn high-entropy alloy increased from 490 MPa to 800 MPa, with the strain rates varying from 900 s−<sup>1</sup> to 4600 s−1. The flow stress curves are likely smoothed when the specimens are deformed at the strain rate of 0.001 s−<sup>1</sup> at an ambient temperature (298 K). However, the serrations appeared on the flow stress curves of the specimens deformed at high strain rates (e.g., 900 s<sup>−</sup>1, 1700 s−1, and 4600 s−1). Furthermore, with the increasing of the strain rates, the serrations on the flow stress curves became more serious.

**Figure 4.** Compressive true stress–strain curves of the NiCrFeCoMn high-entropy alloy at different strain rates.

The EBSD technique was used to investigate the microstructure and the micro-orientation of a NiCrFeCoMn high-entropy alloy. Figure 5a–c are the electron backscattered diffraction images of the as-received specimen, the specimen deformed at the strain rate of 0.001 s−1, and the specimen deformed at the strain rate of 4600 s<sup>−</sup>1. First, the noise of the images was reduced. Then, to Figure 5b,c were added a layer of full Euler angles after adding a layer of band contrast. The size of the Kuwahara filter was 3-pixel points × 3-pixel points and the smoothing angle was 5◦. Among them, high-angle boundaries (60◦) were marked with black lines. It is evident that the alloy in the as-cast condition consisted of grains on the order of micrometers in size. A visual comparison of Figure 5b,c suggests a higher density of high-angle boundaries in the specimen deformed at a high strain rate of 4600 s−<sup>1</sup> compared to the specimen deformed at a low strain rate of 0.001 s<sup>−</sup>1. Further, the intensive deformation bands were generated in the specimen deformed at a high strain rate of 4600 s<sup>−</sup>1, as shown in Figure 5c. Therefore, the obvious serrations in flow stress curves of the specimen deformed at high strain rates are caused by the intensive deformation bands.

**Figure 5.** Electron backscattered diffraction images. (**a**) the Euler image of the as-received specimen; (**b**) the Euler+BC image of the deformed specimen at the strain rate 1 <sup>×</sup> <sup>10</sup>−<sup>3</sup> <sup>s</sup>−1; (**c**) the Euler+BC image of the deformed specimen at the strain rate 4600 s<sup>−</sup>1.

Bright field electron images taken for the specimens deformed at strain rates of 1700 s−<sup>1</sup> and 4600 s−<sup>1</sup> are shown in Figure 6. Figure 6a,c show the high-density dislocations and deformation bands that were formed in the NiCrFeCoMn high-entropy alloy. These deformation bands were distributed in parallel. The deformation band consisted of nanograins, as shown in Figure 6b,d. Comparing Figure 6a,b and Figure 6c,d, the parallel deformation bands were more clear when the specimen was deformed at relative lower strain rates as shown in Figure 6a,c, and the sizes of the nanograins in the specimens under a strain rate of 1700 s−<sup>1</sup> were larger than those in the specimens deformed at a strain rate of 4600 s<sup>−</sup>1.

**Figure 6.** *Cont.*

**Figure 6.** Bright field electron images showing microstructure in the specimens deformed at high strain rates. (**a**) and (**b**) are for the specimen deformed at the strain rate of about 1700 s−<sup>1</sup> (**c**) and (**d**) are for the specimen deformed at the strain rate of about 4600 s<sup>−</sup>1.

#### **4. Discussion**

#### *4.1. The Constitutive Model and Strain-Rate Sensitivity*

The main dynamic constitutive equations are the Johnson–Cook and Zerilli–Armstrong plastic models. Among them, the Johnson–Cook model is the most widely used for its relative simple expression. It can be represented as follows [27,28]:

$$\sigma = (\text{A} + \text{B}\varepsilon^{\text{n}}) \left[ 1 + \text{C} \ln \left( \frac{\dot{\varepsilon}}{\dot{\varepsilon}\_0} \right) \right] \left[ 1 - \left( \frac{T - T\_r}{T\_m - T\_r} \right)^{\text{m}} \right], \tag{4}$$

where A, B, and C are material constants, σ and *ε* are the flow stress and the equivalent plastic strain respectively, . *ε* and *T* are the equivalent plastic strain rate and the experimental temperature respectively, and *Tr* and *Tm* are the reference temperature (usually room temperature) and the melting point, respectively.

The dislocation mechanism is important for studying the plastic deformation of metallic materials under dynamic deformation. Therefore, the Zerilli–Armstrong plastic model improves the Johnson–Cook model on the basis of dislocation mechanism [29]. In this work, the Zerilli–Armstrong plastic model was used for predicting the strain rate flow behavior of the as-cast NiCrFeCoMn high-entropy alloy. It can be represented as follows:

$$\sigma = \mathbf{C}\_0 + \mathbf{C}\_1 \times \boldsymbol{\varepsilon}^P \times \exp\left[-\mathbf{C}\_2 T + \mathbf{C}\_3 T \ln\left(\frac{\dot{\varepsilon}}{\dot{\varepsilon}\_0}\right)\right],\tag{5}$$

where C0, C1, C2, C3, and P are material constants. Note that T is 298 K.

Taking initial values of the Johnson–Cook model as follows: A = 1, B = 4, C = 100, *n* = 1, then the parameter values and the constitutive relation of stress-strain with strain rate could be obtained by using the MATLAB program (Version 7.0). Therefore, the constitutive equation based on the Johnson–Cook plastic model can be obtained as follows:

$$
\sigma = \left( 0.6039 + 4.81 \epsilon^1 \right) \left[ 1 + 121.6 \ln \left( \frac{\dot{\varepsilon}}{\dot{\varepsilon}\_0} \right) \right]. \tag{6}
$$

As above, taking initial values of the Zerilli–Armstrong model as follows: C0 = 400, C1 = 9 × 105, C2 = 0.01, C3 = 0.001, P = 0.5, then the constitutive equation based on the Zerilli-Armstrong plastic model can be obtained as follows:

$$\sigma = 473.3 + 9.3 \times 10^5 \times \epsilon^{0.5} \times \exp\left[-0.0415T + 0.0026T \ln\left(\frac{\dot{\varepsilon}}{\dot{\varepsilon}\_0}\right)\right].\tag{7}$$

Figure 7 shows the comparison of the calculated results obtained from Equations (6) and (7) and the experimental data of the NiCrFeCoMn high-entropy alloy specimens. It can be seen that the results predicted by the Zerilli–Armstrong plastic model are in better agreement with the experimental results.

**Figure 7.** Comparison between the experimental results and the stress calculated by the Johnson–Cook and the Zerilli–Armstrong plastic models at 298K.

Figure 8 shows the yield strength vs. strain rate curves of the as-cast NiCrFeCoMn high-entropy alloy as a function of strain rate at ambient temperature. The strain-rate sensitivity is defined as follows:

$$\mathrm{Im} = \frac{d(\log \sigma)}{d(\log \dot{\varepsilon})}.\tag{8}$$

Notice that the yield strength distributes from 200 MPa to 800 MPa and the slope of the tangent for the curve is increasing with the increase of strain rates. Therefore, the as-cast NiCrFeCoMn high-entropy alloy has distinguished strain-rate sensitivity at high strain rates.

**Figure 8.** The yield strength vs. strain rate curves of the NiCrFeCoMn high-entropy alloy [30,31].

He et al. [32] studied the stress exponent (the reciprocal of strain rate sensitivity) of the NiCrFeCoMn high-entropy alloy deformed under strain rates less than 10−<sup>2</sup> s−1. He found that the stress exponent was in a positive relationship with the strain rates. When the strain rate ranged from 3.205 × <sup>10</sup>−<sup>5</sup> to 8.013 × <sup>10</sup>−<sup>4</sup> <sup>s</sup><sup>−</sup>1, the stress exponent increased simultaneously. Moon et al. [33] also studied the strain rate sensitivity at the elevated temperature and the cryogenic temperature, respectively. The results showed that the flow stress at 77 K was higher than that at room temperature. Therefore, the strain rate sensitivity of the flow stress at RT was higher than that at 77 K under strain rates less than 1 × <sup>10</sup>−<sup>2</sup> <sup>s</sup><sup>−</sup>1. The following formula can explain this phenomenon:

$$
\Delta V^\* = \sqrt{3} \text{k}T \frac{\partial \ln \dot{\varepsilon}}{\partial x},
\tag{9}
$$

where k is the Boltzmann constant and *V\** is activation volume.

The NiCrFeCoMn high-entropy alloy had positive activation volume at strain rates less than <sup>1</sup> × <sup>10</sup>−<sup>2</sup> <sup>s</sup>−1. However, when the specimen deformed under high strain rates (beyond 1 × 103 <sup>s</sup>−1), the deformation time was very short. There was not enough time available for thermal energy to help dislocations overcome the barriers. The NiCrFeCoMn high-entropy alloy had a different deformation mechanism at dynamic loading. According to the literature [22], the large jump in yield strength at high strain rates is probably due to the phonon drag effect on the motion of dislocations. The phonon drag phenomenon becomes very effective during plastic deformation at high strain rates [22]. A phonon is an elastic lattice vibration propagating in a crystal [26]. The viscous drag generating from the interaction of the dislocations has negative impacts on the deformation progress. The drag effects by the phonons can be ignored under quasi-static conditions for low dislocation velocities. However, high dislocation velocities under high strain rates would enhance the phonon drag effects greatly and lead to phonons scattering, which also hinders dislocation movement. Therefore, the dynamic deformations lead to a much higher strain rate dependence of the flow stress than those under quasi-static conditions.

#### *4.2. Mechanism for the Serration Behavior*

High-entropy alloys present serrations on the flow stress curves during the plastic deformation, often at a normal strain rate around 1 × <sup>10</sup>−<sup>4</sup> <sup>s</sup>−<sup>1</sup> or at low temperatures. Serration behavior exhibits in the stress-strain curves of the Al0.5CoCrCuFeNi high-entropy alloy at 7 K, 7.5 K, and 9 K at a strain rate of 4 × <sup>10</sup>−<sup>4</sup> <sup>s</sup>−<sup>1</sup> [19]. The Al5Cr12Fe35Mn28Ni20 high-entropy alloy exhibits typical serration behaviors at the elevated temperatures of 573 K and 673 K, with a strain rate of 1 × <sup>10</sup>−<sup>4</sup> <sup>s</sup>−<sup>1</sup> [34]. Several microstructure mechanisms are proposed to explain the serration behavior, e.g., the Portevin-Le Chatelier (PLC) effect. Serration behavior was supposed to be associated with Cottrell atmosphere interaction with moving dislocations, slip bands, and dynamic strain aging. In Figure 4, it can be seen that the stress-strain curves of the as-cast NiCrFeCoMn high-entropy alloy deformed at a strain rate of 1 × <sup>10</sup>−<sup>3</sup> <sup>s</sup>−<sup>1</sup> show no serrations, and marked serrations are present on the stress-strain curves of specimens deformed at strain rates above 900 s−1. With increasing strain rates, amplitudes of the serrations on the stress-strain curves become much larger. Investigations on the microstructure in the as-cast NiCrFeCoMn high-entropy alloy show that the high density dislocations and the deformation bands are generated in the specimens deformed at high strain rates, shown in Figures 5 and 6. The Portevin-Le Chatelier effect, i.e., Cottrell atmosphere interaction with moving a simple dislocation structure, may not be the main reason for the serration behavior of the as-cast NiCrFeCoMn high-entropy alloy becoming deformed at dynamic loadings. On the other hand, if the dynamic deformation becomes more serious and the value of the strain rate increases, deformation bands are generated in the as-cast NiCrFeCoMn high-entropy alloy and the amplitudes of the serrations on the stress-strain curves become much larger. Therefore, large amounts of the deformation bands, leading to the serration behaviors, play an important role for the mechanical properties of the high-entropy alloy.

#### **5. Conclusions**

The as-cast equiatomic NiCrFeCoMn high-entropy alloy has a simple FCC crystallographic structure. Serration behavior is observed in the stress- strain curves of the as-cast NiCrFeCoMn high-entropy alloy deformed at dynamic loadings. The yield strength of the high-entropy alloy, which distributes from 490 to 800 MPa, presents a positive relationship with the strain rates. The Johnson–Cook plastic model and the Zerilli–Armstrong plastic model of the NiCrFeCoMn high-entropy alloy are obtained. However, the results predicted by Zerilli–Armstrong correspond better with the experimental results. The serration behavior of the NiCrFeCoMn high-entropy alloy at a high strain rate is sensitive to the strain rates. The high density of deformation bands plays an important role in the deformation behavior and mechanical properties of the as-cast NiCrFeCoMn high-entropy alloy deformed at dynamic loadings.

**Author Contributions:** Conceptualization, B.W. and X.Z.; methodology, X.Y. and X.H.; validation, B.W. and Y.Z.; formal analysis, B.W., X.H., C.W. and X.Y.; investigation, B.W., X.H. and X.Y.; resources, B.W.; data curation, B.W., X.H., C.W. and X.Y.; writing—original draft preparation, B.W. and X.Y.; writing—review and editing, B.W. and Y.Z.; visualization, C.W.; supervision, B.W. and Y.Z..; project administration, B.W. and Y.Z.; funding acquisition, B.W. and Y.Z.

**Funding:** This research was funded by National Natural Science of China, grant number 51771231, and by State Key Laboratory of Powder Metallurgy, Central South University, Changsha, China, grant number 20181106.

**Acknowledgments:** Authors wish to express their most sincere gratitude to M.A. Meyers at University of California, San Diego, and Yong Liu at Central South University for their advice and help. Authors would like to express their sincere thanks to Yang Wang and Yu Wang at University of Science and Technology of China, and Xiang Zan at Hefei University of Technology for dynamic testing.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**


© 2018 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

### *Article* **Unconventional Deformation Behaviours of Nanoscaled High-Entropy Alloys**

**Yeqiang Bu 1, Shenyou Peng 2, Shiwei Wu 3, Yujie Wei 2, Gang Wang 3, Jiabin Liu 1,\* and Hongtao Wang 4,\***


Received: 7 September 2018; Accepted: 28 September 2018; Published: 11 October 2018

**Abstract:** The bulk high-entropy alloys (HEAs) exhibit similar deformation behaviours as traditional metals. These bulk behaviours are likely an averaging of the behaviours exhibited at the nanoscale. Herein, in situ atomic-scale observation of deformation behaviours in nanoscaled CoCrCuFeNi face-centred cubic (FCC) HEA was performed. The deformation behaviours of this nanoscaled FCC HEA (i.e., nanodisturbances and phase transformations) were distinct from those of nanoscaled traditional FCC metals and corresponding bulk HEA. First-principles calculations revealed an obvious fluctuation of the stacking fault energy and stability difference at the atomic scale in the HEA. The stability difference was highlighted only in the nanoscaled HEA and induced unconventional deformation behaviours. Our work suggests that the nanoscaled HEA may provide more chances to discover the long-expected essential distinction between the HEAs and traditional metals.

**Keywords:** nanoscaled high-entropy alloys; nanodisturbances; phase transformations; atomic-scale unstable

#### **1. Introduction**

Traditional alloys, such as steels and copper alloys, are fabricated based on one or two principle constituent elements. Yeh et al. [1] proposed the concept of high-entropy alloys (HEAs) that provides a novel basis to design new alloys. These HEAs are composed of multi-principle elements at equiatomic or near-equiatomic ratios, distinguishing them from traditional alloys. Consequently, the deformation behaviours of HEAs are believed to be different from traditional alloys [2], but no convincing experiments have yet been reported to show an essential distinction of the plastic deformation behaviours between HEAs and traditional metals. On the contrary, most previous results in bulk HEAs have shown similar scenarios with traditional metals, in which the plastic deformation is primarily carried by dislocations or twins [3–5]. The underlying mechanisms of macroscopic mechanical responses are essentially the collective behaviours of atomic-scale configurations. Therefore, the physical processes during alloy deformation exhibit no remarkable features if the salient atomic configuration details are essentially blurred and only the average effect can be measured and observed [6,7]. As expected, such macroscopic deformation behaviours are generally controlled by only a few key parameters such as the elastic moduli, the stable and unstable stacking fault energies (SFEs), the microstructure parameters and the temperature.

Close observation at the nanoscaled regime where discrete plasticity dominates may uncover the essential features that distinguish the HEAs from traditional alloys. Extensive investigations using in situ high-resolution transmission electron microscopy (HRTEM) have revealed some interesting deformation behaviours of nanoscaled pure metals [8–17]. Size-dependent behaviours have thus been uncovered such as the reversible deformation twinning and detwinning processes found in nanoscaled W samples [12] and the dislocation-originated stacking fault tetrahedra in nanoscaled Au samples [13]. Meanwhile, surface-mediated plasticity deformation behaviours have frequently been observed such as the partial dislocations emitted from the surface in sub-10 nm-sized Au [14] and the liquid-like deformations in sub-10 nm-sized Ag nanoparticles [15]. It would be interesting to determine whether nanoscaled HEAs behave similarly in the discrete plasticity regime, in contrast to their collective behaviour. The face-centred cubic (FCC) HEA CoCrCuFeNi is a typical HEA which was proposed in the earliest paper about the concept of HEA [1]. Numerous research papers related to HEAs were applied in CoCrCuFeNi from then. Therefore, CoCrCuFeNi was taken as a model HEA to reveal the obscured potential high-entropy effect on plastic deformation behaviours at the nanoscale in this work.

#### **2. Materials and Methods**

The CoCrCuFeNi button ingots were prepared by melting high-purity Co, Cr, Cu, Fe and Ni at equiatomic ratios in a vacuum arc furnace. During specimen preparation, the ingot was first sliced into a rod of 0.25 × 0.25 × 10 mm<sup>3</sup> in dimension. Firstly, we employed a tungsten carbide cutter to make a pair of gaps on the surface of the rod. Then, we used a plier to pull this rod apart in the long dimension and then led to the formation of fresh triangular nano-tips on the fractured surface, which served as the specimens for the in situ TEM experiments (Figure S1 in supplementary). A JEM-2100F field emission TEM equipped with a Nanofactory TEM-scanning tunnelling microscope (STM) sample holder was used in the in situ TEM experiments. Two fractured ends exhibiting nano-tips were mounted on the sample holder, with one at the fixed end of the holder and the other at the piezo-manipulator end (Figure S2 in Supplementary Materials). One nano-tip could thus be driven to touch the other nano-tip on the counter fractured surface, guided by the piezo-actuated nanoscaled manipulator. A nanoscaled welding joint 10–30 nm in size could then be formed instantly via pulsed joule heating. Uniaxial tensile stress could be applied by step-by-step retraction via the nano-manipulator. This method essentially furnished the fabrication, mechanical testing and easily atomic-resolution observation of the nanoscaled HEA. The structure of the nanoscaled HEA fabricated by this method is shown in Figure S3. The crystalline interplanar spacing of the (111) plane is 2.07 Å in the obtained nanoscaled HEA, i.e., its lattice parameter is 3.59 Å. This is consistent with the reported lattice parameter (3.579 Å) of the bulk CoCrCuFeNi [1,18]. This shows that the structure of the nanoscaled HEA is not affected by the process of preparation. Meanwhile, there is no oxide layer on the surface of the nano-tips before the fabrication of the nanoscaled samples, as shown in Figure S4, which indicates that there is no impact of the oxide layer on the deformation behaviours (These nano-tips were exposed to air for only a few minutes.). The energy dispersion spectrum (EDS) mapping results are shown in Figure S5, which indicates that the constituent elements distribute uniformly in the nao-tip (EDS hardly displays the details of composition in several-atoms scale, even the atomic resolution EDS only gives the statistical chemical composition of each atom column in planar view).

The first-principles calculations were performed using the Vienna ab initio simulation package based on density functional theory, wherein all of the HEA samples were relaxed to the energy precision of 0.01 meV. The SFE was calculated with *E*sf = (*E*fault − *E*perfect)/*A*, where *E*perfect and *E*fault are the free energy of perfect and faulted structures, respectively; and *A* is the area of each layer. The structure with one layer fault was obtained by a rigid displacement between two adjacent layers, where the magnitude was equal to that of the Burgers vector, *bp* = 1/6<112>.

#### **3. Results**

#### *3.1. Nanodisturbances*

Figure 1a displays a nanoscaled HEA during in situ straining in the TEM (Vedio S1 in Supplementary Materials), where the loading direction (LD) is around [311] and the beam zone is [011]. The deformed nanoscaled HEA contains several dislocations. The deformation of this nanoscaled HEA is closely related to the behaviours of dislocations. Dislocation cores are labelled in Figure 1 by "⊥", and an enlarged image of a dislocation core is shown in Figure 1b. According to the analysis of the Burgers circuit, these are full dislocations with *b* = 1/2[101]. Some of the dislocations (circled in Figure 1a) appear in pairs and are thus dislocation dipoles, seen in an enlarged image in Figure 1c. Using in situ HRTEM, we could dynamically observe the evolution of these dislocation dipoles. Figure 1d–f show inverse fast Fourier transform (IFFT) images of the one-dimensional {111} plane fringes in the area around a dislocation dipole. At *t* = 20 s, the circled area exhibited a distorted lattice which was induced by high stress in the nanoscaled HEA. At *t* = 20.5 s, a dislocation dipole pair was produced in the distorted area. Under stress driving, the dislocation dipole was observed to expand along the (111) slip plane at *t* = 21 s. This deformation mode could be called nanodisturbance, which has been proposed on body-centred cubic (BCC) "gum metal" by Gutkin et al. [19]. This kind of dislocation dipole nucleation and expansion could act as a mechanism of dislocation multiplication during the deformation process of the nanoscaled HEA.

**Figure 1.** The nanodisturbances in the nanoscaled high-entropy alloy (HEA). (**a**) deformed nanoscaled HEA sample containing several dislocations (Beam // [011], loading direction (LD) ≈ [311]); (**b**) analysis of the Burgers vector; (**c**) enlarged image of the dislocation dipole; (**d**–**f**) Inverse fast Fourier transform (IFFT) images of the one-dimensional {111} plane fringes showing the formation and expansion of the dislocation dipole.

#### *3.2. Phase Transformation*

Figure 2a–c shows the tensile process of a nanoscaled HEA sample (Video S2 in Supplementary Materials), where the LD is around the [311] and beam zone is [011]. Figure 2e–g show enlarged images corresponding to the areas in red squares in Figure 2a–c, respectively. As shown in Figure 2e, the crystal lattice exhibits clear characteristics of a FCC structure; i.e., the angle between two close-packed planes in the {111} family is 70.5◦. As the nanoscaled HEA continues to be stretched, the angles between two {111} planes reduce to 64◦ at *t* = 104 s (Figure 2f), and then to 60◦ at *t* = 384 s (Figure 2g). The angle

of 64◦ indicates that the lattice structure deviates from the original FCC structure notably. The angle of 60◦ represents a typical BCC structure with a [111] zone. Figure 2d plots the angles between the close-packed planes in the red square areas of Figure 2a–c as a function of time under stress, where the angle exhibits a slow and successive transition from ~70◦ to ~60◦. The change of the angle corresponds to the transition that the initial FCC lattice transforms to BCC lattice during in situ tension. The resultant orientation relationship between the FCC and BCC agrees with the K–S relationship, i.e., [011]FCC // [111]BCC, (111)FCC // (110)BCC. The transformation process could be explained by K–S model. The FCC lattice sheared 19.5◦ along the <112>FCC direction on the {111}FCC plane and sheared 10.5◦ along the <110>FCC direction on the {112}FCC plane. Therefore, the stress-induced FCC→BCC transition realized and the BCC lattice formed. The slow and successive transformation was believed to relate with the high lattice friction in HEAs [20–22], the transformation dislocations slip and lattice shear more slowly in HEAs compared to the traditional abrupt Martensitic transformation. Such stress-induced FCC→BCC transition has a good reproducibility, and Figure S6 shows another example.

**Figure 2.** The phase transformation from face-centred cubic (FCC) to body-centred cubic (BCC) in the nanoscaled HEA. (**a**–**c**) elongation process of the nanoscaled sample (Beam // [011], LD ≈ [311]); (**d**) variation of the angle between two close-packed planes in the red square area during in situ tension; (**e**–**g**) High-resolution transmission electron microscopy (HRTEM) images of the red square zones in (**a**–**c**), respectively.

#### *3.3. Fluctuation of Stacking Fault Energy*

The SFE is one of the most significant parameters determining the deformation behaviours of alloys, and is closely related to phase transformation and structural stability. Herein, we employed first-principles calculations to determine the SFE of this HEA, and the results of 52 independent SFE calculations are shown in Figure 3a. These results exhibit a fluctuant distribution of the SFE in HEA, where the SFE value covers a wide range and even is negative. The negative SFE values indicate that some of the HEA atomic configurations are unstable. The instability seems to be strongly correlated with the non-uniform distribution of atoms. As an example, we analyze a typical calculation sample illustrated in Figure 3b and find that the atoms in the dashed boxes (also seen in Figure 3c,d) are not a uniform distribution of all kinds of elements in CoCrFeNiCu. Some areas have more Co and Cr atoms and less Cu atoms (Figure 3c), which leads to a negative SFE of −24 mJ/m2. Some areas have more Cu atoms and less Co and Cr atoms (Figure 3d), which leads to a SFE as high as 109 mJ/m2. We conclude

that the elemental inhomogeneity at the atomic scale leads to SFE difference in local, and the fluctuant distribution of the SFE induces a stability difference at atomic-scale. However, such fluctuant SFE is averaged in the bulk HEAs, the atomic-scale stability difference is also blurred. The small-scaled sample size may highlight such atomic details.

**Figure 3.** The fluctuant distribution of the stacking fault energies (SFEs). (**a**) the SFEs of 52 independent calculations; (**b**) typical structure of a calculation sample; (**c**) atomic configuration producing a negative SFE of <sup>−</sup>24 mJ/m<sup>2</sup> from the upper dashed box in (**b**); (**d**) atomic configuration producing a high SFE of 109 mJ/m2 from the lower dashed box in (**b**).

#### **4. Discussion**

In such in situ HRTEM experiments, the effects of electron beam on the nanoscaled samples should be verified and eliminated. Although the CoCrCuFeNi HEA was verified to be stable under the severe electron irradiation [23], all in situ experiments are still performed under the weak electron beam for minimizing the effect of electron irradiation. Meanwhile, our verification experiments (see Supplementary Figure S6) and theoretical estimation (see Appendix A) confirm the negligible influence of electron irradiation on the in situ experiments.

Nanodisturbances could be an effective mechanism for dislocation multiplication, and the process of nanodisturbances evolving into dislocation dipoles has been observed in the BCC gum metal [19,24]. However, there has not yet been any experimental observation of this novel dislocation-generating mechanism for FCC structure metals. Some theoretical calculations have investigated the nanodisturbance phenomenon in Au and Cu nanowires [25,26], indicating that the nanodisturbance deformation mode could dominate over traditional dislocation generation at high stresses and 0 K. Obviously, the temperature condition of 0 K was not satisfied in this work. However, we still observed the nanodisturbance deformation mode in FCC HEAs for the first time. The high-level stress in nanoscaled samples and the intrinsic features of HEAs both play significant roles on the occurrence of the nanodisturbances. On the one hand, we observed nanodisturbance in the nanoscaled FCC HEA at the relative loose condition (at room tempreture). On the other hand, nanodisturbance is hardly observed in the bulk HEAs because of the relative low stress level in the bulk HEA. Furthermore, we believe that it is the nanoscaled size that triggers the emergence of the intrinsic features of the HEA.

The previous works showed that there are no phase transformation during the deformation process [27]. However, herein we observed the FCC→BCC transformation in nanoscaled CoCrCuFeNi HEA. A similar transformation from FCC to body-centered tetragonal (BCT) has been observed in nanoscaled fractured Au [14,28], where the phase transformation therein was considered to be stimulated by the relaxation of surface stress. The surface stress has an inverse relationship with the sample size, and thus surface stress of a nanoscaled sample is sufficiently high to stimulate a transformation. However, in this study, the phase transformation in the nanoscaled HEA occurred before fracturing, and thus the surface stress had not been completely released. Therefore, the cause of phase transformation in the nanoscaled HEA is not solely surface stress, but an important role may also be played by the unstable nature of nanoscaled HEA. The combined effect of high surface stress and an unstable nature, therefore, stimulates the occurrence of the phase transformation.

Both nanodisturbances and phase transformations are not regular deformation behaviours in bulk CoCrCuFeNi HEA. These unconventional deformation behaviours observed in nanoscaled HEA are believed to be related to atomic configuration details present at such a small scale. Our first-principles calculations could well illuminate the unconventional deformation behaviours in such nanscaled HEAs. The first-principles results show that the elemental inhomogeneity at the atomic scale leads to SFE difference in local, and the fluctuant distribution of the SFE induces a stability difference of atomic-scale HEA. However, the entire structure averaged in the bulk HEA possesses a constant SFE. The stability difference can be outlined and plays a dominant role only when the sample dimensions reach the nanoscale. At that time, nanoscaled HEA exhibits deformation behaviours different with bulk counterparts. We further speculate that the nanoscaled HEA provides more of a chance to discover the long-expected essential distinction between the HEAs and traditional metals.

Besides the FCC HEAs, the BCC and hexagonal close-packed (HCP) HEAs may also possess some distinct characteristic physical properties but are blurred in bulk. It is worth investigating the other nanoscaled HEAs with various structures in future research to reveal the essential distinction between the HEAs and traditional metals.

#### **5. Conclusions**

In summary, we employed in situ HRTEM to investigate the deformation behaviours of nanoscaled HEA. Unconventional deformation behaviours (i.e., nanodisturbances and phase transformations) were observed in the nanoscaled HEA. The first-principles calculations revealed obvious fluctuant distribution of the SFE at atomic scale, which was resulted from the elemental inhomogeneity. The SFE fluctuation leaded to stability difference at the atomic scale, which plays a dominant role in the deformation of the nanoscale sample but tiny roles in bulk counterparts. The nanoscaled HEA provided a chance to highlight the stability difference and therefore exhibited unconventional deformation behaviours. Our investigations reveal some HEA features and are significant for understanding the nature of HEA.

**Supplementary Materials:** The following are available online at http://www.mdpi.com/1099-4300/20/10/ 778/s1, Figure S1: The nano-tips on the fractured surface, Figure S2: Two teared parts with nano-tips are mounted on a Nanofactory transmission electron microscope (TEM)-scanning tunnelling microscope (STM) TEM holder, Figure S3: The structure keep unchanged during the preparation process, Figure S4: The high-resolution TEM image of the nano-tip, Figure S5: The verification experiments about the effects of electron irradiation, Video S1: Nanodisturbances deformation mode in the nanoscaled HEA, Video S2: Phase transformation in the nanoscaled HEA.

**Author Contributions:** H.W. and J.L. initiated this research project. Y.B. performed the TEM experiments. S.P. and Y.W. performed the first-principles calculations. S.W. and G.W. synthesized the initial samples. H.W., J.L. and Y.B. analyzed the data and wrote the manuscript.

**Funding:** This research was funded by the National Natural Science Foundation of China (Nos. 11572281, 11725210 and 11672355), and the Fundamental Research Funds for the Central Universities (No. 2018XZZX001-05), and the Zhejiang Provincial Science & Technology Program of China (2017C31076).

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **Appendix A. Theoretical Estimation of the Effects of Electron Irradiation**

The effects of electron irradiation could be divided into two aspects, i.e., knock-on displacement and heating effect. We could not observe any surface diffusion or atoms hopping occurs in the absence of external force. Therefore, the knock-on displacement driven by electron irradiation has no influence on the nanoscaled samples. For the heating effect, the estimation was based on the Fisher's model [29] and the Bethe–Bloch equation [30]. Herein, the nanoscaled samples are all with a certain thickness (~10 nm), and exposed at the weak beam intensity. Fisher has proposed a model to estimate the temperature increase induced by electron irradiation [29]:

$$
\Delta T = \frac{I}{4\pi K e} (\frac{\Delta E}{d}) (1 + 2 \ln \frac{b}{r\_0}),
\tag{1}
$$

where Δ*T* is the maximum temperature rise heated by electron irradiation, *K* is the thermal conductivity of the sample, *I* is the beam current, Δ*E* is the total energy loss per electron in a sample of thickness *d*, *b* is the radius of the heat sink, and *r0* is the beam radius. The term <sup>Δ</sup>*<sup>E</sup> <sup>d</sup>* could be approximately equal to d*E* <sup>d</sup>*<sup>x</sup>* (where *x* is the position in thickness), which could be calculated by the Bethe–Bloch equation [30], as follows:

$$-\frac{d\underline{E}}{dx} = \frac{2\pi Z\rho (e^2/4\pi a\_0)^2}{m^2} \left\{ \ln \left[ \frac{E(E+mz^2)\beta^2}{2l\_e^2 m^2} \right] + (1-\beta^2) - (1-\sqrt{1-\beta^2}+\beta^2)\ln 2 + \frac{1}{8}(1-\sqrt{1-\beta^2})^2 \right\},\tag{2}$$

where *Z* is the atomic number of the samples, *ρ* is the atomic density, *ε*<sup>0</sup> is the vacuum dielectric constant, *m* is the electron rest mass, *ν* is the electron velocity, *c* is the speed of light, *E* is the electron energy, *Ie* is the average excitation energy of electrons in the target, and *β* = *ν*/*c*.

We take the *Z*ave = (*Z*Co+ *Z*Cr+ *Z*Cu+ *Z*Fe+ *Z*Ni)/5 = 26.8 as the atomic number of this HEA samples containing Co, Cr, Cu, Fe and Ni five principle elements. *ρ* = *ρ*mass/((*m*Co/atom+ *m*Cr/atom+ *<sup>m</sup>*Cu/atom <sup>+</sup> *<sup>m</sup>*Fe/atom+ *<sup>m</sup>*Ni/atom)/5) = 7.231 × 1027 m−3. The thermal conductivity of this HEA is estimated to be *<sup>K</sup>*= 16.2 W·m−1·K−1. Given *<sup>I</sup>* = 4.8 nA, *<sup>e</sup>* = 1.6 × <sup>10</sup>−<sup>19</sup> C, *<sup>b</sup>* = 1.5 mm, *<sup>r</sup>*<sup>0</sup> = 200 nm, *<sup>ε</sup>*<sup>0</sup> = 8.85 × <sup>10</sup>−<sup>12</sup> <sup>F</sup>·m<sup>−</sup>1, *<sup>m</sup>* = 9.3 × <sup>10</sup>−<sup>31</sup> kg, *<sup>ν</sup>* = 2.0837×10<sup>8</sup> <sup>m</sup>·s−1, *<sup>c</sup>* = 3.0 ×10<sup>8</sup> <sup>m</sup>·s−1, *<sup>E</sup>* = 200 KeV, *Ie* = 8.8*Z* = 235.8 eV, the estimated results of the maximum temperature increase of the HEA sample is 0.061 K. Therefore, the heating effect induced by electron irradiation could be neglected.

#### **References**


© 2018 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

### *Article* **Small-Scale Plastic Deformation of Nanocrystalline High Entropy Alloy**

**Sanghita Mridha 1, Mageshwari Komarasamy 1, Sanjit Bhowmick 2, Rajiv S. Mishra <sup>1</sup> and Sundeep Mukherjee 1,\***


**\*** Correspondence: sundeep.mukherjee@unt.edu; Tel.: +1-940-565-4170

Received: 22 October 2018; Accepted: 16 November 2018; Published: 20 November 2018

**Abstract:** High entropy alloys (HEAs) have attracted widespread interest due to their unique properties at many different length-scales. Here, we report the fabrication of nanocrystalline (NC) Al0.1CoCrFeNi high entropy alloy and subsequent small-scale plastic deformation behavior via nano-pillar compression tests. Exceptional strength was realized for the NC HEA compared to pure Ni of similar grain sizes. Grain boundary mediated deformation mechanisms led to high strain rate sensitivity of flow stress in the nanocrystalline HEA.

**Keywords:** nanocrystalline materials; high entropy alloy; sputtering; deformation and fracture; strain rate sensitivity

#### **1. Introduction**

High entropy alloys (HEAs) represent an alloy design paradigm of combining five or more elements in equiatomic or near-equiatomic proportions [1,2]. In certain compositions, high configurational entropy suppresses intermetallic compound formation and leads to single-phase solid solution [3]. HEAs have attracted widespread interest due to their intriguing physical and mechanical properties [3]. Some of the appealing properties include exceptional ductility [3], outstanding thermal stability [4], and cryogenic fracture toughness [5]. Bulk of the research on HEAs have focused on alloy development [3], phase stability [6], and mechanical behavior of coarse grained (CG) and fine-grained systems [3,7,8]. But there are limited reports on nanocrystalline (NC) HEAs and their small-scale deformation behavior [9,10]. NC metals typically show very high strength [11] and good fatigue resistance [12]. Body-centered cubic NbMoTaW refractory NC HEA exhibited exceptional strength at small scales and ductility [9]. Furthermore, NC HEA retained yield strength (YS) of over 5 GPa up to 600 ◦C denoting an exceptional nano-structural stability [10]. Similar studies for face-centered cubic (FCC) systems could provide insights into their deformation mechanisms at reduced length-scale and pave the way for new application domains towards low-cost, durable, ductile, and strong FCC HEAs.

Al0.1CoCrFeNi HEA is a canonical example of a FCC single phase multi-principal element alloy whose mechanical properties have been widely reported. Komarasamy et al. [13] examined the work hardening mechanisms in coarse-grained (CG) (a few mm) and fine-grained (FG) (~3–14 μm) Al0.1CoCrFeNi HEA and concluded that both the conditions exhibited deformation twinning mediated plasticity. Following that, Choudhuri et al. [14] investigated the plastic deformation mechanisms of the same alloy in CG and FG conditions using transmission electron microscopy. After quasi-static tensile testing, both the microstructures exhibited nanoscale (~2 nm) twins signifying twinning assisted plastic deformation. Wu et al. [15] also noted a similar behavior with deformation twins in both CG and FG materials. Furthermore, an increase in average twin spacing and a reduction in twin thickness was

observed in FG condition as compared with CG material. Kumar et al. [16] examined the high strain rate compression behavior of CG Al0.1CoCrFeNi HEA and observed a similar work hardening behavior as quasi-static compression. Yu et al. [17] investigated the deformation mechanism in Al0.1CoCrFeNi HEA subjected to high pressure torsion (HPT). Deformation via dislocation slip was noted at low strains while deformation twinning was activated at large plastic strains. Furthermore, a large strain rate sensitivity of 0.035 was obtained signaling grain boundary related deformation mechanism in the HPT condition. In Al0.1CoCrFeNi HEA, Feng et al. [18] introduced high density of nano-twins and stacking faults, and investigated the mechanical behavior at small length scales. The pillars exhibited a compressive strength of 4.0 GPa with 15% compressive ductility. The exceptional mechanical properties were attributed to the stability of stacking faults and its effective hindrance to dislocation motion.

In this paper, we report on the nano-mechanical behavior of NC Al0.1CoCrFeNi HEA (average grain size ~40 nm) synthesized using magnetron sputtering technique. Stress–strain response was obtained by nano-pillar compression, concurrent with direct observation of their deformation behavior inside a scanning electron microscope (SEM). Different strain rates were used for strain rate sensitivity (*m*) analysis.

#### **2. Materials and Methods**

Target alloy of composition Al0.1CoCrFeNi was prepared using high purity elements (99.99%) and arc melted in an Ar atmosphere. A thin film of the alloy was deposited by magnetron sputtering technique (AJA International, Scituate, MA, USA) on a silicon substrate at room temperature (RT), with the base and process pressure maintained at ~3 × <sup>10</sup>−<sup>6</sup> torr, and ~5 × <sup>10</sup>−<sup>3</sup> torr, respectively. An Ar atmosphere (flow rate 10 sccm) was used and applied power was 75 W. Nano-pillars were synthesized by milling the thin film using Focused Ion Beam (FIB) (FEI). A concentric circular pattern was used to mill out the nano-pillar. Gallium ion beam with current of 10 pA, and operating voltage of 30 kV was used for milling and the pillars had an average diameter of ~450 nm. ImageJ software was used to determine the grain size. In situ compression tests were performed on the nano-pillars with a SEM equipped with PicoIndenter PI 85 (Bruker Nano Surfaces, Minneapolis, MN, USA) using a 2 μm diameter diamond flat punch. Uniaxial compression tests were conducted at strain rates of 1.2 × <sup>10</sup><sup>−</sup>1, 1.9 × <sup>10</sup><sup>−</sup>2, and 7.5 × <sup>10</sup>−<sup>3</sup> <sup>s</sup><sup>−</sup>1.

#### **3. Results and Discussion**

Figure 1a,b shows the X-ray diffraction (XRD) patterns for Al0.1CoCrFeNi HEA thin film and target, respectively, demonstrating a single phase FCC structure. SEM image of the thin film (Figure 1c) shows uniformly distributed nano-sized grains with an average grain size of ~40 ± 5 nm. SEM images of NC nano-pillar HEA before and after compression tests are shown in Figure 2. The pillar diameter was around 450 nm (Figure 2a). Figure 2b shows the nano-pillar with the diamond flat punch just before the compression test. Uniform plastic deformation of the nano-pillar with no evidence of buckling was observed (Figure 2c). Furthermore, crack propagation parallel to the loading direction can be observed in Figure 2c.

**Figure 1.** X-Ray diffraction (XRD) patterns of (**a**) Al0.1CoCrFeNi thin film and (**b**) Al0.1CoCrFeNi target, showing peaks corresponding to the face-centered cubic (FCC) phase; (**c**) high-magnification scanning electron microscope (SEM) image of the thin film showing nano-sized grains with an average grain size of ~40 ± 5 nm.

**Figure 2.** SEM images of the (**a**) nano-pillar with diameter 450 nm (**b**) diamond punch and the nano-pillar, and (**c**) nano-pillar after the compression test.

Compressive engineering stress-strain curve at a strain rate of 7.5 × <sup>10</sup>−<sup>3</sup> <sup>s</sup>−<sup>1</sup> is shown in Figure 3a. Inset images of the pillar at various deformation intervals clearly show that the plastic deformation was uniform without buckling of the nano-pillar. YS of nano-pillar Al0.1CoCrFeNi HEA at 7.5 × <sup>10</sup>−<sup>3</sup> <sup>s</sup>−<sup>1</sup> strain rate was ~3829 MPa. YS of the CG material of the same composition was ~190 MPa indicating a 20 fold increase in strength for the NC HEA. Al0.1CoCrFeNi forms a single phase alloy with FCC crystal structure without any secondary phases. Therefore, only grain size strengthening contribution was investigated. To that end, YS versus *d*−1/2 correlation (where *d* represents grain size) for various HEAs, pure Ni, and the current study are shown in Figure 3b. Hall–Petch relation for various HEAs and Ni are shown in Figure 3b, based on the following equation:

$$
\sigma\_{YS} = \sigma\_{\bullet} + kd^{-1/2} \tag{1}
$$

where, *σ*<sup>o</sup> is the strength of the material for infinitely large grain size also called lattice friction stress, *k* is the Hall-Petch coefficient, and *d* is the grain size. The Hall–Petch equation with coefficients, *σ*<sup>o</sup> and *k*, for all the conditions are given in the bottom inset of Figure 3b. Hall–Petch coefficients based on two independent HEA investigations were used to calculate the expected grain size strengthening

contribution for the current condition with ~50 nm grain size. The calculated YS based on Otto et al.'s [19] and Nilesh et al.'s [8] Hall–Petch coefficients were 2255 and 1849 MPa, respectively. Remarkably, the obtained YS in the current investigation was ~1500 MPa higher than the predicted strength values. This may be attributed to the shift in deformation mode from dislocation-controlled to grain boundary mediated plastic deformation. For the same grain size, HEAs exhibited strength values two fold higher than that of pure Ni [20]. In addition to Hall–Petch and grain boundary mediated deformation mechanism, size effect may also dominate the YS of NC Al0.1CoCrFeNi HEA, which is out of the scope of the current paper. In CoCrCuFeNi HEA, Zhang et al. [21] investigated size-dependent YS based on micro-/nano-pillar uniaxial compression tests and noted that the HEA did exhibit size-dependent mechanical properties.

**Figure 3.** (**a**) Compressive engineering stress–strain curve at strain rate of 7.5 <sup>×</sup> <sup>10</sup>−<sup>3</sup> <sup>s</sup>−<sup>1</sup> showing YS of 3829 MPa. Insets show the *in situ* image of pillar at various stages; (**b**) YS vs. *d*−1/2 plot with Hall–Petch equation fit for Ni and various HEAs.

Effect of strain rate on yield strength and strain rate sensitivity calculation are shown in Figure 4a,b, respectively. Stress-strain curves for 1.9 × <sup>10</sup>−<sup>2</sup> <sup>s</sup>−<sup>1</sup> and 1.2 × <sup>10</sup>−<sup>1</sup> <sup>s</sup>−<sup>1</sup> strain rates were shifted along the strain axis for a clear representation. With the increase in strain rate, yield strength of the NC nano-pillar Al0.1CoCrFeNi HEA increased. This indicates a positive strain rate sensitivity of flow stress at room temperature which was observed in other investigations as well [22]. Strain rate sensitivity (*m*) is defined as:

$$m = \frac{\partial \ln \sigma}{\partial \ln \varepsilon} \tag{2}$$

where, σ is the flow stress and ε is the strain rate. Following this equation, slope of 1% flow stress and strain rate in logarithmic scale yielded a strain rate sensitivity of 0.08 as presented in Figure 4b. Comparison with *m* values reported in literature for other materials is given as inset in Figure 4b. As can be clearly seen, *m* value increases with the reduction in grain size. For example, CG and NC copper exhibited *m* values of 0.009 and 0.06, respectively [11]. Furthermore, *m* value of CG HEA was higher than CG conventional metals/alloys, which was attributed to fluctuating lattice energy controlled deformation mechanism. In the current investigation, due to the nano-scale grain size, there is high possibility that dislocation controlled processes were suppressed and grain-boundary mediated processes were activated. Furthermore, *m* value of 0.08 for NC HEA would translate into lower apparent activation volume of dislocation as compared to NC copper with *m* value of 0.06. Therefore, in addition to differences in CG material due to lattice distortion controlled dislocation activity, current results suggest that grain boundary mediated plastic deformation was influenced by the inherent lattice distortion of HEA.

**Figure 4.** (**a**) Compressive engineering stress-strain plot at different strain rates showing the increase in tensile strength with increase in strain rate; (**b**) flow stress at 1% offset strain versus strain rate plot in logarithmic scale to calculate strain rate sensitivity (*m*).

An important implication of the current investigation is the viability of nano-crystalline HEAs for use in high-strength applications without expensive refractory elements such as Mo, Ta, W, and Nb. In fact, body centered cubic NbMoTaW HEA pillar of 1 μm diameter exhibited yield strength similar to the current Al0.1CoCrFeNi face centered cubic nano-pillar [9]. Therefore, the current nanostructured alloy is a cost-effective alternative towards achieving ultra-strong and ductile wires for small-scale applications.

#### **4. Conclusions**

In conclusion, exceptional strength was seen for nano-crystalline Al0.1CoCrFeNi HEA similar to refractory HEAs of comparable length scales. This was attributed to grain boundary mediated plastic deformation processes. The strain rate sensitivity was higher than conventional NC material implying an even lower activation volume of dislocations.

**Author Contributions:** Conceptualization, S.M.; Formal analysis, S.B. and R.S.M.; Investigation, S.M. and M.K.; Supervision, S.M.; Writing—original draft, S.M., M.K. and S.M.

**Funding:** This research received no external funding.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**


© 2018 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

*Article*
