**Applications of High-Pressure Technology for High-Entropy Alloys: A Review**

**Wanqing Dong 1, Zheng Zhou 1, Mengdi Zhang 1, Yimo Ma 1, Pengfei Yu 1,\*, Peter K. Liaw <sup>2</sup> and Gong Li 1,\***


Received: 23 June 2019; Accepted: 26 July 2019; Published: 8 August 2019

**Abstract:** High-entropy alloys are a new type of material developed in recent years. It breaks the traditional alloy-design conventions and has many excellent properties. High-pressure treatment is an effective means to change the structures and properties of metal materials. The pressure can effectively vary the distance and interaction between molecules or atoms, so as to change the bonding mode, and form high-pressure phases. These new material states often have different structures and characteristics, compared to untreated metal materials. At present, high-pressure technology is an effective method to prepare alloys with unique properties, and there are many techniques that can achieve high pressures. The most commonly used methods include high-pressure torsion, large cavity presses and diamond-anvil-cell presses. The materials show many unique properties under high pressures which do not exist under normal conditions, providing a new approach for the in-depth study of materials. In this paper, high-pressure (HP) technologies applied to high-entropy alloys (HEAs) are reviewed, and some possible ways to develop good properties of HEAs using HP as fabrication are introduced. Moreover, the studies of HEAs under high pressures are summarized, in order to deepen the basic understanding of HEAs under high pressures, which provides the theoretical basis for the application of high-entropy alloys.

**Keywords:** high-entropy; high pressure; high pressure torsion; diamond anvil cells

#### **1. Introduction**

Emerging in recent years, high-entropy alloys (HEAs) are newly developed alloys, with many outstanding properties which have broken the design concept of traditional alloys, and have a variety of principal elements and special crystal structures. Although the composition of the high-entropy alloys varies, the phase composition is very simple. Usually only one or two solid-solution phases are detected by X-ray diffraction, and rare intermetallic compounds are formed. Since there is no principal element in HEAs, the performance of the alloy is affected by the combined influence of the constituent elements. As a result, the performance of HEAs is somewhat unpredictable. The structure and stability of the alloy has crucial influence on its properties. Therefore, it is necessary to study the structures, and the stability of the structure of the different HEAs under various conditions. HEAs are solid solutions composed of many elements which can maintain their stable structures under normal temperatures and atmospheric pressure. However, during long-term or high-temperature annealing, the new phase will be created, which will affect the performance of the alloy [1–3]. The pressure, temperature, and chemical composition are the basic thermodynamic factors that determine the state of a substance. The use and control of temperature and chemical composition are almost synchronous with the development of human civilization. However, due to the limitations of technical conditions, the application of pressure has just begun. High-pressure technology is a relatively-young and emerging discipline, but most of the condensed matter in the universe is under high pressure. The research on high pressures relies heavily on the experimental techniques, and each progression of the method has led to a significant expansion of our basic understanding of material behavior under high pressures. The progress in high-pressure experimental technology has directly promoted the development of the high-pressure science and provided an advanced method for frontier subjects. It has become an important field in modern scientific research. There are few studies on the structures of HEAs under high pressures. Interactions within the materials will change under high pressures and induce the generation of high-pressure phase transitions, thus becoming a new material with special properties. The materials will show many unique properties under high pressures, ones which do not exist under normal conditions, providing a subject for the in-depth study of materials. The behavior of HEAs under high pressures is a potential research direction for the future.

#### **2. High-Entropy Alloys (HEAs)**

#### *2.1. Concept*

HEAs are kinds of alloys developed in recent years. HEAs are loosely defined as solid-solution alloys that contain more than five principal elements in an equal or near-equal atomic percent (at. %) [4]. These kinds of alloys were defined by Yeh et al. [4] in 2004 as HEAs, and in the same year named by Cantor [5] et al. as the multi-component alloy. The concept of HEAs is based on the development of bulk amorphous alloys in the 1990s, when people were looking for alloys with ultra-high glass-forming ability. According to the well-known confusion principle, the more components of the alloy, the higher the chaos of the liquid alloy. That is, with a high entropy of mixing, it is easy for the alloy to retain the structure of the melt, thus forming a disordered amorphous structure. Figure 1 shows a panoramic view of the materials that humans have used over the past 10,000 years. The picture gives different types of materials, from ceramics to metals, polymers, and most recently composites, the same age at which HEA production is noted. With an increasing understanding of HEAs, the requirements for the content of each element and the number of elements have gradually loosened in the definition of HEA. At present, quaternary equiatomic or part of non-equiatomic quaternary alloys are also defined as HEAs [6–8], and some literature now refers to multi-component alloys containing small amounts of intermetallic compounds as HEAs [9,10].

**Figure 1.** Historical evolution of engineering materials [11].

#### *2.2. Four Core E*ff*ects*

Yeh [12] summarized the four core effects in HEAs. Those were: (1) High-entropy effects; (2) sluggish diffusion; (3) severe lattice distortion; and (4) cocktail effects.

#### 2.2.1. High-Entropy Effects

This is the most important characteristic of the HEA; i.e., the formation of intermetallic compounds or complex phases is inhibited due to the high entropy of the HEAs, tending to form solid-solution phases [4,5,13,14]. In other words, the high entropy produced by multiple principal elements can inhibit the generation of intermetallic compounds.

#### 2.2.2. Sluggish Diffusion

In traditional alloys, there is a high probability that a less prevalent atom is surrounded by the main element's atoms, which are called the solute atoms. It means that during the atom diffusion, the interaction with the surrounding atoms is essentially constant. However, the atoms that surround an atom in the HEAs are diverse, leading to different activation energies required by various lattice sites during atomic diffusion in HEAs [15]. In comparison with traditional alloys, more energy is needed during the atomic diffusion for HEAs. Although diffusion activation varies greatly with the elements, the trend of the increased diffusion activation energy can be obtained from low entropy to high entropy.

The sluggish diffusion theoretically suppresses the grain growth in HEAs [4] and explains the formation of nano-sized precipitates. On the other hand, the sluggish diffusion increases the phase stability of HEAs, especially at high temperatures; the phase stability and high-temperature strength even exceed some Ni-based superalloys [16].

#### 2.2.3. Severe Lattice Distortion

In HEAs, each atom is surrounded by other kinds of atoms. Due to different compositional atomic sizes in HEAs, that feature can lead to the severe lattice distortion. The severe lattice distortion affects the mechanical, physical, and chemical properties of the material. As shown in Figure 2, the probability that each element in the HEA occupies the position of the lattice is the same, and the atomic radii of different atoms will cause severe lattice distortion.

**Figure 2.** Schematic illustration of lattice distortion.

#### 2.2.4. Cocktail Effect

The cocktail effect for alloys was first mentioned by Ranganathan [17] to describe some basic characteristics of the elements affecting the properties of alloys. However, for the HEAs that break through the traditional design concept, the cocktail effect does not mean that the performance of the alloys is simply a superposition of the properties of each component. There are also interactions between different elements that eventually lead to a composite effect in HEAs [12].

#### *2.3. Research Status*

Due to the four core effects of HEAs, many HEAs exhibit mechanical, physical, and chemical properties that are superior to pure metals and other alloys. Many research groups have carried out much research on HEAs. At present, several mature high-entropy systems have been formed, including AlxCoCrCuFeNi [4], CoCrFeMnNi [5], WNbMoTaV [8], and GdTbDyTmLu [18]. The structures and properties can be regulated through a series of elemental additions or changes in the content, such as the addition of Ti [19], Zr [20], B [20], V [21], and/or Mo [22,23]. Some HEAs have the high glass-forming ability, such as PdPtCuNiP [24], ScCaYbMgZn [25], and ZrHfTiCuFe [26]. The structures of the HEAs include all the three simple metal crystal structures: A face-centered-cubic (fcc) structure represented by CoCrFeNi, a body-centered-cubic (bcc) structure represented by TaZrHfNbTi, and a hexagonal-closed-packed (hcp) structure represented by YGdTbDyLu. The deformation mechanisms of hcp-phase metals have been summarized: Either tensile or compressive strain induced twins or basal and prismatic slips [27–29].

In addition to the excellent mechanical properties of the HEAs at room temperature, the HEAs also have good performance under extreme conditions, such as high and low temperatures. The VNbMoTaW alloy developed by Senkov et al. maintains a yield strength of 800 MPa at 600 to 800 ◦C, and the yield strength is still more than 400 MPa at 1600 ◦C [30]. The CoCrFeMnNi single-phase fcc alloy demonstrates the characteristics of "lower temperature and more toughness," and the fracture toughness value is greater than 200 MPa·m1/<sup>2</sup> at a low temperature (77 K) [31]. For Co1.5CrFeNi1.5Ti and Al0.2Co1.5CrFeNi1.5Ti HEAs, the wear resistance is at least twice that of conventional wear-resistance steels, such as SUJ and SKH51 steel under similar hardness conditions [32]. Moreover, the mechanism of the phase transition in HEAs has received great attention from researchers. At present, the transformation-induced plastic deformation and polymorphism have been studied [7,33]. However, the phase-transition kinetics of the HEAs still need to be further studied. The atomic mechanism and thermodynamic principle of the HEA phase transition are still in dispute [34,35]. Surface energy at the interface may affect the solidification and diffusion of solids, such as precipitation, resulting in phase transitions [36].

#### **3. High Pressure**

The rapid development of the high-pressure technology and the interpenetration with other technologies have become a hot topic in science today. High pressure is an important means to control the properties of the material, and is a decisive variable in modern scientific research [37]. Bundy synthesized diamond with technologies employing high pressures and high temperatures, which paved the way for the high-pressure technology for the application of new materials [38].

According to the different high-pressure loading methods or technologies, the high-pressure experimental technology can be divided into the static high-pressure loading technology and dynamic high-pressure loading technology. The duration of the traditional, dynamic high-pressure experiments is very short, typically no more than a few milliseconds, and can reach very high pressures and temperatures. However, the defects are also obvious: The time span is too short to be effectively detected and accurately controlled, and the large amount of extra heat will affect the expected experimental results. In contrast, the static high-pressure technique can perform the non-destructive research on the material under fixed pressure conditions and can obtain good data. Static and dynamic high pressures are complementary to each other, which enhances our basic understanding of the

material structure and performance under different high pressures. The dynamic high-pressure loading technology is an experimental technique for subjecting samples to transient high-pressure and high-temperature environments by the shock waves or high-speed physical impact generated from an explosion, or other means. At present, there are many experimental techniques that can achieve dynamic high pressures, which can be simply divided into high-speed dynamic loading and low-speed dynamic loading. The methods implementing the former mainly include the underground nuclear explosion, magnetic flux compression, rapid expansion, or an explosion of air currents. These methods are characterized by fast loading (can realize nanosecond loading), and high-pressure limits. But it is difficult to accurately control the pressure range, relying heavily on experience and repeated experiments. Many high-speed dynamic-loading methods are still in development and inappropriate for research under normal laboratory conditions. In contrast, the experimental techniques of low-speed dynamic loading are relatively mature and have mostly been used in experiments. These loading techniques have a large time span (from hours to milliseconds or even microseconds), high repeatability, high-pressure loading accuracy, and wide adaptation (which can be applied to room temperature, high temperature, low temperature, magnetic field, and other conditions), which can be combined with other test methods and characteristics. It is of great help to study the structures and properties of materials under nonequilibrium thermodynamic conditions, which is an important direction in the high-pressure research.

The static high-pressure loading method is a technique for obtaining high-pressure experimental conditions using a static compression method. The sample is slowly compressed by means of the external mechanical loading device. Since the compression process is slow enough, the heat generated in the process can be fully exchanged with the external environment. Hence, the static high-pressure loading process is an isothermal compression process. At present, the equipment to achieve the static high pressure mainly includes large cavity presses and diamond-anvil-cell presses. Currently, the highest pressure of the former is up to hundreds of thousands GPa, which is the main equipment for the synthesis of high-pressure materials, and it has a unique advantage in the use of high-temperature and high-pressure, synthetic block, superhard materials. Compared to the former, the diamond-anvil cell (DAC) is small in size, requires few samples, and has low experimental costs. Multiple in-situ experiments (including X-ray diffraction, Raman spectrum, etc.) can be performed under a variety of experimental conditions, and higher-pressure limits can be achieved. The hexahedron-anvil-press instrument is the most widely-used static high-pressure device. Figure 3 is the schematic of the instrument. The instrument is composed of six anvil cells, extruding a square-pressure transmitting medium, and the sample is wrapped in the medium. As shown in the figure, the machine applies the pressure from six directions to the center simultaneously to achieve the static equilibrium [39,40].

**Figure 3.** Schematic of the hexahedron anvil press.

The DAC, another device that can achieve the static high-pressure loading, it is smaller in size, and Figure 4 shows the principle of the DAC. DACs are composed of two opposing diamond anvils that squeeze the samples between them to create hydrostatic/non-hydrostatic pressures [41]. In 1977, Bruas combined the DAC technology with synchrotron-radiation X-ray diffraction for the first time [42], which greatly promoted the development of the high-pressure science. In 2015, Dubrovinsky et al. used nano-diamonds in combination with a two-stage pressurized device to obtain a high static pressure of 770 GPa [43]. The core of the DAC is to use the extremely-high hardness of the diamond, with two diamonds pressing on top of each other, to generate high pressures.

The high-pressure torsion (HPT) process was first proposed by Bridgman in 1935 [44]. With the development of microscopy technology, it was not until the 1980s that researchers discovered that ultra-fine crystalline structure could be prepared by high-pressure torsion technology [45]. The principle of HPT is exhibited in Figure 5. The sample is placed between the two anvils and subjected to extremely high pressures of hundreds of GPa. When there is a relative rotation between the two anvils driven by external forces, the friction between the sample and the anvils drives the sample to rotate and causes the shear deformation of the sample. Although the sample is subjected to the large strain plastic deformation, it will not rupture, due to the high pressure. By means of HPT, the size of crystal grains can be significantly reduced, and a dense nano-bulk material can be prepared effectively.

In high-pressure scientific research, the structure-phase transitions, which combine pressure or temperature and pressure as driving forces, can be divided into two basic types, namely reconstructive phase transitions and displacive phase transitions. The classification is based on whether the chemical bonds that form the periodic grid are destroyed after the phase transformation. In the process of the reconstructive phase transition, the main chemical bonds are broken and recombined to form a new structure. There is no clear orientation relationship with the previous phase in crystallography. All the reconstructive phase transitions belong to the first-order phase transition, and the discontinuity of the volume change during the phase transition is obvious. Since there are dynamic barriers between the equilibrium pressures of the adjacent two phases, it often leads to a hysteresis effect in the phase transition. The phase-transition pressure point in the pressurization process is greater than the pressure-relief process. The secondary bonds may break when a displacive phase transition occurs, but the main chemical bonds will not break, usually due to the displacement of the atoms or the tilt of the polyhedral structure. Under a high pressure, many phase transitions are displacive phase transitions, and the degree of the discontinuous volume change is very small. Most of them still belong to the first-order phase transition.

**Figure 4.** Schematic of the diamond anvil cell (DAC).

**Figure 5.** Schematic illustration showing the principles of high-pressure torsion (HPT).

#### **4. HEAs under High Pressure**

#### *4.1. Dynamic High Pressure*

The HEA is a new kind of alloy between the traditional alloy and the amorphous alloy. The deformation mechanism shows different characteristics from the traditional alloys. Therefore, it is necessary to study the structure and phase-transition process of HEAs under high pressures. The microstructures of HEAs under high pressures is also one of the hot topics in the alloy research. Huang et al. [46] developed a facile, two-step carbothermal shock (CTS) method that employs flash heating and cooling (temperature of 2000 K, shock duration of 55 ms, and ramp rates in the order of 105 K/s) of metal precursors on oxygenated-carbon support to produce HEA nanoparticles (HEA-NPs) with up to eight metallic elements. Figure 6 shows the scanning transmission electron microscopy (STEM) elemental maps for PtPdRhRuCe HEA-NPs. The aforementioned is a method of synthesizing HEAs using the high pressure produced by the thermal shock. Nanoparticles are useful in a wide range of applications. Huang et al. [46] developed a method for making HEA nanoparticles, and the "carbothermal shock synthesis" can be tuned to select for the nanoparticle size as well as final structure. These carbothermal shock (CTS) capabilities facilitate a new research area for the materials discovery and optimization.

**Figure 6.** Scanning transmission electron microscope (STEM) elemental maps for PtPdRhRuCe high-entropy-alloy nanoparticles (HEA-NPs), reproduced from [46], with permission from authors.

#### *4.2. Diamond Anvil Cells*

The duration of traditional dynamic high-pressure experiments is very short, generally no more than a few milliseconds. Therefore, it is difficult to effectively detect and accurately control, and the large amount of the extra heat will affect the expected experimental results. Thus, the experiment of placing HEAs under high pressures usually adopts a static high-pressure technology. Yu et al. [47] used DAC to pressurize the rare-earth HoDyYGdTb HEA and studied the phase transition under high pressures. Figure 7 shows the pressure-volume relationship of different phases of the HoDyYGdTb HEA measured at room temperature, and the red line indicates the fitting result using the Birch-Murnaghan equation of the state in [48]. The sample was prepared by arc-melting, then scraped, and loaded into the T301-stainless-steel gasket hole with a diameter of 180 μm. The specimen was pressurized up to 60.1 GPa. The HEA is shown to follow the trivalent rare-earth crystal structure sequence of hcp → Sm-type → dhcp → dfcc, which correlated the s → d charge transfer of the HEA. The bulk modulus and atomic volume of the rare-earth HEA agree extremely well with the calculated values with the "additivity law." The pressure-included phase transformation among bcc, fcc, and hcp phases in transition metals were reported previously. What is most noteworthy is that Au changes the structure of HCP under extreme conditions [49]. Zhang et al. [50] studied the phase transition in Ni-based HEAs under high pressures. For high-pressure experiments, the sample was loaded in the chamber, which was indented from the Re gasket with a pair of diamond anvils with a thickness of 40 μm. The single-phase CoCrFeNi HEA alloy was pressurized to observe the phase transition. A pressure-induced fcc-hcp phase transition was found in the CoCrFeNi HEA at the pressure of 13.5 GPa and at an ambient temperature. The hcp structure is recoverable when the pressure is released. The phase transformation is very sluggish and did not finish at 39 GPa

The HEA is essentially an alloy-design concept, and there is no requirement on which elements must be used. Therefore, it also provides a broad scope for the study of HEAs. The transition elements are the most frequently-used components of HEAs. As the major element in our planet, Fe is the most interesting element to be studied, and the phase transformations (including from bcc to hcp structures) under high pressures or high temperature have already been found [51–53]. In ambient conditions, cobalt is in the hcp structure, and a pressure-induced phase transition to the fcc phase was found [54]. Figure 8 shows this pressure-induced phase transition. However, no pressure-included phase transition was found in Ni up to 260 GPa [55]. [50].

Polymorphism is widely observed in many materials, which describes the occurrence of different lattice structures in a crystalline material, and is a potential research direction in the future. The polymorphism in the CoCrFeMnNi HEA is reported in [33]. By employing in situ, high-pressure, synchrotron radiation X-ray diffraction, the polymorphic transition from fcc to hcp structures in the CoCrFeMnNi HEA is observed. The hydrostatic pressure was up to 41 GPa using a DAC. The CoCrFeMnNi HEA has an fcc single–phase structure and remains stable up to 19.5 GPa. When the pressure reaches 22.1 GPa, new peaks appeared, indicating that the phase transition has occurred. All the new peaks revealed an fcc-to-hcp transition under high pressures. During the decompression, the phase transition was irreversible. Huang et al. [56] studied the deviatoric deformation kinetics in the CoCrFeMnNi HEA under hydrostatic compression. The HEA was subjected to a hydrostatic pressure of 20 GPa via a DAC. The main significance of this study was to provide another perspective for studying the deformation mode of the HEA system through high-pressure experiments [56].

**Figure 7.** The measured pressure-volume relationship for various phases of the HoDyYGdTb high-entropy alloy (HEA) at up to 60.1 GPa, at room temperature, data from [47].

Gong Li et al. [57] studied the pressure-volume relationship of CoCrFeNiAlCu HEA using in-situ high-pressure energy-dispersive X-ray diffraction with synchrotron radiation at high pressures, and the results show that the CoCrFeNiAlCu HEA keeps a stable fcc + bcc structure in the experimental pressure ranges from 0 to 24 GPa. The equation of the state of the HEA determined by the calculation of the radial distribution function in the non-phase-transitional case is:

$$-\Delta V/V\_0 = 2.7P - 0.256P^2 + 0.012P^3 - 2.928 \times 10^{-4}P^4 + 2.907 \times 10^{-6}P^5$$

where *V*<sup>0</sup> is the volume at zero pressure, (Δ*V*/*V*0) is the relative volume change.

Cheng et al. [58] studied an ordered, bcc-structured (B2 phase) AlCoCrFeNi HEA using in situ, synchrotron radiation X-ray diffraction up to 42 GPa and non-in transmission electron microscopy. Pressure-induced polymorphic transitions (PIPT) to potentially disordered phase were observed. Yusenko et al. [59] studied the temperature and pressure stability for a hcp Ir0.19Os0.22Re0.21Rh0.20Ru0.19

HEA. The sample was loaded in a DAC cell and did not result in phase transition with a maximum pressure of 45 GPa. Ahmad et al. [60] performed in-situ high-pressure and high-temperature XRD measurements on bcc-Hf25Nb25Zr25Ti25, fcc-Ni20Co20Fe20Mn20Cr20 and hcp-Re25Ru25Co25Fe25 HEAs; all of the HEAs remained stable and no phase transition was observed.

**Figure 8.** CoCrFeMnNi XRD patterns during compression and decompression, reproduced from [33], with permission from authors.

#### *4.3. High-Pressure Torsion*

Due to the severe lattice distortion effect and different chemical bonds of their constituent elements, the plastic deformation mechanisms of HEA could be different from that of conventional alloys. The high-pressure torsion (HPT) method is currently used mainly as a severe plastic deformation (SPD) technique for grain refinement, and it has taken a long time to synthesize a metastable phase by HPT processing. High-pressure torsion is generally used to prepare non-porous bulk ultrafine grain samples, which can generate large plastic deformation of the sample through shear stress (γ = 10–100) [61–64]. At present, there have been reports on high-entropy alloys treated by high pressure torsion. These reports are mainly focused on the face centered cubic HEAs. Tang [65] obtained the nano-scale Al0.3CoCrFeNi HEA by HPT and studied the strengthening mechanism of the annealing process. Schuh [66] used high pressure torsion to treat the CoCrFeMnNi HEA; the grain can be refined to 50 nm, and the strength and hardness can be increased to 1950 MPa and 520 HV, respectively. After high-pressure torsion of the face centered cubic high-entropy alloy, the grain gradually refines as the strain increases. As shown in Figure 9, the plastic-deformation mechanism mainly includes dislocation slip and twinning.

**Figure 9.** Microstructural evolution in HPT disks investigated with SEM using back-scattered electron contrast, reproduced from [66], with permission from authors.

Yu et al. showed that the plastic-deformation mechanisms of the single-phase fcc Al0.1CoCrFeNi HEA induced by HPT at room temperature [67]. The sample was prepared by arc-melting and then casted as a plate by vacuum induction. The plate was hot-isostatic pressed (HIPed) at 1473 K, and 100 MPa, for 4 h, and then placed in a horizontal tube furnace at 1423 K for 50 h. The samples were compressed at 6 GPa through 1 and 2 turns with the rotation speed of 1 rpm. Processing by HPT produces a very substantial grain refinement in HEAs. The deformation mechanisms include the dislocation slip and twinning at room temperature. As shown in Figure 10, the average Vickers microhardness of the HIPed sample is 135 Hv. After HPT for 1 revolution, the microhardness reaches a saturation of about 482 Hv at the edge of the sample.

A deformation-induced phase transformation in a single-phase fcc Co20Cr26Fe20Mn20Ni14 HEA during the cryogenic HPT was reported [68]. The sample was prepared by arc-melting and then homogenizing by heat treatment at 1050 ◦C for 24 h; the sample was cold-rolled after the heat treatment, followed by annealing at 1050 ◦C for 1 h with water quenching. The HPT processing was performed at 77 K in the liquid nitrogen (cryo-HPT) and at room temperature (300 K-HPT), the samples were compressed at 5 GPa for 5 revolutions. The thermodynamic calculations indicated that the hcp phase showed a higher stability than the fcc phase. The cryo-HPT providing an extra driving force for the fcc to hcp phase transformation of the Co20Cr26Fe20Mn20Ni14 alloy at 77 K. The XRD pattern of the annealed sample shows the FCC phase and the r phase. In contrast, there was no r phase in the XRD pattern of the 300 K-HPT or cryo-HPT treated samples. Only the fcc peak was found. The absence of the r-phase peak after HPT treatment may have been due to grain refinement and residual strain after HPT treatment, which causes the main fcc peak to broaden. Unlike other conditions, the XRD pattern of the low temperature and high pressure treated alloy contains a well-developed hcp peak. This gives us reason to assume that, unlike 300 K-HPT, low temperature HPT induces a phase transition of fcc-to-hcp [68]. Furthermore, the microstructure and thermal stability of the nanocrystalline CoCrFeMnNi HEA after HPT were reported, and the results indicated that the grain size was refined, along with an unusual increase of the strength [66]. The hardness of Al0.3CoCrFeNi was significantly increased processed by HPT [65].

**Figure 10.** Vickers microhardness plotted against the distance from the center for the sample, data from [67].

#### *4.4. Hexahedron Anvils Press*

HEAs with high hardness prepared by high-pressure sintering were reported by Yu et al. [69]. The equiatomic CoCrFeNiCu and CoCrFeNiMn HEAs were prepared by high-pressure sintering (HPS). The elemental powders were developed in a planetary ball mill. Then the powders were sintered in the hexahedron anvil press at 1273 K and 5 GPa for 15 min. A graphite tube was taken as the heating device and the pyrophyllite as the pressure-transmitting medium. The structures and mechanical properties were carefully investigated. It revealed that the structure of the HEA powders have a main fcc phase and a minor bcc phase. After high-pressure sintering, the phase transition from the bcc to fcc structures occurred, and exhibited a simple fcc solid-solution structure. The hardness of the CoCrFeNiCu HEA increased from 133 HV to 494 HV by HPS, and the hardness of the CoCrFeNiMn increased from 300 HV to 587 HV. The increase of hardness is the reason for the decrease of the grain size. The grain size is about 100 nm after HPS. It reveals that the HPS is an effective way to design excellent HEAs.

#### **5. Future Work**

HEAs are considered to be one of the three breakthroughs in alloying theory in recent decades (the other two are bulk metallic glass and metal rubber). There are huge developments underway for this unique design concept. The high mixing entropy makes it have an extensive application potential. Since the discovery of HEAs, it has been about two decades. Only in recent years, it has received sufficient attention from researchers and developed rapidly. There is still a great amount of work to be done. The high-pressure technology has been more and more widely used in the fields of fabricating the new material preparation and changing the structures and organizations of materials, such as synthesizing diamonds with the ultra-high-pressure technology, producing ultra-hard and antifriction materials, etc. The constant renewal of the high-pressure technology will enable the development of excellent and more promising materials. Additionally, high pressure science has obvious frontier characteristics, which brings forth many new opportunities and challenges for the development of science and technology. National security is a potentially important driving force of the international attention, in addition to mere intellectual interest and economic interests. HEAs will be an important

materials development route for the field of the high-pressure science in the future. The properties of HEAs change when high pressure act on the alloys. The most significant change in the alloys, due to the high pressure, is the phase change during the high-pressure crystallization. The materials show many unique properties under high pressures. The research of high entropy alloy under high pressure has already established a certain foundation, but the industrial applications have not matched the expected results. It is necessary to further broaden the application field, which will play a more important role in future theoretical improvement and new material preparation. With HEAs, as emerging alloys with a great potential for the future development, their structures and properties under high pressures need further studies.

#### **6. Conclusions**

HEAs as a new class of alloys have been attracting more and more attention in recent years. With the unique design concept, HEAs show wide application potentials. High pressure is an external condition that is important for structural changes and phase transitions. The unique phenomenon of HEAs under high pressures gives us a new understanding of HEAs, providing a new way forward for the in-depth study of HEAs. Some single-phase HEAs, such as CoCrFeNi, rare-earth HoDyYGdTb, and CoCrFeMnNi phase transition under high pressure, but bcc-Hf25Nb25Zr25Ti25, fcc-Ni20Co20Fe20Mn20Cr20, and hcp-Re25Ru25Co25Fe25 remaine stable and no phase transition is observed under high pressure. These unique structural changes and phase transitions of HEAs under high pressure show wide application potentials. As a new research trend, there is a growing interest in the structure and performance of HEAs under high pressures. The studies of HEAs under high pressures are summarized in this paper, with a hope to deepen the fundamental understanding of HEAs under high pressures.

**Author Contributions:** W.D. performed the data analyses and wrote the manuscript; Z.Z. helped perform the analysis with constructive discussions; M.Z. and Y.M. helped document retrieval; P.K.L. and P.Y. performed the manuscript review; G.L. contributed to the conception of the study.

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

**Acknowledgments:** One of the authors (Gong Li) acknowledges the National Science Foundation of China (Grant No. 11674274). Pengfei Yu acknowledges the National Natural Science Funds of China (Grant No. 51601166). Peter K. Liaw very much appreciates the support of the U.S. Army Research Office Projects (W911NF-13-1-0438 and W911NF-19-2-0049) with the program managers, M.P. Bakas, S. N. Mathaudhu, and D.M. Stepp. Peter K. Liaw thanks the support from the National Science Foundation (DMR-1611180 and 1809640) with the program directors, G. Shiflet and D. Farkas.

**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* **Tribological Behavior of As-Cast and Aged AlCoCrFeNi2.1 CCA**

#### **Fevzi Kafexhiu \*, Bojan Podgornik and Darja Feizpour**

Institute of Metals and Technology, Lepi pot 11, 1000 Ljubljana, Slovenia; bojan.podgornik@imt.si (B.P.); darja.feizpour@imt.si (D.F.)

**\*** Correspondence: fevzi.kafexhiu@imt.si; Tel.: +386-1-4701-931

Received: 31 December 2019; Accepted: 28 January 2020; Published: 1 February 2020

**Abstract:** In the present study, wear behavior as a function of aging time was evaluated for the AlCoCrFeNi2.1 eutectic complex, concentrated alloy (CCA) consisting of B2 (BCC), and L12 (FCC) lamellae in the as-cast state. By aging the material at 800 ◦C up to 500 h, precipitation of a fine, evenly dispersed micro-phase inside the L12 takes place. From 500 h to 1000 h of aging, precipitates coarsen by the Ostwald ripening mechanism. Reciprocating wear tests were characterized by a prevailing abrasive wear mechanism, while adhesive and delamination wear components change with aging conditions. The L12 phase with lower hardness in the as-cast material preferentially deformed during the wear test, which was not the case after aging the material, i.e., with the presence of precipitates. Aging-induced changes show a similar trend for the coefficient of friction and L12 + precipitates phase fraction, whereas changes in specific wear rate are in a good agreement with changes in B2 phase fraction. In general, aging the AlCoCrFeNi2.1 CCA at 800 ◦C up to 500 h decreases its coefficient of friction due to reduced adhesive wear component and enhances its wear performance through precipitation strengthening.

**Keywords:** AlCoCrFeNi2.1; CCA; HEA; aging; precipitates; wear; tribology

#### **1. Introduction**

Multi-principal element alloys (MPEAs) have become an attractive topic in the materials science community because of their great potential in discovering and developing new materials of scientific significance and practical benefit. In contrast to conventional alloys, these materials usually contain at least five elements in equimolar or near-equimolar proportions, which in the case of fulfilling the Hume-Rothery rules can have microstructures of single-phase solid solution with high configurational entropy of mixing—high-entropy alloys (HEAs), or they can have complex microstructures that contain multiple phases of solid solution and/or intermetallic compounds—complex, concentrated alloys (CCAs). The research interest on CCAs is ever increasing as they offer an excellent combination of advantages of single-phase solid solution in HEAs, and secondary-phase strengthening effects of well-established alloys. This combination may as well diminish the strength-ductility competition, which is a well-known issue in conventional alloys.

However, the application of these materials in structural engineering requires, among others, a good understanding of their surface degradation mechanisms including corrosion, erosion, and wear behavior [1]. In this respect, some research has been carried out by different authors on different HEAs and CCAs as reviewed by Ayyagari et al. [2]. Some highlights worth mentioning include the work of Wu et al. [3] who found out that by increasing aluminum content of the Al*x*CoCrCuFeNi, both the volume fraction of the BCC phase and the hardness value increase and thus the wear coefficient decreases. On the other hand, Tong et al. [4] reported that the wear resistance of the Al*x*CoCrCuFeNi was similar to that of ferrous alloys at the same hardness level. Both these authors correlated the high wear resistance to the higher hardness coming from the solid solution strengthening of single-phase HEAs. Hsu et al. [5] in their work discovered that the major wear mechanism of the AlCoCrFe*x*Mo0.5Ni HEA is abrasion. Also, by performing the oxidation test at the pin/disk interface flash temperature, 500 ◦C, they came to the conclusion that the oxidation rate of Fe2.0 markedly exceeds that of Fe1.5, indicating more oxides abrade the surface, resulting in lower wear resistance. Oxidative wear was also encountered by Du et al. [6] who studied the tribological behavior of the Al0.25CoCrFeNi HEA with a simple FCC phase and hardness of 260 HV from room temperature to 600 ◦C. They found out that below 300 ◦C, with increasing temperature, the wear rate increased due to high temperature softening. The wear rate remained stabilized above 300 ◦C due to the anti-wear effect of the oxidation film on the contact interface. The dominant wear mechanism of HEA changed from abrasive wear at room temperature to delamination wear at 200 ◦C, then delamination wear and oxidative wear at 300 ◦C and became oxidative above 300 ◦C. Excellent anti-oxidation property and resistance to thermal softening were reported for the Co1.5CrFeNi1.5Ti and Al0.2Co1.5CrFeNi1.5Ti alloys, which are the main reasons for the outstanding wear resistance, which is at least two times better than that of conventional wear-resistant steels with similar hardness, such as SUJ2 and SKH51 [7].

Löbel et al. [8] investigated the wear behavior of AlCoCrFeNiTi0.5 HEA produced by powder metallurgy under reciprocating wear conditions from room temperature to 900 ◦C. They found out that with increasing temperature up to 650 ◦C, initially, a slight decrease in wear resistance occurred, whereas a further increase in test temperature resulted in a distinct increase in wear resistance and a decrease in coefficient of friction. Their investigations prove the suitability of the AlCoCrFeNiTi0.5 HEA for high-temperature applications, as the formation of protective oxides improves the wear performance.

A novel AlCoCrFeNi2.1 eutectic CCA firstly studied by Lu et al. [9], is an alloy with promising properties in as-cast condition [10,11] due to the contribution of both ductile FCC (L12) and harder BCC (B2) phases resulting in a combination of good strength and ductility. The alloy has also shown to have excellent thermomechanical processing capability by severe cold rolling [12–15] or rotary friction welding [16]. Excellent work on surface wear and corrosion behavior of the AlCoCrFeNi2.1 has been performed by Hasannaeimi et al. [17], where a transition from adhesive to oxidative wear was observed as the duration of reciprocating wear test increased.

The purpose of the present research is to study the tribological behavior of AlCoCrFeNi2.1 eutectic CCA as a function of microstructure evolution by aging the material at 800 ◦C for 100, 500, and 1000 h, both for the sake of fundamental scientific understanding and the application worthiness of this particular alloy. Although the precipitation is the main process taking place during aging at 800 ◦C, the focus of the present research will be only on the effect these precipitates have on the wear behavior of the AlCoCrFeNi2.1. Detailed analysis of precipitate kinetics, thermodynamics, and structure/morphology will be published elsewhere.

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

The eutectic AlCoCrFeNi2.1 CCA was synthesized by vacuum induction melting and ingot casting using commercially available elements with a purity of 99.9%. Before melting, the furnace was evacuated and a subsequent argon gas atmosphere of 300 mbar was created. The molten material was kept for at least 20 min at 1500 ◦C under induction current, which provided sufficient agitation to ensure the homogeneous distribution of elements in the melt. The melt was cast inside the furnace (Ar atmosphere) in a cast-iron crucible and further cooled down in the air. In order to determine the melting point of the alloy and to find out what reactions might take place during the heating/cooling process, differential scanning calorimetry (DSC) analysis using STA 449 C Jupiter Thermo-microbalance (Netzsch-Gerätebau GmbH, Selb, Germany) was subsequently performed in the temperature range of 30–1425 ◦C under the dynamic atmosphere of Ar (15 mL min<sup>−</sup>1), heating and cooling rate of 20 K min<sup>−</sup>1, using an Al2O3 crucible with approximately 219.6 mg of sample material. As indicated by the DSC curves in Figure 1, besides melting and solidification peaks with onset temperatures 1343.9 ◦C and 1346.3 ◦C, respectively, a transformation reaction revealed in form of weak peaks can be seen on the

heating and cooling curves with onset temperatures around 800 ◦C and 822 ◦C, respectively. Therefore, in order to be able to characterize the phase transformation developing around 800 ◦C, dilatometry technique using Bähr DL 805A/D dilatometer (TA Instruments, Inc, New Castle, DE, USA) was used to heat a cylindrical sample with a diameter of 4 mm and a length of 10 mm to a temperature of 800 ◦C at which it was held for 20 h, then quenched by streaming N2 gas to ensure a controlled cooling rate of 10 K s−1. Afterward, the sample was longitudinally cut in half and a metallographic specimen was prepared by hot mounting in Bakelite, then mechanically ground with silicon carbide emery paper from a grade 180 down to 1200, and finally polished with 3 μm and 1 μm abrasive diamond suspensions. Scanning electron microscopy (SEM; JSM-6500F, Jeol, Tokyo, Japan) with back-scattered (BSE) detector and energy dispersive X-ray spectroscopy (EDS) were used for imaging and chemical composition determination. In addition, a lamella for transmission electron microscopy (TEM; JEM-2100 HR, Jeol, Tokyo, Japan, operated at 200 kV) was prepared for detailed microstructural characterization. A TEM lamella was first prepared by cutting a small sample of around 3 mm length, 1 mm height, and 1 mm thickness from a bulk specimen, then coarse and fine ground and polished to around 100 μm thickness using SiC papers from 800 to 4000 grit sizes, and additionally thinned by argon ion-slicing (IonSlicer, EM-09100IS, Jeol, Tokyo, Japan) with ion milling at 6 kV for around 6 h and around 15 min of fine ion milling at 2 kV to reach electron transparency. Scanning transmission electron microscopy (STEM) unit with a bright-field (BF) detector and EDS (JED-2300T, Jeol, Tokyo, Japan) were used for chemical composition determination and elemental mapping. Based on the characterization results, it was decided for the present work to isothermally age the material at 800 ◦C in a simple furnace (air atmosphere) at three different durations, 100, 500, and 1000 h and subsequently quench it in water.

**Figure 1.** Heating and cooling curves from DSC analysis on the AlCoCrFeNi2.1 CCA.

On as-cast and aged material, metallographic samples of cuboid shape with dimensions 20 × 20 × 10 mm3 were machined, ground and polished applying the same method as described above. SEM-BSE imaging and EDS analyses on at least 5 randomly chosen regions at both 1k and 3k magnifications were acquired. For more representative analysis, images acquired at 1k magnification were used for quantitative evaluation of phase fractions, whereas those acquired at 3k were used for quantitative evaluation of precipitates. Images were digitally analyzed using FIJI (ImageJ, ver. 1.52p, Bethesda, MD, USA) [18,19] with an appropriate size filter and color threshold, which enabled a separate analysis of the three distinct phases. As a result, the surface area of the darker phase, as well as surface area and distribution (*x* and *y* coordinates) of each precipitate could be obtained. Having these data and a known surface area of BSE images, the area fraction of all phases could be easily determined, where average values with standard deviations were derived out of five analyzed images. Furthermore, on the metallographic samples (cuboid 20 <sup>×</sup> <sup>20</sup> <sup>×</sup> 10 mm3), HV10 hardness measurements were performed at room temperature with at least three indentations and a holding time of 14 s using Instron Tukon 2100B instrument (Buehler-Illinois Tool Works (ITW), Lake Bluff, IL, USA). In addition, three repetitions

of reciprocating ball-on-plate wear tests using hardened DIN 100Cr6 bearing steel ball with a diameter of 20 mm as a counter-body were performed. The hardened bearing steel ball was used due to its high hardness (~700 HV) thus concentrating the major amount of wear on the investigated alloy's surface, as well as due to the similarity to many machine component applications with the prevailing metal-metal contact. Furthermore, the use of the steel counter-body enables the study of not only abrasive but also adhesive wear mechanisms.

Wear tests were performed using the in-house designed ball-on-plate reciprocating sliding device (Figure 2), as typically found in many tribological investigations, with a stationary ball (counter-body) being loaded against a moving flat-surface specimen (cuboid 20 <sup>×</sup> <sup>20</sup> <sup>×</sup> 10 mm3). Tests were performed at ambient temperature conditions (21 ± 2 ◦C) with a sliding frequency of 15 Hz at a stroke of 4 mm, resulting in a maximal sliding velocity of 0.12 m s−<sup>1</sup> and a total duration of 833 s, which corresponds to a 100 m of sliding. The normal load of 20 N was applied, corresponding to 1 GPa of mean Hertzian contact pressure. All tests were performed in dry sliding conditions.

**Figure 2.** Ball-on-plate configuration used for wear tests: (**a**) Schematics with test parameters; (**b**) Configuration in the actual testing rig; (**c**) The device used for wear tests.

Worn specimens and counter-bodies were characterized using the 3D surface measurement system (Alicona InfiniteFocus, Alicona Imaging GmbH, Raaba, Austria), being able this way to accurately measure the worn-out material volume *V* (mm3) and calculate the specific wear rate coefficient (*k*) according to the equation

$$k = V \cdot F^{-1} \cdot \mathbf{s}^{-1} \tag{1}$$

where *F* (N) is the maximum load applied, and *s* (m) is the total sliding distance.

Finally, SEM imaging and EDS analyses were performed on wear tracks in order to shed some light on wear mechanisms.

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

#### *3.1. Microstructure and Hardness*

The microstructure of the as-cast alloy with typical lamellar/dendritic morphology is shown in the backscattered SEM image in Figure 3a. Precipitation of a fine darker phase inside L12 (FCC) (light gray area) after 100, 500, and 1000 h of aging at 800 ◦C is shown in backscattered SEM images in Figure 3b–d, respectively. Besides precipitation and a slight coarsening/decomposition of the B2 lamellae, there are no major morphological changes of the lamellar/dendritic structure with aging. However, elemental concentrations within both phases slightly change with aging time, as shown in Figure 4.

Figure 4 shows the change in at% concentration of elements in all three phases, obtained by EDS analysis. A decrease of Al and Ni concentrations and an increase of the concentrations of Co, Cr, and Fe in the L12 (FCC) phase (Figure 4a) is more pronounced in the first 100 h of aging at 800 ◦C, as the precipitation kinetics is the fastest in this time span. The opposite can be seen in the B2 (BCC) phase (Figure 4b). This is in good agreement with changes in phase fraction and precipitation process (Figure 5), where precipitates size and area fraction increase at the expense of the area fraction of L12

phase, while the B2 area fraction also shows a slight increase. Furthermore, there is a noticeable decrease in at% concentration of Co, Cr, and Fe in the precipitates after 1000 h of aging at 800 ◦C (Figure 4c) and an increase of Al and Ni concentrations. Note also the similarity of chemical composition between the B2 (BCC) phase after aging (Figure 4b) and the precipitates phase after 1000 h of aging (Figure 4c). This is a clear indication of similarity of B2 and precipitates phase also from the viewpoint of the crystal lattice structure, which needs to be confirmed by additional analysis using XRD, selected area diffraction in TEM, or even high-resolution TEM for a detailed analysis of the interface between precipitates and L12 phase. This is beyond the scope of the present study and will be investigated separately.

**Figure 3.** SEM-BSE images of the microstructure of: (**a**) As-cast AlCoCrFeNi2.1 CCA with B2 (BCC) lamellae—dark grey and L12 (FCC) phase—light grey; Precipitation and coarsening of fine B2-like phase inside L12 after aging at 800 ◦C for: (**b**) 100 h; (**c**) 500 h; (**d**) 1000 h.

**Figure 4.** Changes of elemental concentration in at% with aging at 800 ◦C in: (**a**) L12 (FCC) phase; (**b**) B2 (BCC) phase; (**c**) Precipitates.

**Figure 5.** Microstructure evolution and hardness change as a function of aging at 800 ◦C: (**a**) Phase fractions and Vickers hardness; (**b**) Precipitates size and their number density.

Changes in phase fraction and Vickers hardness with aging are shown in Figure 5a, whereas precipitate size and number density are given in Figure 5b. The initial phase fractions of B2 and L12 phases is around 0.3 and 0.7, respectively. After the first 100 h of aging, the precipitates phase fraction (0.1) increases at the expense of the L12 phase that drops to 0.57, while B2 phase fraction also shows a slight increase (0.33). Precipitates average size expressed as equivalent circle diameter (ECD) at this stage is 0.37 ± 0.1 μm. From this point up to the 1000 h of aging, the L12 phase fraction remains virtually unchanged, whereas the phase fraction of precipitates slightly fluctuates between 0.15 and 0.12 in accordance with the B2 phase after 500 and 1000 h of aging, respectively. However, this does not mean that the L12 phase has reached its solid solution equilibrium state because the precipitation from the L12 phase continues even after 100 h of aging. The precipitate size continues to increase to 0.42 ± 0.1 μm and 0.62 ± 0.2 μm after 500 and 1000 h of aging, respectively. The precipitation rate is the highest in the first 100 h of aging, as represented by the number of precipitates per unit area or number density of precipitates, which at this stage is around 1.05 <sup>±</sup> 0.13 <sup>μ</sup>m<sup>−</sup>2. The precipitation continues up to 500 h of aging but a with much lower rate, as the existing precipitates at this stage continue growing. From this point up to 1000 h of aging, the number of precipitates drops down to 0.37 <sup>±</sup> 0.04 <sup>μ</sup>m−2. Since at the same time precipitates size increases while their number density decreases, it means that by aging the AlCoCrFeNi2.1 at 800 ◦C from 500 h onward, the Ostwald ripening process develops, where larger particles coarsen at the expense of dissolving smaller ones.

The slight increase in hardness of AlCoCrFeNi2.1 after the first 100 h of aging at 800 ◦C can be attributed to precipitation of the fine B2-like phase at the expense of the softer L12 phase, the fraction of which decreases at this stage (Figure 5a). From this point up to 1000 h of aging, the hardness decreases linearly back to the initial value (as-cast condition), which can be attributed to the precipitate coarsening within the softer L12 phase and continuous impoverishing of the solid solution from solute atoms such as Al and Ni, which diffuse towards the precipitates and the B2 phase.

STEM-EDS elemental analysis and mapping shown in Figure 6b and summarized in Table 1, indicate the different composition of the B2 phase in spectra 1 and 2, and the L12 phase in spectrum 3 [20]. Spectrum 2 in Table 1 and the elemental map for Cr, at the lower right corner shows, increased Cr signal, which is coming from the Cr-rich nano-precipitates, as reported by Gao et al. [11]. Similar to the B2 phase (spectra 1 and 2), Al and Ni concentration in precipitates (spectra 4, 5, and 6) is higher as compared to the L12 phase (spectrum 3). A detailed analysis of these precipitates from the viewpoint of their crystal lattice structure, kinetics, thermodynamics of precipitation, etc., is beyond the scope of the present work, therefore it will be studied separately elsewhere.

**Figure 6.** Electron microscopy characterization of the dilatometry sample (20 h at 800 ◦C): (**a**) SEM-BSE image showing partial precipitation; (**b**) STEM bright-field image of precipitates with EDS point analyses and mapping.

(**a**) (**b**)


**Table 1.** Elements concentration in at% from EDS point analyses in Figure 6.

#### *3.2. Wear Tests*

A representative example of 1 out of 3 wear tracks of tested samples is shown in Figure 7 in form of 3D pseudo-color depth images, 2D optical images, and cross-sectional depth profiles of all three wear tracks taken at the same locations shown with the red line across the representative wear tracks in Figure 7a–d. Visual differences in shape or length/width/depth of wear tracks depending on materials condition (aging time) are minor, however, detailed volume measurements could reveal their difference.

The coefficient of friction (COF) shown in Figure 8 is characterized by both short- and long-range fluctuations. The short-range (high frequency) fluctuations can be attributed to the difference in coefficient of friction between the harder B2 phase and the softer L12 phase. This is supported by the findings of Hasannaeimi et al. [17] who evaluated a small-scale phase-specific scratch behavior and found the variation in COF across the scratch line, where the softer FCC phase shows lower COF as compared to the harder BCC one. Long-range fluctuations, however, could be coming from the different morphology of the new contacting surface, which is continuously uncovered by wearing out the top-most layer of the material in dry sliding contact.

The trend at which the average coefficient of friction of the as-cast AlCoCrFeNi2.1 CCA changes with aging time as shown in Figure 9, is in a good correlation with the changing trend of L12 + precipitates phase fraction. This suggests that the plastic deformation of the softer L12 matrix with precipitates results in higher COF than the harder B2 phase. In general, by aging the AlCoCrFeNi2.1 CCA at 800 ◦C, COF of decreases with time.

A trend opposite to the one in Figure 9 is shown in Figure 10 for the specific wear rate of the specimen and counter-body, as well as the B2 phase fraction. In as-cast material, the wear rate of counter-body is lower compared to the wear rate of the material. After 100 h of aging at 800 ◦C, the wear rates of both specimen and counter-body increase and become almost equal. At this point, the B2 phase fraction and hardness also increase slightly, whereas COF decreases. After 500 h of aging, the specific wear rate of the specimen decreases almost three times more than the wear rate of counter-body (Figure 10), whereas the COF remains virtually unchanged (Figure 9). At this point, there is a slight decrease in B2 fraction and a slight increase in precipitate size, fraction, and number density (Figure 5). With further aging up to 1000 h, the wear rate of specimen rises almost to the value of the as-cast state but still remains lower than the wear rate of the counter-body. From this analysis, it can be concluded that adequate aging of AlCoCrFeNi2.1 CCA at 800 ◦C not only decreases the coefficient of friction but also improves the material's wear performance.

**Figure 7.** Pseudo-color 3D depth images, optical images, and depth profiles of three wear tracks in (**a**) As-cast material; Material aged at 800 ◦C for (**b**) 100 h; (**c**) 500 h; (**d**) 1000 h.

**Figure 8.** Coefficient of friction throughout the wear test on as-cast AlCoCrFeNi2.1 and after aging at 800 ◦C at three different durations.

**Figure 9.** Coefficient of friction and FCC (L12) + precipitates phase fraction of as-cast AlCoCrFeNi2.1 and after aging at 800 ◦C at three different durations.

SEM-BSE images of the wear track of as-cast AlCoCrFeNi2.1 CCA shown in Figure 11a reveals the presence of patches of re-deposited material caused by adhesive wear, erosion grooves from abrasive wear, and severe disintegration and some delamination of the lamellar morphology inside the wear track. An SEM-BSE image at higher magnification in Figure 11b indicates extensive plastic deformation through the planar slip mechanism of the L12 phase. When the planar slip lines in the L12 phase reach the semi-coherent interface between the latter and B2 phases, a step-like shape is formed, as shown in Figure 11b.

Inside the wear track, both phases deform uniformly and no inter-phase detaching or void formation is observed. Gao et al. [11] reported that the L12 phase can accommodate several arrays of parallel mobile dislocations and deform by planar slip. They also reported that the B2 phase in the eutectic fails in a brittle manner, while the L12 phase shows ductility and necking, leading to dual-mode fracture in this alloy. Hasannaeimi et al. [17] reported that during wear test of the AlCoCrFeNi2.1, L12 and B2 phases deformed simultaneously, while the B2 phase accommodates medium density of dislocations. This was attributed to the 3D back-stress acting on the L12 phase which can maintain synchronous deformation in heterogeneous systems. This back stress is further enhanced by semi-coherent boundaries between B2 and L12 phases, and the lower fraction of B2 lamellae. The accumulative effect of these conditions resulted in the activation of dislocation in the brittle B2 phase facilitated by a high density of dislocation pile-up at the phase boundaries, and this modified the brittle behavior of B2 phase to accommodate deformation.

**Figure 10.** The specific wear rate of specimens and counter-body, and BCC (B2) phase fraction of as-cast AlCoCrFeNi2.1 and after aging at 800 ◦C at three different durations.

**Figure 11.** SEM-BSE images of wear tracks of AlCoCrFeNi2.1 CCA in as-the cast state: (**a**) Wear track edge with patches of re-deposited material; (**b**) Slip lines in the vicinity of wear track.

SEM-EDS analyses in Figure 12a summarized in Table 2, indicate that inside the wear track, the main wear mechanism is abrasion with little patches of oxidized redeposited material (self-adhesion) and some surface delamination. Figure 12b shows EDS analyses of the piled-up oxides of counter-body and self-adhered material, as summarized in Table 3, where Si and Mn, which are present in the counter-body material (DIN 100Cr6) are also detected.

**Figure 12.** SEM-BSE images with EDS analyses on the wear track of AlCoCrFeNi2.1 in the as-cast state: (**a**) Inside the wear track; (**b**) At the edge of the wear track.

Similar oxide pile-up phenomenon but in reduced amount is present in all durations of the aged material, as shown in Figure 13a for 100 h of aging. The difference between the aged material and the one in as-cast condition is the fraction of L12 and B2 phases, resulting in different hardness level, as well as the presence, size, and number density of precipitates that strengthened the L12 phase, resulting in no plastic deformation by planar slip, as seen in the as-cast material. This is shown in Figure 13b for 100 h of aging and is similar to the material aged for 500 and 1000 h. Differences observed in tribological behavior obtained by aging can be ascribed to these changes, causing alterations in wear mechanism. In the case of as-cast alloy without precipitates, due to lower B2 phase fraction as compared to the L12, abrasive wear is predominantly combined with adhesive wear, where delamination is largely prevented by dislocation pile-up and planar slip at the phase boundaries [11], thus modifying the brittle behavior of B2 phase [17]. Intensified adhesive wear component results in higher friction but low counter-body wear. After 100 h of aging, the adhesive wear component is reduced due to lower L12 phase fraction and increased hardness, also indicated by a small drop in friction. However, increased B2 phase fraction, absence of planar slip, and precipitation strengthening result in increased brittleness and intensified delamination. This amplifies the formation of hard wear particles, which remain in the reciprocating sliding contact and lead to increased wear of the material and counter-body (Figure 10). Prolonged aging time (500 h) results in a slightly reduced B2 phase and increased L12 phase fractions but a high level of strengthening with precipitates (Figure 5b), thus providing the best combination of high hardness and reduced brittleness with minimal wear. In this case, the main wear mechanism is abrasive wear with minimized components of adhesion and delamination. Furthermore, over-aging for 1000 h leads to a decrease in the number density of precipitates and an increase in their size, which results in decreased hardness, as shown in Figure 5. Lower hardness means increased abrasive wear, with a smaller number of large precipitates representing reduced resistance to sliding and a thus lower coefficient of friction (Figure 9).

**Figure 13.** SEM-BSE images of the wear track edge after 100 h of aging at 800 ◦C: (**a**) Oxide pile-up at the edge of the wear track; (**b**) Edge of the wear track showing no deformation by planar slip inside the L12 phase.


**Table 2.** Elements concentration in at% from EDS point analyses in Figure 12a.


**Table 3.** Elements concentration in at% from EDS point analyses in Figure 12b.

#### **4. Conclusions**

In the present study, wear behavior as a function of aging time was evaluated for the AlCoCrFeNi2.1 eutectic complex, concentrated alloy consisting of B2 (BCC) lamellae, and L12 (FCC) phase in as-cast state, as well as fine evenly distributed precipitates inside the L12 phase after aging at 800 ◦C for 100, 500, and 1000 h. Between 0 and 500 h of aging, both precipitates number and size increase, while from 500 h to 1000 h of aging, precipitates coarsen by Ostwald ripening mechanism. Abrasive wear prevailed in reciprocating wear analysis, with adhesive wear and delamination component changing depending on the aging conditions. The L12 phase without precipitates and lower hardness in the as-cast material preferentially deformed during the wear test, resulting in intensified adhesive wear component but minimum delamination, hindered by planar slip between B2 and L12 phases. This was not the case after aging the material, i.e., with the presence of precipitates. In general, aging the AlCoCrFeNi2.1 alloy at 800 ◦C decreases its coefficient of friction due to reduced adhesive wear component and enhances its wear performance through precipitation strengthening. However, under-aging (100 h) results in increased material brittleness and thus increased delamination, while over-aging (1000 h) results in precipitates coarsening, and decreased hardness and abrasive wear resistance. The best performance, combining high hardness and reduced brittleness with minimal wear, is achieved with an intermediate aging of 500 h.

**Author Contributions:** Conceptualization, F.K. and B.P.; Methodology, F.K.; Formal analysis, F.K.; Investigation, F.K. and D.F.; Writing—original draft preparation, F.K.; Writing—review and editing, B.P. and D.F.; Visualization, F.K.; Project administration, F.K.; Funding acquisition, F.K. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research was funded by Javna Agencija za Raziskovalno Dejavnost Republike Slovenije, grant number Z2-9220.

**Acknowledgments:** Acknowledgements go to the Institute of Metals and Technology (IMT) in Ljubljana where research was performed. Also, authors are thankful to all who were involved in the experimental part of this research.

**Conflicts of Interest:** The authors declare no conflict of interest. The funders had no role in the design of the study; in the collection, analyses, or interpretation of data; in the writing of the manuscript, or in the decision to publish the results.

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© 2020 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*
