**1. Introduction**

In recent years, numerous metallic alloys have been studied that contain multiple principal and equimolar elements, typically with five or more major constituents in the range of 5–35 at. %. This new approach to designing such alloys with multiple principal elements has led to the emergence of high-entropy alloys (HEAs), which were proposed by Yeh et al. [1] and Cantor et al. [2], and have more recently been denoted as complex concentrated alloys (CCAs) [3,4]. Unlike conventional alloys, core effects are present, such as high entropy, lattice distortion, sluggish diffusion, and cocktail effects [1–4]. The highentropy effect results from the higher configurational entropy in these alloys compared to conventional alloys. Due to the high entropy of mixing, these alloys may favor the formation of simple solid solutions and prevent the generation of hard but brittle intermetallic compounds. As a result, these alloys can exhibit superior wear, oxidation and corrosion resistance, as well as high-temperature strength [3–6]. Moreover, recent studies have indicated that Al0.5CoCrCuFeNiTix, AlxCoCrCuFeNi and AlCoCrFeNiTix HEAs exhibit a variety of microstructures and mechanical properties, with face-centered cubic (FCC), body-centered cubic (BCC) structures, or a mixture of both [7,8].

According to studies on the tribological performance of HEAs, composition is a key factor in wear resistance. The effect of Al addition on the FeCoCrNiAlx HEA was examined by Liu et al. [9], who concluded that the wear mechanisms changed from previously

**Citation:** Kadhim, D.F.; Koricherla, M.V.; Scharf, T.W. Room and Elevated Temperature Sliding Friction and Wear Behavior of Al0.3CoFeCrNi and Al0.3CuFeCrNi2 High Entropy Alloys. *Crystals* **2023**, *13*, 609. https://doi.org/10.3390/ cryst13040609

Academic Editor: Jacek Cwik ´

Received: 17 February 2023 Revised: 26 March 2023 Accepted: 30 March 2023 Published: 2 April 2023

**Copyright:** © 2023 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 (https:// creativecommons.org/licenses/by/ 4.0/).

mixed wear modes of abrasive, adhesive, and oxidative wear to a mixture of abrasive and oxidative wear. Such a change shows that the FeCoCrNiAlx HEA coating's wear resistance was significantly enhanced by the addition of Al. In a tribological study of Al0.65CoCrFeNi, Miao et al. [10] found that only abrasive wear features could be seen on the remelted layer, whereas adhesive wear predominated on the substrate. As a result, the average friction coefficient and wear rate of the remelted layer were reduced by 23% and 80%, respectively, compared to the substrate. Wu et al. [11] reported on the adhesive wear behavior of AlxCoCrCuFeNi alloys, showing that increasing the aluminum content resulted in the HEA structure transforming from FCC to BCC phases, subsequently raising the hardness value and lowering the wear rate. In addition, a low Al percentage has been found to favor a single-phase FCC lattice in the well-studied AlxCoCrFeNi HEA, whereas a greater Al fraction results in a BCC phase [12].

The well-studied AlxCoCrFeNi HEA system microstructures have been shown to change from having a single FCC phase to a single BCC phase with increasing Al content, with the Al0.3CoCrFeNi HEA being the only one to form a solid solution with an FCC structure. This particular HEA will serve as a baseline alloy in the present study. According to several studies, this HEA exhibits good mechanical properties such as high plasticity, work-hardening capacity, and a balance between cryogenic strength and ductility [13–16]. Due to the inherent property of the FCC structure, Al0.3CoCrFeNi HEA has a relatively low strength at ambient temperature, with a room-temperature yield strength between 150 and 350 MPa [13]; its melting point is at 1870 K [14]. The mechanical properties of Al0.3CoCrFeNi HEAs have been significantly improved through the application of thermo-mechanical processing techniques, such as cold rolling and subsequent annealing. For example, with 90% cold rolling and 550 ◦C annealing, Gwalani et al. [15] produced a more complex microstructure with hierarchical features of ultra-fine grains, fine-scale B2 and σ precipitates, and nanoclusters that resulted in the tensile yield strength significantly increasing from 160 to 1800 MPa. The fully recrystallized states along with grain refinement strengthening and precipitation strengthening were the focus of earlier investigations into strengthening FCC Al0.3CoCrFeNi HEAs. Jiao et al. [16] used instrumented nanoindentation to examine the mechanical characteristics of Al0.3CoCrFeNi and AlCoCrFeNi HEAs over a wide range of loading rates to determine an excellent combination of strength and ductility.

Based on the above studies, Al0.3CoFeCrNi HEAs exhibit a good balance of mechanical properties, so their effects on tribological properties are of interest. However, there have been limited systematic investigations on the room-temperature friction and wear behavior and mechanisms of single FCC-phase Al0.3CoFeCrNi, and no studies on Al0.3CuFeCrNi2 HEA. Additionally, the tribological behavior and mechanistic studies are unknown for these alloys at elevated temperatures, at which alternative structural alloys are of interest to replace bearing steels, such 440C and 52100, which oxidize and form iron oxide deleterious phases. It is at elevated temperatures such as 300 ◦C, used in this study, that mechanical and tribological processes begin to occur in metallic alloys, such as thermal softening and tribochemical (oxidation) reactions, respectively. Therefore, the objective of the present study was to investigate the influence and mechanisms of Ni content on the tribological properties and corresponding tribofilm evolution of Al0.3CoFeCrNi and Al0.3CuFeCrNi2 HEAs, both at room temperature and during 300 ◦C sliding.

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

Two alloys with different Ni content and containing Co or Cu were studied: Al0.3CoFeCrNi (A1) and Al0.3CuFeCrNi2 (A2). Alloy ingots were prepared by arc-melting and casting. The mixtures of the alloying elements with purities higher than 99.5 wt.% were melted in an argon atmosphere several times to improve the chemical homogeneity of the ingots in the liquid state. Both alloys were cold rolled to a 20% reduction in thickness with ten passes and solutionized at 1150 ◦C for 30 min, followed by water-quenching for homogenization. For microstructural and property investigation, the samples were ground and polished by standard metallographic techniques to dimensions of approximately

19 mm (length) × 14 mm (width) × 7 mm (thick) with mass of ~14 g. Surface microstructural characterization was carried out in a field emission gun scanning electron microscope (SEM) using a FEI Nova 200 dual-beam SEM. The SEM is also equipped with energy dispersive x-ray spectroscopy (EDS) for chemical analysis of the alloys before and after sliding (wear track maps). Room and elevated temperature (300 ◦C after the sliding tests) microhardness measurements were performed using a Shimadzu Vickers hardness indenter with a normal load of 9.807 N at a hold time of 10 s. A total of ten measurements were recorded for both alloys at spacings of ~1 mm apart.

The friction behavior of both alloys was Investigated using a Falex ISC-200 pin-on-disk tribometer, following the ASTM G99 standard. The sliding coefficient of friction (COF) was measured at room and elevated (300 ◦C) temperature in lab air (40% relative humidity). The current study was limited to 300 ◦C sliding since elevated temperature studies require a significant amount of time. The HEAs were tested in unidirectional sliding against a Si3N4 ball counterface (3.175 mm diameter) with hardness of 22 GPa to avoid ball wear and transfer to the wear tracks. The sliding speed was 8.5 mm/s for all tests, with a normal load of 0.25 N. Based on these values, the initial maximum Hertzian contact stress is ~0.6 GPa, which was chosen to be below the yield strength of these alloys. The total sliding distance was 200 m for all tests that took about 6 h. At least three measurements were made for each HEA for repeatability purposes. After each test, an optical microscope was used to image the worn surfaces of the HEAs and the Si3N4 counterfaces. A stylus surface profilometer (Veeco Dektak 150 Profilometer) was used to measure wear track depths. At least eight profilometry traces were taken across each wear scar to obtain the cross sectional worn area. The wear factor/rate was calculated as the removed volume loss divided by the applied load and the total sliding distance. The volume loss can be calculated by multiplying the area of the worn surface by the circumference of the circular wear track, assuming uniform wear. Crystal structures were identified with an X-ray diffractometer (Rigaku Ultima III) under radiation conditions of 30 kV, 20 mA, a CuKα anode, and a scanning speed of 2 degrees/minute. Representative wear surfaces were analyzed using SEM and EDS to acquire both secondary electron (SE) and backscatter electron (BSE) images as well as elemental wear maps, respectively. In addition, a Raman spectrometer (Thermo Electron Almega XR) was used to determine tribo-chemical phases on the wear surfaces using a 532 nm laser wavelength.

## **3. Results and Discussion**
