*3.3. Friction and Wear Behavior*

Table 3 lists the averaged steady-state COF and wear factor values for A1 and A2 alloys during RT and 300 ◦C sliding. The RT COF values of ~0.8 for both alloys during steady-state sliding are relatively high friction values that are consistent with Si3N4 sliding on Ni and other metallic materials [27,28]. However, these values are considerably higher than other HEAs with hard secondary phases; for example, COF values of ~0.3 were observed in HEA Al0.25Ti0.75CoCrFeNi [29], since harder alloys exhibit a smaller real area of sliding contact, thereby leading to lower frictional forces. The high RT friction behavior of both alloys is due to the single FCC phase being relatively soft, which is corroborated by the relatively low hardness values listed in Table 2. During sliding at 300 ◦C, the averaged COF values decreased for both alloys, but considerably more in the A2 alloy, which will be further discussed in the next section.

**Table 3.** Average steady-state COF values and wear factors (mm3/N·m) with standard deviations of A1 and A2 alloys at RT and 300 ◦C.


Representative cross-sectional wear track depth and width profiles are shown in Figure 3 for A1 and A2 alloys acquired after sliding at RT and 300 ◦C. It is evident that during RT sliding, the alloys exhibit a smaller worn area compared to sliding at 300 ◦C, despite the higher COF values. The cross sectional worn areas were used to calculate the corresponding wear factors listed in Table 3. The wear factors listed in Table 3 agree well with the RT COF trends, i.e., there are similar wear factors for both alloys at RT. However, during 300 ◦C sliding, there is a slightly lower wear factor for the A2 alloy, since there was no thermal softening, as evidenced by the increase in hardness. Compared to RT sliding wear factors, both alloys have higher wear factors, suggesting that in addition to hardness, there are other factors active that will be discussed in the next section. Since both HEAs exhibit single FCC-phase structures without the presence of typical secondary phases present in structural alloys, e.g., hard intermetallic or carbide phases, the wear factors are about an order of magnitude higher. For example, the aforementioned Al0.25Ti0.75CoCrFeNi BCC HEA has hard intermetallic L21 and χ phases responsible for lower sliding wear rates in the order of 1 × <sup>10</sup>−<sup>5</sup> mm3/N·m [29]. However, a similar single FCC-phase HEA Al0.1CoCrFeNi to that of A1 alloy, but with lower Al content and slightly softer, exhibits a higher RT sliding wear factor of 1.9 × <sup>10</sup>−<sup>4</sup> mm3/N·m [30]. Based on the above, single FCC-phase HEAs do not provide adequate wear resistance compared to HEAs with hard secondary phases, making them similar to other structural alloys such as bearing steels with chromium carbide precipitates. Only when hard secondary phases are present do HEAs such as Al0.25Ti0.75CoCrFeNi have comparable wear rates to those measured for chromium carbide bearing steel 440C, values of ~1 × <sup>10</sup>−<sup>5</sup> mm3/N·m [29].

**Figure 3.** Representative cross-sectional wear track depths and widths of A1 and A2 alloys after RT and 300 ◦C sliding.

#### *3.4. Friction and Wear Mechanisms*

Figure 4 shows representative SEM images and EDS maps of the A1 alloy wear track after RT sliding. The wear track shows signs of abrasive wear, i.e., microabrasion with grooves along the sliding direction. There is also oxidative wear based on the dark features in the BSE images (meaning a higher atomic number contrast) that coincides with the oxygen EDS map in Figure 4. Based on the elemental maps, there is not one particular metal that shows a preference for oxidative wear. This suggests that there is a mixed metal oxide tribofilm on the wear surface. These flattened oxide patches inside the wear track are indicative of a surface fatigue wear mode that results in metallic oxide wear fragments delaminating from the surface. With repeated sliding, the wear fragments can either be ejected from the sliding contact or become entrapped beneath the Si3N4 counterface, with the latter pathway contributing to the formation of micro-grooves by a three-body abrasive wear mode.

**Figure 4.** BSE and SE images and corresponding EDS elemental maps of the A1 alloy wear track after RT sliding.

Figure 5 shows representative SEM images and EDS maps of the A2 alloy wear track after RT sliding. Similar to the A1 alloy, there is micro-ploughing/micro-grooving inside the wear track, indicating a micro-abrasive wear mode with striations running parallel to the sliding direction and across the entire wear track length. There are also oxidative/surface fatigue wear modes present, although the wear track width is slightly smaller, and there is slightly less micro-abrasion/metal oxide tribofilm covering the wear track. This accounts for the slightly lower RT COF and wear factor values listed in Table 3 for the A2 alloy.

**Figure 5.** BSE and SE images and corresponding EDS elemental maps of the A2 alloy wear track after RT sliding.

Figure 6 shows SEM images and EDS maps of the A1 alloy wear track after 300 ◦C sliding. It is evident there is an increased amount of oxidative wear compared to RT sliding for this alloy, based on the BSE image and corresponding oxygen EDS wear map covering almost the entire wear track. Furthermore, it appears that this oxide tribofilm is also a mix of all the metallic elements, based on the EDS maps that do not show a preference for any particular metal oxide phase. Despite the lower COF due to the lower interfacial shear strength oxide tribofilm, the wear track width has increased in size compared to the RT wear track. This is supported by the increased wear factor listed in Table 3, due to thermal softening for this A1 alloy, which is based on the lower hardness value of HV1 = 144. More severe adhesive wear also occurs in the wear track shown by the SEM images in Figure 6, since these wear tracks are covered with a compact tribofilm (often referred to as an oxide glaze layer), indicating a change in wear mechanisms. Abrasive grooves along the sliding

direction are still visible on the wear track in areas not covered by the oxidized tribofilm. In addition, there is evidence of fragmented oxide wear debris that further acts similar to abrasive particles to accelerate the wear process, resulting in the high wear factor. Therefore, while the low interfacial shear strength oxide tribofilm provides low COF values, it not protective at this elevated sliding temperature compared to RT sliding.

**Figure 6.** BSE and SE images and corresponding EDS elemental maps of the A1 alloy wear track after 300 ◦C sliding.

In contrast, the A2 alloy exhibits increased hardness, up to HV1 = 188 at 300 ◦C, which results in a slightly smaller wear track, as shown in Figure 7, and thus a smaller frictional area of contact during sliding. Hence, the A2 alloy has the lowest COF of ~0.29 and a slightly lower wear factor of 9.8 × <sup>10</sup>−<sup>5</sup> mm3/N·m compared to the A1 alloy at 300 ◦C, but this is still higher compared to RT sliding. Hardening in the A2 alloy is likely due to the increased amount of Ni resulting in the formation of a NiO scale, although verification would be needed with cross-sectional microscopy and micro-indentation. The wear track shown in Figure 7 exhibits slightly less oxidative wear compared to the A1 alloy, based on the images and EDS oxygen map. In addition, the fragmented oxide wear debris act as micro-abrasive particles, resulting in increased wear. Like the A1 alloy, the low interfacial shear strength oxide tribofilm results in low COF values, but overall does not provide wear protection at this elevated sliding temperature compared to RT sliding. Similarl to other alloys such as the CoCr-based alloy (Haynes 25), at temperatures greater than 200 ◦C, the friction coefficient is reduced due to the formation of a protective oxide layer, which minimizes the adhesion between the two contacting surfaces [31]. When the temperature becomes greater than 150 ◦C, the sintering rates of the metal oxide films increase, which leads to the formation of glazes (Co and Cr oxides) that can supply prolonged protection against friction and reduce the friction coefficient [32]. In addition, low unidirectional sliding friction coefficients for Co-based alloys at temperatures greater than 200 ◦C can be attributed to the formation of thermally stable oxide glazes on the pin surface, which cause low friction and wear [28].

The Si3N4 counterfaces showed some abrasive wear with extruded wear debris at RT, while at 300 ◦C, the counterfaces exhibited similar features, along with some adhered metal oxide transfer films from the wear tracks. A similar study revealed there is no accumulation of third bodies or transfer films adhered to the Si3N4 balls after RT sliding for HEAs Al0.1CoCrFeNi and CoCrFeMnNi [30]. In order to better determine the tribochemcial oxide phases, Raman spectroscopy was performed inside the four wear tracks shown in Figures 4–7. Figure 8 shows representative Raman spectra for the alloys after RT and 300 ◦C sliding. It is evident there are several oxide phases, both binary oxides and multi-element solid solution oxides, present in both alloys, including CoO, Co3O4, Cr2O3, NiO, and multielement Cr2O4 and Fe2O4 tribochemical phases. These metal oxides tribochemically form at intermediate and higher temperatures via tribo-sintering the compact oxide tribofilm, which has been shown to lower the friction coefficients [28,33]. These oxide phases are in good agreement with the Raman spectra acquired inside FeCrNi medium-entropy alloy and CoCrFeNi HEA wear tracks, respectively [33,34]. The higher-Ni content A2 alloy at 300 ◦C exhibited a higher intensity peak of NiO and Cr2O3 in the tribofilm, which could be responsible for its lower friction and wear. Future studies will explore sliding temperatures higher than 300 ◦C to determine if there are more protective high-temperature oxide phases in both alloy tribofilms that result in low interfacial shear strength for friction reduction, while simultaneously providing low wear. It is possible that at higher sliding temperatures than 300 ◦C, the thermal softening process is counteracted by the formation of more protective tribochemical oxide phases that act as solid lubricants, thereby lowering friction and wear.

**Figure 7.** BSE and SE images and corresponding EDS elemental maps of the A2 alloy wear track after 300 ◦C sliding.

**Figure 8.** Raman spectra acquired inside A1 alloy and A2 alloy wear tracks after RT and 300 ◦C sliding.

#### **4. Summary and Conclusions**

Two FCC single-phase HEAs alloys with different Ni contents and either Co or Cu were studied: Al0.3CoFeCrNi (A1) and Al0.3CuFeCrNi2 (A2). For the A1 alloy with lower Ni content, micro-indentation and sliding wear tests revealed that the hardness decreased, resulting in thermal softening and a higher wear factor during 300 ◦C sliding. In contrast, the higher Ni content A2 alloy exhibited increasing hardness and subsequently a slightly lower wear factor. Mechanistic wear studies showed this was due to the oxidative wear, with the formation of low interfacial shear strength tribofilms that covered the wear tracks. Raman spectra determined that the A2 alloy at 300 ◦C exhibits a higher intensity peak of

NiO and Cr2O3 in the oxide tribofilm, which is likely responsible for lowering both the COF and wear factor. However, compared to RT sliding, both alloys provide no wear protection during 300 ◦C sliding, most likely due to thermal softening in the tribofilms. Thus, these single FCC-phase HEAs provide no further benefit to wear resistance at elevated temperatures, with similar implications likely for other such single FCC-phase HEAs. Lastly, these and other single-phase HEAs without hard secondary phases are no better than current bearing steels.

**Author Contributions:** Conceptualization, D.F.K. and T.W.S.; methodology, D.F.K., M.V.K. and T.W.S.; formal analysis, D.F.K., M.V.K. and T.W.S.; investigation, D.F.K., M.V.K. and T.W.S.; data curation, D.F.K. and M.V.K.; writing—original draft preparation, D.F.K. and T.W.S.; writing—review and editing, D.F.K. and T.W.S.; supervision, T.W.S. All authors have read and agreed to the published version of the manuscript.

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

**Data Availability Statement:** All data generated or analyzed during this study are included in this published article.

**Acknowledgments:** We thank Bharat Gwalani for providing the two alloys, and the UNT Materials Research Facility (MRF) for access to SEM/EDS and XRD.

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