Tribological behavior of Ti(C, N)-TiB2 composite cermets using FeCoCrNiAl high entropy alloys as binder over a wide range of temperatures

The Ti(C, N)-TiB 2 composite cermets with different binders (HEAs or Ni-Co) were fabricated by mechanical alloying and vacuum hot-pressing sintering. Wear resistance of two composite cermets at elevated temperatures was studied. Wear mechanism was characterized by a combination of scanning electron microscopy and energy dispersive spectroscopy. Experimental results indicated that HEAs binder composite cermets possessed excellent wear resistance comparing with Ni-Co binder composite cermets. At lower temperatures, no obvious difference was observed in worn surfaces of two cermets. Abrasive wear mechanism was dominant wear mechanism. At greater than 600 °C, oxidative wear and adhesive wear were found to be dominant wear mechanism. The wear rate of HEAs binder composite cermets was 11.8%, 17%, 39.25%, and 46.7% lower than that of Ni-Co binder composite cermets at 200 ℃ , 400 ℃ , 600 ℃ , and 800 ℃ , respectively. Enhanced wear performance of Ti(C, N)-TiB 2 -HEAs composite cermets is attributed to relatively high hardness and toughness, as well as excellent high-temperature softening resistance and oxidation resistance of HEAs. The microstructure and scar The The chemical compounds on the wear were by X-ray photoelectron


Introduction
WC-Co cemented carbide is widely used in machining, metallurgy, aerospace and other elds, attributing to its high hardness, strength, wear resistance, and corrosion resistance [1][2][3][4]. However, with decreasing global tungsten resources and the increase in cobalt price in recent years, alternative materials for cemented carbide are desperately needed for some aspects. After years of development, Ti(C, N)-based cermet materials have received extensive attention attributing to their high hardness, excellent wear resistance, thermal shock resistance, and high-temperature oxidation resistance. Cemented carbides have been successfully replaced in terms of nishing and semi-nishing as cutting tools [5,6].
Usually, the cutting temperature of cermet cutting tools can reach in the range 600-1000 °C, and this working condition is quite severe during the dry nishing and semi-nishing process [7]. Under this situation, cermet tools with excellent thermomechanical properties, thermal shock resistance, and hightemperature resistance friction and softening resistance are required. Therefore, as a high-temperature material, they must have excellent resistance to high temperature. High-temperature anti-wear properties have become more important than other mechanical properties such as toughness and mechanical shock resistance indicators for evaluating tool materials [8][9][10]. In recent years, the friction and wear properties of Ti(C, N) based cermets have been studied by many researchers [11][12][13][14]. Verma et al. [11] found the coe cient of friction (COF) of Ti(C, N)-based cermets with or without TaC as 0.3 to 1.1. Stewart et al. [12] reported that the addition of Mo 2 C had a detrimental effect on the sliding wear response of the Ti(C, N) based cermets. Yang et al. [13,14] investigated the high-temperature friction and wear properties of TiN-TiB 2 ceramic and found that lubricious oxidized products had a favorable effect on ceramic at high temperature. However, these studies mainly focused on the addition phase, and the effect of the binder phase of Ti(C, N)-based cermets on the friction and wear properties at high temperature has been scarcely investigated. In addition, the Fe, Co, and Ni traditional binders have high-temperature softening and poor oxidation resistance characteristics. Consequently, exploring new binders is important. High entropy alloy (HEA) is a new type of multi-component alloy with many excellent properties such as high strength, hardness, wear resistance and thermal stability [15][16]. It has been used to replace traditional adhesives to improve the high-temperature performance of cermets, which has gradually attracted researchers' attention, resulting in some outstanding ndings. [17][18][19][20][21]. However, the effect of HEAs binder on the wear properties of Ti(C, N) based cermets at high temperature is still unclear. FeCoCrNiAl HEAs possess excellent high-temperature performance [18,22]. Therefore, in this study, the FeCoCrNiAl HEAs were selected as the binder compared to the traditional Ni-Co binder.
The focus of this study was to research the high-temperature friction and wear behavior of Ti(C, N)-TiB 2 composite cermets with HEAs or Ni-Co as the binder. The COF and speci c wear rate were investigated, and wear mechanisms were analyzed. This study will provide a useful reference for the selection of cermet binders in the future.

Experimental
The cermet in this experiment was prepared by the powder metallurgy method. The compositions (wt.%) of the Ti(C, N)-based cermet in this study were Ti(C, N)-TiB 2 -(10 wt.%) HEAs or (10 wt.%) Ni-Co. The Ti(C, N) and TiB 2 mass ratio was 3:1. Commercially available Fe, Co, Cr, Ni, Al, Ti(C, N) and TiB 2 were supplied by Shuitian Technology Co., Ltd., Shanghai, China. Associated purities and particle sizes are shown in Table 1 The Archimedes principle is used to measure the bulk density of sintered samples. The Vickers hardness tester (model HVST-10, China) was used to test the hardness under a load of 98 N. The three-point bending method was adopted to test the bending strength through an electronic universal tester (China WD-10). The indentation method was used to test fracture toughness and was calculated using Palmqvist-type cracks around the indentation according to the Niihara formula represented by Eq. (1) [23]: Where H is the Vickers hardness, E is the elastic modulus (480 GPa), a is the indentation half length, l is the crack length, and Φ is the shape factor (= 3).
The friction and wear were tested using a high-temperature vacuum friction and wear testing machine (Model HVT-1000 Lanzhou, China). With regard to counter body material, the choice of cemented carbide (WC-6Co) with Ø4 mm (Linwei Steel Ball Co., Ltd., Hangzhou, China) was mainly attributed to the cemented carbide regularly used in all crushing & grinding operations and partially was replaced by Ti(C, N) based cermets in cutting tool. Under these circumstances, cemented carbide as a counter-grinding material can have a significant contrast with the research material. And the components and properties of WC-6Co ball are listed in Table 2. Specimens with Ø30×2 mm 3 dimensions were cut from the as-sintered cylindrical bars, ground, polished, and dried. The test sliding velocity, load, and sliding time were 450 rpm, 5 N, and 20 min, respectively. The tests were carried out at 200, 400, 600, and 800 °C. Wear rates were calculated by Eq. (2).
where Wr is the volume wear rate (mm 3 /Nm), is the volume loss of the material, S is the sliding distance, and F is the applied normal force. S is given by where v, r, and t are the rotational speed, track diameter, and time, respectively.
High temperature vacuum hardness tester (Model HVT-1000 Lanzhou, China) is used to test the thermal hardness of cermets under a load of 196 N. SEM was used to measure the diagonal length of the indentation. Then, the average of the two diagonal lines was calculated and the appropriate hardness was compared with the Vickers hardness standard table. The final hardness is taken from the average of five points.
The phase constituents of these wear surfaces were analyzed by X-ray diffraction (XRD) (D8 Advance, Germany). The microstructure of the specimens and wear scar was observed by scanning electron microscopy (SEM) (FEI NovaNanoSEM450, America). The chemical composition was analyzed by energy dispersive spectrometry (EDS) (Model Link-ISIS, Oxford, England). The chemical compounds on the wear surface were determined by X-ray photoelectron spectroscopy (XPS).

Microstructure and mechanical properties
The microstructures and fracture morphology of the sintered CBH and CBN composite cermets are shown in Fig. 1. Based on previous study [24], in Fig. 1(a) and (b), the gray matrix and dark gray massive particles are Ti (C, N) and TiB 2 . The white phase distributed between TiB 2 and Ti(C, N) is HEAs ( Fig. 1(a)) or Ni-Co ( Fig. 1(b)). In Fig. 1(a), the black phase is Al 2 O 3 or pores. The CBH composite cermets (Fig. 1(c)) mainly exhibited transgranular fracture mode, which could consume more energy to improve their toughness. However, the CBN composite cermets ( Fig. 1(d)) mainly have intergranular fracture, indicating that the binder and ceramic phase bind poorly. The density and mechanical properties are listed in Table 3. Clearly, the CBH composite cermets have a density of 4.992 g/cm 3 and exhibit relatively excellent mechanical properties with a Vickers hardness of 1977.3 HV 10 , fracture toughness of 7.9 MPa m 1/2 , and a bending strength of 727 MPa at room temperature. indicating that the COF values of the two composite materials have the same tendency with temperature. At the beginning of the experiment, the friction coefficient gradually increases and then reaches a steady state. In general, this process can be divided into two stages:

Friction and wear characteristics
wearing-in stage and stable wear, which are similar to the literature [25]. At the beginning, the interface is very rough, and the contact area at the micro protrusions is small, so it will be worn away, which will cause the contact point to have a cold welding effect. High shear forces are required to cut the welds, which will rapidly increase the coefficient of friction at the beginning of the running-in phase [25]. With increasing friction time, the friction coefficient became relatively stable and is mainly attributed to the contact area achieving running-in stage, and the relatively rough contact surface is replaced by a relatively smooth friction layer. With increasing testing temperature, the COF decreases first and then increases. The CBH and CBN had the highest COF of 0.36 and 0.41, respectively, at 200 °C (Fig. 3), and the minimum is 0.27 and 0.32, respectively, at 400 °C (Fig. 3). In addition, the COF increased to 0.29 and 0.37 at 600 °C for the CBH and CBN, respectively (Fig. 3), attributed to the high temperature, which softens the sample and grinding ball and forms a soft oxide layer on the sample surface, resulting in slight adhesive wear. The COF reached 0.23 and 0.35 at 800 °C (Fig. 3). This might be due to tribo-oxidation wear and adhesive wear. Relatively low COF values were obtained compared to the literature data [13]. The difference in the friction coefficient indicates that the binder (HEAs and Ni-Co) has an important effect on the COF of the material. The mechanism of the effect of the binder is discussed in detail below. Fig. 4 shows the effect of temperature and wear rate of two cermet composites, indicating that CBH has a lower wear rate than CBN at different test temperatures. With increasing test temperature, the wear rate increases first and then decreases. The relatively higher wear rate of the CBN is mainly because of the low strength of the Ni-Co binder, easily producing the wear debris. CBH composites have a relatively low wear rate, which might be the effect of HEAs binders. FeCoCrNiAl HEAs have excellent high-temperature softening resistance [26], endowing the CBH composites with excellent mechanical properties at high temperature. Table 3 shows that the CBH composites exhibited better hardness, strength, and toughness than those of the CBN composites. In general, the wear is related to the hardness and plastic deformation of materials. Eq. (3) shows that the wear volume of the material is inversely proportional to the hardness and toughness, explaining the CBH composites having a lower wear rate than that of the CBN composites [27,28].
where V is the volume wear rate, α is the material-independent constant, K C is the toughness, P is the normal load, E is the Young's modulus of materials, and H is the hardness. The test temperature plays an important role in evolution of friction and wear of materials. With increasing test temperature to 600 °C, the wear rate of the two materials reduced dramatically, and this might be relative to the change in the wear mechanism and the decreased hardness of the coupled grinding ball. Fig. 5 shows the wear surface morphology of the CBH and CBN composites at different temperatures, indicating that the testing temperature has a significant effect on the wear scar morphology of the composite, especially at >600 °C. The difference in the wear surface is not obvious at 200 °C and 400 °C, all of which are formed by the repeated action of abrasive wear, forming of a large number of shallow grooved scratches and hard particle debris of compaction. The wear mechanism is mainly abrasive wear. With increasing test temperature, the worn surface of the two materials changed obviously. This may be the effect of different binders. For the CBH composites, at 600 °C, there are two different wear scar morphologies in the middle and boundary of the wear scar. In the middle is a dense plow, which is characteristic of the abrasive wear. At the boundary, scattered wear debris and a small amount of adhesive were observed. This may be attributed to the softening of the binder and the grinding ball or the relatively soft and highly adherent oxides (TiO 2 , WO 3 , etc.), mutually transferring the tribo-layers. At 600 °C, the wear surface of the CBN composite did not show any adhesive, but large ploughs and debris appeared. This may be because the abrasive wear was produced between the larger convex hard particles on the grinding ball and the softer surface of the sample surface [29]. As the test temperature rises to 800 °C, the wear scar surface of the CBH composite has a relatively shallow and flat plow marks (Fig. 5(d)). At high temperature, the surface of the sample and the grinding ball oxidized in the oxygen environment. During the wear process, some relatively soft and highly adherent oxides (TiO 2 , WO 3 , etc) are repeatedly extruded, stretched, and finally spread on the wear surface of the sample to form a dense friction layer. Moreover, HEAs also have excellent oxidation resistance [30,31]. The above features will effectively prevent the infiltration of oxygen and the further oxidation of the substrate, thus may reduce the wear rate of CBH composites. However, the wear surface of the CBN composite still exhibits ploughs and debris of abrasive wear at 800 °C (Fig. 5(f)) despite a dense tribo-layer of particles. The CBN composite has a relatively poor density, more easily infiltrating oxygen, resulting in a thicker oxide layer. In addition, the relatively low strength and toughness make the grains easier to pull out during high-temperature wear, leading to a poor wear performance of the CBN composite.

Discussion
Wear mechanism at high temperature (600 and 800 °C) Based on the above analysis of the wear surface morphology, the wear mechanism of the material at 600 and 800 °C is discussed in detail below. In general, the wear mechanisms of cermet materials are typically brittle fracture, grain extraction, and tribo-oxidation wear, as determined by mechanical wear [32,33]. However, with increasing testing temperature, the proportion of tribo-oxidation wear will greatly increase because of the oxygen environment. Figure 7 shows the high magni cation images of the wear morphology and the grinding ball wear morphology of both composites at 600 and 800 °C. Obviously, a small amount of pits and debris are observed on the surface of the CBH composite, most of which is compacted at 600 °C ( Fig. 6(a)). However, the surface of the CBN composite was covered with a large amount of small debris. The EDS analysis (regions 1 and 5 in Figs. 6(a) and (e) Table 4) showed the presence of oxygen, indicating that tribo-oxidation wear also occurred during the wear process. Figure 6(b) shows the wear morphology of the CBH composite grinding ball at 600 °C. A small amount of adhesive was observed on the edge of the grinding ball wear scar. The combined analysis results of Fig. 5(c) and EDS analysis (regions 2 and 6 in Figs. 6(b) and (f) Table 4) indicate that the adhesive is mainly a friction layer, which is transferred from the substrate. However, the surface of the grinding ball of Fig. 6(f) has more adhesives than that observed in Fig. 6(b). These characteristics indicated that the CBN composite has strong adhesive wear at 600 °C, attributed to relatively higher high-temperature hardness of CBH than CBN and excellent high-temperature softening resistance of HEA (Fig. 7). In general, hardness is related to the dislocation slip system in the material [34]. Under high temperature conditions, dislocations are rst formed in the binder, and then the dislocations enter the hard phase. Finally, the overall plastic deformation occurs in the stressed area of the cermet [35]. The HEAs binder is prone to form twins relative to the elemental metal binder, because of its low stacking fault energy. Under these conditions, the twins will hinder the movement of dislocations, improving the hardness of ceramic materials [36]. When the material shows a slight adhesive wear, the adhesion of the friction pair contact area is lower than the shear strength of the two materials. At this time, the friction coe cient increases, the wear rate decreases, and the material transfer is not obvious. This is consistent with the adhesive wear of the CBH composites at 600 °C. However, when the adhesive wear becomes strong, the adhesion of the friction pair contact area is higher than the shear strength of the softer material in the friction pair. The damage occurs in the surface layer of the softer material near the contact surface; therefore, the soft material is transferred to the surface of the hard material. At this time, the friction coe cient is similar to slight adhesive wear, but the relative wear rate is aggravated. This is consistent with the adhesive wear of the CBN composites at 600 °C.
The EDS analysis (regions 3 in Fig. 6(c) and Table 4) shows that the wear surface forms a dense oxide layer. In this case, the tribo-oxidation wear is the main wear mechanism. In addition, the lowest wear rate of the CBH composites indicates that the dense oxide layer has a bene cial effect on the wear properties of the material. At high temperatures, both the ball and the material soften, and the surface is oxidized. In this case, the friction mainly occurs between the softening layer and the oxide layer. The plastic deformation of the friction layer results in the formation of a dense friction surface at the friction interface, and the FeCoCrNiAl HEAs have excellent high-temperature oxidation resistance [18], which can prevent further oxidation of the matrix and effectively reduces the wear. Regarding the CBN composite, the friction surface is relatively rough. The EDS analysis (regions 7 and 8 in Figs. 6(g) and (h) Table 4) shows that the substrate and grinding ball wear surface have mutual transfer of elements. In addition, not only the plow marks, but also the extraction of grains is observed. These features demonstrate that the CBN composite exhibited tribo-oxidation wear, adhesive wear, and abrasive wear at 800 °C, mainly because of their low hardness at 800 °C, causing strong wear. Figure 8 shows the depth of the wear of the two materials at 800 °C, with CBN composite exhibiting deeper wear than that of the CBH composite.
To further con rm the high-temperature wear mechanism, Ti and W elements on the worn surfaces of CBH and CBN composite cermets at 600 and 800 °C were analyzed by XPS, as shown in Fig. 9. For the CBH composite cermets, as shown in Figs. 9(a) and (c), the main peak of Ti2p of the worn surface appears at 458.4 ± 0.1 eV, which are assigned to the TiO 2 [37]. The analysis of W elements shows that the peaks also include WC besides tungsten oxide [38]. However, the peaks of CBN composite cermets did not show any obvious change. The results indicated that the worn surface mainly composed of oxides of Ti and W. In addition, the presence of the WC peak further demonstrates the occurrence of adhesive wear of the CBH composite cermets at 600 and 800 °C.

Conclusions
The wear resistance of Ti(C, N)-TiB 2 composite cermets for different binders (HEAs and Ni/Co) at hightemperature against WC-6Co was studied. The conclusions of this study are summarized as: 1. The Ti(C, N)-TiB 2 composite cermets with HEAs binder had better wear resistance than that of the composite with traditional Ni/Co binder. And as the test temperature increases, the abrasion resistance of CBH composite cermets increases. For example, the wear rate of CBH composite cermets was 11.8%, 17%, 39.25%, and 46.7% lower than that of CBN composite cermets at 200℃, 400℃, 600℃, and 800℃, respectively. 2. The wear mechanisms of the composite cermets are affected by the binder and test temperature. At 200 °C and 400 °C, the abrasive wear plays important roles. The abrasive wear and tribo-oxidation wear are the dominant wear mechanisms of the composite cermets, and mild adhesive wear was observed at 600 °C. 3. At 800 °C, the dense oxide layer of the wear surface prevents further tribo-oxidation wear, provides excellent lubrication effect, and greatly decreases the COF of the CBH composite cermets. The tribooxidation wear is the dominant wear mechanism. However, the wear mechanisms of the CBN composite cermets also include slight abrasive wear besides tribo-oxidation wear.

Con ict Of Interest
The authors declare that they have no con ict of interest. Table 2 Chemical composition, density and mechanical properties of counter body material of WC-Co. Table 3 Relevant properties of the Ti(C, N)-TiB 2 -FeCoCrNiAl HEAs/Ni-Co composite cermets. Table 4 Results of EDS analysis for the worn surfaces of the sample and the WC/Co ball in Fig. 6.   In uence of HEA binder and Ni/Co binder on wear coe cient of omposite cermet.

Figure 4
In uence of HEAs binder and Ni/Co binder on wear rate of omposite cermet.

Figure 8
The two-dimension pro le of the wear tracks of CBH (a) and CBN (b) composites at 800 ℃.