High-temperature oxidation resistance, mechanical and wear resistance properties of Ti(C,N)-based cermets with Al0.3CoCrFeNi high-entropy alloy as a metal binder

Journal Pre-proof High-temperature oxidation resistance, mechanical and wear resistance properties of Ti(C,N)-based cermets with Al0.3CoCrFeNi high-entropy alloy as a metal binder Yihang Fang, Nan Chen, Guoping Du, Mengxian Zhang, Xianrui Zhao, Hu Cheng, Jianbo Wu PII:

S0925-8388(19)33732-6

DOI:

https://doi.org/10.1016/j.jallcom.2019.152486

Reference:

JALCOM 152486

To appear in:

Journal of Alloys and Compounds

Received Date: 18 June 2019 Revised Date:

6 September 2019

Accepted Date: 27 September 2019

Please cite this article as: Y. Fang, N. Chen, G. Du, M. Zhang, X. Zhao, H. Cheng, J. Wu, Hightemperature oxidation resistance, mechanical and wear resistance properties of Ti(C,N)-based cermets with Al0.3CoCrFeNi high-entropy alloy as a metal binder, Journal of Alloys and Compounds (2019), doi: https://doi.org/10.1016/j.jallcom.2019.152486. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2019 Published by Elsevier B.V.

High-temperature oxidation resistance, mechanical and wear resistance properties of Ti(C,N)-based cermets with Al0.3CoCrFeNi high-entropy alloy as a metal binder Yihang Fanga,b, Nan Chena, Guoping Dua∗1, Mengxian Zhangc, Xianrui Zhaod,e, Hu Chengb∗2, Jianbo Wub a

b

School of Materials Science and Engineering, Nanchang University, Nanchang 330031, China

Zhejiang Provincial Key Laboratory for Cutting Tools, Taizhou University, Taizhou 318000, China c School of Mechanical and Materials Engineering, Jiujiang University, Jiujiang 332005, China d

e

Jiangsu Maritime Inst, Dept Naval Architecture & Ocean Engn, Nanjing 211170, China

Nanjing University of Aeronautics and Astronautics, Jiangsu Key Laboratory of Precision and Micro-Manufacturing Technology

Abstract Al0.3CoCrFeNi high-entropy alloy (HEA) prepared by mechanical alloying was used as a metal binder for the Ti(C,N)-based cermets by hot-pressing sintering. Room-temperature microstructure and high-temperature oxidation resistance, mechanical and wear resistance properties of the Ti(C,N)-based cermets with HEA metal binder (cermets TA) were comprehensively investigated and compared with the Ti(C,N)-based cermets containing the conventional Ni-Co metal binder (cermets TN). Different from the Ni-Co binder, the HEA metal binder not only largely enhanced the hardness and fracture toughness of Ti(C,N)-based cermets at room temperature, but also retained notably higher hardness (Vickers hardness 1137 HV), fracture toughness (6.46 MPa·m1/2) and flexural strength (761 MPa) at 1000 °C. The mechanism for such an excellent retention of mechanical properties at high-temperature by the cermets TA is attributed to the hindrance of their slip systems and the higher oxidation resistance of HEA metal binder. Furthermore, the hardened cermets TA exhibited better ∗1 ∗2

Corresponding author, E-mail: [email protected] Corresponding author, E-mail: [email protected]

friction and wearing properties at high-temperature. Keywords: Ti(C,N)-based cermets; Al0.3CoCrFeNi high-entropy alloy; Metal binder; High-temperature mechanical properties; Wear resistance

1. Introduction Ti(C,N)-based cermet has many advantages including high hardness, good chemical stability and thermal properties, outstanding wear resistance and small friction coefficient against steels and Ni-based alloys. It therefore has been used in many fields such as wear resistance parts, die materials and high-speed cutting tools for many years [1,2]. At present, the main drawback of the Ti(C,N)-based cermets for their applications in high-speed cutting/machining industry is their inadequate strength-toughness at high temperature, which is prone to flaking or chipping during the process of service [3,4]. In order to address these difficulties, different metal binders have been used to enhance the strength-toughness of Ti(C,N)-based cermets at high temperature [1,5-8]. Generally, the metal binders can be classified into two groups: intermetallic compounds and metal alloys. For intermetallic compounds, Ni3Al metal binder is most widely used. Qi et al. [5] reported that the oxidation resistance of the cermets has been improved by adding 15 wt. % Ni3Al metal binder. For the metal binders, Ni-Co is most widely used. This is because it not only reduces the sintering temperature of the Ti(C,N)-based cermets, but also stabilizes the hard phase and meanwhile refines the grains in the cermets. Xu et al. [6] obtained good mechanical properties for the Ti(C,N)-based cermets by adding 25 wt. % Ni-Co metal binder. It is well known that the high-temperature mechanical property of cutting tool materials is one of the critical factors in assessing their capability in high-speed cutting/machining [9]. Nevertheless, it has been reported that these conventional metal binders, especially at high contents (greater than 10 wt. %), are detrimental to the high-temperature mechanical properties of Ti(C,N)-based cermets [6]. Hence, it is necessary to develop a new type of metal binder which can render

outstanding high-temperature mechanical properties for the Ti(C,N)-based cermets with a lower content of binder. However, little research has been done to address the mechanical properties of Ti(C,N)-based cermets at high temperature. Recently, high entropy alloys (HEA) have been proposed to replace the conventional metal binders for Ti(C,N)-based cermets, and they are believed to be conducive to sintering of cermets and meanwhile provide a good wettability with the Ti(C,N) grains [10-12]. These alloys are solid solutions consisting of at least five major metal elements, of which the atomic percentages are equal or close to the same, and have simple BCC or FCC structures. As one of the typical HEA systems, AlCoCrFeNi has been intensively researched due to its excellent malleability, resistance to wear, thermodynamic stability, and oxidation resistance at high temperature. In fact, these high-temperature properties perfectly match the working requirements for the cermets in high speed cutting/machining operations if it is used as metal binder [13,14]. However, most of the published research has been only focused on the microstructures and oxidation resistance of the HEA metal binder itself [15,16]. Up to now, little work has been done to investigate the high-temperature mechanical properties of the cermets containing HEA metal binders. In this work, the HEA system Al0.3CoCrFeNi was selected and prepared by mechanical alloying, and it was added to the Ti(C,N)-based cermets as a metal binder to fabricate the cermets TA. For comparison, the conventional Ni-Co metal binder was also used to prepare the cermets TN. Herein, the atomic ratio of Al is lowered to 0.3 for the Al0.3CoCrFeNi HEA, which has been found to have a higher softening temperature and hardness than the AlCoCrFeNi HEA [10]. As a result, the Al0.3CoCrFeNi HEA as metal binder in the

Ti(C,N)-based cermets is expected to have better high-temperature performance. The room-temperature microstructures, high-temperature oxidation resistances, mechanical properties and wear resistances of the cermets TA and TN were comprehensively investigated. This work will provide a new approach to design cermets-based cutting tools for the high-speed cutting/machining industry.

2. Experimental 2.1 Materials and preparation procedure The raw materials included Ti(C,N) (1 ~ 4 µm), WC (< 1 µm), TaC (< 1 µm), Mo2C and metal binder of Co and Ni (< 1 µm), all from Shanghai Shui Tian Materials Technology Co., Shanghai, China, 99% purity. The Al0.3CoCrFeNi HEA was prepared from Al, Co, Cr, Fe and Ni metal powders (purity ＞99.9%, average grain size ~ 40 µm, Shanghai Shui Tian Materials Technology Co., Shanghai, China) using the mechanical alloying (MA) method. The MA preparation details of the Al0.3CoCrFeNi HEA were as follows. The Al, Co, Cr, Fe and Ni metal powders with a molar ratio of 0.3 : 1 : 1 : 1 : 1 were weighed and mixed. The MA process was carried out in a stainless steel vial filled with high purity argon. The grinding media were stainless steel balls with a media to material ratio of 15 : 1, and the grinding speed was 200 rpm. N-heptane was added to the grinding vial to avoid excessive cold welding and oxidation of the metal powders. After 25 h of grinding, the collected powders were manually ground in argon atmosphere for 30 min. The compositions (wt. %) of the cermets TA and TN are listed in Table 1. In order to improve the homogeneous distribution of the mixed powders, ultrasonic dispersion ball milling aid was used in this study. Firstly, the powder was dispersed in ethanol for 10 min by

ultrasonic vibration and mechanical stirring, and a well-mixed powder was obtained. Secondly, powder mixtures were wet-milled with WC-Co balls (ball-to-powder ratio at 10 : 1) in an ethanol bath for 24 h. The ball milled slurry of the powder mixtures was then dried in a vacuum oven at 100 °C. The dried powder mixtures were manually ground in argon atmosphere and sieved through a 100-mesh sieve. Finally, the powder mixtures were sintered under uniaxial pressure of 8 MPa at 1400 °C for 1 h by the hot-pressing technology in vacuum to obtain the cermets TA and TN. After the temperature dropped to room temperature, the sintered cylindrical samples having a diameter of 30 mm and a height of 5 mm were obtained. At the same time, the sample disks were cut into bars using an electrical discharge machine and polished to the corresponding final sizes for characterization. 2.2 Characterization Samples (size 3 x 2 x 2 mm) for the oxidation test were cut from the sintered cylinder rods, and they were ground and polished with 2.5 µm diamond slurry. Their oxidation resistance was evaluated by their increased mass after an isothermal processing in air at different temperatures (800, 900, 1000, 1050 1100 °C) for a holding time of 4 h, and three samples were tested for each of the processing temperatures. The heating rate was 10 °C/min. Before and after the oxidation experiment, the samples were weighed by a balance with a sensitivity of 0.01 mg in a closed environment. The relative density of each sintered sample was calculated from its theoretical and measured densities using Archimedes’ method. The hardness of the cermets was evaluated at room temperature and 1000 °C at a load of 200 N using a Vickers hardness testing machine (HVT-1000, Lanzhou, China) at a fixed dwell time (10 s) under a vacuum environment. Test procedures of their mechanical properties

followed the national standards of China. All samples were mechanically polished before the hardness test. The fracture toughness and flexural strength of the cermets were measured at 1000 °C in air using the 8801 high-temperature comprehensive mechanical testing system (INSTRON 8801, Instron, United Kingdom). The fracture toughness of cermets was measured by single-edge notched bending method with a loading velocity of 0.02 mm/min and a span of 30 mm. The samples for the fracture toughness test had a size of 2 x 5 x 35 mm and notches (0.1 mm wide, 3 mm deep). The flexural strength of the cermets was tested using the three-point bending test method, and the sample size was 3 x 4 x 35 mm. The values of all of these high-temperature properties of the cermets were averaged from four samples. Their room temperature mechanical properties were tested in a similar way except that the testing temperature was at 20 °C. The high-temperature dry sliding friction and wear tests were performed on high-temperature friction and wear testing machine (Model GHT-1000E, Lanzhou, China) against WC-Co carbide ball. The samples used in wear and abrasion test were machined to circular patch specimen (dimension 26 x 5 mm) and mush-room pin with 4 mm friction surface to compose friction pairs. The pin-on-disk mode was chosen with the friction speed of 200 rpm, and a load of 4.9 N (restriction by test equipment) was applied. The testing temperature and total sliding time were 900 °C and 10 min, respectively. The phase structures of these specimens were investigated using X-ray diffraction (XRD) with copper Kα radiation (Model D2, λ = 0.15406 nm, Bruker, Germany). Scanning electron microscopy (SEM, Model S-4800, Hitachi, Japan) or backscattered electron (PDBSE) mode with energy dispersive spectrometry (EDS, Model Link-ISIS, Oxford, England) and transmission electron microscopy (TEM, FEI, TF20, America) were used to observe the

microstructures of the surfaces and fracture surfaces. The polished sintered samples were hot-etched at 1000 °C for 15 minutes in static air for the observation of their grain morphologies. The thermal gravimetric and differential scanning calorimetry evaluations of the metal powders and sintered cermets were performed (TG-DSC, NETZSCH, STA 449F3, Germany).

3. Results and discussion 3.1 Microstructural characteristics XRD analysis was conducted to study whether there was a phase transition in the two sintered cermets (Fig. 1). As indicated by the XRD patterns in Fig. 1, only the Ti(C,N) phase and the corresponding metal binders were present in the cermets TN and TA. In particular, the metal binders of Ni-Co and HEA were face-centered cubic (FCC) solid solutions (Fig. 1). Furthermore, as shown in Fig. 1, the diffraction peaks of the Ni-Co binder phase in cermets TN slightly shifted to the left (lower 2θ value) of the pure Ni-Co diffraction peak, indicating that the solid solution binder phase had an increased lattice parameter after sintering, but the HEA binder phase was the opposite, with a decreased lattice constant. The larger lattice parameter of the Ni-Co solid solution binder phase could be ascribed to the larger radius of W and Mo replacing the Ni-Co atoms. The reduced lattice constant of the cermets TA could be a result of high stress due to dislocations and defects in the HEA binder [17], which was nanocrystalline as stated in the next paragraph. Figs. 2 and 3 show the microstructures of the metal binder powders of Ni-Co and HEA and the sintered cermets TN and TA. As shown in Fig. 2a, the Ni-Co binder powders had an irregular agglomeration structure with a particle size of about 1.35 µm. The as-prepared

Al0.3CoCrFeNi HEA binder powders (Fig. 2b), however, were dispersed elliptical particles with an average particle size about 24.7 µm. Furthermore, as shown in the TEM image in Fig. 2c, the Al0.3CoCrFeNi HEA particles were composed of nanocrystals with a size of about 18 nm. The small crystal size could enhance the strength of cermets [18]. The electron diffraction pattern (Fig. 2d) suggests that the as-prepared Al0.3CoCrFeNi HEA had a polycrystalline structure with strong FCC and weak BCC diffraction patterns. However, in the XRD pattern (Fig. 1c), no BCC diffraction peaks were present. This should be because the BCC phase in the HEA was too minor to be detected by XRD. The SEM images in Fig. 3a and b show the microstructures of the polished cermets TN and TA. Both of the two cermets were quite compact as no pores were observed, and their relative densities at room temperature were quite close as shown in Table 2, 99.1% and 98.8% for cermets TN and TA, respectively. Furthermore, as shown in the insets of Fig. 3a and b, the typical core-rim structures embedded into the metal binder skeletons were readily observed. It is noted that the rim structure of the black core had white inner rim and gray outer rim in the cermets TN. Previous research has shown that the inner ring was a (Ti,W,Mo)(C,N) solid solution rich in W and Mo formed by partial diffusion of the core structure [19]. Correspondingly, the outer rim was a (Ti,W,Mo)(C,N) solid solution of barren W and Mo which was formed by dissolution and precipitation. Compared with the cermets TN (Fig. 3a), the cermets TA (Fig. 3b) had a completely different core-rim microstructure. As shown in Fig. 3b, the presence of the (Ti,W,Mo)(C,N) solid solution was not obvious enough to determine whether it was the outer rim or inner rim. In addition, for the cermets TN (Fig. 3a), the thickness of the rim was uneven and did not surround the Ti(C,N) hard phase, and some of

them even exceeded 0.5 µm. It is well known that, in the Ti(C,N)-based cermet, the rim phase improves the wettability of the metal binder to the Ti(C,N) hard phase, allowing the binder phase to bond well with the hard phase. Besides, the relatively complete rim phase can increase the interfacial bonding strength between the hard phase and the binder phase, so that the crack does not easily spread along with the phase interface, thereby improving the mechanical properties. Therefore, the dissolution and precipitation process between binder powders and rim phase plays an important role in influencing the properties of cermets, particularly their mechanical properties [20, 21]. DSC test was conducted to evaluate the solubility of Ti(C,N) in the Ni-Co and HEA metal binders. As indicated in Fig. 4, the start and end solid-liquid transition temperatures of the Ni-Co binder were 1190 ℃ and 1348 ℃, respectively, while for the HEA binder they were 1269 ℃ and 1348 ℃, respectively. It means that the Ni-Co binder had longer solid-liquid transition time than the HEA binder. It is well-known that the metals easily form a low-melting eutectic liquid phase at high temperature and promotes the plastic flow of the raw material particles resulting in particle rearrangement. Both the earlier liquid phase formation temperature and the longer liquefaction time favour the wetting of the Ti(C,N) hard phase. In addition, it also facilitates the dissolution of carbides, which facilitates the formation of a complete and uniform rim phase. Thus, the Ni-Co binder had a higher solubility of Ti(C,N), and the cermets TN is expected to have relatively better room temperature mechanical properties than the cermets TA. This is also verified from the subsequent results on their mechanical properties. 3.2 Oxidation resistance

Thermogravimetric analysis was performed on the binder powders and the hot-pressed cermets TN and TA in N2-O2 mixture to investigate their oxidation resistance properties. As shown in Fig. 5, both the binder powders and cermets exhibited an initial slow oxidation weight gain at first, followed by a rapid mass increase along with the rise of temperature. Finally, the trend of oxidation weight gain slowed into a relatively stable stage. In addition, it can also be seen that the initial oxidation temperature of the Ni-Co binder powders was approximately 300 ℃ (Fig. 5a), and this was far earlier than the HEA powders which had the initial oxidation temperature at about 1000 ℃ (Fig. 5b). As shown in Fig. 5c and d, the hot-pressed cermets TN and TA exhibited a similar trend, but the latter had an initial oxidation temperature at about 1200 ℃, while the former only had about 1000 ℃. To evaluate the oxidation scale and depth of oxide layers at high temperature, the polished sections of cermets TN and TA after oxidation at 1000 °C for 4 h in static air were observed using SEM. The oxidation of both the cermets TN and TA exhibited a multi-layered structure. Fig. 6a and b shows the morphological characteristics from the exposed surfaces toward the unoxidized inner regions. In the cermets TN, a number of defects in the form of pores or microcracks were formed in the oxidation layer on surface. This suggests that the inward oxygen transport process through microcracks or porosities results in the poor oxidation resistance of the cermets TN. Moreover, the thickness of the oxidation scale was about 3.43 µm in the cermets TN, while the cermets TA had only about 1.73 µm of oxidation thickness. The oxide layer of the cermets TN was loose and porous, while the oxide layer of the cermets TA was found to be compact, stable, and continuously adherent to the surfaces. These results support the higher oxidation resistance for the cermets TA over the cermets TN.

In order to further study the oxidation resistance of the cermets TN and TA, the mass gains per unit area of cermets TN and TA were compared after an isothermal oxidation process at 1000 ℃ up to 4 h in air. As shown in Fig. 7, the oxidation mass gain curves of the cermets TN and TA were similar to the parabolic-like law in the early stage, but the oxidation tended to be linear along with the increasing oxidation time. When the oxidation time was 90 min, the mass gains per unit area of the cermets TN were almost 2.5 times that of cermets TA. In fact, the characteristics of cermets oxidation resistance were related to the anti-oxidation metal binder phase. According to Ref. [22], there are four vital effects for HEA, which are: (1) high entropy effect; (2) sluggish diffusion effect; (3) severe lattice distortion effect and (4) cocktail effects. The effects of (1), (2), and (3) could slow down the oxygen atom diffusion rate and inhibit the formation of oxides. Based on the experimental data (Figs. 5-7), the cermets TA possessed superior high-temperature oxidation resistance, and this should be a result of the excellent oxidation resistance of the multi-component HEA binder phase. Their phase structures at different oxidation temperatures were also investigated. The XRD patterns for the cermets TN and TA before and after the 4 h oxidation test at different temperatures of 800, 900, 1000 and 1100 °C are shown in Fig. 8. As shown in the XRD patterns for the cermets TN in Fig. 8A, the main diffraction peaks correspond to the Ti(C,N), TiO2, CoO, NiWO4 and Ni3TiO5 phases at 800 ℃. For the cermets TA (Fig. 8B), some phases of TiO2, CoO, NiWO4 and Ni3TiO5 in the cermets TN were also present. However, additional oxide phases such as Fe2O3, Cr2WO6, Al2TiO5 and Cr2O5 were observed in the cermets TA, but there was no Ti(C,N) phase. Due to the loose and porous structure of the oxide layer of the

cermets TN (Fig. 6a), the partial spallation of the external oxide surface was also found in the specimen exposed at 800 ℃, and thus the Ti(C,N) phase could be observed. The NiWO4 could be produced from a reaction between NiO and WO3, and the Ni3TiO5 might be from a reaction between NiO and TiO2. The NiWO4 was probably from NiO and WO3, while TiMoO5 was likely from TiO2 and MoO3. The oxide phase of Cr2WO6 in the cermets TA would be produced from a reaction between Cr2O3 and WO3, while Al2TiO5 was likely from a reaction between Al2O3 and TiO2, which is why there were no diffraction peaks of Cr2O3 and Al2O3. As the temperature continues to rise, the most significant change is the gradual enhancement of the diffraction peaks of TiO2 in cermets TN, while the changes of cermets TA was not obvious. When oxidized at 1100 ℃, the cermets TN primarily contained the phases of TiO2, Mo2C, TiMoO5, and Ni3TiO5, while the cermets TA mainly consisted of the phases of Cr2O5, Fe2O3 TiO2 and Ni3TiO5. The TiO2 was mainly formed by the diffusion of oxygen atoms into the interior and reacting with Ti(C,N). The antioxidation activity of the cermets TA could effectively prevent the diffusion of oxygen, thereby reducing the formation of TiO2. Analysis from the point of oxides view, it could be attributed to two aspects: one the one hand, since the thermal expansion coefficients of Fe2O3 and Cr2O5 were similar to the Ti(C,N) matrix, the presence of Fe2O3 and Cr2O5 could improve the anti-flaking capability of the oxide film for the cermets TA [23]. On the other hand, Fe2O3 and Cr2O5 formed by oxidation in HEA had strong anti-oxidation ability and could improve the stability of oxide layer and reduce the diffusion of oxygen and the formation of TiO2 in cermets TA [24]. According to the analysis above, the HEA binders rendered the cermets TA with a much better oxidation resistance than the cermets TN.

3.3 Mechanical properties at room and high-temperatures The mechanical properties of the cermets TN and TA at room and high temperatures (1000 ℃) are listed in Table 2. All of the mechanical properties of the cermets TN and TA largely decreased when the temperature increased from room temperature and 1000 ℃. For the cermets TN, the hardness, flexural strength, fracture toughness and elastic modulus decreased from 1939 HV to 894 HV, 1488 MPa to 826 MPa, 7.84 MPa·m1/2 to 5.55 MPa·m1/2 and 294.7 GPa to 188.5 GPa, respectively. The decrease of the hardness, flexural strength, fracture toughness and elastic modulus was about 54 %, 45 %, 29 % and 36 %, respectively. For the cermets TA, the hardness, flexural strength, fracture toughness and elastic modulus decreased from 2054 HV to 1137 HV, 968 MPa to 761 MPa, 8.92 MPa·m1/2 to 6.46 MPa·m1/2 and 296.8 GPa to 189.7 GPa, respectively. The decrease of the hardness, flexural strength, fracture toughness and elastic modulus was about 45 %, 21 %, 28 % and 36 %, respectively. By comparison, when the temperature increased from room temperature to 1000 ℃, the decrease extents of the hardness and flexural strength of the cermets TA were notably smaller than the cermets TN. This indicates that the HEA binders largely improved the high-temperature mechanical properties of cermets. In cutting applications, hardness is one of the most critical properties of cutting tools at high temperature [25]. The hardness of cermets TA at 1000 ℃ was 1137 HV, which is significantly higher than 1000 HV and maintained about 55 % of the room-temperature hardness. In general, the influence of temperature on hardness is usually related with the dislocation slip system within the materials [26,27]. Under high-temperature conditions, dislocations are first formed in the metal binder when applied stress acts on cermets. Then the

dislocations will accumulate at the rim of (Ti,W,Mo)(C,N) solid solution grain boundaries until they move into the (Ti,W,Mo)(C,N) solid solution and Ti(C,N) hard phase. Eventually, the overall plastic deformation of cermets in the stressed region is formed [28]. For the HEA metal binder, the stacking fault energy is relatively low, which enhances the formation of twins. The stacking fault energy of dislocated dislocations is relatively high, which results in the hindrance of their motion in the most active dislocation slip, improving the hardness of ceramic materials [29-31]. This effect could also be confirmed by Fig. 9 of the indentation surface after hardness test at 1000℃. For the cermets TN (Fig. 9a), the metal phase was significantly softened, and this is prone to dislocation slip. On the contrary, the HEA grains in the cermets TA (Fig. 9b) were irregular with quite a few folds, which can hinder the dislocation slip in the cermets TA grains [32]. Based on the above analysis, the hardness of the cermets TA is expected to be higher. This is consistent with the hardness results at 1000 ℃ (Table 2). Relative to a large decrease of their hardness at 1000 ℃, the fracture toughness of both the cermets TN and TA only dropped by about 28%. This suggests that the cermets could also maintain a good fracture toughness at high temperature, due primarily to the residual compressive stress originated by the thermal expansion of oxides at high temperature. The residual compressive stress could increase resistance to crack propagation and promote crack deflection, and it therefore enhances the fracture toughness. The reasons for higher fracture toughness of cermets TA may be attributed to its larger plastic deformation (Fig. 9). This plastic deformation was mainly caused by intragranular slip and grain boundary slip, thus causing a significant change in grain morphology. The plastic zone at the crack tip increased,

resulting in a decrease in the stress intensity factor at the crack tip of the plastic zone when the cermets were plastically deformed at high temperature. Accordingly, the more fracture energy was weakened by grain deformation and grain boundary slip. As mentioned above, the fracture toughness of cermets TA with HEA addition was higher than that of cermets TN at 1000 ℃. Although the flexural strength of the cermets TN was higher than that of the cermets TA at room temperature (Table 2), their high-temperature flexural strength was very close (826 MPa and 761 MPa for TN and TA, respectively, Table 2). According to the Griffth fracture theory, brittle fracture under static condition can be expressed by the following formula: ଶாఊ

σୡ = ට గ௔

(2)

where σ is the fracture energy, E is the effective elastic modulus, γ is the surface energy density of fracture, ܽ is the half length of crack. For cermets materials, the decrease of flexural strength at high temperature is mainly due to the substantial reduction of elastic modulus, oxidation effect and weakened grain boundaries. From Table 2, it is clear that the elastic modulus of the two cermets decreased by about 31 %. At the temperature of 1000 ℃, as shown in Fig. 10, voids and microcracks appeared on grain and grain boundary, which were caused by serious oxidation damage. Combined with analysis of Fig. 7 and 8, a conclusion can be drawn that the Ti(C,N) and (Ti,W,Mo)(C,N) grains and the metal binder distributed around the peripheral structure of core-rim were destroyed by oxidation during the heating and holding process. The above-mentioned voids and microcracks would cause local stress concentration, which would cause new cracks to occur, and the cracks would gradually propagate the substrate material. In addition, the free energy of atoms at the grain boundaries

would be increased, which could weaken the attraction between the particles, causing a reduction in the surface energy of the fracture while releasing residual stress. These factors worked together and resulted in a decrease in the flexural strength of the cermets TN and TA at 1000 ℃. Fig. 11 shows the SEM micrographs and EDS spectra of cermets TN and TA surfaces after thermally etched at 1000 ℃ for 15 min in static air. The EDS spectra indicated the existence of oxides on both the cermets TN and TA. For the cermets TN (Fig. 11a), a large number of oxides has been formed on the surface of the grains and grain boundaries. On the contrary, for the cermets TA (Fig. 11a), only a small amount of oxides was observed. Hence, the cermets TA with HEA metal binder had better oxidation resistance than that of cermets TN with Ni-Co metal binder. 3.4 Wear resistance at high-temperature For cutting tool application, the friction and wear characteristics represent the basis for the limitations of cutting material performance and durability [32]. The wear resistance test was carried out at 900 ℃. Their room temperature wear resistance test was not done because their hardness was too high for the test. The friction coefficient of both cermets versus time is presented in Fig. 12. In general, the dry friction process can be divided into two stages: wearing-in stage and stable wear, and the two stages were quite different. However, the wearing-in stage were both short and less than 15 s, while cermets TN was shorter (only 6 s). In the early part of sliding between friction pairs, the interface was rough and had a small contact area at the microprotrusions, which would be abraded so that the contact point created a cold soldering effect. Thus, a high shear force was needed to cut the welded point, and this is why a rapid raise of friction coefficient appeared at the beginning of the wearing-in stage

(Fig. 12). The average friction coefficient of the cermets TN was about 0.21, which was higher than the average friction coefficient of the cermets TA of about 0.13. This can be attributed to the higher high-temperature hardness for the cermets TA than TN (Table 2). Both cermets undergo abrasive wear process, in which the particles are wearied off from matrix materials and friction pairs act as abrasive. The addition of HEA can effectively improve hardness and reduce particles peeling and wear, thereby it could reduce the friction coefficient. It should be noted that Ti(C,N) and (Ti,W,Mo)(C,N) in the air can be oxidized to TiO2 when cutting. The existence of TiO2 at the friction interfaces can increase the wear resistance and decrease the friction coefficient [33, 34] In order to investigate the wear mechanism during wear test at 1000 ℃, it is necessary to investigate the microstructures of the worn surface of the cermets TN and TA after the wear test at 900 ℃. As shown in Fig. 13a and b, the groove widths of the cermets TN and TA were 172 µm and 49 µm, respectively, indicating that the cermets TA possessed better wear resistance. Additionally, the wear tracks of both cermets showed less accumulation of particles and deep grooves or indentations (Fig. 13c and d), indicating that the cold welding was obvious. This suggests that the dominant wear mechanism is adhesive wear mechanism [35]. Nevertheless, the cermets TN displayed more cold welding phenomenon with accumulation of debris either in or around the wear tracks, indicating a higher wear rate. The cermets TA, however, had lower wear volume and therefore a better wear resistance property.

4 Conclusions Two types of Ti(C,N)-based cermets with the conventional Ni-Co metal binder and the

Al0.3CoCrFeNi HEA metal binder were successfully prepared. The influence of these two types of metal binders on the cermets’ microstructures, oxidation resistance, mechanical properties at room and high-temperatures and wear resistance at high-temperature were systematically investigated. The following conclusions were drawn: (1) The cermets TN displayed typical core-rim structure embedded into a metal binder skeleton. The cermets TA had a different core-rim structure with an uneven thickness of the rim. (2) The cermets TA exhibited excellent high-temperature mechanical properties, which were generally better than the cermets TN, including a Vickers hardness of 1137 HV, the fracture toughness of 6.46 MPa·m1/2 and flexural strength of 761 MPa at 1000 ℃. The better high-temperature properties of the cermets TA than the cermets TN are attributed to the stronger hindrance of the dislocation slip systems and higher oxidation resistance of the HEA metal binder over the conventional Ni-Co metal binder. (3) The friction coefficient of cermets TA was smaller than that of cermets TN due to its higher hardness. The cermets TA exhibited better wear resistance than the cermets TN. Wear mechanism of cermets was dominated by the adhesive wear mechanism.

Acknowledgments This work was sponsored in part by the National Natural Science Foundation of China (21571095, 51362020), the Taizhou science and technology project (1802gy05), the Zhejiang Provincial basic Public Welfare Research Project (LY17E050003, 2017C31118), and the Open Foundation of Jiangsu Key Laboratory of Precision, and Micro-Manufacturing Technology.

References [1] K.Vasanthakumar, S. Ghosh, N. Koundinya, S. Ramaprabhu, S. RaoBakshi, Synthesis and mechanical properties of TiCx and Ti(C,N) reinforced Titanium matrix in situ composites by reactive spark plasma sintering, Mater. Sci. Eng. A. 759, (2019) 30-39. https://doi.org/10.1016/j.msea.2019.05.021. [2] P. Alvaredo, M. Dios, B. Ferrari, E. Gordo, Understanding of wetting and solubility behavior of Fe binder on Ti(C,N) cermets, J. Alloy. Compd. 770 (2019) 17-25. https://doi.org/10.1016/j.jallcom.2018.07.243. [3] M. Dios, I. Kraleva, Z. González, P. Alvaredo, B. Ferrari, E. Gordo, R. Bermejo, Mechanical characterization of Ti(C,N)-based cermets fabricated through different colloidal

processing

routes,

J.

Alloy.

Compd.

732

(2018)

806-817.

https://doi.org/10.1016/j.jallcom.2017.10. 274. [4] W.C. Wan, J. Xiong, M.X. Liang, Effects of secondary carbides on the microstructure, mechanical properties and erosive wear of Ti(C,N)-based cermets, Ceram. Int. 43 (2017) 944-952. https://doi.org/10.1016/j.ceramint.2016.10.024. [5] K.F. Qi, M. Yang, S.H. Li, J.B. Liu, T.J. Li, J.L. Ye, Microstructure and oxidation behavior of Ti(C,N)-based cermets with in situ synthesized Ni3Al phase, Int. J. Refract. Met. H. 73 (2018) 157-161. https://doi.org/10.1016/j.ijrmhm.2018.02. 008. [6] Q.Z. Xu, X. Ai, J. Zhao, F. Gong, J.M. Pang, Y.T. Wang, Effects of metal binder on the microstructure and mechanical properties of Ti(C,N)-based cermets, J. Alloy. Compd. 644 (2015) 663-672. https://doi.org/10.1016/j.jallcom.2015.05.

059. [7] G. Zhu, Y. Liu, J.W. Ye, Fabrication and properties of Ti(C,N)-based cermets with multi-component AlCoCrFeNi high-entropy alloys binder, Mater. Lett. 113 (2013) 80-82. https://doi.org/10.1016/j.matlet.2013.08.087. [8] H.W. Xiong, Y.X. Wu, X.P. Gan, Z.Y. Li, K.C. Zhou, Strengthened Ti(C, N)-based cermets using high-energy ball-milled NiTiC binders: Microstructure and mechanical properties, Ceram. Int. 44 (16) (2018) 19113-19122. https://doi.org/ 10.1016/j.ceramint.2018.06.080. [9] M.A. Moore, F.S. King, Abrasive wear of brittle solids, Wear, 60 (1980) 123-140. https://doi.org/10.1016/0043-1648(80)90253-7. [10] H. Cheng, X.Q. Liu, Q.H. Tang, W.G. Wang, X.H. Yan, P.Q. Dai, Microstructure and mechanical properties of FeCoCrNiMnAlx high-entropy alloys prepared by mechanical alloying and hot-pressed sintering, J. Alloy. Compd. 775 (2019) 742-751. https://doi.org/ 10.1016/j.jallcom.2018.10.168. [11] Q. Ye, G. Yang, B. Yang, Effect of aging on microstructure and property of AlCoCrFeMo0.05Ni2 high entropy alloy, Mat. Sci. Eng. A. 760 (2019) 1-6. https://doi.org/10.1016/j.msea.2019.05.058. [12] Z. Fu, R. Koc, Processing and characterization of TiB2-TiNiFeCrCoAl high-entropy alloy composite, J. Am. Ceram. Soc. 100 (2017) 2803-2813. https://doi.org/10.1111/jace.14814. [13] V. Ferrari, W. Wolf, G. Zepon, F.G. Coury, M.J. Kaufman, C. Bolfarini, C.S. Kiminami, W.J. Botta, Effect of boron addition on the solidification sequence and

microstructure of AlCoCrFeNi alloys, J. Alloy. Compd. 775 (2019) 1235-1243. https://doi.org/10.1016/j.jallcom.2018.10.268. [14] P.D. Niu, R.D. Li, T.C. Yuan, S.Y. Zhu, C. Chen, M.B. Wang, L.Huang, Microstructures and properties of an equimolar AlCoCrFeNi high entropy alloy printed

by selective

laser

melting,

Intermetallics.

104

(2019)

24-32.

https://doi.org/10.1016/j.intermet.2018.10.018. [15] G. Zhu, Y. Liu, J.W. Ye, Early high-temperature oxidation behavior of Ti(C,N)-based cermets with multi-component AlCoCrFeNi high-entropy alloy binder, Int. J. Refract. Met. H. 44 (2014) 35-41. https://doi.org/10.1016 /j.ijrmhm.2014.01.005. [16] W. Ji, J.Y. Zhang, W.M. Wang, H. Wang, F. Zhang, Y.C. Wang, Z.Y. Fu, Fabrication and properties of TiB2-based cermets by spark plasma sinteringwith CoCrFeNiTiAl high-entropy alloy as sintering aid, J. Eur. Ceram. Soc. 35 (2015) 879-886. https://doi.org/10.1016 /j.jeurceramsoc.2014.10.024 [17] R.M. Pohan, B. Gwalani, J Lee, T. Alam, J.Y. Hwang, H.J. Ryu, R. Banerjee, S.H. Hong, Microstructures and mechanical properties of mechanically alloyed and spark plasma sintered Al0.3CoCrFeMnNi high entropy alloy, Mater. Chem. Phys. 210 (2018) 62-70. https://doi.org/10.1016/ j.matchemphys.2017.09.013. [18] Z. Islam, B.M. Wang, K. Hattar, H.J. Gao, A. Haque, Departing from the mutual exclusiveness of strength and ductility in nanocrystalline metals with vacancy induced plasticity, Scripta. Mater. 157 (2018) 39-43. https://doi.org/10.1016/ j.scriptamat.2018.07. 037.

[19] J.W. Kim, S.Y. Ahn, S. Kang, Effect of the complete solid-solution phase on the microstructure of Ti(CN)-based cermet, Int. J. Refract. Met. H. 27 (2009) 224-228. https://doi.org/10.1016/j.ijrmhm.2013.01.001. [20] Z.Z. Fu, R. Koc, Ultrafine TiB2-TiNiFeCrCoAl high-entropy alloy composite with enhanced mechanical properties, Mat. Sci. Eng. A. 702 (2017) 184-188. https://doi.org/10.1016/j.msea.2017.07.008. [21] J. Zackrisson, U. Rolander, H. O. Andren. Development of cermet microstructures during sintering, Metall. Mater. Trans. A, 32A(1) (2001) 85-94. https://link.springer.com/article/10.1007/s11661-001-0104-z [22] J. Chen, X.Y. Zhou, W.L. Wang, B. Liu, Y.K. Lv, W. Yang, D.P. Xu, Y. Liu, A review on fundamental of high entropy alloys with promising high-temperature properties, J. Alloy. Compd. 760 (2018) 15-30. https://doi.org/10.1016/j.jallcom. 2019.01.386. [23] L. Chen, Z. Zhou, Z. Tan, High temperature oxidation behavior of Al0.6CrFeCoNi and Al0.6CrFeCoNiSi0.3 high entropy alloys, J. Alloy. Compd. 764 (2018) 845-852. https://doi.org/10.1016/j.jallcom.2018.06.036. [24] N.S. Bharasi, M.G. Pujar, C. Mallika, U.K. Mudali, Corrosion and Passive Film Formation Studies on Modified 9Cr-1Mo Steel in Different Sodium Hydroxide Concentrations at Room Temperature and in Boiling Condition, T. Insian. I. Metals. 70 (8) (2017) 1953-1963. https://doi.org/10.1007/s12666-016-0958-9. [25] A.Z.A. Azhar, H. Mohamad, M.M. Ratnam, Z.A. Ahmad, The effects of MgO addition on microstructure, mechanical properties and wear performance of

zirconia-toughened alumina cutting inserts, J. Alloys Compd. 497 (2010) 316-320. https://doi.org/10.1016/j.ceramint.2014.11.080. [26] C.P. Alpert, H.M. Chan, S.J. Bennison, B. Lawn, Temperature dependence of hardness of alumina-based ceramics, J. Am. Ceram. Soc. 71 (1988) 371-373. https://doi.org/10.1111/j.1151-2916.1988.tb06394.x. [27] Y.V. Milman, S.I. Chugunova, I.V. Goncharova, T. Chudoba, W. Lojkowski, W. Gooch, Temperature dependence of hardness in silicon-carbide ceramics with different porosity, Int. J. Refract. Met. H. 17 (1999) 361-368. https://doi.org/ 10.1016/S0263-4368(99)00022-0. [28] S. Bolognini, G. Feusier, D. Mari, T. Viatte, W. Benoit, TiMoCN-based cermets: high-temperature deformation, Int. J. Refract. Met. H. 21 (2003) 19-29. https://doi.org/10.1016/S0263-4368(02)00091-4. [29] R.J. Arsenault, N. Shi, Dislocation generation due to differences between the coefficients of thermal expansion, Mater. Sci. Eng. 81 (1986) 175-187. https://doi.org/10.1016/0025-5416(86)90261-2. [30] D. Dunand, A. Mortensen, Reinforced silver chloride as a model material for the study of dislocations in metal matrix composites, Mater. Sci. Eng. A. 144 (1991) 179-188. https://doi.org/10.1016/0921-5093(91)90224-B. [31] Ł. Rogal, D. Kalita, A. Tarasek, P. Bobrowski, F. Czerwinski, Effect of SiC nano-particles on microstructure and mechanical properties of the CoCrFeMnNi high entropy alloy, J. Alloys Compd. 708 (2017) 344-352. https://doi.org/10.1016/ j.jallcom.2017.02.274.

[32] Z.B. Yin, C.Z. Huang, B. Zou, H.L. Liu, H.T. Zhu, J. Wang, High temperature mechanical properties of Al2O3/TiC micro-nano-composite ceramic tool materials, Cera. Int. 39 (2013) 8877-8883. https://doi.org/10.1016/j.ceramint.2013.04.081. [33] Q.Z. Wang, F. Zhou, K. Chen, M.L.Wang, T. Qian, Friction and wear properties of TiCN coatings sliding against SiC and steel balls in air and water, Thin. Solid. Films. 519 (2011) 4830-4841. https://doi.org/10.1016/j.tsf.2011.01.038. [34] F. Zhou, K. Adachi, K. Kato, Wear-mechanism map of amorphous carbon nitride coatings sliding against silicon carbide balls in water, Surf. Coat. Technol. 200 (2006) 4909-4917. https://doi.org/10.1016/j.surfcoat.2005.04.041. [35] K. Feng, Y. Chen, P.S. Deng, Y.Y. Li, H.X. Zhao, F.G. Lu,R.F. Li, J. Huang, Z.G. Li, Improved high-temperature hardness and wear resistance of Inconel625 coatings fabricated by laser cladding, J. Mater. Process. Tech. 243 (2017) 82-91. https://doi.org/10.1016/j.jmatprotec.2016.12. 001.

Figure and table captions Fig. 1 XRD patterns for the binder powders and hot-pressed cermets at 1400 ℃ with 1 h: (a) Ni-Co powders; (b) cermets TN; (c) HEA powders; (d) cermets TA Fig. 2 Microscopic structure of the binder powders: (a) SEM image of the Ni-Co powders; (b) SEM image of the HEA powders; (c) TEM image of the HEA powders; (d) TEM diffraction spot of the HEA powders Fig. 3 Microscopic structure of the sintered samples obtained at 1400 ℃ for 1h: (a) PDBSE of the cermets TN; (b) PDBSE of the cermets TA Fig. 4 DSC curves of the milling powders of TN and TA Fig. 5 Thermogravimetric curve for the binder powders and milling powders in N2-O2 mixture (N2: 78 vol%, O2: 22 vol%): (a) Ni-Co powders; (b) HEA powders; (c) cermets TN; (d): cermets TA Fig. 6 Cross-sectional morphologies of cermets TN and TA after isothermal oxidation at 1000 ℃ for 4 h in static air: (a) PDBSE micrographs from a polished section of cermets TN; (b) PDBSE micrographs from a polished section of cermets TA Fig. 7 Mass gains curve per unit surface area relative to the oxidation time of the cermets TN and TA in air 1000 ℃ Fig. 8 XRD spectra of the oxide scales on cermets TN(A) and TA(B) before and after isothermal oxidation at 800-1100 °C for 4 h in static air (♦ Ti(C,N); ∇ Ni-Co; ∨ FCC; ♣ TiO2; ∀ NiWO4; ⇓ Ni3TiO5; ⊕ CoO; ⊗ Mo2C; ♥ TiMoO5; • WO3; ◊ Fe2O3; ∅ Cr2WO6; ∩ Al2TiO5; θ Cr2O5) Fig. 9 SEM micrographs of indentation surface after hardness test at 1000℃ (a):

cermets TN; (b): cermets TA Fig. 10 SEM micrographs after high flexural strength test at 1000℃ (a): the cermets TN; (b): the cermets TA Fig. 11 SEM micrographs and EDS spectra of cermets TN and TA surfaces thermally etched under 1000℃ in static air (a): the cermets TN; (b): the cermets TA Fig. 12 Friction coefficient of cermets TN and TA under 900℃ (a): the cermets TN; (b): the cermets TA Fig. 13 SEM images on the worn surfaces morphology under 900℃ (a): the cermets TN; (b): the cermets TA; (c): Magnification of cermets TN; (d) Magnification of cermets TA Table 1 Compositions of cermets TA and TN Table 2 Mechanical properties of cermets TN and TA at room and high temperatures.

Fig. 1

Fig. 2

Fig. 3

Fig. 4

Fig. 5

Fig. 6

Fig. 7

Fig. 8

Fig. 9

Fig. 10

Fig. 11

Fig. 12

Fig. 13

Table 1

Cermets TA TN

Components (wt. %) Ti(C,N)

WC

TaC

Mo2C

Ni

Co

Al0.3CoCrFeNi

70 70

10 10

5 5

5 5

0 8

0 2

10 0

Table 2

Relative

Vickers

Flexural

Fracture

Elastic

density

hardness

strength

toughness

Modulus

(%)

(HV)

(MPa)

(MPa·m1/2)

GPa

Room

TN

99.1

1939

1488

7.84

294.7

temperture

TA

98.8

2054

968

8.92

296.8

TN

894

826

5.55

188.5

TA

1137

761

6.46

189.7

1000

Highlights •

High entropy alloy Al0.3CoCrFeNi was used as the metal binder for the Ti(C,N)-based cermets.

•

Room-temperature microstructure and high-temperature properties of the Ti(C,N)-based cermets with the Al0.3CoCrFeNi and the conventional Ni-Co metal binder were compared.

•

Compare with the conventional Ni-Co metal binder, Al0.3CoCrFeNi remarkably improved the oxidation resistance, Vickers hardness, fracture toughness and wear resistance properties of the Ti(C,N)-based cermets at high-temperature.

•

Mechanisms for their high-temperature properties were proposed and discussed.