Microstructure, mechanical properties and tribological performance of CoCrFeNi high entropy alloy matrix self-lubricating composite

Microstructure, mechanical properties and tribological performance of CoCrFeNi high entropy alloy matrix self-lubricating composite

Materials and Design 114 (2017) 253–263 Contents lists available at ScienceDirect Materials and Design journal homepage: www.elsevier.com/locate/mat...

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Materials and Design 114 (2017) 253–263

Contents lists available at ScienceDirect

Materials and Design journal homepage: www.elsevier.com/locate/matdes

Microstructure, mechanical properties and tribological performance of CoCrFeNi high entropy alloy matrix self-lubricating composite Aijun Zhang a,b, Jiesheng Han a, Bo Su a, Pengde Li a, Junhu Meng a,⁎ a b

State Key Laboratory of Solid Lubrication, Lanzhou Institute of Chemical Physics, Chinese Academy of Sciences, Lanzhou, 730000, P.R. China University of Chinese Academy of Sciences, Beijing, 100049, P.R. China

H I G H L I G H T S

G R A P H I C A L

A B S T R A C T

• A high entropy alloy matrix self-lubricating composite was firstly reported. • The composite was successfully prepared by spark plasma sintering. • The composite has good mechanical and tribological properties.

a r t i c l e

i n f o

Article history: Received 30 September 2016 Received in revised form 16 November 2016 Accepted 19 November 2016 Available online 23 November 2016 Keywords: High entropy alloy Self-lubricating composite Microstructure Mechanical properties Tribological properties

a b s t r a c t A CoCrFeNi high entropy alloy matrix self-lubricating composite was prepared by spark plasma sintering using a mixture of CoCrFeNi high entropy alloy powder, nickel-coated graphite powder and nickel-coated MoS2 powder. The composite consisted of four phases: face-centered cubic phase of the high entropy alloy, graphite, MoS2 and nickel. Graphite and MoS2 existed in the composite were mainly due to the nickel coatings onto the lubricants and the unique advantage of spark plasma sintering. The hardness, yield strength, compressive strength and fracture toughness of the composite were 271 HV, 610 MPa, 921 MPa and 14.3 MPa·m0.5, respectively. The good mechanical properties of the composite were mainly derived from the CoCrFeNi high entropy alloy matrix. The composite exhibited excellent tribological properties from room temperature to 800 °C. From room temperature to medium temperatures, the reduction of friction coefficients and wear rates of the composite was attributed to the synergistic lubricating effect of graphite and MoS2. At high temperatures, it was found that various metal oxides formed on the composite surfaces played a key role in improving the tribological properties. © 2016 Elsevier Ltd. All rights reserved.

1. Introduction The tribological properties of sliding parts determine the efficiency, reliability and durability of the mechanical equipment that served in ⁎ Corresponding author. E-mail address: [email protected] (J. Meng).

http://dx.doi.org/10.1016/j.matdes.2016.11.072 0264-1275/© 2016 Elsevier Ltd. All rights reserved.

extreme conditions such as high temperature, cryogenic, vacuum, high-speed and heavy-load [1,2]. However, it is difficult to use conventional oil or grease lubricating systems in those extreme conditions, especially at high temperatures [3–5]. The metal matrix self-lubricating composites are highly desirable for high temperature tribological applications since their remarkable properties [3,4]. The selection of matrix of the metal matrix self-lubricating composites is an important factor

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Fig. 1. Typical SEM secondary electron (SE) images of the atomized CoCrFeNi HEA powders (low magnification image (a) and high magnification image (b)), the nickel-coated graphite powders (c) and the nickel-coated MoS2 powders (d).

according to the working conditions [4]. Over the past decades, Nibased alloys have been widely used as matrix materials due to their good mechanical properties and thermal stability [3–5]. Usually, in order to achieve good tribological properties in a wide range of temperatures, a large amount of various solid lubricants are added into Nibased alloys [3–5]. However, the addition of solid lubricants is inevitably detrimental to the mechanical properties of the composites, which may cause serious deformation, fracture or collapse in the process of applications [3,5]. Therefore, it is still a challenge to achieve good balance between mechanical properties and tribological properties in the metal matrix self-lubricating composites. In the last decade, high entropy alloys (HEAs) have shown considerable promise from both a scientific and an application perspective since they have many excellent properties [6,7]. HEAs are multi-component mixtures of element in equal or near-equal concentrations, where the high entropy of mixing benefits the formation of solid solution phases with simple structures and suppresses the formation of numerous intermetallic phases, avoiding the disadvantages of conventional multi-component alloys [6–8]. Among the huge number of HEAs, the CoCrFeNi HEA is one of the most simple alloy systems and has been widely studied [6,9,10]. Compared with the conventional Ni-based alloys, the CoCrFeNi HEA has more excellent mechanical properties and better high temperature stability [9–12]. Therefore, the CoCrFeNi HEA is a promising candidate for the matrix of metal matrix self-lubricating composites used in extreme conditions. However, until now, the design strategy of using HEAs as matrix of self-lubricating composites has rarely been reported.

Table 1 Chemical compositions of the raw powders. Powders

Composition (wt.%)

Atomized CoCrFeNi HEA powders Nickel-coated graphite powders Nickel-coated MoS2 powders

HEA N 99.5% 25% graphite 24% MoS2

b0.5% impurities 75% Ni b0.3% impurities 76% Ni b0.5% impurities

The objective of this study is to design and prepare HEAs matrix selflubricating composites and to explore their friction and wear mechanisms at high temperatures. Here, we report on a CoCrFeNi HEA matrix self-lubricating composite with excellent mechanical and tribological properties that was successfully prepared by spark plasma sintering (SPS). In the composite, molybdenum disulfide (MoS2) and graphite were used as lubricants since they have a synergistic effect of lubrication in a wide range of temperatures [5]. The graphite and MoS2 were coated by nickel, which could decrease the reaction and decomposition of the lubricants and improve the bonding strength between the HEA matrix and the lubricants. For comparison, the bulk CoCrFeNi HEA was also prepared by SPS, and its mechanical and tribological properties were investigated as well. 2. Materials and methods The bulk CoCrFeNi HEA was prepared by SPS from argon atomized CoCrFeNi HEA powders. The atomized HEA powders were filled in a graphite die and sintered in a SPS-20 T-10 furnace (Shanghai Chen Hua Technology Co., China) under vacuum. In order to obtain the HEA with high relative density, the samples were heated from room temperature (RT) to 1150 °C at a heating rate of 100 °C/min, and then were sintered at 1150 °C for 20 min under a pressure of 30 MPa, followed by furnace cooling. The CoCrFeNi HEA matrix composite was fabricated by SPS using a mixture of atomized CoCrFeNi HEA powder, nickel-coated graphite powder and nickel-coated MoS2 powder. The contents of graphite and MoS2 in the composite were 5 wt.% and 8 wt.%, respectively. In order to suppress the reaction between the lubricants and the HEA matrix, and avoid the decomposition of the lubricants during sintering, the fast SPS heating rate (150 °C/min) and short holding time (3 min) at 1150 °C were used during the preparation of the composite. The phase constitutions of the samples were characterized by X-ray diffraction (XRD, D/MAX-2400) with 40 kV operating voltage and Cu Kα radiation at a scanning rate of 5 deg./min. The microstructures of the samples were examined by scanning electron microscope (SEM, JSM5600) and the chemical compositions were investigated by energy

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Fig. 2. XRD patterns (a) and magnified main peak (b) of the atomized CoCrFeNi HEA powders, the SPSed CoCrFeNi HEA and the HEA matrix composite.

dispersive spectroscopy (EDS, Oxford instrument). The AZtec-Energy software that was installed on the EDS instrument was employed for elements contents analysis. The elements contents in the present study were the average values that were calculated from at least ten measurements. The volume fractions of the Cr-rich phase in the samples were determined by calculating the areas of the phase in SEM back scatter electron (BSE) images. At least ten different SEM-BSE images were measured with the same magnification to obtain the average values. The density of the samples was measured by Archimedes method. The hardness of the samples was measured by a HV-1000 type Vicker's hardness instrument using a load of 300 N. The microhardness of the samples was also measured by a MH-5-VM microhardness tester using a load of 2 N. Both the hardness and the microhardness measurements were repeated at least ten times to obtain the mean values. The compressive tests were performed on a CMT5202 materials testing machine (Shenzhen Sans Material Test Instrument Co., China) with a strain rate of 1.6 × 10−3 s−1. The compressive specimens with nominal dimension of Φ 5 mm × 10 mm were cut from the SPSed samples. The RT fracture toughness KIc of the samples was determined by three-point bending

tests using single edge notch binding (SENB) specimens according to ASTM E399–12. The SENB specimens were 24 mm × 5 mm × 2.5 mm in dimensions, and a straight notch of about 2.5 mm depth was prepared at the mid-length of the specimens. The tests were executed with a span of 20 mm and a nominal cross-head speed of 0.1 mm/min. Both the compressive tests and fracture toughness tests were carried out at least three times and the average values were reported. The dry friction and wear tests from RT to 800 °C in air were performed on a ball-on-disk high-temperature tribometer (HT-1000, China). The ball was made of Si3N4 with a size of Φ 6 mm, and the disk was made of the SPSed CoCrFeNi HEA and the HEA matrix self-lubricating composites with a size of Φ 25 mm × 5 mm. The test temperatures were RT, 200, 400, 600 and 800 °C, and the loads were 5 N. The sliding speed was 0.28 m/s, and the testing time was 30 min. Each test was repeated at least three times. After tribological testing, the worn surfaces were examined by SEM and EDS. The final wear volume of the specimens was measured by a surface profilometer (Micro-XAM3D, USA). In order to obtain the variations of phase structure in the

Fig. 3. Typical SEM-BSE image and corresponding EDS elemental mapping of Co, Cr, Fe and Ni of the SPSed CoCrFeNi HEA.

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Fig. 4. Typical SEM-BSE image and corresponding EDS elemental mapping of Co, Cr, Fe, Ni, S, Mo and C of the HEA matrix composite.

worn surfaces, the specimens were analyzed by a Raman spectrometer (Renishaw inVia, UK) using 633 nm wave length laser light. The worn surfaces of the samples tested at 800 °C were examined by a microbeam XRD diffractometer (Bruker, XRD) to further identify the oxides.

Table 2 Mechanical properties of the SPSed CoCrFeNi HEA and the HEA matrix composite. Materials

SPSed CoCrFeNi HEA

HEA matrix composite

Density (g/cm3) Hardness (HV) Yield strength (MPa) Compressive strength (MPa) Ultimate plasticity strain (%) Fracture toughness (MPa∙m1/2)

8.15 ± 0.03 238 ± 11 387 ± 16 N2495 (not fractured) N51 (not fractured) 24.8 ± 1.2

7.39 ± 0.05 271 ± 14 610 ± 13 921 ± 16 7.5 ± 1.2 14.3 ± 1.7

3. Results and discussion 3.1. Material characterization Morphologies of the atomized CoCrFeNi HEA powders, the nickelcoated graphite powders and the nickel-coated MoS2 powders are shown in Fig. 1, and the chemical compositions of the raw powders are listed in Table 1. The atomized HEA powders (Fig. 1(a)) were spherical in shape with an average diameter b70 μm. The high magnified SEM secondary electron (SE) image (Fig. 1(b)) shows that the surface of the HEA powders consisted of small cellular structures at submicron scales, due to the higher cooling rate and solidified in a shorter time during the process of argon atomization [13]. The nickel-coated graphite powders (Fig. 1(c)) and the nickel-coated MoS2 powders (Fig. 1(d)) had irregular shape, and the average diameters were b80 μm and 65 μm, respectively.

Fig. 5. High magnified SEM-BSE micrographs of the atomized CoCrFeNi HEA powder (a), the SPSed CoCrFeNi HEA (b) and the HEA matrix of the composite (c).

A. Zhang et al. / Materials and Design 114 (2017) 253–263 Table 3 Composition of the Cr-rich FCC phase in the atomized HEA powder, the SPSed CoCrFeNi HEA and the HEA matrix of the composite. Composition of Cr-rich phase (at.%) Materials

Co

Cr

Fe

Ni

CoCrFeNi HEA powder SPSed CoCrFeNi HEA HEA matrix of the composite

16.6 ± 2.9 12.6 ± 1.7 14.9 ± 1.6

51.1 ± 4.1 62.7 ± 2.7 54.5 ± 3.2

17.1 ± 3.5 12.5 ± 1.4 15.4 ± 1.8

15.2 ± 2.7 12.2 ± 1.2 15.2 ± 1.6

Table 4 Microhardness and volume fraction of the Cr-rich phase of the atomized CoCrFeNi HEA powder, the SPSed CoCrFeNi HEA and the HEA matrix of the composite.

Materials CoCrFeNi HEA powder SPSed CoCrFeNi HEA HEA matrix of the composite

Microhardness (HV)

Volume fraction of Cr-rich phase (%)

395 ± 21 231 ± 23 377 ± 17

35.7 ± 4.1 13.5 ± 1.7 22.6 ± 3.3

Fig. 2 shows the XRD patterns of the atomized CoCrFeNi HEA powders, the SPSed CoCrFeNi HEA and the HEA matrix composite. It can be found that there was only a single face-centered cubic (FCC) solid solution phase in the XRD pattern of the atomized HEA powders. The lattice constant of the FCC phase was 3.567 Å, which was calculated from the principal peaks of the XRD data. The SPSed HEA also consisted of a single FCC phase and its lattice constant was similar to that of the atomized HEA powders. The HEA's FCC phase, graphite phase, MoS2 phase and Ni phase could be obviously observed in the XRD pattern of the

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HEA matrix composite, and no other impurities could be found. The existence of the lubricants in the composite was mainly due to the rapidly SPS process which effectively avoided the decomposition and reaction of the lubricants during sintering. The peaks of Ni (as shown in Fig. 2(b)) appeared in the XRD pattern of the composite were attributed to the Ni coatings on the surface of the lubricants. The Ni coatings also played a key role in reducing the reaction and decomposition of the lubricants during SPS. Since the combined effect of the rapidly SPS process and the Ni coatings on the surface of the lubricants, the formation of impurities in the composites was effectively avoided. Fig. 3 shows the typical SEM micrograph of the SPSed CoCrFeNi HEA imaged by BSE and EDS elemental mapping of Co, Cr, Fe and Ni. It can be found that there was no compositional macro-segregation in the microstructure of the SPSed HEA. Pores and primary particle boundaries were rarely found in the microstructure. It was mainly due to the high sintering temperature and the long dwelling time during SPS which were beneficial for densifying the HEA sample [14]. Fig. 4 shows the typical SEM-BSE image of the HEA matrix composite and EDS elemental mapping of Co, Cr, Fe, Ni, S, Mo and C. The HEA powders distributed uniformly in the composite and they connected with each other or were bonded by the nickel-coated lubricants. The pores and voids were rarely observed in the microstructure of the composite, which suggested that the composite was close to full densification. Elemental distributions of Ni, C, Mo and S indicated that the nickel-coated graphite and the nickel-coated MoS2 distributed uniformly in the composite. The Ni coatings of the lubricants effectively suppressed the reaction between the lubricants and the HEA powders, and also improved the bonding strength between the lubricants and the HEA powders. Consequently, the composite may have excellent mechanical and tribological properties over a wide temperature range according to Refs [15,16].

Fig. 6. Typical fracture surfaces (SEM\ \SE images) of the SPSed CoCrFeNi HEA (a) and the HEA matrix composite (b).

Fig. 7. The variation of friction coefficients (a) and wear rates (b) of the SPSed CoCrFeNi HEA and the HEA matrix composite with temperatures.

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Fig. 8. Typical worn surface morphologies of the SPSed CoCrFeNi HEA: (a), (c), (e), (g) and (i) are the low magnified SEM\ \SE images of the specimens tested at RT, 200, 400, 600 and 800 °C, respectively, and the (b), (d), (f), (h) and (j) are the corresponding high magnified images of (a), (c), (e), (g) and (i), respectively.

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3.2. Mechanical properties In order to evaluate the mechanical properties of the SPSed CoCrFeNi HEA and the HEA matrix composite, the density, microhardness, compressive behavior and fracture toughness were tested, and the results are listed in Table 2. The density of the SPSed HEA was 8.15 g/cm3, which was similar to the results reported in Refs [17,18]., indicating that the high density HEA sample had been prepared by SPS successfully. The HEA matrix composite had a lower density (7.39 g/cm3) than the CoCrFeNi HEA since the added lubricants had low densities. As shown in Table 2, the RT yield strength (σy), compressive strength (σmax) and ultimate plasticity strain (εp) of the SPSed CoCrFeNi HEA were 387 MPa, N 2495 MPa and N51% (not fractured), respectively. The SPSed HEA showed excellent ductility since the FCC phase had more slip systems and higher crystal symmetry [6–9]. The σy, σmax and εp of the HEA matrix composite were 610 MPa, 920 MPa and 7.5%, respectively. Compared with the SPSed HEA, the composite exhibited lower σmax and εp. It may be due to the solid-lubricants with low strength interrupted the continuity of the composite and acted as a potential source of crack nucleation [19]. However, the hardness and the σy of the composite were obviously higher than that of the SPSed HEA. For further investigation on this abnormal phenomenon, the microstructures of the atomized CoCrFeNi HEA powder, the SPSed HEA and the composite's HEA matrix were investigated using high-magnified SEMBSE images, and the results are shown in Fig. 5. As indicated by arrows in Fig. 5, the dark gray areas (located both at grain boundaries and the interior of grains) in the microstructures were the Cr-rich FCC phase, which was identified by EDS. The compositions of the Cr-rich FCC phase are listed in Table 3. Many researchers have reported that the Cr-rich FCC phase exists in the CoCrFeNi HEA [17,18]. It is hard to detect the Cr-rich FCC phase by conventional XRD, due to the lattice constants of the Cr-rich FCC phase and the HEA's major FCC phase were very similar [17,18]. It has also been reported that the Crrich phase plays a key role in improving the strength of the CoCrFeNiMn HEA as a result of precipitation strengthening [20]. The microhardness and the volume fraction of the Cr-rich phase of the atomized HEA powder, the SPSed HEA and the HEA matrix of the composite are summarized in Table 4. It can be seen that the microhardness and the volume fraction of the Cr-rich phase had a similar variation trend. Therefore, it can be concluded that the Cr-rich phase also played a key role in improving the strength of the CoCrFeNi HEA. The atomized CoCrFeNi HEA powder had the highest volume fraction of the Cr-rich phase which led to the highest microhardness among these three materials (see Table 4). However, for the SPSed HEA and the HEA matrix of the composite, the volume fraction of the Cr-rich phase decreased significantly after SPS. In addition, comparing the SPSed HEA (sintered at 1150 °C for 20 min) and the HEA matrix of the composite (sintered at 1150 °C for 3 min), it was found that the volume fraction of the Crrich phase decreased obviously with increasing sintering time. The

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decrease of the volume fraction of the Cr-rich phase may be attributed to the SPS process promoting the dissolution of the Cr-rich phase. This phenomenon may be similar to the solution heat treatment promoting the dissolution of the γ' strengthening precipitates in Ni-based superalloys [21,22]. Therefore, compared with the SPSed HEA, the higher hardness and σy of the HEA matrix composite were mainly due to the precipitation strengthening effect, which was caused by the higher volume fraction of the Cr-rich phase in the composite's HEA matrix. In other words, the good mechanical properties of the composite were mainly derived from the HEA matrix. The RT fracture toughness (KIc) of the SPSed CoCrFeNi HEA and the HEA matrix composite were 24.8 MPa·m1/2 and 14.3 MPa·m1/2, respectively. In order to analyze the fracture mechanisms, the fracture surfaces were observed by SEM and the typical SEM\\SE images are shown in Fig. 6. A large number of dimples appeared in the fractured SPSed HEA surface (Fig. 6(a)), indicating that intense plastic deformation occurred in the HEA before fracture. This was mainly due to the HEA with FCC phase had excellent ductility and toughness [6,23]. There were a few pores and other metallurgical defects in the fractured surface of the SPSed HEA. These defects may lead to stress concentration during fracture toughness testing. This was may be the main reason why the SPSed HEA had a lower fracture toughness than the values in Refs [6,24]. As shown in Fig. 6(b), a large amount of dimples could also be found in the fractured surface of the HEA matrix composite, suggesting the HEA powders bonded well. Meanwhile, the fractured graphite and MoS2 with lamellar structures could be obviously found in the fracture surface of the composites. This was mainly due to the cracks were easily propagated in these lubricants since they had lower strength than the HEA matrix. Therefore, compared with the SPSed HEA, the lower KIc of the composite was mainly due to the soft graphite and MoS2 phases which interrupted the continuity of the hardened HEA matrix and acted as a potential source of crack nucleation. 3.3. Friction and wear behaviors The variation of friction coefficients and wear rates of the SPSed CoCrFeNi HEA and the HEA matrix composite with temperatures are shown in Fig. 7(a) and (b), respectively. It can be seen that the SPSed HEA had high friction coefficients at RT, 200, and 400 °C. As the temperatures increased to 600 and 800 °C, the friction coefficients significantly decreased to 0.37 and 0.28, respectively. The addition of lubricants significantly decreased the friction coefficients at RT, 200, and 400 °C for the HEA matrix composite. It was mainly due to graphite and MoS2 had a synergistic lubricating effect in the temperature range [5]. At 600 and 800 °C, the composite also exhibited low friction coefficients. As shown in Fig. 7(b), the SPSed HEA had high wear rates at RT and 200 °C, while the wear rates decreased significantly with increasing

Table 5 EDS results of worn surfaces of the SPSed CoCrFeNi HEA. Chemical composition (at.%) Test temperature (°C) RT 200 400 600 800 a b

Area in specimen a

Inside Outside b Inside Outside Inside Outside Inside Outside Inside Outside

This was the inside of worn surface. This was the outside of worn surface.

Co 21.7 23.7 18.3 20.2 17.9 20.3 13.6 16.8 11.3 17.2

Cr ± ± ± ± ± ± ± ± ± ±

1.1 1.2 1.1 2.1 1.3 0.8 1.9 1.3 1.3 1.7

20.6 25.3 19.1 22.6 16.9 19.3 11.8 18.5 10.4 17.5

Fe ± ± ± ± ± ± ± ± ± ±

0.9 1.3 1.2 1.8 1.4 1.4 1.3 1.1 0.5 2.3

21.0 21.2 18.5 21.9 18.1 21.1 14.1 18.1 10.8 18.1

± ± ± ± ± ± ± ± ± ±

1.3 1.8 0.8 0.9 0.8 1.2 1.5 1.2 0.9 0.9

Ni

O

21.4 ± 0.8 23.4 ± 0.7 20.4 ± 0.5 24.0 ± 0.9 16.0 ± 1.9 18.9 ± 0.7 13.4 ± 0.8 18.1 ± 1.4 9.9 ± 1.4 14.9 ± 1.3

15.3 ± 1.2 6.4 ± 1.1 23.7 ± 1.3 11.3 ± 1.4 31.1 ± 1.2 20.4 ± 1.5 47.1 ± 1.4 28.5 ± 2.8 57.6 ± 2.3 32.3 ± 2.8

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Fig. 9. Typical worn surface morphologies of the HEA matrix composite: (a), (c), (e), (g) and (i) are the low magnified SEM\ \SE images of the specimens tested at RT, 200, 400, 600 and 800 °C, respectively, and the (b), (d), (f), (h) and (j) are the corresponding high magnified images of (a), (c), (e), (g) and (i), respectively.

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Table 6 EDS results of worn surfaces of the HEA matrix composite. Chemical composition (at.%) Test temperature (°C)

Area in specimen

Co

RT

Inside a Outside b Inside Outside Inside Outside Inside Outside Inside Outside

13.3 18.6 15.1 17.5 11.2 17.0 12.3 13.5 10.5 13.2

200 400 600 800 a b

± ± ± ± ± ± ± ± ± ±

0.9 1.4 0.9 1.3 1.4 1.4 2.1 1.7 0.9 1.6

Cr

Fe

Ni

15.1 ± 1.7 16.3 ± 1.6 13.4 ± 1.1 17.7 ± 1.9 12.5 ± 0.9 16.0 ± 1.9 11.2 ± 0.9 14.6 ± 1.4 9.6 ± 1.6 12.4 ± 1.2

17.2 ± 0.9 20.4 ± 1.2 12.1 ± 0.5 17.9 ± 0.7 13.2 ± 1.1 15.0 ± 1.3 10.5 ± 1.2 15.7 ± 0.6 9.7 ± 1.1 12.0 ± 0.5

22.7 27.1 19.3 26.1 18.2 24.4 15.5 21.9 13.9 17.4

± ± ± ± ± ± ± ± ± ±

1.2 1.1 2.7 3.8 0.8 1.1 1.2 1.3 1.6 1.7

C

S

12.3 ± 1.3 5.1 ± 1.1 15.1 ± 0.8 5.6 ± 0.4 13.2 ± 1.4 4.3 ± 0.2 3.1 ± 0.5 4.1 ± 0.3 1.2 ± 0.4 6.2 ± 0.1

1.6 0.8 1.7 1.2 1.5 1.3 0 0.1 0 0.2

Mo ± ± ± ± ± ±

0.3 0.2 0.2 0.3 0.2 0.1

± 0.04 ± 0.07

0.9 0.3 0.8 0.3 1.1 0.8 0.7 0.5 0.6 0.7

O ± ± ± ± ± ± ± ± ± ±

0.1 0.1 0.2 0.1 0.2 0.2 0.2 0.1 0.1 0.1

16.9 11.4 22.5 13.7 29.1 21.2 46.7 29.6 54.5 37.9

± ± ± ± ± ± ± ± ± ±

0.7 1.3 2.2 1.6 1.8 3.1 2.4 2.0 2.5 2.7

This was the inside of worn surface. This was the outside of worn surface.

temperature. At 600 and 800 °C, the wear rates of the SPSed HEA were as low as 10−5 mm3/N·m. The HEA matrix composite always had low wear rates (10−5 mm3/N.m) from RT to 800 °C (see Fig. 7(b)). Compared with the SPSed HEA, as well as some reported self-lubricating materials such as Ni-based composites [25], Co-based composites [4] and Fe-based composites [26], it can be seen that the HEA matrix composite had excellent anti-wear and friction reducing abilities over a wide temperature range. In order to investigate the wear mechanism, the worn surfaces of the SPSed CoCrFeNi HEA and the HEA matrix composite were analyzed by SEM and EDS. As shown in Fig. 8, at RT, 200, and 400 °C, there were abundant parallel grooves and wear particles on the worn surfaces of the HEA, revealing that the wear mechanism was the typical abrasive wear. The grooves in the worn surface were produced by the micro-cutting action and the furrow action of the micro-asperities on the counterparts. The wear particles were the chips or debris produced by microcutting during friction test. Meanwhile, adhesive wear also occurred, evidenced by transfer of material with the formation of craters and patches due to delamination. The high friction coefficients and wear rates of the SPSed HEA from RT to 400 °C were mainly due to the low hardness and strength which led the HEA to be easily furrowed and cut by the counterpart. Comparing the elements contents between the inside and the outside of the worn surfaces in the SPSed HEA (see Table 5), it can be found that the oxygen contents in the worn surfaces were obviously higher than the outside of the worn surfaces for all the specimens tested at different temperatures. This indicated that friction-induced oxidation had occurred during the process of friction. As shown in Fig. 10(a), the Raman results also confirmed that various metal oxides existed in the worn surfaces. Furthermore, the friction-induced oxidation was improved by high temperatures since the oxygen contents of the worn surfaces increased with increasing test temperatures. After testing at 600 and 800 °C, a smooth glaze layer appeared on the worn surfaces of the SPSed HEA specimens (shown in Fig. 8(g)-(j)). The micro-beam XRD spectrum (Fig. 10(c)) and the Raman results confirmed that the glaze layers were composed of various metal oxides such as Cr2O3, Fe3O4, etc., which were similar to the results in Refs [27,28]. The glaze layers on the worn surfaces were very important for improving the HEA's anti-wear and friction reducing abilities at high temperatures. As shown in Fig. 9, the worn surfaces of the HEA matrix composite tested at RT, 200 and 400 °C were much smoother, and the wear widths were narrower than the SPSed CoCrFeNi HEA, suggesting that the abrasive and adhesive wear were significantly reduced by the addition of the lubricants. The EDS results (Table 6) and the Raman results (Fig. 10(b)) showed that the graphite and MoS2 gathered on the worn surfaces in the temperature range from RT to 400 °C, indicating that the solid lubricants played a key role in improving the tribological properties of the composite in this temperature range. Moreover, the EDS and Raman

results showed that the friction-induced oxidation also occurred in the composite. In Table 6, it can be seen that the oxygen contents increased significantly, while the contents of C and S decreased obviously with increasing temperature. The increase of oxygen contents was mainly caused by high temperature oxidation of the specimens, and the decrease of C and S contents may be attributed to the decomposition of the solid lubricants at high temperatures. After testing at 600 and 800 °C, the smooth glaze layers (Fig. 9(g)-(j)) also appeared on the worn surfaces of the composite specimens. According to the EDS results (Table 6), Raman results (Fig. 10(b)) and XRD spectrum (Fig. 10(d)), it can be concluded that the glaze layers mainly consisted of Cr2O3 , Fe3O4, etc., which were similar to the oxides formed on the SPSed HEA's worn surfaces after testing at high temperatures. The graphite and MoS2 were not be detected by Raman on the worn surface of the sample that was tested at 800 °C. This may be due to the decomposition of the solid lubricants at high temperature. Therefore, the metal oxides on the worn surfaces played a crucial role in improving the anti-wear and friction reducing abilities at high temperatures. In other words, the graphite and MoS2 were responsible for the lubrication of the composite at low temperatures (RT-400 °C), and the metal oxides were crucial for improving tribological properties at high temperatures (400–800 °C). The combination of solid lubricants and various metal oxides formed at high temperatures made the composite has good tribological properties over a wide range of temperatures. 4. Conclusions A CoCrFeNi HEA matrix self-lubricating composite was firstly designed and then was successfully prepared by SPS using a mixture of atomized CoCrFeNi HEA powder, nickel-coated graphite powder and nickel-coated MoS2 powder. The composite consisted of four main phases: HEA's FCC phase, graphite, MoS2 and nickel. The graphite and MoS2 uniformly distributed in the composite, and they could be successfully survived in the composite mainly due to the nickel coatings on the surfaces of the lubricants and the unique advantages of the SPS technology. The composite exhibited excellent mechanical properties, which were mainly derived from the CoCrFeNi HEA matrix. Due to the synergetic lubricating effect of the lubricants and the various oxides that formed on the worn surfaces at high temperatures, the composite had excellent self-lubrication and wear-resistance in a wide range of temperatures. The high-performance CoCrFeNi HEA matrix self-lubricating composite may have broad application prospects in the mechanical equipment which are used in the extreme conditions in the future. Acknowledgments This work was funded by the Hundred Talent Program of the Chinese Academy of Sciences (Junhu Meng).

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Fig. 10. Raman spectrum of the SPSed CoCrFeNi HEA (a) and the HEA matrix composite (b) after testing at RT, 400 and 800 °C, and the XRD patterns of the worn surfaces (after testing at 800 °C) of the SPSed CoCrFeNi HEA (c) and the HEA matrix composite (d), respectively.

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