Sliding wear of CoCrNi medium-entropy alloy at elevated temperatures: Wear mechanism transition and subsurface microstructure evolution

Journal Pre-proof Sliding wear of CoCrNi medium-entropy alloy at elevated temperatures: Wear mechanism transition and subsurface microstructure evolution Shuai Pan, Cancan Zhao, Pengbo Wei, Fuzeng Ren PII:

S0043-1648(19)31324-9

DOI:

https://doi.org/10.1016/j.wear.2019.203108

Reference:

WEA 203108

To appear in:

Wear

Received Date: 3 September 2019 Revised Date:

25 October 2019

Accepted Date: 26 October 2019

Please cite this article as: S. Pan, C. Zhao, P. Wei, F. Ren, Sliding wear of CoCrNi medium-entropy alloy at elevated temperatures: Wear mechanism transition and subsurface microstructure evolution, Wear (2019), doi: https://doi.org/10.1016/j.wear.2019.203108. 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.

Graphical Abstract

Sliding Wear of CoCrNi Medium-Entropy Alloy at Elevated Temperatures: Wear Mechanism Transition and Subsurface Microstructure Evolution Shuai Pana, b, Cancan Zhaoa, Pengbo Weia, c, Fuzeng Rena,* a

Department of Materials Science and Engineering, Southern University of Science

and Technology, Shenzhen, Guangdong 518055, China b

Department of Materials Science and Engineering, Harbin Institute of Technology,

Harbin 150001, PR China c

Department of Mechanical and Aerospace Engineering, The Hong Kong University

of Science and Technology, Clear Water Bay, Kowloon, Hong Kong, China ∗Corresponding author. E-mail address: [email protected]

Abstract To gain a comprehensive understanding of the wear mechanism and the combined effect of low stacking fault energy (SFE) and external elevated temperature on sliding-induced plastic deformation of CoCrNi medium-entropy alloy (MEA), herein, we report on the wear response of fine-grained CoCrNi MEA against Inconel alloy 718 counterparts between room temperature (RT) and 300 °C, with particular focus

on

the wear

mechanism

transition

and

sliding-induced

subsurface

microstructure evolution. The results show that the hardness of the MEA and coefficients of friction (CoFs) start to decrease at 200 °C, but wear rates monotonously decrease with rising temperatures. The wear mode changes from 1

abrasive wear at RT to oxidative and adhesive wear at 200 °C. Between RT and 150 °C, stacking faults and deformation twins play a significant role in the formation of gradient subsurface microstructure. The improved wear resistance is mainly attributed to the thermal softening of the mating material and the increased contribution of adhesive wear. However, at 200 °C and above, the reduced wear rates and CoFs are associated with the formation of glaze layer. The present findings provide insights into understanding the wear mechanism and sliding-induced deformation of metallic alloys with low SFE at elevated temperatures.

Keywords:

CoCrNi,

medium-entropy

alloy,

microstructure

2

wear

mechanism,

subsurface

1. Introduction Sliding wear of metallic contacts is a complex process involving in plastic deformation underneath the sliding surface, fracture, materials transfer, mechanical mixing, and even chemical reactions with environment [1]. Moreover, sliding wear-induced plastic deformation shows distinct features from the traditional bulk deformation modes, attributing to large strain gradients [2], high strain rates and strain rate gradients in the subsurface region [3, 4]. At elevated temperatures, the contact surfaces may also experience thermal softening [5] and tribo-chemical reactions which lead to the formation of the ‘glaze’ layer and consequently affect wear performance [6]. A thorough understanding of such a complex process requires detailed characterization of the worn surface, the generated wear debris, and particularly the subsurface microstructure [7-9]. High-entropy alloys (HEAs) with multi-principal elements, are a new class of emerging advanced materials with novel alloy design concept [10, 11] and have demonstrated attractive properties that are not attainable in conventional alloys with one principal element [12]. One of the interesting property is that the HEAs shows improved wear resistance as compared to the traditional alloys with the same level of hardness [13, 14]. In the past fifteen years, sliding wear behavior of HEAs has thus attracted a lot of research interest. Early investigations mainly focused on exploring the effects of addition of alloying elements on the phase, microstructure, hardness and thus, the wear performance of HEAs [13, 15, 16]. Moreover, HEAs have also been considered potential candidates for high-temperature structural applications, due to

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the sluggish diffusion effect, excellent thermal stability, and good oxidation resistance [17]. This motivates researchers to explore the wear behavior of HEAs at elevated temperatures. The complex mechanical and thermal environment interplay at elevated temperatures makes the wear mechanisms distinct from those at room temperature (RT). For instance, sliding wear of CoCrFeNi-based HEAs at elevated temperatures has shown that the wear mechanism changes from abrasive wear at RT to the oxidation and delamination wear at high temperatures due to the formation of oxide films [18-21]. In comparison with HEAs, some medium-entropy alloys (MEAs) have been found to exhibit even better mechanical properties [22, 23]. This has been demonstrated by comparative studies between a series of single-phase face-centered cubic (fcc) MEAs comprising the elements in the CoCrFeMnNi HEA, which show that there has been no systematic correlation between mechanical properties and number of principal elements or configurational entropy [24, 25]. CoCrNi, a representative MEA with a very low stacking fault energy (SFE < 5 mJ/m2) [26], has demonstrated excellent tensile properties and fracture toughness at cryogenic temperatures [22], owing to the twinning induced plasticity effect [27], formation of 3D twin structure during deformation [28], and fcc to hexagonal close-packed (hcp) phase transformation at lager tensile strain [29] or under high strain rate [30]. Despite that significant achievements have been made in understanding the basic mechanical properties and deformation mechanisms at cryogenic temperatures [22, 27, 28, 31], limited resources have been directed towards the wear behavior and the

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sliding-induced subsurface microstructure evolution at elevated temperatures. The only available literature reports that addition of molybdenum (Mo) could increase the wear resistance of CoCrNi MEA through the formation of σ and µ phases [32]. To gain a comprehensive understanding of the wear mechanism and the combined effect of low SFE and external elevated temperature on sliding-induced plastic deformation of CoCrNi MEA, in the present study, we report on the wear response of fine-grained CoCrNi MEA against the typical high-temperature alloy Inconel 718 counterparts between RT (25 ± 1 °C) and 300 °C, with particular focus on the wear mechanism transition and sliding-induced subsurface microstructure evolution.

2. Experimental methods 2.1. Fabrication of the CoCrNi MEA The CoCrNi MEA was produced via a combination of mechanical alloying and spark plasma sintering (SPS). Commercially pure cobalt (1.6 µm, 99.8%, Alfa Aesar), chromium (< 10 µm, 99.2%, Alfa Aesar) and nickel (-325 mesh, 99.8%, Alfa Aesar) powders with an equiatomic ratio were subjected to high energy ball milling using a SPEX 8000D mill at ambient temperature in an argon glove box for predefined time intervals. The 12 h-ball milled powder was then packed into a graphite die with an inner diameter of 10 mm and then consolidated by SPS (SPS-211Lx, Fujidempa Kogyo. Co., ltd, Japan) at 1050 °C and 50 MPa for 5 min.

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2.2. Hardness measurements and pin-on-disk wear testing Since wear empirically correlates material’s hardness, the Vickers hardness values of both the CoCrNi MEA and the mating Inconel alloy 718 disks at the predefined wear temperatures were measured using a high-temperature Vickers diamond pyramidal indenter (Archimedes Industrial Technology Co., Ltd., London, United Kingdom) under a load of 2 kgf for 10 s. At least five individual measurements were performed at each temperature and the average values with standard deviations are provided. Prior to wear tests, the CoCrNi MEA was cut into 3 mm in diameter and 6 mm in length from the bulk cylindrical samples using electrical discharge machining. Both contact surfaces were mechanically polished with SiC papers down to 1200 grit. For direct comparison of the wear behavior of CoCrNi MEA with the conventional Inconel-series superalloys, the conventional high-temperature alloy Inconel 718 with hardness of ~ 350 HV was selected as the mating material. Pin-on-disk wear tests (Anton-Parr High Temperature Tribometer, Austria) in accordance with ASTM G99– 17 were performed in air under a load of 5 N (with the nominal contact pressure of ~ 0.71 MPa) at a constant sliding velocity of 0.1 m/s with a total sliding distance of 1000 m between RT and 300 °C. Wear rates were calculated by direct measurement of weight loss of the pins to accuracy of ± 0.1 mg. Three independent tests were run for each temperature, and average wear rates and coefficients of friction (CoFs) with standard deviations are provided.

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2.3. Materials characterization The phases of the powders, the as-sintered alloy and the generated wear debris were identified by X-ray diffraction (XRD) recorded by a diffractometer in a 2θ range from 40° to 100° using a Cu-Kα radiation (Rigaku Smartlab-9 kW, 45 kV, 200 mA) with a step size of 0.02° and speed duration time of 10° min−1. The morphology, elemental distribution and grain size of the as-sintered CoCrNi MEA were characterized by scanning electron microscopy (SEM; TESCAN MIRA 3, Czech Republic) equipped with energy dispersive X-ray spectroscopy (EDX) and electron backscatter diffraction (EBSD). Surface morphology and composition of the worn pins and disks were examined by SEM and EDX. X-ray photoelectron spectroscopy (XPS, Thermo K-Alpha+, America) was used to identify the chemical composition of the glaze layer formed at 200 °C and 300 °C. Phase, morphology and chemical composition of the generated wear debris were characterized by XRD, SEM and EDX, respectively. Cross-sectional subsurface microstructures of the pins after wear were characterized by SEM, EBSD, transmission electron microscopy (TEM), high-resolution TEM (HRTEM), and high-angle annular dark-field scanning TEM (HAADF-STEM; Thermo ScientificTM TalosTM F200i, FEI) with attached EDX operated at 200 kV. All TEM samples were prepared by focused ion beam (FIB; Helios NanoLab™ 600i, FEI, USA) milling using the site-specific standard lift-out technique. The directions normal to the sliding surface, along the sliding direction and perpendicular to the sliding direction in the sliding plane were defined as ND, SD and TD, respectively.

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3. Results 3.1 Phase and microstructure of the CoCrNi MEA Fig. 1a shows XRD patterns of starting powder mixture and the powders after ball milling for different time intervals. After 3 h of ball milling, peaks corresponding to hcp-Co are absent while those of bcc-Cr are still visible, suggesting that hcp-Co has dissolved into the fcc-Ni phase. Differing from the previous study [33] which reported that a certain amount of Cr-like bcc phase remained even after 35 h of ball milling, here, when the milling time extends to 12 h, a single-phase solid-solution with an fcc structure is formed. This should be attributed to the high-energy ball mixer used in our study. The as-sintered alloy with relative density exceeding 98% preserves the single fcc phase (Fig. 1b) with intense and narrow diffraction peaks owing to grain growth and stress release during SPS. Lattice parameter of the fcc phase is measured to be ~ 3.564 Å. The backscattered electron (BSE) image (Fig. 2a1) and corresponding EDX elemental maps (Fig. 2a2-a4) further confirm the single-phase structure and uniform distribution of the constituent elements in the alloy. EBSD inverse pole figure (IPF) map (Fig. 2b) shows no pronounced preferred orientation of the grains. The corresponding phase map (Fig. 2c) also confirms the single fcc structure of the alloy. Statistical measurements (Fig. 2d) show that the grain size ranges from 200 nm to 4.5 µm, with an average of 1.0 µm. Actually, over 60% of the grains are smaller than 1 µm.

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3.2 Hardness Fig. 3 presents the hardness varying with temperature of both the CoCrNi MEA and the Inconel alloy 718 disks. For the CoCrNi MEA, the measured hardness remains stable from RT to 150 °C, with the value of ~ 380 HV. Then, the hardness values start to decrease with rising temperatures, from 374 (± 28) HV at 200 °C to 350 (± 36) HV at 300 °C. In contrast, the hardness values of Inconel alloy 718 disks monotonously decreases from 349 (± 3) HV at RT to 319 (± 5) HV at 300 °C, suggesting that the present CoCrNi MEA has better thermal softening resistance than the Inconel alloy 718 between RT and 150 °C. The relatively large fluctuations of the hardness values of the CoCrNi MEA should be attributed to the multi-leveled grain size microstructure.

3.3 Coefficients of friction and wear rates CoFs and wear rates of the CoCrNi MEA varying with temperatures are presented in Fig. 4. The wear rates monotonously decreases with the temperature, from 1.15 (± 0.01) × 10‒3 mm3/N·m at RT to 8.16 (± 0.02) ×10‒5 mm3/N·m at 300 °C, approximately 14 times reduction. The CoF evolution with sliding distances at different temperatures (Fig. S1) suggests that the tribo-pairs reach steady state wear upon sliding distance over about 200 m. The CoFs are relatively stable upon wear from RT to 150 °C, in the range of 0.64-0.66. However, when the temperature goes higher, the CoF starts to decrease, with 0.62 at 200 °C to 0.59 at 300 °C.

9

3.3 Wear debris The generated wear debris at different temperatures shows no significant difference in morphology (Fig. S2). All are in the form of big flakes with size ranging from several tens microns to several hundred microns. EDX analyses reveal that the wear debris is sourced from both the Inconel alloy 718 alloy disks and CoCrNi MEA pins, denoted as fcc1 and fcc2 phases, respectively. However, XRD patterns show phase evolution of the wear debris generated during wear from RT to 300 °C (Fig. 5). The deconvolution of the (111) reflections (Fig. S3) clearly shows the variation of the relative contents of the two fcc phases with the wear temperatures. Between RT and 150 °C, the wear debris is predominantly sourced from the CoCrNi MEA pin. However, at 200 °C, the relative content of fcc1 in the debris significantly increases, almost equal to that of fcc2. Interestingly, the relative content of fcc1 in the wear debris further increases with increasing temperature. At 300 °C, the XRD pattern of the debris shows almost a single fcc1 phase, suggesting that the debris is mainly sourced from the mating material. Combined with CoFs and the phase component evolutions with temperatures, we can deduce that a wear mechanism transition occurs at the critical temperature of 200 °C. Thus, the following analyses would focus on two representative temperature ranges: (i) RT – 150 °C and (ii) 200 °C – 300 °C.

3.4 Worn surfaces

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Fig. 6 shows a comparison of typical surface morphology and composition of pins after sliding wear at RT, 150 °C, 200 °C and 300 °C. Here, it should be mentioned that since the Inconel alloy 718 disk also contains ~ 52 wt.% Ni and ~ 20 wt.% Cr, Co and Fe are traced to distinguish the pin and disk materials, respectively. At RT, many parallel grooves along SD were observed on the worn surface (Fig. 6a1), characteristics of abrasive wear. This is consistent with the morphology of wear debris (Fig. S2a1). The uniform distribution of Co with a trace amount of Fe and O (Fig. 6a2-a4) on the worn surface suggests almost no material transfer from the disk to the pin surface at RT. Upon wear at 150 °C, besides the grooves, smearing and fracture of materials were also observed (Fig. 6b1). A small amount of disk material was transferred onto the worn pin surface, as confirmed by the richness of Fe in some local surface regions (Fig. 6b2-b3). Within this temperature range, the wear mechanism does not change significantly. From 200 °C to 300 °C, the grooves diminished and the concentrations of Fe and O were found to become much higher in the surface, suggesting that more disk material was transferred to the pins. The wear mode changes to adhesive wear as well as oxidative wear. To better understand the wear mechanism transition and the tribo-oxidation process, surface morphology and composition of the wear tracks on the Inconel alloy 718 disks at varying temperatures (Fig. S4) were also analyzed. At RT and 150 °C, some parallel grooves along with scattered patches and pits were observed. A minor amount of Co is present, indicating a small amount of material transfer from the CoCrNi MEA pin onto the mating disk. (Fig. S4a2-a4). At 200 °C, a small amount of

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pin material adhered to the disk surface and the transferred materials were oxidized (Fig. S4c2and c4). Upon wear at 300 °C, a tribo-oxidation layer was formed on the disk surface (Fig. S4d1). In contrast to high concentrations of Fe and O, a very limited amount of Co was detected, revealing that the oxidation layer was mainly sourced from the mating disk (Fig. S4d2-d4).

3.5 Cross-sectional microstructure of the CoCrNi MEA In order to understand the wear mechanism transition at elevated temperatures, we have performed detailed characterization of the subsurface microstructures of CoCrNi MEA pins. We first used SEM coupled with EDX elemental maps and EBSD to give a broad overview of the subsurface microstructures, and then used analytical TEM to give a close observation on gradient microstructures. The ND-SD cross-sectional microstructure (Fig. 7a1-a4) of worn pins at 150 °C shows no formation of glaze layer, suggesting pure metallic contacts of the tribo-pairs. However, at 200 °C, a glaze layer with a thickness of 5-6 µm was formed (Fig. 7b1). EDX elemental maps (Fig. 7b2-b4) show that the glaze layer contains Co, Fe and O, sourced from both the pins and the mating disks. When temperature is further increased to be 300 °C, besides with a larger thickness (7-8 µm), the glaze layer is found to be deficient in Co but rich in Fe and O (Fig. 7c1-c4). Quantitative EDX analysis of the chemical composition of the glaze layer at 200 °C, 250 °C and 300 °C is included in Table S1. With the rising temperature, the content of Co in the glaze layer apparently decreased. In contrast, Ni, Cr, Fe and O significantly increased. The

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composition change reveals that the glaze layer formed at high wear temperatures is mainly sourced from mating material. To provide deep insights into the chemical state of Co, Cr, Ni and O in the glaze layers, XPS analysis of the worn CoCrNi MEA pin surfaces at 200 °C and 300 °C was conducted. A comparison of the high-resolution XPS spectra of Co 2p, Cr 2p, Ni 2p and O 1s of the glaze layer formed at 200 °C and 300 °C is presented in Fig. 8. The results show that at 200 °C, the Co is in the form of predominantly unoxidized metal state and a small amount of Co3O4; the Cr2O3 and unoxidized Cr have almost same intensity; the Ni remains mainly unoxidized, but with a minor amount of NiO. However, at 300 °C, the total amount of Co in the glaze layer significantly reduced, as evidenced from the significant reduction in intensity, in agreement with above EDX quantitative analysis. The Co is more in the form of oxides. Besides Co3O4, a small amount of Co2O3 is also present. The Cr is almost fully oxidized. Only a trace amount of Cr remains unoxidized. More interestingly, NiCr2O4 is the dominant oxide formed at 300 °C, due to the reaction of oxides at this temperature. Fig. 9 shows EBSD IPF and corresponding Kernel Average Misorientation (KAM) maps of ND-SD cross-sections of CoCrNi MEA pins after wear tests at RT, 100 °C, and 150 °C. The grains in the topmost regions (presented by black pixels) can hardly be resolved by EBSD due to severe plastic deformation and the small grain size. Below the unidentified regions, the grains are heavily deformed along SD and refined (as marked by the yellow box in Fig. 9 a, d and g). The average grain sizes (d) of these deformed layers at RT, 100 °C and 150 °C are measured to be ~ 630 nm, 680

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nm and 600 nm, respectively, which are considerably smaller than their adjacent undeformed base material. This layer at a specific wear temperature is thus referred to as deformed ultrafine-grained layer (DUGL). The DUGLs generated after wear at RT, 100 °C and 150 °C have the thickness of 14 µm, 12 µm and 7 µm, respectively. Deformation twins are found in this layer (as marked in white dashed circles). The grains have weak textures, as indicated by the pole figures (Fig. 9c, f and i). The KAM maps (Fig. 9b, e and h) also reveal that the wear-induced strain level reduced from RT to 150 °C. To facilitate a direct comparison with the gradient microstructure of the worn pins at wear temperature of 150 °C, The ND-SD cross-sectional grain structures beneath the glaze layers of the pins after wear tests at 200 °C, 250 °C and 300 °C were also analyzed by EBSD (Fig. 10). Except the topmost glaze layers, the sliding-induced ND-SD cross-sectional microstructures at 200 °C and 250 °C also gradient microstructure. Interestingly, more severe plastic deformation is found at 300 °C and the plastically deformed grains are almost parallel to SD, despite with the presence of glaze layer. Quantitative KAM distributions at varying temperatures are summarized in Fig. 11. From RT to 150 °C, the peak positions shift to smaller KAM angles and the peaks became higher and more concentrated, indicating gradual decrease in the degree of deformation. However, an inverse trend was observed in the temperature range of 200 °C to 300 °C, where the peak position shift to larger KAM angles and relative frequency of all peaks decreased, suggesting more severe deformation at higher wear temperatures.

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To obtain more detailed subsurface microstructure information, especially those failed to be resolved by EBSD after wear tests between RT and 150 °C, and to extract the chemistry, crystallography, and microstructure information in the glaze layer at the transition temperature of 200 °C, TEM and HAADF-STEM analyses have also been performed. Fig. 12 presents the subsurface microstructure of the CoCrNi MEA pin after sliding wear at RT. A nanocrystalline layer (NL) (d = ~ 60 nm) with a thickness of 1 µm was found in the topmost region (Fig. 12a and b1). Selected area electron diffraction (SAED) pattern (Fig. 12b2) shows continuous rings of a single fcc phase, indicating the random orientations of the grains. Below the NL is the above mentioned DUGL with a mixed deformed structure. In this layer, a high density of dislocation cells, nanometer-scale deformation twins (Fig. 12c1) and a large amount of stacking faults are present (Fig. 12d1). The non-uniform intensity of the SAED patterns suggests weak texture (Fig. 12c2), in agreement with the pole figures presented in Fig. 9. Upon wear at 150 °C, the formed subsurface microstructure shows no significant difference from that at RT, except reduced thickness of plastic deformation region and the nanocrystalline layer (~ 700 nm) (Fig. 13). Fig. 14 presents a detailed characterization of the microstructure and composition of the glaze layer formed at 200 °C. The glaze layer can be further divided into two sub-regions according to their microstructure and composition features. The top region with a depth of 2-3 µm consists of nanocrystalline grains (Fig. 14a and b1) with relatively uniform distribution of Co, Cr, Ni, Fe and O (Fig. 14b2-b6). EDX quantitative analysis shows that the atomic ratios of Ni and Cr are over 90% and

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40% higher than that of Co, suggesting that this region contains a significant amount of transferred disk material (Inconel alloy 718). SAED pattern (Fig. 14f) further confirms that this nanocomposite region consists of randomly oriented nanocrystalline (Co, Cr, Ni, Fe)-based oxides, CoCrNi MEA and transferred Inconel alloy 718. The bottom region, with the depth extending to several microns, is a vortex-like mechanical mixture consisting of long fragments of the transferred disk material (bright contrast), the CoCrNi MEA (gray contrast) and a small amount spherical oxides (black contrast) (Fig. 14a & c1). EDX elemental maps (Fig. 14c1- c6) and quantitative analysis (Fig. 14d) further reveal that the CoCrNi MEA is rich in oxygen but almost no oxygen was detected in the long fragments of disk material, suggesting that CoCrNi MEA is more easily oxidized than the Inconel alloy 718. In addition, micro-pores and cracks are also found in this region.

4. Discussion This work has systematically investigated dry sliding wear behaviors of CoCrNi MEA against Inconel alloy 718 disk from between RT and 300 °C, with particular focus on wear mechanism and the sliding-induced gradient microstructure evolution. The keys findings are: (i) at wear temperatures between RT and 150 °C, sliding wear induces the formation of gradient microstructure in MEA, and the elevating temperature mitigates the sliding-induced plastic deformation; and (ii) a compacted glaze layer starts to form at 200 °C, and consequently leads to the wear mode transition and the reduction of CoF and wear rates. Thus, the following discussion

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would mainly focus on the formation of the gradient microstructure and the evolution of wear mechanisms.

4.1 Formation of gradient microstructure Upon wear between RT and 150 °C, the gradient strain and strain rate imposed by sliding wear can refine the near surface microstructure of the CoCrNi MEA and the degree of refinement decreases along with the depth away from the surface. The formation of nanocrystalline layer (with d = ~ 60 nm) of the CoCrNi MEA pins suggests that the strain during sliding wear is sufficiently high for microstructural refinement. The gradient microstructures are similar to the previous findings in the dynamic plastic deformation (DPD)-treated Cu-Al alloys [34] and ‘hard turning’ -processed CoCrNi MEA [31]. Previous works have demonstrated that the formation of gradient microstructure induced by surface mechanical attribution or grinding treatment [35, 36], DPD [34] or high pressure torsion (HPT) [37] are intrinsically rooted in two fundamental deformation mechanisms: dislocation slip and deformation twinning, while the dominant mode for a specific metal/alloy strongly correlates with its SFE and initial grain size, as well as external loading conditions [38]. For coarse-grained alloys with low SFE, deformation twinning governs the formation of gradient structure and microstructural refinement. In this work, the CoCrNi MEA has a heterogeneous grain structure with grain sizes ranging from 200 nm to 4.5 µm, and an average of 1.0 µm. Over 60% of the grains are smaller than 1 µm. Thus, the effect

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of grain size on the plastic deformation and microstructure refinement should also be considered. During sliding wear at RT, planar slip is activated at the low strain level (the bottom region of the DUGL in Fig. 9) owing to the limitation of dislocation movement, which is caused by the lowing SFE and short-range order (SRO) effect of the CoCrNi MEA [12, 39, 40]. The full dislocation dissociates into partials to form stacking faults as the twin nuclei [41]. The low SFE provides a low threshold stress of twinning formation and facilitates the transformation from stacking faults to deformation twins [27, 29]. However, the threshold stress to form deformation twins in ultrafine-grained or nano-grained material can be written as [42]:

τtwin =

Gb dm

(1)

where G is the shear modulus, b is the magnitude of Burgers vector for partial dislocation, d is the grain size and m is the shear factor. Also, the Hall-Petch slope of twinning is much larger than that of the slip of full dislocations, indicating a strong dependence of grain-size on twinning stress [43]. So, the critical twin stress of small grains is higher than that of large ones. Hence, at the low strain level, deformation twins can only be found in large grains (of micron-scale) in the CoCrNi MEA. For those ultrafine grains (< 1 µm, account for ~ 60%), dislocations are emitted and trapped at grain boundaries, rather than grain interior [44]. As the strain increases upon approaching the sliding surface, for large grains, deformation twins pass through the whole grains and many twin bundles form. As a result, deformation twins with the increasing density subdivide original coarse grains into nanoscale twin/matrix (T/M) 18

lamellae, which are indicated by red arrows in the Fig. 15a. Then, dislocations formed between the T/M lamellae arrange themselves into interconnecting boundaries to minimize the further increasing strain energy (Fig. 15b) [45]. Secondary twins will also be induced in this process. The interactions between dislocations and deformation twins increase the misorientations of dislocation boundaries and primary grains eventually evolve into subgrains or new nanograins. For the ultrafine grains, the increasing strain will stimulate dislocations trapped at grain boundaries and form dislocation cell structures (as indicated by yellow arrows in Fig. 15a). When the local stress exceeds the critical stress required for twin nucleation, nano-scale deformation twins will also form in the ultrafine grains (Fig. 15c and d). This is also same to the new formed nanograins generated from original coarse grains. With further increasing strain (near the NL), the density of deformation twins further increases. The dislocations will also pile up at the twin boundaries and stacking faults, which make the original atomically flat coherent twin boundaries and stacking faults evolve into curved incoherent high-angle grain boundaries [37]. When the grain size decreases below 100 nm, stacking faults and nano-twins are still found in the grains in the nanocrystalline layer (Fig. 12d1), but the strain level at the sliding surface is not sufficient to further promote the multiplication of dislocations and the formation of nanoscale twins at this grain size level (d = 60 nm). Regarding the influence of increasing temperature of sliding wear tests on the plastic deformation and microstructure refinement, it can be explained by the parameter Z (Zener-Hollomon parameter) [46]:

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InZ = ln ε& +

Q RT

(2)

where ε& is the strain rate, Q is the activation energy for diffusion, R is the gas constant, and T is the temperature. Due to the stable CoF and hardness of CoCrNi MEA, the strain rate can be supposed to be the same between RT and 150 °C. The decrease of the value of lnZ at elevated temperature will result in the reduction of the slip stress of deformation [47]. Thus, the dislocation slip would become more significant at elevated temperature at the low strain level (at the bottom of DUGL). Also, it should be noted that the threshold for nanotwin formation in the CoCrNi MEA increases with temperature, as observed during plastic deformation at RT than cryogenic temperature [29, 31]. This may explain the reduced depth of plastic deformation region from RT to 150 °C, but low SFE alloys are not sensitive to lnZ like high- or medium- SFE alloys at the high strain level (in the NL and the top region of DUGL), and therefore deformation twins are still the dominant mechanism of microstructure refinement during sliding wear at 150 °C. Regarding the evolution of the subsurface microstructure between 200 °C and 300 °C, the sliding-induced plastic deformation in the subsurface region is more severe with the elevating temperature, despite with the presence of the glaze layer (Figs. 10 and 11b). This should be attributed to thermal softening of the CoCrNi MEA. However, the effect of thermal softening, the glaze layer and oxidation reaction during sliding wear on the deformation microstructure still needs to be further systematically investigated in future.

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4.2 The evolution of wear mechanism Between RT and 150 °C, one interesting observation is that the hardness and CoFs of the CoCrNi MEA remain unchanged, suggesting the similar wear mechanism, but the wear rates continuously decrease with the rising temperature. Since no oxygen is found in the wear debris and no glaze layer is formed on the worn surfaces of the pins (Figs. 5-7), the metal-to-metal sliding contact explains the relatively higher CoFs in this temperature range (Fig. 4b). The wear rate reduction with rising temperature seems in contradiction with previous observations in sliding wear of nickel-based alloys [48] and other HEAs [18, 20, 49] at elevated temperatures, all of which show that the wear rate increases with the rising temperature before the formation of glaze layer, due to thermal softening of alloys [20, 50-52]. This nominally unexpected result is due to the complex nature of the present tribo-pairs. The CoCrNi MEA pins show stable hardness, but the thermal softening is present in the mating Inconel alloy 718 disks (as revealed by their reduced hardness in Fig. 3), which promotes the adhesion and bonding of the transferred CoCrNi patches to the mating disks during rubbing of the contacting surfaces (Fig. S4b1). The transferred CoCrNi MEA thus reduces the direct asperity contacts between the pin and the disk surfaces, and could thus contribute to the reduction of the wear rates and the strain accumulation in the subsurface regions of the pins. In addition, despite with dominant wear mode of abrasion, the role of adhesive wear becomes more significant at higher temperature (Fig. S4b1-b4), which can also help to reduce wear rate.

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The oxidation reactions during wear between 200 °C and 300 °C play significant roles. The wear mode changes from abrasive wear between RT and 150 °C to oxidative wear and adhesive wear at 200 °C, due to the formation of a compacted glaze layer in the topmost surface region of the CoCrNi MEA pins. The glaze layer prevents direct metal-to-metal contacts of the tribo-pairs and consequently reduces both the wear rate and CoF. Recently, several works have also highlighted the importance of the formation of glaze layer on the wear scars of HEAs during high temperature wear [18-20]. The glaze layer in this work consists of two sub-regions. The formation of the top cohesive and compacted nanocomposite region (Fig. 14) can be explained by the extreme surface plastic deformations during sliding wear [1, 53]. The formed oxides are mainly Cr2O3 and Co3O4. Viat et al. [51] investigated the contribution of main alloying elements (Co, Cr, Ni) in Haynes 25 alloy in the formation of glaze layer during fretting wear. They reported that the oxidation and sintering processes were mainly controlled by the cobalt element, due to the formation of Co3O4-based oxides, such as (Ni, Cr)xCo3-xO4. The Co3O4 oxide has a better sinterability than other oxides, such as Cr2O3, and NiO, because of the highest self-diffusion coefficient of cobalt [51]. The bottom part of glaze layer is composed of a vortex-like mechanical mixture of the pin and the disk materials along with spherical nanoscale oxides (Fig. 14c1-c6 and g). The EDX analysis (Fig. 14d) suggests that CoCrNi MEA gets oxidized more easily than the Inconel alloy 718 at 200 °C. This could be explained by the Gibbs free energy of the oxides. The Gibbs free energy values of CoO, Co3O4, Cr2O3 and NiO at 500 K (227 °C) are -198.9 kJ, -705.9 kJ,

22

-991.0 kJ, and -194.7 kJ, respectively [54]. The Gibbs free energy values to form Cr2O3 and Co3O4 are much lower than those of CoO and NiO. Thus, thermodynamically, oxygen prefers to react with CoCrNi MEA when it diffuses into the glaze layer and thus, resulting in the formation of Cr2O3 and Co3O4 at 200 °C. With increasing temperature, more Inconel alloy 718-sourced debris were generated (Figs. 5 and S3) due to the thermal softening of the disks (Fig. 3) and thus, the content of Co in the glaze layer decreased and conversely the contents of Ni, Cr and Fe increased (Table S1). This also suggests that CoCrNi MEA has better wear resistance than Inconel alloy 718 above 250 °C. The presence of more micro-pores and cracks in the glaze layer (Fig. 7 a1-c1) should be attributed to the insufficient sintering of the Niand Cr-based oxides in this temperature range.

5. Conclusions A systematical investigation of the dry sliding wear behaviors of CoCrNi MEA with low SFE against Inconel alloy 718 between RT and 300 °C was conducted. The following major conclusions were drawn: (1) The spark plasma-sintered CoCrNi MEA has a single fcc structure with heterogeneous distribution of grain size, ranging from 200 nm to 4 µm with an average of 1.0 µm and over 60% grains smaller than 1 µm. (2) The hardness of the CoCrNi MEA and CoF start to decrease at 200 °C, but wear rates monotonously decrease with rising temperatures. The wear mechanism

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changes from abrasive wear between RT and 150 °C to adhesive and oxidative wear at 200 °C. (3) Between RT and 150 °C, without the presence of glaze layer, a gradient microstructure with a high density of dislocation cells, nanoscale deformation twins and a large amount of stacking faults was formed in the CoCrNi MEA. Stacking faults and deformation twins play a significant role for the formation of the gradient microstructure. The decreased wear rates in this temperature range should be mainly attributed to the thermal softening of the mating disk and the increased contribution of adhesive wear. (4) A glaze layer starts to form at 200 °C, which consists of two sub-regions: (i) the top one has a nanocomposite structure consisting of randomly oriented nanocrystalline oxides and the contacting components, and (ii) the bottom one is a vortex-like mechanical mixture of the long fragments of oxygen-rich CoCrNi MEA and the transferred disk material along with a small amount spherical oxides. The presence of glaze layer is responsible for the reduced wear rates and CoFs at high temperature.

Acknowledgements This work was financially supported by the Fundamental Research Program of Shenzhen (Grant Nos. JCYJ20170307110418960 and JCYJ20170412153039309), and Guangdong Innovative & Entrepreneurial Research Team Program (No. 2016ZT06C279), China. This work was also supported by the Pico Center at

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SUSTech that receives support from Presidential fund and Development and Reform Commission of Shenzhen Municipality, China.

Appendix A. Supplementary Material Evolution of CoFs as a function of sliding distance, characterization of the wear debris, deconvolution of the (111) reflections of the XRD patterns of the wear debris, surface morphology and EDX elemental maps of the mating disks, and chemical composition of the glaze layers are included.

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Figures

Fig. 1 Phase evolution of the Co, Cr, and Ni powders during ball milling and SPS. (a) XRD patterns of Co, Cr, and Ni powders after ball milling for time intervals; (b) XRD pattern of the as-sintered CoCrNi MEA.

1

Fig. 2 Elemental distribution and microstructure of the as-sintered bulk CoCrNi MEA. (a1) BSE image; (a2-a4) are corresponding EDX elemental maps of Co, Cr, and Ni, respectively; (b) EBSD IPF map; (c) phase map corresponding to (b); and (d) grain size distribution.

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Fig. 3 The measured Vickers hardness versus temperature of CoCrNi MEA and Inconel alloy 718.

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Fig. 4 Wear rates (a) and coefficients of friction (b) of CoCrNi MEA upon sliding wear against Inconel alloy 718 at varying temperatures.

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Fig. 5 XRD patterns of the generated debris during wear at different temperatures, in comparison with those of CoCrNi MEA and Inconel alloy 718.

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Fig. 6 Surface morphology and EDX elemental maps of the worn CoCrNi MEA pins after wear tests at different temperatures.

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Fig. 7 SEM images and corresponding EDX elemental maps of ND-SD cross-sections of the CoCrNi MEA pins after sliding wear at different temperatures.

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Fig. 8 XPS analysis of the glaze layer formed on top of the worn CoCrNi pin surface after sliding at 200 °C and 300 °C. (a)-(d) are high-resolution XPS spectra Co 2p, Cr 2p, Ni 2p, and O 1s, respectively.

8

Fig. 9 ND-SD cross-sectional EBSD IPF maps (a, d, and g) and the corresponding KAM maps (b, e, and h) of CoCrNi MEA pins after sliding wear at RT, 100 °C, and 150 °C with pole maps (c, f, and i) of regions marked by yellow box. The white dashed circles in (a), (d) and (g) denote the regions containg deformation twins.

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Fig. 10 ND-SD cross-sectional SEM images, EBSD IPFs, and corresponding KAM maps of CoCrNi MEA pins after sliding wear at 200 °C, 250 °C, and 300 °C.

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Fig. 11 KAM distributions of the CoCrNi MEA after sliding wear at varying temperatures.

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Fig. 12 ND-SD cross-sectional TEM analysis of CoCrNi MEA after sliding wear at RT. (a) a bright field TEM image; (b1) a high magnification bright field TEM image of region 1 in (a); (b2) electron diffraction pattern of the selected area in (b1); (c1) a high magnification bright field TEM image of region 2 in (a); (c2) electron diffraction pattern of the selected area in (c1); (d1) a high-resolution TEM image from the nanocrystalline layer; (d2) and (d3) are corresponding FFT patterns of the matrix and SFs selected in (d1).

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Fig. 13 ND-SD cross-sectional TEM analysis of CoCrNi MEA pin after sliding wear at 150 °C. (a) a BF-TEM image; (b1) a high magnification BF-TEM image; (b-c) corresponding SAED patterns of region 1 and 2 in (a), respectively; (d) a HRTEM image; (e) corresponding FFT pattern of (d).

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Fig. 14 ND-SD cross-sectional TEM analysis of the glaze layer formed at 200 °C. (a) a HAADF-TEM image; (b1) a high magnification HAADF-TEM image of region 1; (b2-b6) corresponding EDX elemental mappings of (b1); (c1) a HAADF-TEM image at the depth about 5 µm; (c2-c6) corresponding EDX elemental mappings of (c1); (d) chemical compositions of region 1 in (b1) and area 2 and 3 in (c1); (e-f) a BF-TEM image near the worn surface and corresponding SAED pattern; (g-h) a BF-TEM image at the depth about 5.5 µm and corresponding SAED pattern.

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Fig. 15 (a-b) TEM image of deformed ultrafine-grained layer ; (c) high-resolution TEM image obtained from the nanocrystalline layer and (d) inverse fast Fourier transform (iFFT) and FFT (inset) images of the selected area in (c).

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Highlights

• Wear mechanism and sliding-induced plastic deformation of CoCrNi alloy was investigated. • Sliding wear tests were performed between 25 °C and 300 °C. • The hardness and CoFs start to decrease at 200 °C but wear rates keep decreasing. • The wear mode changes from abrasive wear to oxidative and adhesive wear at 200 °C. • The sliding-induced gradient microstructure and the glaze layer were characterized.

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Declaration of interests The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.