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Article

Influence of Plasma Arc Current on the Friction and Wear Properties of CoCrFeNiMn High Entropy Alloy Coatings Prepared on CGI through Plasma Transfer Arc Cladding

by
Peihu Gao
1,2,*,
Ruitao Fu
1,
Jilin Liu
3,
Baiyang Chen
1,
Bo Zhang
1,
Daming Zhao
1,
Zhong Yang
1,2,
Yongchun Guo
1,2,
Minxian Liang
1,2,
Jianping Li
1,2,*,
Wei Wang
4,
Zhiyi Yan
4 and
Lina Zhang
4
1
School of Materials and Chemical Engineering, Xi’an Technological University, Xi’an 710021, China
2
Shaanxi Province Engineering Research Centre of Aluminum/Magnesium Light Alloy and Composites, Xi’an 710021, China
3
Representative Office in Datong of PLA, Datong 037036, China
4
Shanxi Diesel Engine Co., Ltd., Datong 037036, China
*
Authors to whom correspondence should be addressed.
Coatings 2022, 12(5), 633; https://doi.org/10.3390/coatings12050633
Submission received: 28 March 2022 / Revised: 20 April 2022 / Accepted: 3 May 2022 / Published: 5 May 2022
(This article belongs to the Special Issue Structural, Mechanical and Tribological Properties of Hard Coatings)
Browse Figures
Versions Notes

Abstract

:
High-entropy alloys receive more attention for high strength, good ductility as well as good wear resistance. In this work, CoCrFeNiMn high-entropy alloy (HEA) coatings were deposited on compacted graphite iron through plasma transfer arc at different currents. The microstructure and wear properties of the CoCrFeNiMn HEA coatings were investigated. The coatings are composed of single phase with FCC structure. The CoCrFeNiMn HEA coating had the highest microhardness of 394 ± 21.6 HV0.2 and the lowest wear mass loss when the plasma current was 65 A. All of the HEA coatings had higher friction coefficients than that of the substrate. There were adhesive, abrasive and oxidation wear forms in the HEA coatings with the wear couple of N80 alloy. The HEA coating presented higher friction coefficient and better wear resistance than compacted graphite iron.

1. Introduction

High-entropy alloy (HEA), a new type of multi-principal metallic materials, usually contains 5–13 major elements with equimolar or nearly equimolar ratios of each element, which has attracted more and more attention due to their unique microstructures and properties [1,2,3,4]. The number of continuous phases in HEAs is much lower than that predicted according to the Gibbs phase law for HEAs’ relatively low mixing enthalpy and high entropy, resulting in low Gibbs free energy of the HEAs system [2]. The low Gibbs free energy stabilizes the simple solid solution in the alloy and inhibits the formation of intermetallic compounds [5]. High entropy alloys actually exhibit good mutual solubility and are capable of forming simple face-centered-cubic (FCC), body-centered-cubic (BCC) or hexagonal-close-packed (HCP) solid solution phases [6,7,8], as well as some intermetallic phases [9] or even amorphous phases [10]. The microstructural and compositional characteristics of high entropy alloys will have four main effects, including high mixed entropy, severe lattice distortion, slow diffusion and cocktail effects, which endow high-entropy alloys with many excellent properties, such as ultra-high strength [11,12], high ductility [13], excellent catalytic properties [14], excellent irradiation resistance [15], high corrosion resistance [16] and excellent wear resistance [17].
In general, there is a relationship between the material’s hardness and its wear resistance. HEAs have higher hardness and better wear resistance than conventional alloys [18]. Parisa Moazzena et al. [19] added Zr to the NiCoCrFe alloy so that the micro-hardness and nano-hardness of NiCoCrFe alloy increased from 682 ± 7 and 672 ± 7 to 828 ± 10 and 845 ± 10 Vickers, respectively. Moreover, the addition of Zr led to significant increasing in wear resistance and decreasing the coefficient of friction. In the work of Cheenepalli Nagarjunaa et al. [20], an equiatomic CoCrFeMnNi high entropy alloy (HEA) was fabricated by a rapid solidified gas atomization process. The hardness of HEA bulks increased from 270 ± 10 to 450 ± 10 Hv with increasing milling time, while the lowest coefficient of friction 0.283 and specific wear rate 1.03 × 10−5 mm3/Nm were obtained for the 60 min milled HEA. In Mohamed AliHassana‘s work [21], the effect of copper coated particles on the properties of CoCrFeNi and AlCoCrFeNi high entropy alloys (HEAs) was investigated. It decreased from 189.1 to 134.5 HV for Cux(CoCrFeNi)1−x and from 403 to 191 HV for Cux(AlCoCrFeNi)1−x HEAs by the addition of the nano Cu. In addition, wear rate was increased gradually by the addition of the nano Cu. Furthermore, the use of HEA coatings for repairing and remanufacturing can significantly reduce material consumption, achieving energy savings and enabling secondary use of mechanical parts. Methods of preparing HEA coatings include magnetron sputtering [22], thermal spraying [23,24], laser melting [25] and so on. Qiu et al. [26] prepared the Al2CrFeNiCoCuTix high-entropy alloy coatings through laser cladding and found that the relative wear resistance of the high-entropy alloy coatings was greatly improved with increased titanium content. The hardness and modulus of ZrNbTaTiW films fabricated by multi-target magnetron sputtering process by Feng et al. [27] were in the range of 5.9–11.5 GPa and 134.4–190.4 GPa, respectively. Chen et al. [28] investigated the formation of σ phase by adding molybdenum to FeCoCrNi coatings manufactured through tungsten inert gas cladding, which increased the microhardness and wear properties of the coatings. Chen et al. [29] prepared the high entropy Al0.6TiCrFeCoNi coating through HVOF spray with dense microstructure, two BCC phases with similar lattice parameters, high microhardness of 789.54 HV0.2 and good fracture toughness of 8.4 MPa·m1/2. Wu et al. [30] investigated the effect of molybdenum on the microstructure and mechanical properties of laser-cladded Al2CrFeNiMox coatings and found that the wear resistance of the coatings was greatly improved as compared to that of stainless steel. Cai et al. [31] studied the effect of alloying elements on the dilution rate and found that excess Fe in the original powder increased the dilution rate, which led to a decrease in the mixing entropy of the coating but not to the degree of forming new phases in the coating. The difference in atomic radii among the elements could lead to large lattice distortion in the FCC lattice. The lattice distortion and microhardness of the FexCoCrNi coating increased with the increase of Fe element content. In addition, the excessive amount of Fe element would weaken the self-passivation ability and reduce the high temperature oxidation resistance of the coating.
There are many research works focusing on the preparation method. Magnetron sputtering was suitable to prepare thin films. Meanwhile, the bonding force between the film and the substrate was relatively poor. Plasma transfer arc cladding had higher energy density than thermal spray, but lower energy density than laser cladding. Plasma cladding technology has the advantages of concentrated heat, short action time, low thermal impact, good metallurgical bonding and high joint strength, low cost and low requirement for substrate surface treatment as compared to other methods [32,33]. In this work, plasma cladding technology was adopted to prepare HEA coatings.
Compacted graphite iron (CGI) was often used to manufacture heavy engine parts, for example, cylinder head, for its comprehensive mechanical properties [34]. While, valve seats in cylinder heads would experience severe wear during long-term service at high temperature, which would inhibit the effective working of the engine further.
Therefore, in this work a CoCrFeNiMn high entropy coating was prepared on CGI through plasma transfer arc cladding to protect CGI from wear. The effect of plasma arc current on the friction and wear properties of CoCrFeNiMn HEA coatings were investigated. The plasma arc currents to prepare CoCrFeNiMn HEA coatings on CGI were optimized.

2. Materials and Methods

2.1. Preparation of the HEA Coatings

The feedstock was an equiatomic CoCrFeNiMn HEA powder (Beijing Avimetal Powder Metallurgy Technology Co., Ltd. Beijing, China). The feedstocks distributed from 15 to 53 μm with a mean size of 42 μm. The chemical composition of the feedstock is listed in Table 1. The powder had a near-equal atomic percent. CoCrFeNiMn HEA coatings were deposited on the CGI with a size of 100 mm × 100 mm × 10 mm through a plasma transfer arc system (DML-V03BD, Shanghai Duomu Mechanical Co., Ltd., Shanghai, China). Table 2 shows the chemical composition of the compacted graphite iron. Before coating deposition, the CGI substrate was cleaned with acetone. Based on a previous study, the plasma melting parameters selected in this study are as follows [35]. During the plasma cladding, argon was used as protective gas, plasma gas and powder feeding gas. The plasma torch was kept 10 mm from the torch exit to the substrate and the travelling speed was kept at 10 mm/min. The ion gas flow and the protection gas flow were 2.0 L/min and 8.0 L/min, respectively. The powder feeding rate was kept at 10 r/min. The lap width of the multiple cladding was 4 mm. The plasma arc currents were 50 A, 55 A, 60 A and 65 A.

2.2. Characterization of the HEA Coatings

The powder and coatings were charactered by scanning electron microscopy (SEM, VEGA II-XMU, TESCAN, Bron, Czech Republic) with EDX (7718, Oxford Instruments, Abingdon, UK). The sample was etched with a 10% nitric acid alcohol solution for 3 s. The phases of powders and coatings were analyzed by X-ray diffraction (D2, Bruker, Billerica, MA, USA) using Cu Kα radiation. The X-ray tube voltage was 20 kV. The diffraction angles of 2 θ were from 20 to 80° at a speed of 2°/min with a step of 0.02°. The micro-hardness of the coating was measured by a Vickers micro-hardness tester (HV-5, Taiming, Shanghai, China) under a load of 200 gf for a loading duration of 30 s.

2.3. Wear Test

The multi-pass claddings were prepared on CGI with a size of Φ 30 mm × 4 mm. The prepared friction wear specimens have a coating thickness of 2 mm and a substrate thickness of 4 mm. The coatings were ground by grinding machine to a certain surface roughness. The friction coefficient and the wear of the coatings was tested through a pin-on-disk friction and wear tester (HT1000, Lanzhou Zhongke Kaihua Technology Development Co., Ltd., Lanzhou, China). The wear test was conducted at room temperature with a load of 10 N at a rotational speed of 300 rad/min for a test time of 30 min. The N80 bulk was selected and cut into cylindrical pin with a size of Φ 3 mm × 10 mm serving as a counterpart for the HEA coatings in the friction and wear test. The stability of the wear process was analyzed by the friction coefficient. The samples were weighed before and after wear test through an electronic balance with the accuracy of 0.0001 g, which formed the wear mass losses. After wear tests, the wear morphologies were characterized through SEM and 3D confocal microscope (VK-X3000, Keyence, Osaka, Japan).

3. Results and Discussions

3.1. Microstructure of the HEA Feedstock and Coatings

Figure 1 shows the morphology and cross-sectional microstructure of the CoCrFeNiMn HEA feedstock. It had a spherical morphology, which was related to the gas atomization process. The feedstock’s spherical morphology was helpful in powder feeding. Its mean size was about 42 μm. Figure 2 shows the X-ray diffraction pattern of the powder. The powder had a single phase with face centered cubic (FCC) structure. Figure 3 shows the microstructure of the compacted graphite iron, consisting of worm-like graphite as well as ferrite.
Figure 4 shows the microstructure of the CoCrFeNiMn HEA coatings prepared by plasma transfer arc cladding under a plasma arc current of 60 A. As shown in Figure 4a, there was a strong metallurgical bonding between the HEA coating and the compacted graphite iron without cracks. Fine needle-like martensite was formed near the compacted graphite iron, and ledeburite presented in the bonding area. The surface of the compacted graphite iron melted quickly during the plasma melting process. Some parts of the molten areas cooled rapidly enough to form martensite, the others cooled relatively slowly and mixed with the HEA melt to form ledeburite. The middle part of the coating, as shown in Figure 4b, consisted of columnar grain and equiaxed grain, which was related to the temperature changes in the overlapped region with the circulation of heating and cooling during the multi-pass fusion process. According to the solidification thermodynamic theory and heat dissipation direction, the solidification rate (SR) tended to be zero and the value of GL/SR tended to infinity due to the very large temperature gradient (GL) between the melt layer and the liquid alloy in the substrate’s top region. Therefore, the nucleation rate of the grains was much higher than the growth rate, which would result in the formation of equiaxed grain near the substrate. As the liquid-solid interface moved toward the surface, the actual GL in the liquid metal decreased, SR increased and the GL/SR ratio decreased, and the grain morphology changed from equiaxed grain to columnar grains in the middle of the coating, as shown in Figure 5b. As the heat dissipation in the middle part of the coating was slower, the columnar grain region generated more widely. As the solidification proceeded, the temperature gradient decreased accordingly. Therefore, the cross-sectional microstructure of the molten layer changed from columnar to equiaxed grains, as shown in Figure 4c [36]. The top region of the coating was mainly composed of equiaxed grains. All of the equiaxed grains in the top zone of the cladding had their growth direction pointing to the edge of the coating, which meant that the growth directions tended to be parallel. The growth directions were not unique. The direction of heat dissipation at the top of the coating was not confined after the cladding finished, which resulted in the growth of grains without unique directions. It indicated that the temperature gradient had a great influence on the growth direction of grains.
Figure 5 shows the cross-sectional distributions of elements in the CoCrFeNiMn HEA coatings at different plasma arc currents of 50 A, 55 A, 60 A and 65 A. Point analysis results of elements in dendrite (DR) and interdendrite (ID) regions of the central part in the CoCrFeNiMn HEA coatings at different plasma arc currents are listed in Table 3. Cr and Mn were enriched in the interdendrite (ID) region. Fe distributed uniformly. Ni and Mn were more uniformly distributed and enriched in part of the dendrite (DR). The point analysis results of elements confirmed that the weight percentages of Ni and Mn were higher than those at the grain boundaries. The dendrite (DR) region had a higher concentration of Co and Ni, while larger amounts of Cr and Mn were in the interdendrite (ID) region. The iron contents of the claddings were all higher than that of the raw feedstock due to the loss of elements during melting at high temperatures and the dilution on the surface of the compacted graphite iron substrate [37]. Co and Mn were found to be higher along the boundaries between DR than those within them, which was similar to the results of M. Laurent-Brocq et al. [38], G.A. Salishchev et al. [39] and in agreement with the results of the mapping analysis. The presence of high content of Fe in the dendrites, which diffused out from the compacted graphite iron during the cladding process, led to solid solution strengthening of the dendrites. On the other hand, the contents of Co, Ni, Cr and Mn in the grains and grain boundaries were lower than those in the powder. The loss of powder elements during the rapid melting of HEA powders might be related to the lower saturation vapor pressure [40]. At the grain boundaries of the coating, there were 2.51 wt.% oxygen, which was higher than that in powders without significant oxygen. The oxidation at the grain boundaries could be mainly attributed to the oxidation of the powder surfaces during the melting process of HEA powder.

3.2. Phases

Figure 6 shows the X-ray diffraction patterns of CoCrFeNiMn high entropy alloy coatings prepared by plasma transfer arc cladding at different plasma currents. It was found that CoCrFeNiMn high-entropy alloy coatings had three main characteristic peaks at 43.23°, 52.69° and 73.60°. According to the research results of D.Y. Lin et al. [41] and PDF cards (PDF card No. 33-0397), the characteristic peaks were calibrated to be (111), (200) and (220) lattice planes in face centered cubic (FCC) solid solutions, respectively. The HEA coatings consisted of a single-phase face-centered cubic structure. The formation of the single-phase face-centered cubic structure was attributed to that the alloy elements where all adjacent elements in the fourth period and their atomic radii were closed [42]. According to the thermodynamic relationship, the Gibbs free energy (ΔG mix = ΔH mix − TΔS) changed, and there was a large negative mixing among the elements of the system due to the equiatomic ratio of the mixed system, where the enthalpy changed and entropy change were mutually constrained [43]. At high temperature, the formation of common intermetallic compounds was suppressed due to the high mixing entropy of the cobalt-manganese high-entropy alloy. Therefore, the Gibbs free energy changes were small and the system tended to be thermodynamically stable [37]. No oxide peaks were found in the XRD patterns because the Mn element in the powder had a strong deoxidizing and slagging effect during the plasma cladding process and argon gas also served to prevent oxidation [30].
Table 4 shows the lattice parameters and dendritic crystal width of the CoCrFeNiMn HEA powder and coatings prepared at different plasma arc currents calculated from the X-ray diffraction patterns through MDI Jade software. The grain size was measured by Image-pro Plus software on SEM images of the specimens. Table 4 shows the grain width of dendritic structure at different plasma currents. With the increase of plasma current, the width of dendrites in the coating became small, which would produce a fine grain strengthening effect.
With the increase of plasma arc current, the lattice parameters of the CoCrFeNiMn HEA coatings decreased. The four HEA coatings had larger lattice parameters than the powder. With the increase of plasma arc current, the heat input in the plasma cladding process increased greatly, which led to fast heating and cooling simultaneously. The super cooling effect led to great solid solution and finally led to larger lattice parameters in the CoCrFeNiMn HEA coatings than that in the powder. Meanwhile, the cooling times became longer and longer with the increase of the plasma arc current. A dynamic equilibrium process would exist between the solid solution and de-solution, which was affected by both the cooling speed and time. Although the cooling speed increased with the increase of plasma arc current, the cooling times were extended obviously, which brought much more chances to de-solute. Therefore, the lattice constant of the CoCrFeNiMn HEA coatings decreased with the increase of plasma arc current. With the increase of the plasma arc current, the heat input during the plasma cladding process increased and the grain size of the CoCrFeNiMn HEA coating tended to decrease with grain refinement.

3.3. Microhardness

Figure 7 shows the microhardness distribution of the coating in the cross section along the top of the coating to the substrate. The highest microhardness of the coating was 394 ± 21.6 HV0.2 when the plasma arc current was 65 A and the lowest microhardness of the HEA coating was 360 ± 15.5 HV0.2 when the plasma arc current was 60 A. The average hardness of CoCrFeMnNi high-entropy alloy prepared by microbeam plasma arc melting was 158.9 HV0.2 by F.X. Ye et al. [44]. The hardness of FeCoNiCrMn high-entropy alloy coating prepared by plasma spraying was 273 ± 35 HV0.2 by J.K. Xiao et al. [45]. CoCrFeNi high-entropy alloy with a hardness of 181.2 HV0.2 was prepared by A.J. Zhang et al. [46]. With the increase of plasma arc current, the heat input increased, which led to superheating and affected the size of solidification structure by the way of affecting the degree of undercooling and finally led to the coatings’ grains refining and acted as fine grain reinforcement. The mean microhardnesses of the coatings prepared at different plasma currents were all higher than the reported works [44,45]. Meanwhile, the coating’s hardness was higher than that of the dense compacted graphite iron substrate due to the strong solid solution strengthening effect. The hardness at the top of the coating was higher than that of the inner part due to the slow solidification rate during melting and the presence of small equiaxed crystals at the top edge of the coating, which had a fine crystal strengthening effect. The microhardness of both the bonded and heat-affected zones were much higher than that of the HEA coating and the compacted graphite iron substrate for the formation of martensite in the bonded and heat-affected zones. The highest microhardness in the bonded zone varied between 430 and 690 HV0.2, and that in the heat affected zone varied between 430 and 630 HV0.2. In the bonding zone, there were comixed CoCrFeNiMn HEA solid solutions, ledeburite and fine acicular martensite formed near the compacted graphite iron for some parts remelted at very high heating and cooling speeds on the surface of the compacted graphite iron during plasma transfer arc cladding, which led to a great increase of the microhardness in the bonding zone. Meanwhile, in the heat affected zone, the pearlite in the compacted graphite iron substrate was heated and cooled quickly and changed to martensite with high hardness, which had been reported in references [35,40].

3.4. Wear Resistance

Figure 8 shows the friction coefficient curves and the mean friction coefficient of the CoCrFeNiMn HEA coating at different plasma arc currents. During the break-in period, the friction coefficient increased sharply and then stabilized. The fluctuation in stability was mainly attributed to instrument vibration and measurement accuracy. The initial lower friction coefficient could be attributed to the oxidation of the coating surface, where a thin oxide film existed. When it was destroyed by frictional sliding, the adhesive wear between the coating surface and the counterpart became very severe and the friction coefficient increased sharply [45]. At the same time, the frictional heat promoted the formation of the oxide film on the surface of the coating and the substrate, which slightly reduced the friction coefficient. When the formation of the oxide film and the wear reached a dynamic equilibrium, the friction coefficient remained stable. When the plasma arc current was 55 A, the highest friction coefficient of HEA coating was 0.75. When the plasma arc current was 60 A, the lowest friction coefficient of the coating was 0.5. The friction coefficient of the HEA coatings in this work were all lower than that of the plasma sprayed coatings with the friction coefficient of 0.8, which was reported in J.K Xiao’s work [45]. Meanwhile, the lowest friction coefficient of the HEA coatings prepared by plasma transfer arc cladding in this work was lower than that of the CoCrFeNiMn high-entropy alloy prepared by the spark plasma sintering method with the friction coefficient of 0.62 in Z.M. Guo’s work [47]. The friction coefficients of all four coatings were higher than that of the compacted graphite iron substrate. The friction coefficient of the HEA coating increased slightly with the prolongation of testing time and then gradually reached a constant value throughout the wear test period, which indicated an improvement in its tribological properties. The compacted graphite iron substrate had a lower coefficient of friction than the coating because the compacted graphite iron substrate consisted of vermicular graphite, which had lubricating properties and acted as a lubricant and reduced the coefficient of friction.
Figure 9 shows the wear mass loss of the coating, the substrate and the pins. The HEA coating had the highest wear mass loss when the plasma arc current was 60 A. At 65 A plasma arc current, the wear mass loss of the HEA coating was the lowest. At the same time, all of the wear mass losses of the HEA coatings were lower than the wear mass loss of the substrate. The wear mass loss of the pins was the lowest one at a plasma arc current of 50 A. The highest wear mass loss of the pins was observed for plasma arc currents of 60 A. At 65 A plasma arc current, the HEA coating’s hardness was the highest and the wear mass loss was the lowest. The trend in coating wear resistance was similar to the trend in microhardness. The degree of wear was reconnected to the resistance in the contact zone. As the friction resistance increased, the shear stress on the surface particles of the coating increased, which further led to easy wear of the surface particles. The friction coefficient of the alloy was also related to the phase change and the variation in hardness.
This result was in agreement with Khruschov’s conclusion that the wear resistance of a material was usually proportional to its Vickers hardness [48].
Figure 10 shows the abrasion depth of the coating for the substrate and HEA coatings at different plasma arc currents. When the plasma arc current was 65 A, the abrasion depth of the coating was the shallowest. At 60 A plasma arc current, the abrasion depth of the coating was the deepest. All of the HEA coatings’ abrasion depths were shallower than that of the compacted graphite iron substrate. The trends of abrasion depth of the coating were consistent with that of the microhardness. The HEA coatings had better wear resistance than the CGI substrate. Combined with the friction coefficient, the coating had the best wear resistance when the plasma arc current was 65 A.
Figure 11 shows the wear morphologies of the HEA coatings prepared at different plasma arc currents after wear tests. When the plasma arc current was 50 A, the wear morphology of the coating was relatively flat with parallel and shallow grooves and wear particles on the surface as shown in Figure 11a, which revealed that the wear mechanism was typically abrasive wear. The grooves on the wear surface were produced by the micro-cutting of the micro-convex bodies on the pin surface. The wear particles were chips or debris produced by micro-cutting in the wear test. When the plasma arc current was 55 A, a large number of wide and deep parallel grooves as well as pits and patches appeared on the wear surface of the HEA coating as shown in Figure 11b, which meant that adhesive wear occurred. The wear composed of both abrasive and adhesive wear. When the plasma arc current was 60 A, the wear surface of the coating was rough with obvious flaking and wear particles as shown in Figure 11c. Generally, the wear mechanism of metal materials was mainly influenced by the hardness. The high wear mass loss of the HEA coating at 60 A plasma arc current was mainly due to the relatively low hardness. The spalling phenomenon in the wear scars was due to the micro-cut of the coating by the corresponding friction sub. When the plasma arc current was 65 A, the surface of the HEA coating was flat. Meanwhile, there were a large number of parallel grooves, spalling and wear particles on the wear surface as shown in Figure 11d, which revealed that the wear mechanism was abrasive wear and adhesive wear. Since the friction volume was much smaller than the contact peak, the peak temperature dropped instantly once the contact was removed. The typical local high temperature lasted only a few milliseconds. In this momentum, the oxide film on the coating surface was broken down, which broke the adhesion at the contact peak, and the metal on the surface was torn off, forming abrasive particles and also leading to the formation of flaking on the surface. Some metal particles stuck to another metal surface, which formed adhesive wear. The mechanism of adhesive wear was a cyclic process of adhesion, destruction and re-adhesion. All forms of wear were accompanied by plastic deformation, adhesive wear, grooving and spalling. During wear, the coating layer was subjected to high temperatures and oxygen, thus oxidation particles appeared on the wear surface. Thus, the wear mechanism was adhesive wear, abrasive wear with a small amount of oxidation.
Figure 12a shows the 3D wear morphologies of the HEA coating at the plasma arc current of 55 A. There were many furrow-like abrasive scars and a large number of adhesion characteristics. Figure 12b shows the 3D wear morphologies of the HEA coating at the plasma arc current of 60 A. Figure 12c,d shows the variation of the depth of the abrasion marks obtained from the line scan of the HEA coating at the plasma arc current of 55 A and 60 A. The depth of the abrasion scars of the HEA coating at the plasma arc current of 55 A varied greatly due to the presence of furrow-like abrasion scars. While, the depth of the abrasion scars of the coating at the plasma arc current of 60 A varied less for the reason that the wear form of this coating was mainly adhesion wear. The HEA coatings at the plasma arc currents of 50 A and 65 A both had furrow abrasion and adhesion wear. The HEA coating at the plasma arc current of 55 A had furrow abrasion. The HEA coating at the plasma arc current of 60 A had obvious spalling.
Figure 13 shows the wear scars of the N80 pin corresponded to the coating when the plasma arc current was 55 A. There were both a large number of furrow-like scratches and some adhesion characters. Figure 14 shows the point analysis of elements in the HEA coating prepared at 55 A plasma arc current after the friction and wear test. Three points were selected for the analysis as shown in Figure 14. The point analysis results of elements in the HEA coating prepared at 55 A plasma arc current are shown in Table 5. Oxygen element was found in the composition analysis at all three points, which indicated that oxidation occurred at the wear scars during the wear tests. Point 1 indicated the oxidation of the particles under the action of frictional heat. During friction and wear, these oxidized particles between the friction subsets acted as an abrasive material, which then produced visible grooves on the wear surface (Figure 12b) [49]. Point 3 showed the adhesive wear products with high oxygen content. Due to the presence of Cr active metal in the CoCrFeNiMn HEA coating, the coating having undergone oxidation reaction under the action of high temperature and oxygen during the friction wear test. The oxidation products generated by the coating played an important role in the formation of lubricating film during the friction process. Carbon element was mainly derived from the counter-abrasive pin N80. The chromium content at point 1 was 55.60 wt.%, which was much higher than that of the coating. The particles were the debris or chips generated by the coating due to micro-cutting in the friction test.
According to the wear morphologies and the element analysis of the wear scars, the HEA coatings mainly had the wear mechanism of abrasive wear, adhesive wear with a small amount of oxidation against N80 friction pair.

4. Conclusions

The effects of plasma arc currents on the phase composition, microstructure, microhardness and wear resistance of CoCrFeNiMn HEA coatings on CGI through plasma transfer arc cladding were investigated.
  • The CoCrFeNiMn HEA coating consisted mainly of martensite and ledeburite in the bonding zone, columnar and equiaxed grains in the middle of the coating and equiaxed grains in the top zone of the coating. The HEA coating had a simple face-centered-cubic solid solution structure similar to the powder.
  • The CoCrFeNiMn HEA coating had the highest microhardness of 394 ± 21.6 HV0.2 and the lowest wear mass loss when the plasma arc current was 65 A.
  • The CoCrFeNiMn HEA coating had mean friction coefficients ranging from 0.5 to 0.75, which were all higher than that of the compacted graphite iron substrate.
  • There were abrasive, adhesive and oxidation wear forms in the HEA coatings against N80 friction pair. The HEA coating presented higher friction coefficient and better wear resistance than compacted graphite iron.

Author Contributions

Conceptualization, P.G., R.F., J.L. (Jilin Liu), B.C. and M.L.; methodology, P.G., R.F., J.L. (Jilin Liu), B.C., B.Z., D.Z., W.W. and M.L.; software, P.G., R.F., B.C. and Y.G.; validation, P.G., B.C. and B.Z.; formal analysis, P.G., R.F., B.C. and Z.Y.; investigation P.G., R.F., J.L. (Jilin Liu), B.C., B.Z., D.Z., W.W., Z.Y. (Zhiyi Yan) and L.Z.; data curation, Z.Y. (Zhong Yang); writing—original draft preparation, R.F.; writing—review and editing, P.G.; project administration, J.L. (Jianping Li); funding acquisition, P.G. and J.L. (Jianping Li) All authors have read and agreed to the published version of the manuscript.

Funding

This work was funded by the National Natural Science Foundation of China (51771140), Shaanxi Key Science and Technology Innovation Team (2017KCT-05), The Youth Innovation Team of Shaanxi Universities: Metal Corrosion Protection and Surface Engineering Technology, Shaanxi Provincial Key Research and Development Project (2019ZDLGY05-09), Local Serving Special Scientific Research Projects of Shaanxi Provincial Department of Education (19JC022), Project of Yulin Science and Technology Bureau (2019-121).

Institutional Review Board Statement

Not applicable.

Informed Consent Statement

Not applicable.

Data Availability Statement

Not applicable.

Conflicts of Interest

The authors declare no conflict of interest.

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Figure 1. Microstructure of CoCrFeNiMn powder: (a) global morphology, (b) cross-sectional microstructure.
Figure 1. Microstructure of CoCrFeNiMn powder: (a) global morphology, (b) cross-sectional microstructure.
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Figure 2. XRD pattern of the CoCrFeNiMn powder.
Figure 2. XRD pattern of the CoCrFeNiMn powder.
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Figure 3. Microstructure of the compacted graphite iron.
Figure 3. Microstructure of the compacted graphite iron.
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Figure 4. Cross-sectional microstructure of CoCrFeNiMn coating at 60 A plasma arc current: (a) bonding zone, (b) middle part of the coating, (c) top zone of the coating.
Figure 4. Cross-sectional microstructure of CoCrFeNiMn coating at 60 A plasma arc current: (a) bonding zone, (b) middle part of the coating, (c) top zone of the coating.
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Figure 5. Cross-sectional distribution of elements in the CoCrFeNiMn HEA coatings at different plasma arc currents: (a) 50 A, (b) 55 A, (c) 60 A and (d) 65 A.
Figure 5. Cross-sectional distribution of elements in the CoCrFeNiMn HEA coatings at different plasma arc currents: (a) 50 A, (b) 55 A, (c) 60 A and (d) 65 A.
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Figure 6. XRD patterns of the CoCrFeNiMn HEA coatings.
Figure 6. XRD patterns of the CoCrFeNiMn HEA coatings.
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Figure 7. Microhardness distribution along the top surface to the substrate across the cross-section of the CoCrFeNiMn HEA coatings.
Figure 7. Microhardness distribution along the top surface to the substrate across the cross-section of the CoCrFeNiMn HEA coatings.
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Figure 8. (a) Friction coefficient of the CoCrFeNiMn HEA coating and CGI substrate, (b) the average friction coefficient of the CoCrFeNiMn HEA coating and CGI substrate.
Figure 8. (a) Friction coefficient of the CoCrFeNiMn HEA coating and CGI substrate, (b) the average friction coefficient of the CoCrFeNiMn HEA coating and CGI substrate.
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Figure 9. Wear mass loss of the CoCrFeNiMn HEA coatings, CGI substrate and N80 pin counterpart after wear test.
Figure 9. Wear mass loss of the CoCrFeNiMn HEA coatings, CGI substrate and N80 pin counterpart after wear test.
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Figure 10. The wear scar depth of the substrate and the CoCrFeNiMn HEA coatings at different plasma arc currents.
Figure 10. The wear scar depth of the substrate and the CoCrFeNiMn HEA coatings at different plasma arc currents.
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Figure 11. Wear morphologies of the CoCrFeNiMn HEA coatings after wear test at: (a) 50 A, (b) 55 A, (c) 60 A and (d) 65 A.
Figure 11. Wear morphologies of the CoCrFeNiMn HEA coatings after wear test at: (a) 50 A, (b) 55 A, (c) 60 A and (d) 65 A.
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Figure 12. 3D morphologies of the CoCrFeNiMn HEA coatings after wear test: (a) coating at 55 A, (b) coating at 60 A and wear depth variation through line scanning of the HEA, coating at (c) 55A and (d) 60A.
Figure 12. 3D morphologies of the CoCrFeNiMn HEA coatings after wear test: (a) coating at 55 A, (b) coating at 60 A and wear depth variation through line scanning of the HEA, coating at (c) 55A and (d) 60A.
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Figure 13. Wear scars of N80 counterparts corresponded to the CoCrFeNiMn HEA coating at 55 A plasma arc current after wear test: (a) global morphology, (b) magnified part.
Figure 13. Wear scars of N80 counterparts corresponded to the CoCrFeNiMn HEA coating at 55 A plasma arc current after wear test: (a) global morphology, (b) magnified part.
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Figure 14. Point analysis of elements in the CoCrFeNiMn HEA coating prepared at 55 A plasma arc currents after the friction wear test (a,b) spectra at point 1, (c) spectra at point 2, (d) spectra at point 3.
Figure 14. Point analysis of elements in the CoCrFeNiMn HEA coating prepared at 55 A plasma arc currents after the friction wear test (a,b) spectra at point 1, (c) spectra at point 2, (d) spectra at point 3.
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Table
Table 1. Elemental composition of CoCrFeNiMn HEA powder.
Table 1. Elemental composition of CoCrFeNiMn HEA powder.
ElementsCoCrFeNiMn
Contents/wt.%20.5818.3519.9820.4920.48
Table
Table 2. Elemental composition of the compacted graphite iron.
Table 2. Elemental composition of the compacted graphite iron.
ElementsCSiMnSPFe
Contents/wt.%3.4–3.72.4–3.0≤0.6≤0.6≤0.06Bal
Table
Table 3. Point analysis of elements in the middle part of the CoCrFeNiMn HEA coatings prepared at different plasma arc currents.
Table 3. Point analysis of elements in the middle part of the CoCrFeNiMn HEA coatings prepared at different plasma arc currents.
Weight Percent/wt.%RegionCrMnFeCoNi
Coatings
50 ADendrite2.764.03581.855.5655.79
Interdendrite7.8855.979.544.032.64
55 ADendrite2.1956.7373.856.7956.85
Interdendrite14.9259.13565.236.3654.33
60 ADendrite4.1255.59574.477.9057.805
Interdendrite12.149.2168.4935.8134.343
65 ADendrite3.235.25577.1556.6757.69
Interdendrite12.039.12372.9133.4132.52
Table
Table 4. Lattice parameters and grain width of dendritic structure of the CoCrFeNiMn HEA powder and coatings prepared at different plasma arc currents.
Table 4. Lattice parameters and grain width of dendritic structure of the CoCrFeNiMn HEA powder and coatings prepared at different plasma arc currents.
Plasma Arc Current/ALattice Parameters/ÅDendritic Crystal Width/μm
Powder3.59751 ± 0.0001834 ± 19
50 A3.59857 ± 0.0001318.902 ± 2.101
55 A3.60065 ± 0.0001316.837 ± 1.868
60 A3.60311 ± 0.0003312.878 ± 2.081
65 A3.60063 ± 0.0000710.023 ± 1.075
Table
Table 5. Point analysis of elements in the worn surface of the CoCrFeNiMn HEA coating prepared at 55 A plasma arc current.
Table 5. Point analysis of elements in the worn surface of the CoCrFeNiMn HEA coating prepared at 55 A plasma arc current.
Weight Percent/wt.%OCCrMnFeCoNi
Point
Point 17.7612.8255.603.4512.752.205.41
Point 23.784.5310.4411.6249.598.7910.24
Point 331.413.419.134.6821.204.1426.02
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Gao, P.; Fu, R.; Liu, J.; Chen, B.; Zhang, B.; Zhao, D.; Yang, Z.; Guo, Y.; Liang, M.; Li, J.; et al. Influence of Plasma Arc Current on the Friction and Wear Properties of CoCrFeNiMn High Entropy Alloy Coatings Prepared on CGI through Plasma Transfer Arc Cladding. Coatings 2022, 12, 633. https://doi.org/10.3390/coatings12050633

AMA Style

Gao P, Fu R, Liu J, Chen B, Zhang B, Zhao D, Yang Z, Guo Y, Liang M, Li J, et al. Influence of Plasma Arc Current on the Friction and Wear Properties of CoCrFeNiMn High Entropy Alloy Coatings Prepared on CGI through Plasma Transfer Arc Cladding. Coatings. 2022; 12(5):633. https://doi.org/10.3390/coatings12050633

Chicago/Turabian Style

Gao, Peihu, Ruitao Fu, Jilin Liu, Baiyang Chen, Bo Zhang, Daming Zhao, Zhong Yang, Yongchun Guo, Minxian Liang, Jianping Li, and et al. 2022. "Influence of Plasma Arc Current on the Friction and Wear Properties of CoCrFeNiMn High Entropy Alloy Coatings Prepared on CGI through Plasma Transfer Arc Cladding" Coatings 12, no. 5: 633. https://doi.org/10.3390/coatings12050633

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