Friction and Wear Properties of CoCrFeNiMnSnx High Entropy Alloy Coatings Prepared via Laser Cladding

Abstract

Due to its unique single-phase multivariate alloy characteristics and good low-temperature mechanical properties, CoCrFeNiMn high entropy alloy (HEA) has attracted the interest of many researchers in recent years. In this paper, to improve the wear resistance of Q235 alloy steel surface, CoCrFeNiMnSn_x HEA coatings were prepared on the surface of Q235 steel via laser cladding. X-ray diffractometry, optical microscopy, scanning electron microscopy (SEM), and energy dispersive spectrometry were used to determine the microstructure and chemical composition. The research findings revealed that the CoCrFeNiMn HEA coatings were formed from a single FCC phase. As the Sn content in the coating increased, a new MnNi₂Sn phase formed. Microhardness and friction and wear results showed that when the mole content of Sn was 0.2, the hardness of the CoCrFeNiMn HEA coating was increased by approximately 45%, the friction coefficient decreased by 0.168, and the wear loss decreased by 16.6%. Three-dimensional noncontact morphology and SEM results revealed that the wear mechanisms of CoCrFeNiMn HEA coatings were abrasive wear, delamination wear and a small amount of oxidative wear under dry friction conditions, whereas the friction mechanisms of CoCrFeNiMnSn_0.2 HEA coatings were primarily abrasive wear and oxidative wear.

1. Introduction

The novel concept of high entropy alloy (HEA), defined as alloys composed of five or more alloy elements with the concentration of each at 5 to 35 at.%, represents an innovative alloy design strategy [1]. Because of their unique microstructure, HEAs have superior mechanical properties to traditional alloys [2], including high strength, ductility, and fracture toughness [3,4,5,6,7]. The prototypical CoCrFeNiMn HEA was first successfully prepared and synthesized by Cantor et al. in 2004 Bernard G et al. found that CoCrFeNiMn HEA had the characteristics of high strength and high toughness at low temperature and an extremely high research value [8,9,10], so it has been widely used in the fields of aerospace and shipbuilding. However, CoCrFeNiMn HEA is a single-phase FCC solid solution, whether in the as-cast or heat-treated state, resulting in it having low strength and good plasticity [11,12]. To improve the wear resistance of CoCrFeNiMn HEA based on the FCC phase, one of the primary methods used is to form one or more intermetallic compounds by adding one or more elements to the CoCrFeNiMn HEA.

In recent years, many scientists have performed research on the modification of CoCrFeNiMn HEA by adding elements. Stepanov N D et al. investigated CoCrFeNiMnV_x HEA with the added element V [13]; the results showed that as V increased, the phase structure of the alloy transitioned from FCC to FCC + σ phase. When x = 1, the phase changed to a matrix, the hardness rose from 135 to 636 HV_0.3, and the yield strength rose from 230 to 1660 MPa. N R Wang et al. created (CoCrFeNiMn)₉₀Hf₁₀ HEA powder by incorporating the element Hf [14]. The results showed that (CoCrFeNiMn)₉₀Hf₁₀ HEA powder had a dual solid solution structure and exhibited typical ferromagnetic behavior. The strength of FeCoNiCrMn HEA could be increased by adding the element B to delay grain growth [15]. Klimova et al. discovered that after deformation heat treatment, adding C to FeCoNiCrMn HEA precipitated Cr-rich M₂₃C₆-type carbides in the matrix [16]. The pinning effect caused by carbides strongly restricted grain growth, resulting in an alloy with excellent mechanical properties at room and low temperatures. The same phenomenon was also observed when a trace amount of N was added to FeCoNiCrMn [17]. He J Y et al. studied the addition of the element Al to FeCoNiCrMn HEA [18]. The alloy phase composition changed from FCC to FCC + BCC and then to BCC as the Al content increased. HEA hardness increased from 176 to 538 HV_0.3, and yield strength increased from 209 to 832 MPa, because the slip system of the BCC structure was much less than that of the FCC structure, which improved the alloy’s hardness. The atomic radius of Al was also much larger than that of Co, Cr, Fe, and the other elements, resulting in strong lattice distortion that enhanced strengthening of the solid solution and improved the alloy’s hardness. L Lin et al. [19] studied the microstructure of FeCoCuNiSnx HEAs by adding the element Sn; the results showed that when the proportion of Sn was low, the alloy had a single FCC solution. Furthermore, with increase in Sn content, the microstructure of the alloy changed to an FCC solid–solution matrix with Cu₈₁Sn₂₂ intermetallics. The addition of the element Sn changed the chemical composition of the alloy. Zhu et al. studied the influence of changes in the Sn content on the microstructure of TiZrTaNb HEA [20]. The results showed that the microstructure of the HEA was a single-layer BCC structure. The addition of elements promoted the growth of dendritic structures and accelerated the interdendritic segregation of TiZrTaNb alloys. These results indicate that the element Sn can promote structural change of HEAs. However, the observed effect of Sn on the microstructure and friction and wear properties of CoCrFeNiMn HEA using a laser cladding preparation process is not perfect, which potentially limits the use of CoCrFeNiMn HEA in the future.

In this study, CoCrFeNiMnSn_x HEA coatings were prepared on Q235 alloy steel substrate by adding different quantities of the element Sn and using a laser cladding preparation process. The effect of changing the Sn element content on the microstructure was investigated, and the friction and wear mechanisms of CoCrFeNiMnSnx HEA coatings were explicated.

2. Experimental Method

2.1. Preparation of HEAs Coatings

The substrate was Q235 steel; the surface was polished to 800 mesh to remove oxide and other pollutants before being cleaned with acetone and dried. The laser cladding experiments were performed using a YLS-8000 semiconductor laser cladding system produced by Kawasaki, Japan. The powder material was CoCrFeNiMn HEA powder, as shown in Figure 1. Each elemental powder can be seen to be completely mixed. Figure 1b shows a local magnification of the powder; it can be seen that the sphericity of the powder was complete, with particle sizes between 60 and 100 nm. As shown in Figure 2, scanning the white area of the CoCrFeNiMn HEA powder revealed that Cr powder, Mn powder, and Fe powder accounted for 20.14%, 21.44%, and 18.38% of the total content in the CoCrFeNiMn HEA powder, respectively. Table 1 shows the element distribution of coating powders and Q235 alloy steel with different Sn content (referred to below as Sn₀, Sn_0.1, Sn_0.2, Sn_0.3, and Sn_0.5). The CoCrFeNiMn HEA powder and Sn elemental powder were homogenized for 5 h at 300 r/min in a YXQM-2L vertical planetary ball mill (Changsha Miqi Instrument Equipment Co., Ltd. Changsha, China), then dried in a vacuum drying box at 120 °C for 2 h before use. To reduce the cladding layer’s cracking susceptibility, the substrate was preheated to 200 °C before cladding and the temperature maintained until the end of the experiment. After a series of experimental tests, the parameters were determined to be 1300 W laser power and 1.2 r/min powder feeding speed. To prevent oxidation, high-purity argon gas was used as a shielding gas applied through a coaxial nozzle.

2.2. Structural Analysis

The specimens were sectioned into a suitable size, mounted, polished, and etched via aqua regia for metallographic examination. The microstructure and morphology of the coatings were studied using optical microscopy (OM) with a DM-400C (Shanghai Caikang Optical Instrument Co., Ltd. Shanghai, China) microscope, and scanning electron microscopy (SEM) with an EM30AXP energy dispersive spectrometer (EDS). The scanned photos were imported into Image J software. The phase volume fraction and porosity of each phase in the alloy were calculated five times, with the average value taken. The phases of the coatings were determined using a D8 FOCUS X-ray diffractometer (XRD) from Bruker in Beijing. The radiation source was Cu Kα, operated at 40 kV and 30 mA with a scanning rate of 7°/min, from 20° to 90°.

2.3. Hardness and Friction and Wear Performance Testing

The HEA coating Vickers hardness was determined using an automatic microhardness tester (United States, Buehler CG-200). The experimental parameters were set to 500 g load, 15 s load retention time, and an average value of three measurements was obtained.

The friction and wear experiment on the coatings was carried out using an MPX-3G in a pin-on-disk configuration in wear media. The experiment was conducted at room temperature with the following specific parameters: contact pressure test of 40 N, torque of 1.4 Nm, spindle speed of 500 r/min, wear test of 5 mm, and test revolution number of 10,000. Following the experiment, SEM was used to examine the morphology of the worn surface to determine the sample’s wear mechanism. Simultaneously, the cladding layer was examined using a PS50 three-dimensional noncontact profiler manufactured by the NANOVEA Company in the United States.

3. Results and Discussion

3.1. Microstructure and Phase Structure Analysis of Cladding Layer

Figure 3 shows the X-ray diffraction patterns of the CoCrFeNiMnSn_x HEA coatings. As can be seen, the Sn₀ coating formed a simple FCC structure with diffraction peaks at 43.68°, 51.12°, and 75.42°, which corresponded to the three crystal planes (111), (220), and (222). Because the extremely high entropy value of HEAs reduces the Gibbs free energy of the system, the formation of other compounds was inhibited, and only one FCC phase structure was formed. With increase in Sn content, a weak diffraction peak appeared at 2θ = 42.54°; the phase structure corresponding to this diffraction peak was the second phase of MnNi₂Sn, and the corresponding positions of the other three peaks remained unchanged, indicating that the Sn_0.1, Sn_0.2, and Sn_0.3 coatings were composed of FCC phase and MnNi₂Sn phase. As the Sn content increased, more diffraction peaks corresponding to MnNi₂Sn appeared in the XRD pattern; the corresponding positions were, respectively, 25.14°, 42.54°, and 61.53°, indicating that the crystal structure of the Sn_0.5 coating contained the FCC phase and the MnNi₂Sn phase.

Figure 4 shows the microstructure of the Sn₀ coating under an optical microscope. As shown in Figure 4a, the thickness of the heat-affected zone (HAZ) of the coating was 12 μm, indicating that the HAZ of the Sn₀ coating prepared via laser cladding was small. As can be seen from Figure 4b, the coating formed a dendritic structure. Because of the large temperature gradient at the bottom of the coating, the growth direction of the column dendrites tended to be perpendicular to the surface of the substrate. As shown in Figure 4c, the heat flux in the top area was largely controlled by the movement of the laser beam, which caused the growth direction to be more parallel to the surface of the substrate. A planar crystal structure was formed in the joint area between the coating and the substrate, as shown in Figure 4d, indicating that the coating and the substrate formed a good metallurgical bond.

Figure 5a,b show the EDS line scan area of the Sn₀ alloy coating microstructure and the distribution of each element, respectively. As can be seen, the five elements, Co, Cr, Fe, Ni, and Mn, were quite evenly distributed between the dendrites, with the content of Co, Cr, and Fe elements being relatively high and of Ni and Mn elements being relatively low. There was no significant difference in element content between dendrites, indicating that the Sn₀ alloy coating’s microstructure was a single FCC phase.

Figure 6 shows an SEM photograph of CoCrFeNiMnSn_x HEA coatings. When combined with XRD analysis, it was found that the Sn₀ coating formed a dendritic structure and a simple FCC structure, as shown in Figure 6a. The Sn_0.1 alloy coating (Figure 6b) appeared in black irregular blocks with uniform distribution after Sn element addition. The microstructures of the Sn_0.2 alloy (Figure 6c), Sn_0.3 alloy (Figure 6d), and Sn_0.5 alloy (Figure 6e) showed no obvious significant change, but the size of the black irregular block gradually increased, and the content gradually increased. The EDS of Figure 7b was obtained by scanning areas 1 and 2 of the Sn₀.₂ coating (Figure 7a). As can be seen, the black area (Figure 7b Spectrum 1) contained more Mn and Ni elements than the white area (Figure 7b Spectrum 2), indicating that Mn, Ni, and Sn elements were more concentrated in the black block area. According to XRD, Jade fitting analysis, and two-phase mass fraction, the white skeleton tissue in the Sn_0.2 coating was the FCC phase, and the black massive structure was the second phase MnNi₂Sn.

Table 2 depicts the EDS area scanning results of the chemical compositions (at.%) of each phase of the CoCrFeNiMnSn_x HEA coatings. Regardless of Sn content, Co and Cr elements were distributed in the FCC phase and MnNi₂Sn with more atoms. After adding the Sn element to the alloy, almost all the Sn element was enriched in the MnNi₂Sn phase. The Ni and Mn element content in the MnNi₂Sn was higher than that in the FCC phase, whereas the Fe, Co, and Cr element content in MnNi₂Sn is lower. These results show that the Sn element dissolved easily in the Mn and Ni elements, and hence formed a large proportion in the MnNi₂Sn phase.

3.2. Analysis of Mechanical Properties of CoCrFeNiMnSn_x HEAs Coating

3.2.1. Hardness Analysis of Cladding Layer

Figure 8 shows the microhardness distribution of CoCrFeNiMnSn_x HEA coatings with different Sn content from the cladding layer to the substrate. As can be seen, as the Sn content increased, the alloy hardness increased and then decreased. The Sn₀, Sn_0.1, Sn_0.2, Sn_0.3, and Sn_0.5 coatings exhibited average Vickers hardnesses of 288, 327, 417, 369, and 353 HV_0.5, respectively. The hardness value of the Sn_0.2 coating reached the highest value, which was nearly 45% higher than the average hardness of the Sn₀ alloy coating. This demonstrated that the addition of the Sn element improved the hardness of the alloy coating, but the hardness of the coating decreased with further addition of the Sn element, which was primarily related to the porosity of the coating surface. This is considered further below.

Figure 9 shows the SEM images of the CoCrFeNiMnSn_x HEA coatings. Figure 10b,c show the phase volume fraction and surface porosity of each coating after calculation and processing with Image J software. When comparing the relationship between the change in coating surface hardness (Figure 10a), and the volume fraction of the MnNi₂Sn second phase and surface porosity are compared, as the Sn element content increased, the mass fraction of the MnNi₂Sn phase and porosity of CoCrFeNiMnSn_x HEA coatings gradually increased, whereas the hardness first increased, reaching a maximum at Sn_0.2, and then gradually decreased. There are two primary reasons for the observed increase in hardness. First, Sn was dissolved in the FCC solid solution phase. Because Sn has a much larger atomic radius than the other elements, it inevitably causes distortion in the original FCC lattice, expanding the FCC lattice, and generating a strain field, thereby impeding dislocation movement. Therefore, deformation of the material caused by external pressure is hindered; that is, solid solution strengthening increases the hardness of the alloy. Second, the Sn element promotes the formation of the MnNi₂Sn second phase, and as Sn levels rise, so does the volume fraction of the MnNi₂Sn phase, which aids in the precipitation of the second phase. The decrease in the hardness of the HEA coatings after Sn_0.2 occurred because, after the addition of excessive Sn powder, the content of the molten pool was relatively large after solidification, so that the infiltration of the Sn powder and CoCrFeNiMn HEA powder was insufficient, resulting in increased porosity and decreased compactness. At this time, increasing the mass fraction of the MnNi₂Sn phase could not prevent the hardness of the CoCrFeNiMnSn_x HEA coatings from decreasing, resulting in a decrease in the coating’s overall hardness.

3.2.2. Tribological Analyses of the Cladding Layer

Figure 11 shows the volume losses of the CoCrFeNiMnSn_x HEA cladding layer and Q235 substrate due to wear. The wear volume of the Sn₀ alloy coating (0.7 mg/mm³) was about five times lower than that of the base Q235 alloy steel (3.6 mg/mm³), significantly reducing the surface wear of the Q235 alloy steel. When the content of the Sn element increased from Sn₀ to Sn_0.2, the volume wear of the alloy coating gradually decreased by approximately 0.2 mg/mm³. The volume loss of the alloy coating increased when the Sn element content increased from Sn_0.2 to Sn_0.5. The volume wear was about 0.5 mg/mm³ when the Sn element content reached Sn_0.2. In summary, the Sn alloy coating reduced wear loss and the Sn_0.2 alloy coating was the best wear-resistant coating.

Figure 12 shows the friction coefficient of the coating samples, comprising different materials via the Q235 alloy steel substrate samples and laser cladding. As can be seen, the friction coefficient rapidly rose within 0~200 s and continued to fluctuate within 200~600 s. This was because the CoCrFeNiMnSn_x HEA coatings were in the running-in stage to the worn piece during the early stage of friction and wear, and the surface roughness of the alloy will have risen rapidly due to the falling off of small particles, as will have the friction coefficient. The fine particles on the coating’s surface began to fall off, and the contact area between the alloy coating and the abrasion piece grew larger and more stable, causing the friction coefficient to fluctuate. In the same friction and wear environment, the average friction coefficient of the Sn-containing alloy coatings decreased compared to the average friction coefficient of the Sn₀ coating (Figure 12a) obtained of 0.283. Among them, the Sn_0.1 (Figure 12b), Sn_0.2 (Figure 12c), Sn_0.3 (Figure 12d) and Sn_0.5 (Figure 12e) alloy coating average friction coefficients of the layers were 0.129, 0.115, 0.119, and 0.123, respectively. The friction coefficient of the Sn_0.2 alloy coating was approximately 2.5 times lower than that of the Sn₀ alloy coating. This demonstrated that the addition of Sn could effectively reduce wear loss.

Figure 13 shows the three-dimensional (3D) morphologies of the CoCrFeNiMn HEA coating after adding the Sn element wear under the same load. As can be seen, the wear area of the coating had pear grooves parallel to the grinding direction. This was because the abrasion particles formed during the wear process were subjected to both normal and tangential stresses. The normal stress pressed the abrasion particles into the matrix, whereas the tangential stress caused the abrasion particles to move in a circular motion along with the wear. This eventually led to the formation of the pear ditch. As shown in Figure 13a, pear ditches and fragments appeared on the wear surface of the Sn₀ alloy coating. The flaky debris indicated that the coating experienced periodic layered fracture during the friction and wear process, indicating that the wear mechanisms of the Sn₀ alloy coating were primarily abrasion wear and minor oxidative wear. As shown in Figure 13b, the wear surface ditch of the Sn_0.2 alloy coating increased, indicating that, in the process of friction and wear, the wear damage caused by the high Sn coating surface became larger. Therefore, the wear mechanisms of the Sn_0.2 alloy coating were mainly abrasive wear and oxidation wear.

Figure 14 shows SEM images of the wear scar microstructure of the Sn₀ and Sn_0.2 alloy coatings. As can be seen, there were numerous pear grooves, small holes, and a small number of fragments produced on the surface as a result of friction and wear of the cladding layer. The tangential stress transmitted by the friction pair caused the pear grooves. The debris was due to a certain amount of oxidative wear. The Sn₀ alloy coating (Figure 14a) exhibited regional spalling, which was due to the increase in cyclic alternating stress on the coating surface during the wear process, with a large number of spalling layers appearing on the surface of the alloy. Figure 14b is the enlarged area of the peeling ditch. The main wear mechanisms of the Sn₀ alloy coatings were abrasive wear and delamination wear, accompanied by a small amount of oxidative wear. The surface of the Sn₀.₂ alloy coating (Figure 14c) had more tiny fragments, and there were no peeling layers. Figure 14d is the enlarged area of the fragment. Therefore, the alloy coating’s wear mechanisms were mainly dominated by abrasive wear and oxidative wear.

4. Conclusions

• (1). The CoCrFeNiMn HEA coating showed a uniform FCC phase structure. When Sn was added to the alloy, an Sn-rich MnNi₂Sn structural phase formed. The coatings were made up of two structural phases: FCC and MnNi₂Sn. With increase in Sn content, the content of the MnNi₂Sn phase increased in the coating, and the coatings were composed of the FCC phase and MnNi₂Sn phase.

• (2). The microhardness of CoCrFeNiMnSn_x HEA coatings increased with increase in Sn content. The microhardness of the Sn_0.2 coating reached 417 HV_0.5. The atomic radius of Sn was much larger than that of the other elements when dissolved in FCC solid solution, resulting in solid solution strengthening and improving the hardness of the coating. Second, the Sn promoted the formation of the second phase of MnNi₂Sn, which played a role in strengthening the precipitation of the second phase. The porosity of the CoCrFeNiMnSn_x HEA coating increased as the Sn content increased, whereas the compactness decreased. Porosity was the primary cause of the hardness change in the CoCrFeNiMnSn_x HEA coating, resulting in a decrease in the coating’s overall hardness.

• (3). The wear loss of the Q235 base alloy was 3.6 mg/mm³, while that of the Sn_0.2 alloy coating was only 0.5 mg/mm³. The friction coefficient of the Sn_0.2 alloy coating was 0.115, indicating that the increase in Sn element content improved the surface wear resistance.

Author Contributions

Methodology, J.S.; Writing—original draft, J.S.; Investigation, D.Z. and C.Z.; Project administration, D.Z.; Validation, C.Z. and M.Z.; Data curation, W.S. and M.Z.; Funding acquisition, W.S. and B.J.; Writing—review & editing, D.S. and S.D.; Supervision, D.S.; Visualization, B.J. and S.D.; Formal analysis, X.X. and L.W.; Resources, X.X. and L.W. All authors have read and agreed to the published version of the manuscript.

Institutional Review Board Statement

Not applicable.

Informed Consent Statement

Not applicable.

Data Availability Statement

Not applicable.

Conflicts of Interest

The authors declare that they have no conflict of interest.

Publisher’s Note: MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

Figures and Tables

Figure 1. Electron microscope images of CoCrFeNiMn HEA powder (a) 200 and (b) 500 times.

Figure 2. Energy spectrum of CoCrFeNiMn HEA powder.

Figure 3. XRD patterns of CoCrFeNiMnSnx HEA coatings.

Figure 4. OM images of Sn0 coating: (a) longitudinal section of the coating; (b) bottom area of the coating; (c) topside area of the coating; (d) bonding area between the coating and the substrate.

Figure 5. EDS line scan area (a) and element distribution (b) of Sn0 alloy coating.

Figure 6. Microstructure and morphologies of (a) Sn0, (b) Sn0.1, (c) Sn0.2, (d) Sn0.3, and (e) Sn0.5 alloy coatings.

Figure 7. EDS scanning area (a) and corresponding energy spectrum (b) of Sn0.2 alloy coating.

Figure 8. Cross-sectional microhardness distribution of CoCrFeNiMnSnx HEAs with different Sn content.

Figure 9. SEM of (a) Sn0, (b) Sn0.1, (c) Sn0.2, (d) Sn0.3, (e) Sn0.5 alloy coatings.

Figure 10. (a) Surface microhardness, (b) mass fraction of MnNi2Sn phase and (c) surface porosity of CoCrFeNiMnSnx HEA coatings.

Figure 11. Volume losses of CoCrFeNiMnSnx HEAs coating and Q235 base material due to wear.

Figure 12. Variation curve of average friction coefficients of CoCrFeNiMnSnx HEAs (a) Sn0, (b) Sn0.1, (c) Sn0.2, (d) Sn0.3, (e) Sn0.5, (f) fircion cofficient.

Figure 13. CoCrFeNiMnSnx HEAs coatings (a) Sn0 and (b) Sn0.2 under the same pressure wear 3D morphologies.

Figure 14. SEM morphology of wear scars of CoCrFeNiMnSnx HEA coatings (a) Sn0, (b) Sn0 peeling ditch magnified area, (c) Sn0.2 and (d) Sn0.2 fragment magnified area.

Table 1

Chemical composition of CoCrFeMnNiSn_x and Q235 (wt. %).

Materials	Sn Element Percentage %	Co Element Percentage %	Cr Element Percentage %	Fe Element Percentage %	Ni Element Percentage %	Mn Element Percentage %
Sn0	0	19.66	19.17	19.32	19.38	21.28
Sn0.1	1.96	19.42	19.01	18.56	18.59	20.24
Sn0.2	3.92	18.89	18.45	18.25	18.26	19.62
Sn0.3	5.88	18.15	18.05	17.97	18.01	19.34
Sn0.5	9.80	17.16	17.24	17.02	17.31	18.72
Q235	0	21.28	20.01	20.22	20.01	20.23

Table 2

Chemical composition of each phase of CoCrFeNiMnSn_x HEA coatings (at.%).

Alloy	Phase Structure	Cr	Mn	Fe	Co	Ni	Sn
Sn0	FCC	21.54	19.00	15.59	30.35	13.52	-
Sn0.1	FCCMnNi2Sn	20.0925.8	9.1412.40	23.5412.08	32.2130.14	14.5217.68	0.51.9
Sn0.2	FCCMnNi2Sn	29.1824.26	8.1215.32	17.945.95	30.6825.14	13.1323.44	0.955.98
Sn0.3	FCCMnNi2Sn	30.2927.32	10.1015.26	13.937.84	33.9124.76	10.5714.49	1.208.33
Sn0.5	FCCMnNi2Sn	30.3024.96	9.3012.23	16.376.75	30.6729.33	10.5516.29	2.8110.44

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